Borohydride Synthesis for Development of Porous Zirconium Oxide Nanopowders for Novel Applications

Home , Borohydride

BOROHYDRIDE SYNTHESIS FOR DEVELOPMENT OF POROUS ZIRCONIUM OXIDE NANOPOWDERS FOR NOVEL APPLICATIONS

Nadiya Bihary Nayak

Department of Ceramic Engineering National Institute of Technology Rourkela

BOROHYDRIDE SYNTHESIS FOR DEVELOPMENT OF POROUS ZIRCONIUM OXIDE NANOPOWDERS FOR NOVEL APPLICATIONS

Dissertation submitted in partial fulfillment

of the requirements of the degree of

Doctor of Philosophy

Ceramic Engineering

Nadiya Bihary Nayak

(Roll Number: 510CR101)

based on research carried out

under the supervision of

Prof. Bibhuti B. Nayak

Department of Ceramic Engineering National Institute of Technology Rourkela

January, 2017

Department of Ceramic Engineering National Institute of Technology Rourkela

Date: 12/01/2017 Certificate of Examination

Roll Number: 510CR101 Name: Nadiya Bihary Nayak Title of Dissertation: Borohydride synthesis for development of porous zirconium oxide nanopowders for novel applications

We the below signed, after checking the dissertation mentioned above and the official record book (s) of the student, hereby state our approval of the dissertation submitted in partial fulfillment of the requirements of the degree of Doctor of Philosophy in Department of Ceramic Engineering at National Institute of Technology Rourkela. We are satisfied with the volume, quality, correctness, and originality of the work.

Bibhuti Bhusan Nayak Principal Supervisor

______Swadesh Kumar Pratihar Himanshu Bhushan Sahu Member, DSC Member, DSC

Japes Bera Kanhu Charan Barick Chairman, DSC External Examiner

______Bibhuti Bhusan Nayak Head of the Department

Department of Ceramic Engineering National Institute of Technology Rourkela

Prof. Bibhuti Bhusan Nayak (Ph.D., IIT Bombay) Associate Professor and HOD Department of Ceramic Engineering National Institute of Technology Rourkela PIN: 769 008, INDIA Email: [email protected]; [email protected] Phone: 0661-246-2209/2201 (O); +91-9861422166 (M)

Date: 12/01/2017

Supervisor’s Certificate

This is to certify that the work presented in the dissertation entitled “Borohydride synthesis for development of porous zirconium oxide nanopowders for novel applications” submitted by Nadiya Bihary Nayak, Roll Number 510CR101, is a record of original research carried out by him under my supervision and guidance in partial fulfillment of the requirements of the degree of Doctor of Philosophy in Department of Ceramic Engineering. Neither this dissertation nor any part of it has been submitted earlier for any degree or diploma to any institute or university in India or abroad.

Bibhuti Bhusan Nayak

DEDICATED TO

My elder brother Vakta Charan Nayak

Declaration of Originality

I, Nadiya Bihary Nayak, Roll Number 510CR101 hereby declare that this dissertation entitled “Borohydride synthesis for development of porous zirconium oxide nanopowders for novel applications” presents my original work carried out as a doctoral student of NIT Rourkela and, to the best of my knowledge, contains no material previously published or written by another person, nor any material presented by me for the award of any degree or diploma of NIT Rourkela or any other institution. Any contribution made to this research by others, with whom I have worked at NIT Rourkela or elsewhere, is explicitly acknowledged in the dissertation. Works of other authors cited in this dissertation have been duly acknowledged under the sections “Reference” or “Bibliography”. I have also submitted my original research records to the scrutiny committee for evaluation of my dissertation. I am fully aware that in case of any non-compliance detected in future, the Senate of NIT Rourkela may withdraw the degree awarded to me on the basis of the present dissertation.

Date: 12/01/2017 NIT Rourkela Nadiya Bihary Nayak

Acknowledgments With deep regards and profound respect, I avail this opportunity to express my deep sense of gratitude and indebtedness to my supervisor and Head of the Department of Ceramic Engineering Prof. Bibhuti B. Nayak, Department of Ceramic Engineering, National Institute of Technology, Rourkela, for introducing the present research topic and for his inspiring guidance, constructive criticism and valuable suggestion throughout this research work. It would have not been possible for me to bring out this thesis without his guidance, support and and constant encouragement. Under his guidance, I successfully overcame many difficulties and learned a lot. Despite of his busy schedule, he used to review my thesis progress and give his valuable suggestions. His infinite haunts for finding something new, actually motivated me to work with his noble thoughts. My heartfelt gratitude to Prof. Bibhuti B. Nayak, a person I admire not only for his learning but also for his extended cooperation rendered to me as and when I required. I feel myself extremely lucky to have the company of him. From prologue to epilogue, he is always with me. I owe a deep debt of gratitude to him. Words fail me to express my appreciation to Madam Dr. Aparna Mondal, Department of Chemistry, NIT Rourkela for her motivation, constantly encourage, support, generous care and the homely feelings help to push me and motivate me. Master Krishu with smiling face helps me to reduce my internal depression and motivate me to go ahead. I am highly grateful to all DSC members Prof. Santanu Bhattacharyya, Prof. Japesh Bera, Prof. S. K. Pratihar of the Department of Ceramic Engineering and Prof. Himanshu Bhushan Sahu of the Department of Mining Engineering for their valuable advice, constructive criticism and their discussions around my work. I am also highly grateful to Prof. Santanu Kumar Behera for his valuable advice and discussions during my work. I am also very much grateful to all other faculties of Department of Ceramic Engineering, Department of Chemistry, Department of Physics and Department of Metallurgical and Materials Engineering NIT, Rourkela, for their unconditional support and advice. I thank to all my teachers and professors, from academic and nonacademic levels, who inspire me to be wise and knowledgeable. I am truly indebted to all who have supported me and brought me so far. I gratefully acknowledge Dr. Binod Bihari Nayak, Principal Technical Officer CSIR-IMMT, Bhubaneswar, for making arrangement of adsorption behaviour of porous zirconia nanoparticles. Based on this research work, my supervisor received one sponsored project titled “Fabrication of porous and hollow metal oxide nanostructured materials via gas-bubbles template method for different applications” by DST-Nano Mission. Thanks to DST- Nano Mission for accepting the novel idea and sanctioning the research project for three years. I would like to thank all of my friends especially Subrat, Sarat, Geeta, Smruti, Juga, Bhabani, Abhisekh, Jayrao, Sreenivasulu, Shubham, Swarnima, Rahul, Bappaditya, and Niladri for providing all joyful environment in the lab and helping me out in different ways. I should extend my gratitude to Mr.P.K. Mohanty, Mr.S.K.Sahoo, Mr.L. Dhal, Mr.N.Barik, Mr. S. Chakraborty, Mr. G. Behera, Mr.Aravind, Mr. R. Pattanaik, Mr. U. Sahoo, and Mr. G. Dash all from NIT, Rourkela, for their technical help for the experimentation involved in the Thesis. It would be an immense pleasure for me like to thank my grandparents Late Brajabandhu Nayak and Late Smt. Sumitra Nayak and other family members for their endless love. I also like to thank my eldor brother Bairagi Charan Nayak, elder sister Subhadra Nayak, Banita Nayak and brother-in-law Sanjay Kumar Bariki who bear a great feeling and love towards me. It is also the time to thank my lovely nephew Dibyajit and niece Sama and Banisri who bear a great feeling and love towards me. I feel a deep sense of gratitude for my father Shri Mahesh Prasad Nayak and mother Smt. Susama Nayak who formed a part of my vision and taught me the good things that really matter in life. I am thanking my beloved wife Jyostnamayee Nayak for all the hope, true love, affection, caring, concern and constant encouragement she lends. I would like to thank her family members for giving me love and support. My son Master Naibedya gives me many reasons to smile every day..

Date: 26/07/2016 (Nadiya Bihary Nayak)

Abstract The concept of this Ph.D. research work is to find out a suitable in-situ gas-bubbles evolving aqueous precipitating agent to develop loose or agglomeration free powders of amorphous zirconium hydroxide in the as-synthesized condition and further to develop a moderate temperature stable amorphous or crystalline zirconium oxide nanopowders with porous structure for use in different novel applications. This research work discloses the importance of aqueous sodium borohydride (NaBH4) as a precipitating agent towards the development of moderate temperature stable amorphous and crystalline nature of zirconium oxide with porous structure. The role of boron species as well as in-situ hydrogen (H2) gas-bubbles present in aqueous NaBH4 are presented in this research work for the development of porous zirconium oxide nanopowders. Also, we are reporting, a new concept of reaction mechanism between aqueous ZrOCl2·8H2O and NaBH4 with the help of Fourier Transformation Infra-Red (FTIR) spectroscopy for producing zirconium hydroxide loose powders in the as-synthesized condition. In addition, we are emphasis on advantages of aqueous NaBH4 for the development of porous zirconium oxide nanopowders for removal of toxic ion such Pb (II) for environmental applications. The removal efficiency of Pb (II) as a function of time as well as adsorption kinetic mechanism and regeneration of Pb (II) loaded zirconium oxide sample was studied in detail. Quick adsorption of almost hundred percentages of toxic lead ions from water solution within 15 minutes as well as regeneration study suggests that the porous zirconium oxide can be a potential adsorbent for removal of toxic ions for environmental applications. Further, a novel information on the temperature-mediated phase transformation, pore geometry as well as pore hysteresis transformation of in-born porous zirconium hydroxide nanopowders with the help of X-ray diffraction (XRD), Brunauer–Emmett–Teller (BET) isotherm and Transmission Electron Microscopy (TEM) images was also discussed in this research work.The novel hydrogen (H2) gas-bubbles assisted borohydride synthesis route led to develop thermally stable porous zirconium hydroxide/oxide nanopowders with an adequate pore size, pore volume, and surface area and thus these porous materials are further suggested for promising use in different areas of applications. Moreover, using borohydride synthesis route, the rare earth ions [Eu3+ (5 mol%), Tb3+(5 mol%), and mixture of Eu3+(2 mol%) and Tb3+(5 mol%)] are incorporated in porous zirconium oxide to stabilize t- or c-phase of porous zirconia as well as to develop multi-colour phosphor nanomaterials using various excitation wavelengths ranging from 205 nm to 350 nm. Structure, powder morphology, rare-earth ion distribution, luminescence and sample colour under UV light were studied for these phosphor nanomaterials for suitability in lighting applications. In summary, the borohydride synthesis route using sodium borohydride is found to be a potential precipitating reagent for the development of moderate temperature stable amorphous as well as crystalline t-zirconia with a porous structure for possible applications in the field of adsorption of heavy metal ions as well as rare earth based porous zirconia powders for lighting applications.

Keywords: Zirconium oxide; Porous; Gelation; Precipitation; H2 (hydrogen) gas-bubbles; Sodium borohydride; Adsorption; Phosphor.

Contents Page No Supervisor’s Certificate Dedication Declaration of Originality Acknowledgments Abstract List of Figures List of Tables

Chapter 1: Introduction 1-5 1.1: Background and purpose of selecting zirconium oxide 1 1.2: Background and purpose of selecting sodium borohydride 3 1.3: Advantages of loose and porous zirconium oxide powders 4 1.4: Organization of thesis 5 Chapter 2: Literature review 6-30 2.1: Stabilization of tetragonal zirconium oxide 6 2.2: Materials synthesized using sodium borohydride 9 2.3: Gas-bubbles derived porous nanopowders 11 2.4: Pore morphology of porous nanoparticles 14 2.5: Adsorption of toxic ions using porous powders 17 2.6: Luminescence behaviour of zirconium oxide 20 2.7: Summary of literature 26 2.8: Statement of the problem 27 2.9: Objectives 30 Chapter 3: Experimental work 31-34 3.1: Raw Materials 31 3.2: Powder synthesis 31 3.3: General characterization 31 3.3.1: Thermal 31 3.3.2: Phase analysis 32 3.3.3: Particle morphology 32

3.3.4: Particle and crystallite size 32 3.3.5: Distribution of rare earth ions 32 3.3.6: Fourier transformation infrared spectroscopy (FTIR) 32 3.3.7: BET-isotherm and surface area 33 3.3.8: Adsorption of toxic ions and regeneration study 33 3.3.9: Photoluminescence 34 3.3.10: Colour purity (R/O ratio) 34 3.3.11: CIE-chromaticity coordinate graph 34 3.3.12: Powder colour under UV light source 34 Results and discussion Chapter 4: Motivation for selecting borohydride synthesis route to prepare 34-40 zirconium oxide nanopowders 4.1: Experimental 35 4.2: Results and discussion 36 4.2.1: Thermal 36 4.2.2: Structure 36 4.2.3: Powder morphology 39 4.3: Remarks 40 Chapter 5: Gelation-precipitation reaction mechanism of borohydride 41-48 synthesis to prepare zirconium oxide nanopowders 5.1: Experimental 41 5.2: Results and discussion 41

5.2.1: Gelation-precipitation mechanism between Zr-salt and NaBH4 41 5.2.2: Thermal 47 5.3: Remarks 48 Chapter 6: Enhanced activation energy of crystallization, development of 49-62 porous zirconia and application of porous zirconia for removal of Pb (II) toxic ions 6.1 Experimental 49 6.2 Results and discussion 50 6.2.1: Thermal 50 6.2.2: Structure 51 6.2.3: Powder morphology 52 6.2.4: Pore evolution mechanism 55 6.2.5: Adsorption kinetics mechanism 57 6.2.6: Regeneration of the adsorbent 61 6.3: Remarks 62

Chapter 7: Temperature-mediated phase transformation, pore geometry 63-77 and pore hysteresis transformation of borohydride derived in- born porous zirconium hydroxide nanopowders 7.1: Experimental 63 7.2: Results and discussion 64 7.2.1: Structure 64 7.2.2: Powder morphology and BET-isotherm 65 7.3: Remarks 77

Chapter 8: Phase formation, pore morphology and photoluminescence 78-112 behaviour of rare earth (single-doped of Eu3+ or Tb3+ and co- doped of Eu3+ and Tb3+) based porous zirconia nanopowders 8.1: Experimental 78 8.2: Results and discussion 79 8.2.1: Photoluminescence behaviour of un-doped porous zirconia 79 8.2.2: Remarks 8.2.3: Phase analysis of doped (Eu3+/Tb3+) co-doped (Eu3+-Tb3+) 85 porous zirconia nanopowders 8.2.4: Powder morphology of doped (Eu3+/Tb3+) co-doped (Eu3+- 86 Tb3+) porous zirconia nanopowders 8.2.5: Distribution of rare earth ions in porous zirconia particles 88 8.2.6: Pore morphology of rare earth doped and co-doped porous 92 zirconia 8.2.7: Photoluminescence behaviour of Eu-doped porous zirconia 94 8.2.8: Photoluminescence behaviour of Tb-doped porous zirconia 100 8.2.9: Photoluminescence behaviour of co-doped (Eu-Tb) porous 105 zirconia 8.2.10: Proposed mechanism for showing different colour at 110 different excitation wavelength 8.3: Remarks 112

Conclusions 113-115 Scope for future work 115 References 116-123 CURRICULUM VITAE 124

List of Figures Page No Fig. 2.1: Types of sorption isotherms and hysteresis loops 15 Fig. 4.1: DSC-TG curve of as-prepared zirconia synthesized using (a) 36 NaBH4 (b) NH4OH. Fig. 4.2: XRD pattern of un-washed as-synthesized powders, derived via 37 NaBH4 and NH4OH. XRD pattern of washed sample prepared using NaBH4 shows amorphous in nature.

Fig. 4.3: XRD pattern of calcined zirconia powders, derived via NaBH4. 37 [‘t’ stands for t-phase of zirconia]

Fig. 4.4: XRD pattern of calcined zirconia powders derived from NH4OH. 38 [‘t’ and ‘m’ stands for t- and m-phase of zirconia, respectively.]

Fig. 4.5: TEM micrograph of as-synthesized powder derived from (a) NaBH4 39 and (b) NH4OH.

Fig. 4.6: TEM micrograph of calcined (600°C) powders, derived via (a) NaBH4 40 and (b) NH4OH. Fig. 5.1: FTIR spectra of un-washed gel sample 43 Fig. 5.2: Schematic representation of three-dimensional network of 44 Zr(OH)4 with boron species Fig. 5.3: FTIR spectra of un-washed precipitate sample. 44 Fig. 5.4: Solid pieces of boron complex were phase separated from 45 precipitation Fig. 5.5: FTIR spectra of solid borate sample 46 Fig. 5.6: FTIR spectra of washed precipitate and calcined (800 °C) sample 47 Fig. 5.7: DSC of borohydride derived as-prepared zirconia powders. [a and b 48 stand for gelation and precipitation derived sample prepared using NH4OH; c and d stand for gelation and precipitation derived sample prepared using NaBH4] Fig. 6.1: DSC curves of as-prepared zirconia synthesized through three 50 different routes. Fig. 6.2: XRD patterns of zirconia powder calcined at 600°C. 51 Fig. 6.3: XRD patterns of zirconia powder calcined at 800 °C. 51

Fig. 6.4: TEM micrographs of calcined (800°C) ZrO2 powders, 52 synthesized through (a) gelation, (b) precipitation and (c) constant pH. Fig. 6.5: TEM micrographs of as-synthesized powders calcined at (a) 400 54 °C, (b) 600 °C and (c) 800 °C. Inset of each micrographs show electron diffraction pattern. Fig. 6.6: Higher magnified TEM micrograph (a) and BJH curve (b) of 55 porous zirconium oxide, calcined at 800 °C. Fig. 6.7: Removal percentage of Cr(VI) and Pb(II) with different interval 57 of time.

Fig. 6.8: Pseudo-first order (a) and pseudo-second order (b), Elovich (c), 58 Intra-particle diffusion (d) and Bangham (pore diffusion) (e) kinetic plot for adsorption of Pb (II) by porous zirconium oxide powders. Fig. 6.9: Removal (during adsorption) and recovery (during desorption) 61 percentage of Pb(II) as a function of number of cycles. Fig. 7.1: X-ray diffraction pattern of the as-synthesized as well as calcined 65 samples (a) [Note: ‘t’ stands for tetragonal and ‘m’ stands for monoclinic zirconium oxide] and crystallite size and volume percentage of tetragonal phase as a function of calcination temperature (b). To determine the activation energy (Q) of particle growth, Figure 1 (b) was re-plotted to Figure 1 (c). Fig. 7.2: TEM micrograph (a) of as-synthesized porous zirconium 66 hydroxide and the corresponding electron diffraction pattern in the inset of (a) indicates that the powder was amorphous in nature. BET-isotherm (b) and pore size distribution (c) of the as- synthesized porous zirconium hydroxide. Inset of (c) is the enlarge view of the pore size distribution. Fig. 7.3: TEM micrograph (a) of the as-synthesized powders calcined at 68 500 °C and the corresponding electron diffraction pattern in the inset of (a) indicates that the powder was amorphous in nature. BET-isotherm (b) and pore size distribution (c) of the calcined (500°C) porous powders. Inset of (c) is the enlarge view of the pore size distribution. Fig. 7.4: TEM micrograph (a) of the as-synthesized powders calcined at 70 600 °C and the corresponding electron diffraction pattern in the inset of (a) indicates that the powder was in crystalline in nature. BET-isotherm (b) and pore size distribution (c) of the calcined (600°C) porous zirconium oxide. Inset of (c) is the enlarge view of the pore size distribution. Fig. 7.5: TEM micrograph (a) of the as-synthesized powders calcined at 72 800 °C and the corresponding electron diffraction pattern in the inset of (a) indicates that the powder was in crystalline in nature. Higher magnified TEM image (b) indicates the presence of two types (spherical as well as lamellar) of pore. Inset of (b) indicates the presence of ultra-fine loose particles on the surface of the particle. BET-isotherm (c) and pore size distribution (d) of calcined (800°C) zirconium oxide. Inset of (d) is the enlarge view of the pore size distribution. Fig. 7.6: TEM micrograph (a) of the as-synthesized powders calcined at 74 900 °C. BET-isotherm (b) and inset of (b) shows the pore size distribution. TEM micrograph (c) indicates the non-porous zirconium oxide, calcined at 1000 °C. Fig. 8.1: Kubelka-Munk plot of borohydride derived zirconia. 79 Fig. 8.2: Typical excitation spectra versus wavelength of borohydride 80 synthesized un-doped zirconia.

Fig. 8.3: Photoluminescence spectra of borohydride derived un-doped 81 zirconia powders. Fig. 8.4: Excitation wavelength variation chromaticity diagram for 83 borohydride derived un-doped zirconia (a) and powder colour under UV light source (b) indicate that the sample is nearly violet-blue in colour. Fig. 8.5: XRD patterns of calcined (600 °C and 800 °C) samples of Eu3+- 85 doped (a), Tb3+-doped (b) and Eu3+-Tb3+-co-doped zirconia (c). Fig. 8.6: TEM micrographs of Eu3+-doped (a), Tb3+-doped (b) and Eu3+- 87 Tb3+-co-doped (c) zirconia, calcined at 800 °C. Fig. 8.7 HAADF-STEM images of (a) Eu3+-doped, (b) Tb3+-doped and 88 (c) co-doped (Eu3+- Tb3+) zirconia. Fig. 8.8: EDX line profile of Eu3+-doped porous zirconia particles. Inset 89 represents the HAADF-STEM image of Eu3+-doped porous zirconia. Line scan along the direction denoted by the green line. Dotted line represents that the Eu signal was not present on the surface of the zirconia particle. Fig. 8.9: EDX line profile of Tb3+-doped porous zirconia particles. Inset 90 represents the HAADF-STEM image of Tb3+-doped porous zirconia. Line scan along the direction denoted by the green line. Dotted line represents that the Tb signal was not present on the surface of the zirconia particle. Fig. 8.10: EDX line profile of co-doped (Eu3+-Tb3+) porous zirconia 91 particles. Inset represents the HAADF-STEM image of co-doped porous zirconia. Line scan along the direction denoted by the green line. Dotted line represents that the Eu and Tb signal was not present on the surface of the zirconia particle. Fig. 8.11: EDX elemental mapping (b) and line scan (c) of co-doped (Eu3+- 92 Tb3+) porous zirconia particles from HAADF image (a). Fig. 8.12: BET-isotherm and pore size distribution (inset of Fig. 8.12) of 93 (a) Eu3+, (b) Tb3+ and (c) co-doped (Eu3+-Tb3+) porous zirconia sample, calcined at 800 °C. Fig. 8.13: Excitation spectra of Eu-doped sample calcined at 600 °C. 94 Fig. 8.14: Photoluminescence spectra of calcined (600 °C) Eu3+-doped 95 zirconia powders. Fig. 8.15: Photoluminescence spectra of calcined (800 °C) Eu3+-doped 96 zirconia powders. Fig.8.16: Excitation wavelength dependent R/O ratio of Eu-doped 97 zirconia. Fig. 8.17: Excitation wavelength variation chromaticity diagram for Eu- 99 doped zirconia. Fig. 8.18: Powder colour image of Eu-doped zirconia powder under UV 99 lamp. Fig. 8.19: Photoluminescence spectra of calcined (600 °C) Tb-doped 100 zirconia powders.

Fig. 8.20: Photoluminescence spectra of calcined (800 °C) Tb-doped 102 zirconia powders. Fig. 8.21: Excitation wavelength variation chromaticity diagram for Tb- 104 doped zirconia. Fig. 8.22: Powder colour image of Eu-doped zirconia powder under UV 104 lamp. Fig. 8.23: Photoluminescence spectra of calcined (600 °C) co-doped (Eu- 105 Tb) zirconia. Fig. 8.24: Photoluminescence spectra of calcined (800 °C) co-doped (Eu- 106 Tb) zirconia. Fig.8.25: Excitation wavelength dependent R/O ratio of co-doped (Eu-Tb) 107 zirconia. Fig. 8.26: Excitation wavelength variation chromaticity diagram for Tb- 109 doped zirconia Fig. 8.27: Powder colour image of co-doped (Eu-Tb) zirconia powder 109 under UV lamp. Fig. 8.28: Schematic diagram indicating the penetration depth of near VUV 111 and UV light on the porous zirconia particle (a) and different penetration depth of UV light led to show colour tuning in rare earth based porous zirconia particle. [Black small circles indicate the rare earth ions within the zirconia particle]

List of Tables Page No Table 2.1: The equations of the different kinetic models 18 Table: 2.2: PL behaviour of zirconia powders prepared via different 26 synthesis methods. Table 6.1: Kinetic parameters of Pb (II) using porous zirconium oxide as 60 adsorbent Table 7.1: Textural properties of borohydride derived porous nanopowders 76 Table 8.1: Chromaticity co-ordinates of borohydride synthesized zirconia 84 at different excitation wavelengths Table 8.2: Phase and crystallite size of doped and co-doped zirconia 86 sample Table 8.3: Chromaticity co-ordinates of Eu-doped zirconia obtained at 98 different excitation wavelengths Table 8.4: Chromaticity co-ordinates of Tb-doped zirconia obtained at 103 different excitation wavelengths. Table 8.5: Chromaticity co-ordinates of co-doped (Eu-Tb) zirconia 108 obtained at different excitation wavelengths.

Chapter 1

Introduction

This research work discloses the use of a novel precipitating agent to develop thermally stable zirconium oxide [also known as zirconia (ZrO2)], a transition metal oxide nanomaterial with loose as well as porous structure for use in suitable and novel applications. The background and purpose of selecting zirconium oxide material, background and purpose of selecting sodium borohydride (NaBH4), a novel precipitating agent, advantages of loose and porous zirconium oxide powders and organization of thesis are discussed in this chapter. 1.1: Background and purpose of selecting zirconium oxide

Zirconia is the one of the most studied engineering ceramics and it has intrinsic physical and chemical properties of hardness, wear resistance, low coefficient of friction, elastic modulus, chemical inertness, ionic conductivity, electrical properties, low thermal conductivity, and high melting point that make it attractive as an engineering material [1-4]. Pure zirconia at atmospheric pressure exists in three well-known polymorphs in monoclinic

(m), tetragonal (t) and cubic (c) crystal structure with m-ZrO2 being the equilibrium bulk

[2, 5-8] structure at low temperature . The m-ZrO2 at room temperature undergoes m t and

[2, 5-8] t c phase transitions at 1173C and 2370C respectively . The pure ZrO2 was of very limited interest as a structural or engineering ceramic due to the displacive tetragonal (t) to monoclinic (m) phase transformation, which occurs at ~ 950°C on cooling in pure ZrO2 and is accompanied by a shear strain of ~ 0.16 and a volume expansion of ~ 4 % [1-4]. This change of shape in the transforming volume can result in catastrophic fracture and hence, structural unreliability of fabricated components. However, the stabilized zirconium oxide [i.e. stabilization of t-ZrO2 and c-ZrO2 at room temperature] is the most studied engineering ceramics for use in different applications including high temperature fuel cells, thermal

barrier coatings, oxygen sensors, hard ceramic tools, catalysis, adsorbents, luminescent and biomaterials [9-29]. There are various factors that influencing the stabilization of t or c phase at low temperature. One of the factors that stabilizes the tetragonal or cubic phase of zirconia is the addition of suitable amount of stabilizers. The high temperature technologically important phases of zirconium oxide have been stabilized by doping with appropriate amount of tetravalent dopant such as Si4+, Ce4+ and Ge4+ as well as aliovalent dopants such as Al3+, Y3+, Ca2+, Mg2+ and Na+, which give rise to its functional properties and was used for various applications [30-41]. Several methods have been developed to stabilize the t and or c phase of zirconia at low or moderate temperatures. However, stabilization of t or c –

ZrO2 at low or moderate temperature without addition of any stabilizers is still a challenge. Garvie proposed the crystallite size theory that the stabilization of the tetragonal form could be accounted to crystallite size effect [42]. In addition, Garvie discloses the stabilization factor based on the particle size and found that the particle size for stabilizing t-ZrO2 phase [42] must be  30 nm . Garvie developed the equation i.e (Gt-Gm) + Stt – Smm  0, for critical grain size in pure unconstrained ZrO2 based on surface free energy considerations. Where G is the molar free energy and S is the surface area in the single crystal particle. On the basis of the lower value of the surface energy of t-phase (t) in relation to m-phase (m), Garvie considered that in order to stabilize t-phase at low temperature the above equation must be satisfied. Garvie’s hypothesis has been investigated and proven by numerous researchers [43-46]. Other than this factor, the impact of lattice defects on the stabilization of metastable [30, 31,18, 19] t-ZrO2 has also been investigated by many researchers . Further, smaller crystallite size, having higher surface energy agglomerate during synthesis and grows in an uncontrolled way during calcination, which causes the main difficulty to get a stable pure t- [7, 47, 48] ZrO2 at moderate temperature . Various synthesis methods have been explored by many researchers using various precipitating agents to control the particle size below 30 nm, so that the t-ZrO2 can be stable up to moderate temperature without using any stabilizers [2, 49-51]. Generally, zirconia crystals grow in rapid way through agglomeration and achieve a quick crystallization peak temperature at around 450°C, which lead to form stable pure t-

ZrO2 up to 500°C, when using the common precipitating agent such as NH4OH without adding any surfactants or stabilizers [47, 49, 50]. The crystallization peak temperature or activation energy of crystallization of t-ZrO2 can be enhanced by controlling the crystallite size in a smaller dimension using a suitable precipitating agent. It is also possible to grow zirconia with smaller crystallites in a controlled way during wet-chemical synthesis through an efficient way using a suitable precipitating agent. So, the prime theme of this research

work is to synthesize zirconium oxide nanopowders without adding any stabilizers, and also to control the particle size of zirconium oxide limit to 30 nm and further to stable the tetragonal phase zirconium oxide up to moderate or high temperature. Also, to enhance the activation energy of crystallization of zirconia through an efficient way of synthesis using the suitable precipitating agent, so that smaller crystallites pure t-ZrO2 can be stable up to moderate temperature. In this regard, sodium borohydride (NaBH4), a novel precipitating agent is chosen for the development of ultrafine agglomeration free powders in the as- synthesized condition as well as moderate temperature stable tetragonal zirconium oxide nanopowders. So, the purpose of choosing sodium borohydride (NaBH4) as a novel precipitating agent is described as follows. 1.2: Background and purpose of selecting sodium borohydride

Aqueous sodium borohydride (NaBH4) is well known for its reducing property and well-established for the development of metal or metal-boride nanoparticles through reduction method [52-53]. This reagent is also commonly used for preparing alloys, magnetic, [53, 54, 55] magnetic–nonmagnetic composite materials, and also used for H2 storage . However, Zhu and Manthiram [56] and Manthiram et al. [57] showed that borohydrides can also be used effectively to obtain a range of transition-metal oxides. Thus, Nayak et al. [46] explored the potential of this technique to prepare nanocrystalline ZrO2 powders.

So, this research work discloses the importance of aqueous NaBH4 as a precipitating agent towards development of zirconium oxide nanopowders. It is well-known that in aqueous solution, hydrolysis of NaBH4 proceeds to form two species such as tetrahydral [58] boron [B(OH)4¯] and hydrogen (H2) gas bubbles . The role of tetrahydral boron and hydrogen (H2) gas bubbles during synthesis was not explored in zirconium oxide system.

Further, the boron species present in aqueous NaBH4 play an important role to develop boron free zirconium hydroxide in the as-synthesized condition during synthesis and it was discussed in this research work. In addition, the in-situ H2 gas-bubbles also play an important role for formation of loose or agglomeration free porous powders in the as- synthesized condition and the mechanism for development of loose porous powders are also discussed in this research work. However, the detail reaction mechanism on aqueous metal- salts with aqueous NaBH4 for producing metal-oxide systems particularly, zirconium oxide is not yet reported. For utilizing the advantage of borohydride concept to other transition metal oxide systems, a new concept of reaction mechanism between aqueous ZrOCl2·8H2O and NaBH4 was disclosed in this research work. So, the prime theme of this research work

is to synthesize loose porous zirconium hydroxide in the as-synthesized condition with the help of gas-bubbles assisted borohydride route using sodium borohydride. Further, the advantages of these loose powders of porous zirconium hydroxide/oxide are described as follows. 1.3: Advantages of loose and porous zirconium oxide powders

While synthesizing zirconium oxide powder via wet-chemical route using a common precipitating agent such as ammonium hydroxide (NH4OH), without using any stabilizers or surfactants, an amorphous nature of zirconium hydroxide powders were formed in the as- synthesized condition [7, 47, 48]. In general, when the amorphous zirconium hydroxide powders undergo a calcination process, then amorphous phase was converted to crystalline t-phase of zirconium oxide at nearly 450 °C [7, 47, 48]. Due to sudden particle growth, the t- phase of zirconium oxide was difficult to stable up to moderate temperature [7, 47, 48]. If somehow the transition or crystallization temperature from amorphous to crystalline nature t-zirconium oxide increases to higher than 450 °C, then, the amorphous as well as t-phase of zirconium oxide can be stable up to moderate temperature and it may be useful for various potential applications. So, the use of gas-bubbles evolving aqueous sodium borohydride during borohydride synthesis may develop loose porous zirconium hydroxide powders in the as-synthesized condition. The presence of loose nature may inhibit the mass transfer between loose nanoparticles and may help to restrict the coarsening of particles during calcination [59]. So, the loose porous zirconium hydroxide in amorphous nature may enhance the amorphous state as well as stable tetragonal zirconium oxide up to moderate temperature. In the same time, the as-synthesized loose particles may develop moderate or high temperature stable porous particle during calcination process due to coarsening of particles as well as coalescence of voids. So, it is possible to develop thermally stable loose and porous amorphous as well as crystalline zirconium oxide via gas-bubbles assisted borohydride route. Additionally, the presence of nanoscale pores in amorphous zirconium hydroxide as well as tetragonal zirconium oxide nanoparticles play a vital role to enhance the physical and chemical reactivity of nanoparticles and promising potential applications. The porous nanostructured materials have potential applications including sensors, thermal, mechanical, electrical, environmental, luminescence and medical field [60-67]. The developed porous zirconium oxide powders can efficiently be useful for removal of toxic ions for environmental applications. In addition, if the rare earth ions are incorporated in porous structure, then colour tuning is possible at various excitation wavelength. Also, white

light in t-zirconium oxide may be possible by incorporating combination of suitable rare earth ions. The above two applications of borohydride derived porous zirconium oxide nanopowders are discussed in this research work. 1.4: Organization of thesis

The background and purpose of selecting zirconium oxide and NaBH4 along with advantages of loose and porous powders are discussed in the introduction of Chapter 1. Chapter 2 provides a detail literatures based on stabilization of tetragonal zirconium oxide, synthesis of metal-oxide nanopowders using aqueous NaBH4, gas-bubbles derived porous metal-oxide nanopowders, pore morphology of porous nanoparticles, adsorption of toxic ions using porous powders, and luminescence behavior of zirconium oxide. Statement of the problem and objectives of this research work, which is based on the literature survey, is presented towards end of Chapter 2. In Chapter 3, detail experimental work are described. Motivation for selecting borohydride synthesis route to prepare zirconium oxide nanopowders; Gelation-precipitation reaction mechanism of borohydride synthesis method to prepare zirconium oxide nanopowders; Enhanced activation energy of crystallization, development of porous zirconia and application of porous zirconia for removal of Pb (II) toxic ions; Temperature-mediated phase transformation, pore geometry and pore hysteresis transformation of borohydride derived in-born porous zirconium hydroxide nanopowders and Phase formation, pore morphology and photoluminescence behaviour of doped (Eu3+ or Tb3+) and co-doped (Eu3+-Tb3+) porous zirconia nanopowders are described in Chapter 4 to 8, respectively. Finally, the major outcomes from this Ph.D. research work are presented in the conclusions section.

Chapter 2

Literature review

Based on the theme of this research work, as discussed in introduction section of Chapter 1, a detail literature review on the stabilization of tetragonal zirconium oxide, synthesis of metal-oxide nanopowders using sodium borohydride, gas-bubbles derived porous metal-oxide nanopowders, pore morphology of porous nanoparticles, adsorption of toxic ions using porous powders, and luminescence behaviour of porous zirconium oxide are presented in this literature review section of Chapter 2. 2.1: Stabilization of tetragonal zirconium oxide

Pure and stable tetragonal zirconia (t-ZrO2) with smaller crystallite size, which having an average grain size < 100 nm, has many potential applications in different areas of fields due to their unusual physical, chemical and mechanical properties [1-4]. According to

Garvie, pure t-ZrO2 is stable up to a certain critical size of 30 nm. Without adding any stabilizers or surfactants, the particles are generally agglomerated during synthesis due to its higher surface energy of smaller particles, which led to grow at faster rate during [7, 47-49] calcination . So, it is a great challenge to retain t-ZrO2 within this critical size up to moderate or high temperature [7, 47-49]. Many researchers have adopted different synthesis techniques such as chemical precipitation, hydrothermal, gas-condensation, sonochemical, and sol–gel processes and also by varying pH or adding surfactants using different precipitating agents to avoid the agglomeration process so that pure t-ZrO2 can be obtained at different calcination temperatures. Additionally, the tetragonal zirconia are transformed to the monoclinic zirconia, if the size exceeds 30 nm [47-51], which further supress the applicability of the zirconia material for potential applications. Literatures based on the synthesis and stabilization of ZrO2 with small crystallite size without adding any stabilizer as well as controlling the various processing parameters are discussed as follows.

It was found that the starting precursor material effects the phase transformation behaviour of zirconia powders. In this context, Osendi et al. [31] have used two different precursor materials such as zirconium oxychloride and zirconyl acetate for preparing two different kinds of zirconia powder through precipitation and thermal decomposition method. It was found that the phase transformation from t-zirconia to m-zirconia took place at 700 °C and complete m-zirconia was observed at 1100 °C for the case of the sample prepared using precipitation method. However, the phase transformation from t-zirconia to m- zirconia occurred suddenly at ~ 950 °C for the sample prepared though thermal decomposition method. In both cases, the crystallite size for t-zirconia was less than 30 nm. However, when the average size was above 30 nm, the monoclinic phase of zirconia was observed. Additionally, two different salts of zirconia were also used to prepare zirconia nanopowders by Tahmasebpour et al. [68] through polyacrylamide gel method. While using oxynitrate salt, the as-synthesized zirconia was found to be amorphous in nature up to 300 °C and for a semi-crystalline structure of t-zirconia at 400 °C. When the calcination temperature increased to 600 °C, both m- and t-zirconia were detected. Finally monoclinic phase was obtained at 800 °C. On the other hand, presence of t-zirconia along with m- zirconia was observed even at 400 °C when oxychloride salt was used as precursor. The monoclinic phase was prominent at 600 °C along with t-zirconia and finally pure monoclinic phase was observed at 800 °C. It was also observed that the particle size of sample prepared using oxynitrate salt showed bigger than the sample prepared using oxychloride salt. Moreover, processing routes also effect the stabilization of t-zirconia was observed by Rashad et al. [69]. In this case, conventional precipitation (CP), citrate gel combustion (CGC) and microemulsion refined precipitation (MRP) was explored to prepare zirconia nanopowders. When the powders were prepared through precipitated precursor and the citrate precursor, t-zirconia was stable up to 700 °C and convert to m-zirconia at 1000-1200 °C. However, MRP route showed tetragonal zirconia phase at 500–700 °C and cubic zirconia phase at 1000–1200 °C. During calcination process, the soaking time play a major role for the change in crystallite size of zirconia powders and thus effect the phase transformation behaviour. To validate, Srinivasan et. al. [70] have studied the crystallite size effect on the low-temperature transformation of t-phase ZrO2. In this case, zirconia nanopowders were prepared using aqueous solution of ZrCl4 and NH4OH at a pH of ~ 2.95. After synthesis, the as-synthesized powders were calcined at 500 °C for a soaking period from 15 h to 200 h. It was found that the major fraction of tetragonal zirconia phase was formed at 15 h of calcination, however, a large fraction was transformed to the m-phase after 200 h of heating at 500 °C. Similarly, the phase formation of crystalline t-zirconia was observed at 530 °C for the powder sample,

prepared by Xia et al. [48] using low-temperature vapour-phase hydrolysis of inorganic

ZrCl4. In this case, the powders was found to be smaller in size of around 5.8 nm with a narrow size distribution. The zirconia grain grow accompanied by an increase of monoclinic phase and a decrease of tetragonal phase, when the sample was heated at a higher temperature. The pH of the reaction mixture is also a factor for influence the phase transformation behaviour of zirconia materials. So, Berry et al. [71] have developed zirconia powders from zirconium (IV) acetate solutions via two processes such as boiling under reflux and hydrothermal method. The pH was controlled by adding hydrochloric acid (HCl) (pH- 1) and NH4OH (pH- 10). When the sample was prepared using at higher pH through NH4OH, mixture of t- and m- phase of ZrO2 was observed at 500 °C. When the calcination temperature increases for m 500 °C to 900 °C, the volume percentage of t-phase of zirconia decreases and simultaneously m-phase of zirconia increases. It was further found that the volume percentage of t-zirconia and m-zirconia was 43 and 57, respectively. However, only m-zirconia was observed for the sample prepared using HCl and calcined at 500 °C. It was observed that direct formation of monoclinic zirconia was formed under acidic conditions, while using high pressure associated with the hydrothermal treatment. The particle size of the initial powder also effect the stabilization of t-zirconia which was explained by Shukla et al. [72]. In this paper, sol-gel technique was adopted to prepare two different sized particles such as nanosized (∼20-25 nm) and submicron sized (∼500- 600 nm), monodispersed, spherical zirconia particles without any doping of trivalent impurities. When the zirconia nanocrystallites of size ~20-25 nm were calcined at 400 °C for 2h, it shows tetragonal and monoclinic crystal structure and the monoclinic phase increases with increase in calcination temperature. However, the submicron sized (500-600 nm) spherical zirconia particles showed pure t-zirconia phase up to 600 °C and finally converted to m-zirconia at 800 °C. On the other hand, smaller crystallite size with narrow particle size distribution also stabilizes t-zirconia up to moderate temperature. In this context, Bhagwat and Ramaswamy [73] have used citrate route to develop nano-sized zirconia powder. It was found that a high temperature stability of t-zirconia along with minute amount of m-zriconia was up to 750 °C and after which tetragonal to monoclinic phase transformation was observed till 1200 °C. In this transformation, the particle size increased from 8 nm to 55 nm. The formation of t-zirconia is also possible at lower calcination temperature using novel synthesis method, which was implemented by Mondal et al. [47]. In their study, a monolithic t-ZrO2 nanopowders with an amorphous ZrO(OH)2.xH2O polymer precursor was prepared using NH4OH. The t-zirconia nanopowders were formed on heating the as-

synthesized sample at temperatures as low as 200°C. The particle size grow when the sample was heated at higher temperature and form a mixture of 90 vol % t-zirconia long with 10 vol % m-zirconia at 600 °C. However, pure m-zirconia with a particle size of 22 nm was observed at 800 °C. Similarly, zirconia nanopowders having tetragonal structure was observed at low temperature of 450 °C and it was converted to m-phase at 550 °C, when the sample was prepared in a supercritical carbon dioxide reverse micro-emulsion method [74]. Moreover, a lower temperature stabilization of t-zirconia was also observed at 300 °C and pure monoclinic phase was formed at around 1200 °C for the sample prepared via sonochemical method by Liang et al. [75]. In addition, many researchers have adopted different synthesis techniques with the help of different precipitating reagents and able to stabilize the t-zirconia up to moderate or even higher temperature with the use of different surfactants. Use of surfactants also help to stabilize the pure phase of t-zirconia up to moderate or higher temperature due to formation of agglomeration free powders in the as-synthesized condition as well as slow growth of the particle during calcination process [77, 78]. Other than hydroxide based precipitating agent, a known reducing agent such as hydrazine hydrate (N2H5OH) was also used for preparing zirconia nanopowders, which was studied by Zhu et al. [76]. The presence of major phase of t-zirconia along with minor phase of m-zirconia was observed, when it was prepared via hydrothermal route at 150 °C. The size of the t-zirconia was found to be controlled due to formation of hybrid complex between

N2H4 and zirconium salt. Based on the literatures, it was observed that the reducing agent also develops smaller size particles of zirconium oxide. So, literature based on the various materials synthesized using sodium borohydride is discussed as follows. 2.2: Materials synthesized using sodium borohydride

Aqueous sodium borohydride (NaBH4) is well-known for its reducing property. This reagent is commonly used as a reducing agent specially for preparing metals and or metal- borides [52-53]. This is also used for preparing alloys, magnetic, magnetic-nonmagnetic [53, 54, 55] composite materials and also used for H2 storage . Many researchers have developed different types of materials using sodium borohydride and these are discussed as follows. The reducing nature of sodium borohydride led to develop alloy or metal or metal- boride nano materials by many researchers for different applications. In this context, Ag- Fe-Ni ternary metal nanopowders were prepared by Sterling et al. [79] using sodium borohydride reduction of metal nitrates. When the powders were sintered at 900 °C under hydrogen (H2) atmosphere, fine-grained of Ag-Invar alloy was formed. Similarly, Wang et

al. [80] have prepared three carbon-supported Pt-Ni alloy electro-catalysts with varying Pt:

Ni atom ratios (Pt3-Ni/C, Pt2-Ni/C, Pt-Ni/C) using NaBH4 in glycerol at room temperature for use in cathode catalyst in air cathode microbial fuel cell (MFCs). In addition, uniform spherical shape and good monodispersed cobalt (Co) nanoparticles were prepared using reduction method using sodium borohydride [81]. Also Glavee et al. [82] reported the formation of ultrafine cobalt particles using NaBH4. In another research work, Glavee et al. [83] reported the formation of both Ni and NiO when the reaction was carried out in aqueous [53] medium using NaBH4. Moreover, Nayak et al. have prepared Ni nanoparticles with different shapes (spherical, ellipsoidal, cylindrical, hexagonal, and polyhedral) and having a core–shell structure of Ni–nickel oxide by a combination of chemical and gaseous reduction using sodium borohydride. In addition to metal and metal oxides nanoparticles, metal boride materials had also studied using sodium borohydride. In this context, Legrand [84] et al. attempted to synthesize Ni nanoparticles by chemical reduction using NaBH4 in air and they found a mixture of Ni and Ni-B in the reaction product. In addition, nickel boride [85] also have been prepared by Caputo et al. from the decomposition of NaBH4 promoted by the addition of nickel bromide at different concentrations in a dispersing organic medium, tetrahydrofuran and pentane. Moreover, Wu et al. [86] have synthesized cobalt boride materials from the chemical reaction of aqueous sodium borohydride with cobalt chloride. However, the mechanism of sodium borohydride-cobaltous chloride reduction was discussed by Heinzman and Ganem [87]. A new understanding of the chemistry leading to nanoscale particles of metals, borides, and metal borates using borohydride reductions of metal ions was discussed by Glavee et al. [88]. Furthermore, Co catalysts with different supports were prepared for hydrogen generation from catalytic hydrolysis of alkaline sodium borohydride solution. In this case, -Al2O3 supported Co catalyst has quick [89] response and good durability to the hydrolysis of alkaline NaBH4 solution . It is feasible to use this catalyst in hydrogen generators with stabilized NaBH4 solutions to provide on- site hydrogen for proton exchange membrane fuel cell (PEMFC) systems. In addition, Co– B catalyst was acted as the best catalytic activity and used for hydrogen generation for proton exchange membrane fuel cell (PEMFC) application [86]. In spite of preparing metal, alloy or metal oxide nanoparticles, it is also possible to develop oxide materials with smaller crystallites in a controlled way during wet-chemical synthesis through an efficient way using a suitable precipitating agent like sodium [56] borohydride (NaBH4). First, Zhu and Manthiram have used an ambient temperature reduction of aqueous solutions for the synthesis of tungsten oxide bronzes. Addition of

aqueous sodium borohydride solution into aqueous sodium tunstnate solution at pH = 6.5 result in the formation of a reduced sodium tungsten oxide gel. The gel crystalizes sharply at ~ 440 °C to yield the crystalline sodium tungsten oxide bronze (NaxWO3). Similarly, Manthiram et al. [57] showed that borohydrides can also be used effectively to obtain a range of transition metal oxides such as vanadates, molybdates and tungstates. In the view for preparing oxide based materials using NaBH4, our group developed various types of oxide materials using the sodium borohydride as a precipitating reagent. Different oxide based materials such as encapsulation of Ni–O/Zr–O on an Ni core, size-controlled nanocrystallites of Ni varying in size from 2 to 26 nm distributed in a nonmagnetic matrix of Ni(OH)2–ZrO2 and flake-like particles of tetragonal zirconia (t-ZrO2) have been prepared using sodium borohydride [46, 54, 90]. Other than these materials, our group also developed 3+ phosphor nanomaterials such as YBO3:Eu using sodium borohydride. For the first time, we have reported the use of sodium borohydride (NaBH4) as a boron source as well as a 3+ precipitating agent for successfully preparing high temperature stable YBO3:Eu nanophosphor [91, 92]. The borohydride strategy emphasizes the feasibility of synthesizing 3+ thermally stable Eu -doped YBO3 red phosphor materials, which may be suitable for lighting applications. It was observed that gas-bubbles evolves during borohydride synthesis and it forms porous nanostructured materials. So, a detail literatures based on the gas-bubbles derived porous nanopowders is discussed as follows. 2.3: Gas-bubbles derived porous nanopowders Nanoscale sized pores in nanopowders play a vital role to enhance the physical and chemical reactivity of nanoparticles and promising potential applications in adsorption, catalysis, gas purification, sensors as well as in biological application [93-95]. Several synthesis approaches such as sol-gel [96, 97], precipitation [98], tape casting [99], oil emulsion [100, 101] and solvothermal [102] methods have been employed to develop porous metal oxide including zirconium oxide by using hard or soft templates as well as self-assembly process with the help of specific surfactants to meet the requirement for different applications [103, 104]. The approach using either templates or self-assembly process using surfactants (such as sodium dodecyl sulfate, ammonium lauryl sulfate, sodium lauryl sulfate) for developing porous materials needs complicated pre-preparation and or post-treating processes. For template-based method, building a porous framework, inset of materials into the template and removing the pore-forming framework is a big challenge [105, 106]. In addition, self- assembly process is a comparatively simple but removal of surfactants at the end of the

fabrication is also a big problem. In these above syntheses, development of porous in the as-synthesized condition as well as sustaining the porous nature up to moderate temperature is also difficult [103]. In addition, zirconium oxide in the form of tetragonal (t) at room temperature was considered to be an important ceramic material and the above synthesis methods have been employed for the development of porous structure. However, poor structural stability (i.e. phase transformation of zirconium oxide from tetragonal to monoclinic (m), when the sample was cooled down from high or moderate temperature to ambient condition) and difficult in removal of template led to affect the potential application of porous zirconium oxide [105, 107]. Further, in order to develop a porous structure in nanoparticles, it is necessary to form loose or agglomeration free powders in the as- synthesized condition. But it is too difficult to prepare agglomeration free powders through the wet-chemical route without adding any surfactants or templates, because the highly energetic nuclei particles in solution coagulate with each other to form large agglomerate particles by decreasing their surface energy. So, in this context, a new progress has been made in tailoring the synthesis of agglomerated free or loose particles in the as-synthesized condition through gas-bubbles template mechanism. In this method, the gas-bubbles create numerous nucleation sites throughout the solution during synthesis and help to control the agglomeration of nuclei and thus control the size of the particle in the as-synthesized condition. The agglomeration free smaller size particle may form stable tetragonal zirconium oxide up to moderate temperature [98, 108]. In the same time, the as-synthesized loose particles may develop moderate temperature stable porous particle during calcination process due to coarsening of particles as well as coalescence of voids [109-111]. In gas-bubble template, the gas-bubbles can be generated either by simple blowing a gas into precursor solution or by evolving in-situ gas bubbles through chemical reactions [112, 113]. Gas-bubbles induced porous structure has been emerged as more advantageous than other pore forming methods because of simple room temperature synthesis, clean, template-free, environment friendly and short reaction time [109,114, 115]. Synthesis and assembly of porous materials as well as designing nanoscale pores into various types of materials has attracted intense interest for enhanced performance in the various fields of application including absorption, efficient catalysis, sensors and drug delivery. By carefully designing and controlling of different parameters, it is possible to manipulate the optical and electronic properties of these materials for specific technological applications. Many researchers have developed porous structures of different materials using different type of gas-bubbles template method and these are discussed as follows.

The gas microbubbles of NH3, N2 and H2 generated during the reaction of cobalt salt with hydrazine monohydrate provide the aggregation centre and thus led to develop a hollow

cobalt (Co) mesosphere chainlike structure [116]. This is a simple method in general and can be applied to the synthesis of chainlike structures of hollow nanospheres of many other metals.

Using a simple solution-phase approach in ammonia solution, the evolved N2 gas- [113] bubbles were used to develop macro and mesoporous nitridated titania (TMMN-TiO2) .

The interconnected mesoporous and ordered tubular macroporous channels of TMMN-TiO2 have promising applications for high-performance anode materials. Similarly, N2 gas bubbles acted as aggregation centers during the formation of ZnSe hollow spheres [117]. The hollow nature together with the tunable interior nanocrystal size of these microspheres make them for novel optical, electronic and magnetic applications. Ammonia gas-bubbles can also be produced from high concentration of ammonium acetate and these gas-bubbles were used to fabricate mesoporous Fe3O4 nanoparticles. A monodisperse and mesoporous magnetite nanospheres with particle size of ~100 nm and pore size of 7.6 nm were synthesized through a solvothermal process and it was performed by Dong et al. [118]. Also, by decreasing the amount of ammonium acetate, the simple hollow spheres and tiny aggregated particles could be also prepared facilely.

Sulphur dioxide (SO2) gas bubbles was evolved from a thiosulfate complex, formed by reacting lead nitrate with sodium thiosulfate (Na2S2O32H2O). These gas-bubbles was used to develop PbS hollow sphere quantum dots with strong luminescence properties for early cancer diagnosis [110].

Hydrogen sulphide (H2S) gas bubbles was derived from a thioacetamide solution and when this solution was reacted with zinc nitrate [Zn(NO3)6H2O], it develop well- defined hollow and solid ZnS nanospheres [119]. A porous zirconia having ~ 60nm in particle size and filled with smaller spherical mesopores (size ~5 nm) was formed by directly decomposing a Zr(NO3)45H2O ethanol sol– [120] gel solution . In the decomposition products, some micron-sized ZrO2 blocks that full filled with spherical air pores and this suggests a new optional approach for preparing high- porosity metal oxides. [121] Thiourea can acts as the bubble template to develop nanoporous g-C3N4 (npg-C3N4) . In this case, the gas bubbles produced by thiourea were uncontrollable and untunable, which resulted in formation of npg-C3N4 with different kinds of nanoporous morphology. This material can be used for adsorption and photodegradation application.

In addition to, the hydrogen (H2) gas-bubbles can also be utilized to develop different types of porous materials. The H2 gas bubble dynamic template has many advantages, such as low cost, facile one-step process of formation and elimination of the template, and facile control of 3D porous materials [114, 122]. Many researchers have developed porous structures

of different materials using H2 gas-bubbles template method. A novel one-step electrochemical method was adopted in which it involves a repeated gold oxidation−reduction and intensive hydrogen gas-bubbles were evolved. These gas-bubbles help to form three-dimensional (3D) micro−nano hierarchical porous gold films [123]. Porous films of Pb or Sn can also be prepared by the above method [124]. Particularly, this method is green, convenient, and economical, which enables to fabricate the 3D porous structure from the metal itself. Also, a novel kind of three dimensional (3D) porous micro/nanostructured interconnected (PMNI) metal/metal oxide electrode was successfully [125] fabricated by Chen et al. via a facile H2 gas bubble dynamic template route. Similarly, three-dimensional porous nickel (3D-PN) film with large specific surface area (퐴푠) and high porosity by hydrogen bubble dynamic template (HBDT) method was developed [126]. A hierarchically structured mesoporous nanowall arrays of MnO20.5H2O that were grown on a cathodic substrate by means of water-electrolysis-induced precipitation due to release of hydrogen gas bubbles from the cathode [127]. 2.4: Pore morphology of porous nanoparticles Depending on the pore size of a porous material, it is generally categorized into three categories such as microporous, mesoporous and macroporous based on the International Union of Pure and Applied Chemistry (IUPAC) [128]. The pore diameter of a microporous material is less than 2 nm and the pores are generally narrow-size distribution. The microporous materials are generally used in catalysis, adsorption and separation, ion- exchange, gas purification systems, and gas-storage. The pore diameter of a macroporous material is generally greater than 50 nm. The macroporous materials have broad size distribution and potential applications in adsorption, separation and heterogeneous catalysis. However, the pore size of a mesoporous materials lies in between 2 to 50 nm. The mesoporous materials have potential applications and lots of research are going on for the development of these porous materials for different applications. The pore size and its distribution may be determined using BET-adsorption and desorption isotherm plot. The adsorption isotherms can be used to analyze the pore size distribution, surface area and pore volume of the porous materials. According to IUPAC classification, the isotherm can be classified in six different types, as shown in Fig. 2 (a) [128]. Type I is typically for microporous (pore size <2 nm) materials [129]. It is a single-molecule adsorption type, and it is characterized by continuous increases in the adsorption volume till the relative pressure surpasses a certain value [130]. The type II isotherm is characteristic of the non-porous or macroporous (pore size >50 nm) materials [129]. This type exhibits multi-layered adsorption, capillary filling, and capillary condensation. Under these circumstances, no adsorption saturation point exists on the curve [130]. The type III of adsorption isotherm exhibits a

gradual increase of the adsorption volume with the relative pressure [130]. Type III is for non- porous materials. Type IV and V exhibit a hysteresis loop, i.e. the adsorption and desorption isotherms do not coincide over a certain region of external pressures. However, the type IV isotherm is typically for mesoporous materials [129]. Type V hysteresis loop is a typical sign of a weak fluid-wall interaction. It is less common, but observed with certain porous adsorbents. Type VI is exhibited by non-porous solids.

2.1: Types of sorption isotherms and hysteresis loops [128, 130., 131].

Pore shape can have a regular or an irregular, either cylindrical or ink-bottle or funnel or conical or slit like shapes. Also, pores can be closed (not accessible from the outside), blind (open only at one end), or through (open at both ends). In addition, pores can be isolated or connected to other pores to form a porous network. Pore shape affects the mechanisms of condensation and evaporation during adsorption and desorption process of BET isotherm and four types of hysteresis have been recognized. Based on the adsorption- desorption of BET, the hysteresis loops can be divided into four categories based on the IUPAC guideline, as shown in Figure 2 (b) [130]. H1 and H4 represent two extreme cases, and H2 and H3 are the intermediate situations. The adsorption and desorption branches of H1 are almost vertical and nearly parallel over an appreciable range of gas uptake. The adsorption and desorption branches of H4 are almost horizontal and nearly parallel over a wide range of relative pressure [128-130]. Type H1 hysteresis is characteristic of solids crossed by channels with uniform sizes and shapes (typically cylindrical). Materials of this type tend to have relatively narrow distributions of pore size. Type H2 corresponds to channels with a pore mouth smaller than the pore body (typically ink-bottle-shaped pores). In this case, generally the distributions of pore size radii are wide. Type H3 hysteresis is usually found on solids with a very wide distribution of pore size (slit or panel type pores) and type H4 corresponds to limited amounts of mesopores limited by micropores [128-130]. In this thesis

work, the BET adsorption and desorption curve was analysed based on the consideration of the type of hysteresis. Based on the above discussion, many researchers have correlated the pore morphology of the porous zirconia based material, prepared by different synthesis methods with the hysteresis loop and analyze the porous zirconia based material for suitable applications. The pore size distribution of mesoporous zirconium oxides can be tuned by changing the preparation composition [132]. One of the sol–gel-hydrothermal method, an irregular, closed pores were embedded in the single zirconia crystals. The hysteresis loop of the calcined (550 °C for 5 h) porous zirconia was an H1 type, which is characteristic of solids consisting of particles crossed by nearly cylindrical channels or made by aggregates or agglomerates of spherical particles [133]. A network of more or less organized pseudo- hexagonal pores was observed in the dried (140 °C) material of mesoporous zirconia, prepared by hydrolysing the zirconium propoxide in the presence of anionic surfactants (alkyl-phosphate and -sulfate) [134]. In another case, hexagonally ordered mesostructured surfactant composites based on zirconia have been synthesized by Ciesla et al. [135] using cationic surfactants and zirconium sulfate as the inorganic precursor species. A Super- microporous zirconium phosphonate hybrids showed typical type I isotherms and a tuned pore sizes from small micropore (0.87 nm) to mesopore (2.50 nm) range, depending on the hydrothermal time [136]. A series of well-ordered mesoporous composite frameworks containing zirconium oxide and 12-phosphomolybdic acid compounds has been prepared by Armatas et al. [137] by using the evaporation-induced cooperative assembly method. Additionally, a highly ordered mesoporous structure was formed from the combination of inexpensive and commercially available polymers with metal inorganic salts solubilized in ethanol solution [138]. Similarly, high-quality zirconium oxide material with ordered pore structure was fabricated by Chen et al. [139] using a surfactant-assisted route and a post- synthetic treatment with phosphoric acid. Also, monodisperse porous zirconia (ZrO2) microspheres with nanocrystallized framework, prepared by impregnation of porous polymer microspheres as a novel hard template with zirconia precursors, have also analysed [140] using BET isotherm . In another case, the monolithic macroporous zirconia (ZrO2), [141] derived by Guo et al. showed tetragonal ZrO2 and monoclinic ZrO2 at 400 and 600 °C, respectively, without spoiling the macroporous morphology. In this case, the as-dried and 600 °C heat-treated gels exhibit isotherms of type-IV, while the gels heat-treated at 700 and 800 °C show isotherms of type-I. Additionally, a hysteresis loop in the relative pressure range of 0.4–1.0 for the cage-like ZrO2 and hollow ZrO2 microspheres with high surface area (~384 m2/g) sample indicating the presence of the inhomogeneous mesopores [142].

Mesoporous and tetragonal zirconia (t-ZrO2) from a metal–organic framework

(Universitetet i Oslo: UiO-66) was formed at 500 °C and the isotherm is of type IV based on IUPAC classification, characteristic of the formation of mesoporous materials. The sample exhibits a hysteresis loop of type H2, which is typically featured as “ink-bottle” pores [98]. Zirconia based porous materials have potential applications such as catalysis, adsorption and separation, ion-exchange, gas purification systems, gas-storage, luminescence and biomedical. However, a very few literatures are available for use of only porous zirconia in catalysis application. Additionally, mesoporous sulfated zirconia, mesoporous silica-zirconia, gold supported mesoporous zirconia, clay supported mesoporous zirconia and nickel supported mesoporous zirconia have been used for catalysis applications [93-95]. As this research work is not concentrated on the catalysis applications, so, the literatures based on porous zirconia and its applications in catalysis are not elaborated in this literature review. Rather, this thesis work focuses only on removal of Pb(II) ions and luminescence applications on porous zirconia, so literatures based on adsorption of toxic ions using porous powders and luminescence behaviour of zirconium oxide are discussed in detail as follows. 2.5: Adsorption of toxic ions using porous powders Release of heavy toxic metal ions such as As(III/V), Cu(II), Cd(II), Pb(II),Cr(III/VI), Ni(II), Hg(II) and Zn(II) into the water system create a widespread environmental problem and thus the water containing the heavy metal ions must be treated before use. Various treatment technologies have been employed in the removal of heavy metals from aqueous systems such as electrochemical treatment, chemical precipitation, reverse osmosis, ion exchange, and membrane separation. However, the adsorption method is considered quite attractive in terms of high efficiency, ease of handling, and availability of different adsorbents for removal of toxic ions from aqueous systems. In this context, various adsorbent materials have been studied for their ability to remove toxic ions from water samples. In addition, determination of the rate at which toxic ions removal takes place in the used solid/solution system is one of the crucial factors for the effective design of the sorption system. In order to examine the controlling mechanism of sorption processes, five kinetic models were used to test the experimental data, i.e. the pseudo-first, the pseudo- second order, Elovich equation, intraparticle diffusion and Bangham model [143]. Kinetic models have been used to investigate the mechanism of sorption and potential rate controlling steps, which is helpful for selecting optimum operating conditions for the full- scale batch process. The experimental kinetic data were treated with the models given in Table 2.1. Pseudo-first, the pseudo-second order indicates the adsorption process is surface reaction controlled, with chemisorption type. Elovich model also suggests that the

chemisorption is the main adsorption controlling mechanism. Since the pseudo-first, the pseudo-second order and Elovich kinetic models cannot identify the influence of diffusion on sorption; Weber and Morris’ equation and Bangham’s model were used. Table 2.1: The equations of the different kinetic models Kinetic model Equation [Ref]

퐾 푡 pseudo-first order log⁡(푞 − 푞 ) = 푙표푔푞 − 1 [144] 푒 푡 푒 2.303

pseudo-second-order 푡 1 푡 [145] = 2 + 푞푡 퐾2푞푒 푞푒

[146] Elovich 푙푛 푎푒 푏푒 1 푞푡 = + ln 푡 푏푒 푏푒

1/2 Intra-particle diffusion 푞푡 = 퐾푖푡 + 퐶 [147,

148]

Bangham’s model 퐶 퐾 퐶 [149] log 푙표푔 푖 = log 푏 푠 +  log 푡 퐶푖 − 퐶푠푞푡 2.303푉

Nomenclature: -1 qe and qt is the amount of chromium (mg g ) adsorbed at equilibrium and at time t, respectively. -1 K1 (min ) is the rate constant for pseudo-first order adsorption reaction. -1 -1 K2 is the rate constant (g mg min ) for pseudo-second order reaction. ae is the initial adsorption rate (mg/g min). be is related to the extent of surface coverage and activation energy for chemisorption (g/mg). -1 1/2 Ki is the intra-particle diffusion rate constant (mg g min ). C is the intercept. Ci is the initial concentration (g/L) Cs is the weight of adsorbent used per liter of solution (g/L).  and kb are constants V is the volume of solute (mL)

The adsorption process were performed by a suitable adsorbent and the adsorbent with a porous structure is found to be more efficient for removal of toxic metal ions through adsorption method [150]. In addition, a large specific surface area, pore volume and the unsaturated atoms in the surface of porous materials endow the materials with good adsorption capacity. In some cases, metal oxides at the nanosize level especially aluminum, iron, titanium and cupric based oxides have demonstrated superior performance for arsenic adsorption, because of their large surface areas and preferred surface properties [151]. Some

researchers have used porous transitional metal oxides such as zirconium oxide in water treatment. The zirconia has better ion transfer performance and the concentration of oxygen vacancy because it has both acid sites and basic sites on surface. Zirconia can acts as not only the carrier but the catalyst for some reactions. Introducing a porous structure into the zirconium metal oxides is a promising route to upgrade the properties. In addition, zirconium-based oxides are stable, nontoxic, and not dissolvable in water, and they could be an attractive choice for drinking water purification. The relevant literatures based on adsorption properties of zirconia based nanopowders are described as follows. Zirconia nanopowders can adsorb affectively different mono-dye stuffs (red, yellow, and brown). It was suggested that when the solution pH was between 3.0 and 4.5, the adsorption rate of yellow dye on ZrO2 was more than 60% after 3 h, and the adsorption rate [152] of red and brown dye reached 90% after 2 h . A mesoporous spherical zirconia (ZrO2) with a surface area of 113 m2.g-1 and average pore size of 5.0 nm effectively adsorbs Cs ion [153]. It was also found that with suitably charged functional groups in the hosts or a zirconia based composite materials may release the toxic metals from water. In this context, a novel hybrid nanomaterial by encapsulating ZrO2 nanoparticles into spherical microporous polystyrene (MPS) beads covalently bound with charged sulfonate groups (−SO3 −) exhibited more preferential sorption toward Pb(II) than the simple equivalent mixture of [154] MPS and ZrO2 . In another case, a monodisperse mesoporous zirconium titanium oxide microspheres with high surface areas (up to 413 m2/g) and uniform worm-hole like mesopores (∼ 5.5 nm) and monodisperse grain size (~780 nm) removes Cr (VI) anions from solution with very high adsorption capacities [155]. In addition, polyacrylamide zirconium (IV) iodate, a hybrid inorganic−organic material was able to adsorb Pb(II) ions effectively and follows pseudo-second-order kinetic model. It was also suggests that intraparticle diffusion is not the sole rate-controlling step; some degree of boundary layer diffusion is also involved [156]. A new kind of inorganic−organic hybrid zirconium phosphonate (NTAZP) material with worm-like mesoporous (about 2.7 nm) structure efficiently remove the heavy metal ions (e.g., Pb2+,Cu2+, and Cd2+) [157]. A hierarchically porous ZrP monoliths with size-tunable co-continuous macropores (from 0.5 to 5 mm) with a high reactive surface area (600 m2.g-1) and relatively high mechanical strength (Young’s modulus 320 MPa) was applied to ion adsorption. A simple syringe device inserted tightly with the ZrP monolith as a continuous flow setup was demonstrated to remove various toxic metal ions in aqueous solutions, which shows promising results for water purification [158]. In addition, zirconium

oxide (ZrO2)-carbon nanofibers (ZCN) were used effectively to adsorb antimonite (Sb (III)) and antimonate (Sb (V)) [159]. Zirconium hydroxide and or amorphous zirconia was also used effectively to remove various heavy metal ions by many researchers and found to be a good adsorbent materials for water purification. Also, zirconium (IV) hydroxide was used as ion-exchange material in fine purification of potable water to remove silicon, arsenic, and phosphorus compounds and other impurities. Also it was used in radiochemistry for separation and purification of [160] radionuclides . A type of hydrous zirconium oxide such as ZrO2·nH2O adsorbs Cr (VI) within 60 minutes and showed pseudo-second-order model [161]. Due to its high adsorption capacity, this type of hydrous zirconium oxide has the potential for application to control Cr (VI) pollution. In addition, a porous resin loaded with monoclinic or cubic hydrous zirconium oxide adsorbs efficiently As (III) and As (V) and showed a strong adsorption for As (V) at slightly acidic to neutral pH region, while As (III) was favourably adsorbed at pH around 9 to 10 [162]. In another case, mesoporous hydrous zirconium oxide exhibits higher affinity towards arsenic (V) and showed higher kinetics of ion-exchange and improved thermal stability that allows retention of adsorption properties even after treatment at 500– 550 °C [163]. Moreover, amorphous zirconium oxide nanoparticles synthesized by a simple and low-cost hydrothermal process, and their phosphate removal performance was explored in aqueous environment under various conditions [164]. The adsorption mechanism of phosphate onto amorphous -ZrO2 nanoparticles was based on the surface -OH groups, which played a major role in the phosphate removal. Another application of porous zirconia is in the luminescence field, when it was incorporated with rare earth ions. So, literatures based on luminescence behaviour of zirconium oxide is presented as follows. 2.6: Luminescence behaviour of zirconium oxide Inorganic oxide materials such as zirconium oxide has been extensively modified by the addition of suitable stabilizers, allowing the development of this ceramic material for diverse practical applications starting from structural to optical. The addition of stabilizers such as Y2O3, CaO and MgO and rare-earth (RE) materials to the ZrO2 is known to enhance the structural stability in both elements cubic and tetragonal phases, depending on the dopant amount and processing temperature. Besides to the role of the rare earth ions as stabilizers, their optical activation constitutes an opportunity to explore the zirconia-based material in the luminescence applications [165]. Generally, due to its wide band gap (5.0-5.5 eV), optical transparency, high refractive index, and photochemical stability, ZrO2 nanoparticles have been exploited in luminescence applications. The luminescent materials

have been utilized widely in applications, such as cathode ray tubes, fluorescence lamps, vacuum fluorescent display devices, colour plasma display panels, and electroluminescent flat-panel displays. In addition, luminescence properties of zirconia nanoparticle mainly depends either on the size of the particles or on the presence of defects [166, 167]. The existence of smaller size may affect the photoluminescence (PL) properties of un-doped zirconia. In general, Zr4+ is nonluminous, so PL property is most likely due to small size with defects in the system [168]. Several reports state that zirconia with different nanostructures show PL property due to their morphology with defect structure [169-171]. The photoluminescence properties of zirconia arise due to its smaller size with defect structure that involves the movement and diffusion of atoms and a change of electronic structure [172- 175]. According to electronic band theory, the band gap is dependent on material structure and due to which the band gap is dependence on the zirconia crystal structure [176]. Usually, the experimental band gap energies vary from 4.2 to 5.83 eV for m- ZrO2, from 4.2 to 5.78 [177-179] eV for t-ZrO2, from 4.6 to 6.1 eV for c-ZrO2 and ~ 5.5 eV for amorphous zirconia . In addition, the photoluminescence and the band gap energy are well associated with each other [180]. The photoluminescence band position can be divided into three regions. One region with band peaks below 2 eV ( > 620 nm), the second region is between 2.0 (= 620 nm) and 3.5 eV (= 354 nm), whereas the third above 4.0 eV ( < 310 nm) [181]. Luminescence bands with peak position < 2.0 eV are assumed as due to dopants [mostly rare earth (RE) ions elements], the region 2.0−3.5 eV was described as native defect luminescence and the bands at region > 4.0 eV were interpreted as ZrO2 excitonic [182, 183] luminescence . There have been lot of studies on the luminescent properties of ZrO2, but still there is a lack of clarification for the luminescent mechanisms in ZrO2. In general, luminescence from nanoparticles is thought to be related to small size or defects and/or impurities in the system. In ZrO2 nanoparticles, the involvement of mid gap traps such as surface defects or oxygen vacancies defects make contribution towards luminescence properties [167, 172]. Furthermore, it was suggested that the luminescence properties are dependent on the size, crystallinity and morphology of the nanoparticles which in turn is dependent on the synthesis procedure and reaction conditions [167, 172]. There are different synthesis methods are adopted to synthesize un-doped pure zirconia nanoparticles by several researchers and luminescent properties are well described in many literatures and are described as follows. In one of the microwave assisted chemical method led to develop monoclinic and tetragonal ZrO2 at 400 °C and when it was excited at wavelength of 270 nm, it showed a broad peak in the range of 300 nm to 600 nm along with prominent emission peak at around 414, 475, and 563 nm, with 414 nm. The broad fluorescence band seems to be mostly caused

by the small particle size leading to an inhomogeneous broadening from a distribution of surface or defect states. It was also seen that the fluorescence band position and the band shape stayed nearly the same if the excitation wavelength changes. This indicated that the fluorescence involved the same initial and final states in the excitation wavelengths ranging from 270 to 320 nm [184]. Similarly, a broad PL spectra having three fluorescence emissions at 402 nm, 420 nm, and 459 nm were observed for the calcined (320 °C for 3h) polymer- stabilized tetragonal ZrO2 nanopowders (average size of 2 nm), while using an excitation wavelength of 254 nm [185]. These emissions was due to the involvement of mid-gap trap states, such as surface defects and oxygen vacancies. It seems that the large amounts of surface defects exist on the as-synthesized nano-ZrO2 particles because of their high surface area. A hydrothermal derived zirconia nanostrcutures also show a broad emission band with maximum intensity at around 400 nm, while excited at wavelength of 290 nm and 300 nm depicting the violet emission, which was attributed to the ionized oxygen vacancy in the material [186]. The main source of defects centers was oxygen vacancies/interstitial, Zr vacancies/interstitials, which are expected to exist on the surface of ZrO2 nanostructures due to their high surface area, given to their low dimensional structure. Pechini-type sol-gel process derived t-zirconia powder calcined at 500 °C shows an intense whitish blue emission

(max= 425 nm) with a wide range of excitation (230-400 nm). This emission decreased in intensity after being annealed at 600 °C (t + m-ZrO2) and disappeared at 700 (t + m-ZrO2),

800 (t + m-ZrO2), and 900 °C (m-ZrO2). After, further annealing at 1000 °C (m-ZrO2), a strong blue-green emission appeared again (max= 470 nm). The whitish blue (425 nm) and blue-green (470 nm) emission bands was ascribed to interstitial carbon defects in the [187] tetragonal ZrO2 and oxygen defects in the monoclinic ZrO2, respectively . Different morphology of zirconia powders can also show the photoluminescence behaviour. In this context, a mesoporous ZrO2 with a tetragonal (t) nanocrystalline framework which has narrowly distributed pore size (2–11 nm) with average pore diameter of 5 nm and surface area of 65 m2/g exhibited a photoluminescence band centered at 419 nm under excitation at 285 nm wavelength at room temperature and it was attributed to the ionized oxygen vacancy [170] in t-ZrO2 in the mesoporous structure . Similarly, mesoporous nanocrystalline zirconia prepared by combination of soft-templating and solid–liquid method (CSSL) possesses a wormlike arrangement of mesopores surrounded by tetragonal ZrO2 nanocrystallites at 600 °C. When excited at 240, 250 and 260 nm wavelength, the sample showed emission peaks centered at 392, 394 and 398 nm, respectively in the UV region. It was suggests that the oxygen-vacancy defects in ZrO2 crystal lattices are responsible for the photoluminescence emission of zirconia under UV excitation [188]. The inherent fluorescence was observed for a porous zirconia particles of nearly 380 nm diameter and it was originates from surface

defects possibly coordinatively unsaturated Zr4+ ions [189]. This nanoparticles are suitable for the possibility of drug delivery application. In addition, nanowires of zirconia having an average diameter of around 80 nm and a length of over 10 µm showed the emission peaks are ranging from 330 nm to 450 nm and centered at 388 nm, when the sample was excited at wavelength of 230 nm and 235 nm [190]. Further, The emission spectra obtained are almost independent of the excitation wavelengths used; therefore excitation wavelengths of 230, 235, 240, and 244 nm produce emission peaks at 388, 388, 390, and 384 nm, respectively, in the UV region at room temperature. This result indicates that the PL emission comes from the ZrO2 nanowires, not from other impurities. There are oxygen vacancies in ZrO2 crystals, and the oxygen vacancies can induce the formation of new energy levels in the bandgap.

The strong photoluminescence of the ZrO2 nanowires in the UV region suggests for possible use in luminescent labels, light-emitting molecular substances in nanoscale photoluminescent or nano-optoelectronic devices and as optical memory materials in optical memory systems. Effect of calcination temperature also effect the photoluminescence properties of zirconia nanopowdeers. In one of the zirconia samples, prepared using without and with surfactants i.e. sodium dodecyl sulfate (SDS) and it was seen that the emission spectra of as-prepared nanoparticles (prepared without surfactant) showed a peak at 391 nm, while for nanoparticles annealed at 600 °C the peak is at 423 nm. When the sample was heated at higher temperature, PL intensity decreased due to the phase transition from tetragonal to monoclinic which was accompanied by increase in grain size and decrease in the concentration of oxygen vacancies. However, the PL intensity of the surfactant modified zirconia nanoparticles has been decreased almost by five orders as compared to zirconia nanoparticles prepared without surfactant due to the passivation of surface defects [191]. Monodisperse spherical zirconia particles with a narrow size distribution exhibit broad, intense visible photoluminescence in the range between 350 nm to 650 nm. Also, the emission colours of the ZrO2 samples was tuned from blue to nearly white to dark orange by varying the annealing temperature [192]. The experimental results obtained in the present study reveal that the control of the carbon content by optimizing the annealing temperature is important for the production of ZrO2 phosphors with emission wavelengths tunable over the visible light emission. Further, these phosphors can be potentially used as new environmentally friendly luminescent materials. Recently, much interest was focused on the luminescent properties of rare earth doped zirconia nanoparticles [193]. An exponential interest in the field of optically active trivalent lanthanides (Ln3+) embedded in nanomaterials has been observed, as confirmed by the high number of scientific reports in this subject. Among the mentioned lanthanide ions,

trivalent europium (Eu3+) and terbium (Tb3+) are known to give rise to the intense red and 5 7 green luminescence. The red emission is mainly due to D0,→ FJ(0–4) multiplet transitions 3+ in several insulator and wide band gap semiconductor hosts. A typical Eu : t-ZrO2 consists of four PL band groups, namely 560–600, 600–645, 645–680, and 680–730 nm, in the 5 7 5 7 5 7 5 7 3+ [29] D0,→ F1, D0,→ F2, D0,→ F3, and D0,→ F4 (Eu ) transitions, respectively . The peak intensities of the representative members are in the ratio of 63:100:34:87 [29]. The 5 7 5 7 characteristically weak D0,→ F0 band has no visible value. The D0,→ F1 (592 nm) band, which is a magnetic dipole allowed transition, has a relatively enhanced value. The 5 7 D0,→ F3 (658 nm) band, a forbidden transition by the electric dipole as well as the magnetic dipole, appears to be the weakest band of the spectrum. As a hypersensitive forced electric- 5 7 5 7 dipole transition, the D0,→ F2 (608 nm) band is the predominant group in the D0,→ FJ, J = 04 series. The implication is that Eu3+ occupies sites of presumably distorted symmetries in the Eu3+: zirconia based systems. Similarly, for Tb3+-doped zirconia based green phosphor, the spectra display the characteristic emission of a Tb3+ ion composed of 5 7 5 7 four bands associated with the D4  Fj (j = 3-6) transition, at 488 nm ( D4  F6), 543 nm 5 7 5 7 5 7 [194] ( D4  F5), 584 nm ( D4  F4), and 622 nm ( D4  F3) nm . The luminescence properties Eu3+ or Tb3+ doped zirconia are strongly dependent on the size, crystallinity and morphology. There are several synthesis methods were adopted to synthesize Eu3+ or Tb3+ doped zirconia nanoparticles by several researchers and luminescent properties are studied and some relevant literatures are described as follows. While doping with Eu3+ in zirconia, prepared under nonaqueous conditions using different organic solvents (benzyl alcohol and biphenyl-4-methanol), a tunable emission colour ranging from purplish-pink to greenish-blue was obtained by simply selecting different excitation wavelengths and it was due to the overlapping of the various emission components involved (i.e., the emission of europium(III) transitions, defects in the zirconia and capping ligands) [195]. It was also found that a partial Eu3+  Zr4+ substitution led to stabilize t-zirconia and a wide and intense PL in the visible–near-IR regions (550–730 nm) was observed [29]. Similarly, the structural phase composition and room temperature luminescence properties of terbium doped (at different percentages such as 0, 1, 2, 3, 5, and 10 wt%) zirconium oxide powders obtained by solution combustion synthesis method was studied by López-Romero et al. [196]. Studies of the phase composition revealed that when the Tb3+ increase, the content of the monoclinic phase decreases, while the content of the tetragonal phase increases. Un-doped tetragonal and monoclinic phases exhibited intense green photoluminescence by means of a broad band centered at 480 nm. Presumably that 3+ PL emission was due to structural defects of the host lattice. All ZrO2/Tb samples showed 5 7 four emission bands peaking at 490, 545, 598, and 620 nm corresponding to the D4→ Fj,

(j = 3 – 6), electronic transitions of the Tb3+ ion, when excited by ultraviolet radiation (325 nm). The sample doped with a 2 wt% of the Tb3+ ions showed the maximum green PL emission intensity. The observed green emission was very strong and it was possible to perceive it in normal room light with the naked eye. A transparent Tb doped zirconia (Tb: [194] ZrO2) ceramics that luminesce in the visible was developed by Hardin et al. . The visible luminescence was temperature dependent, yielding samples that have integrated temperatures sensing capabilities. The Tb dopant serves to both stabilize the tetragonal phase of zirconia and for emitting light. The Tb: ZrO2 ceramics have an excellent combination of structural and optical properties, suggesting that the ceramics can be used in a wide range of applications such as temperature sensitive transparent armor, windows, or thermal barriers. In addition to only red and green based phosphors, the use of rare-earth co-doped phosphors represents a fast-growing industry, owing to their applications in white-light- emitting diodes (LEDs), field emission displays (FEDs), and plasma display panels (PDPs). White LEDs have drawn much attention in recent years. In this context, co-doped of rare earth ions such as Eu3+ and Tb3+ in zirconia matrix was studied by researchers and are discussed as follows for potential use in white light application. 3+ Un-doped ZrO2 as well as single-and double-doped ZrO2: M (where M = Tb and Eu3+) nanophosphors were prepared by a simple sonochemical process and a characteristic blue and green emission from Tb3+ ions and red from Eu3+ dopant ions were observed [197]. The Commission Internationale de l’éclairage (CIE) coordinates of the double-doped 3+ 3+ ZrO2:Tb (1.2 %): Eu (0.8 %) nanophosphor lie in the white light region of the chromaticity diagram and show promise as good phosphor materials for new lighting devices. Similarly, the developed ZrO2: 8% rare earth particles by the complex 1 3 polymerization method showed a photoluminescence emission spectra of the type G4 H6 5 7 5 7 (466 nm) from Tm 3þ, D4 F5,6 (495 nm and 550 nm) from Tb 3þ, and D0 F1,2 (597, 619, 656 and 706 nm) from Eu 3þ.The CIE coordinates calculated for particles treated at 600 °C and 800 °C showed values of (x  0.34, y  0.34) and (x  0.31, y  0.34), respectively; according to the CIE diagram, the values represent a points in the white region.

The developed ZrO2: RE materials are promising for applications in new white light- emitting devices. On the contrast, hydrothermal technique as well as polyol route was adopted to prepare cubic zirconia doped with Eu3+ [198, 199]. In this case, the dopant Eu3+ is used to stabilize crystalline phase to cubic and at the same time to get red counterpart of the white light. On monitoring the excitation at 244 nm, a broad peak centered at 486 nm and well known Eu3+ peaks are observed. The peak at 486 nm was attributed to O2- vacancy. 5 7 3+ The peak at 592 nm was attributed to magnetic dipole transition, D0 F1, of Eu . On the

other hand, the peak at 609 nm was attributed to structurally sensitive electric dipole 5 7 3+ transition, D0 F2, of Eu . Relative intensity of the electric (E) and magnetic (M) dipole transitions (i.e. E/M), also called asymmetry ratio, strongly depend on the local symmetry of Eu3+. The asymmetry ratio of the sample is 1.04. This value is lower than that of reported tetragonal phase (1.46) [200]. It suggests that the sample have higher symmetry than tetragonal phase i.e. the sample is cubic. CIE co-ordinate of this nanocrystal (0.32, 0.34) is 3+ closed to that of the ideal white light (0.33, 0.33). It suggests that the ZrO2:Eu nanocrystals synthesized by hydrothermal technique may find applications in simulating the vertical daylight of the Sun. For comparison purpose, PL behaviour of zirconia powders prepared via different synthesis methods are discussed in Table 2.1. Table: 2.2: PL behaviour of zirconia powders prepared via different synthesis methods. Synthesis method Phase of zirconia Photoluminescence behavior [Ref]

Microwave assisted m-ZrO2 and t-ZrO2 Broad peak 300 nm to 600 nm [184] chemical method at 400 °C emission peak at around 414, 475, 563 (at λex=270 nm)

Microwave t-ZrO2 at 320 °C Emission peak at 402 nm, 420 nm, 459 nm [185] when excited at λex = 254 nm Hydrothermal m-ZrO2 at 200 °C Broad emission and maximum intensity at 400 [186] and 250 °C nm, (at λex=290 nm and 300 nm) Pechini type sol-gel t-ZrO2 at 500 °C Maximum peak intensity at 425 nm with a wide [187] range of excitation (230 nm-400 nm) Sol-gel Mesoporous t-ZrO2 Broad peak centered at 419 nm when excited at [170] λex = 285 nm Soft templating and Mesoporous Emission peak intensity at 392, 394 and 398 nm [188] solid-liquid method wormlike t-ZrO2 when excited at 240 nm, 250 nm and 260 nm, respectively.

Template method Nanowire t-ZrO2 Emission peak 330 nm to 450 nm centered at [190] 388 nm when excited at λex = 230 and 235 nm. 3+ Sol-Gel Eu doped t-ZrO2 Emission peak at 592 nm, 608 nm, 630 nm, 658 [29] nm, when excited at 465 nm. 3+ Sol-Gel Eu in (c or t) ZrO2 Tunable emission, purplish-pink to greenish – [195] blue at 285 nm and 393 nm. Solution combustion Tb3+ in (m and t) Bright green (emission bands 490nm, 545 nm, [196] ZrO2 588 nm, 620 nm at 325 nm.) Sono chemical process Eu3+ and Tb3+ doped Blue and green emission for Tb3+; red for Eu3+ [197] 3+ 3+ (m and t) ZrO2 and 1.2 % Tb & 0.8 % Eu lie in white light region. Hydrothermal method Eu3+ doped cubic White light having CIE coordinates x = 0.32, y [198] = 0.34 (at 160 °C) when excited at λex= 244 nm. ZrO2

3+ Polyol method Eu doped ZrO2 Two peaks at 58 nm and 640 nm when excited [199] cubic phase (500 at λex= 267 nm °C)

2.7: Summary of literature Several synthesis methods were adopted for developing zirconia nanopowders using different precipitating agents. However, preparation of agglomeration free as-synthesized zirconium hydroxide or oxides powders by using different wet-chemical routes, without any

addition of stabilizing agents or surfactants is a great challenge. In general, due to smaller crystallite size, it agglomerates during synthesis process and further it grows to a bigger particles during calcination process. It was further suggested that if the particle size increases to 30 nm, it transformed from t-zirconia to m-zirconia and thus supress the applications of t-zirconia in potential areas of fields. Among precipitating agents, the common precipitating reagent such as NH4OH led to prepare amorphous phase of zirconium hydroxide in the as-synthesized condition and led to convert to t-zirconia at around 450 °C and in most cases m-zirconia was formed along with t-zirconia above 500 °C. So, it a hard task to select a suitable precipitating reagent to develop agglomeration free powders in the as-synthesized condition as well as stable the t-phase of zirconia up to moderate or higher temperature. In this context, it was found that aqueous sodium borohydride which evolve in-situ hydrogen gas bubbles may be an appropriate precipitating agent to enhance the phase stability of amorphous as well as t-phase of zirconia. In addition it may also help to form agglomeration free powders due to the presence of gas-bubbles in aqueous NaBH4 during synthesis process. However, there was a few related literatures are available, where it showed that NaBH4 can be applied as a precipitating agent for preparing a transition metal oxides such as zirconium oxide. However, there was no reaction mechanism between

ZrOCl2·8H2O and NaBH4 was available in the literature. In the same time, the gelation- precipitation mechanism was not discovered. It was also observed that using gas-bubbles template method, various porous nanostructure materials have been developed. These agglomeration free powders may also develop porous powder during calcination process. So, use of sodium borohydride on the formation of thermally stable porous t-zirconia was not discussed in any literatures. The porous nature such as pore size, pore shape and pore morphology at different calcination temperature was also not discussed in detail in any literatures. In the same time, the use of the porous zirconia derived from borohydride route may suitable for different novel applications such as adsorption of heavy metal ions. Also, the incorporation of rare earth ions such as Eu3+, Tb3+ or mixture of these two in porous zirconia may lead to develop a phosphor material for lighting applications. 2.8: Statement of the problem

Zirconium oxide or zirconia having tetragonal structure (t-ZrO2) with smaller crystallite or particle size has many potential applications in different fields. Generally, zirconia crystals grow in rapid way through agglomeration process and achieve a quick crystallization peak temperature at around 450°C, which lead to form stable pure t-ZrO2 up

to 500°C, when using the common precipitating agent such as NH4OH without adding any surfactants or stabilizers. The crystallization peak temperature or activation energy of crystallization of t-ZrO2 can be enhanced by controlling the crystallite size in a smaller dimension having agglomeration free nature powders during wet-chemical synthesis through an efficient way using a suitable precipitating agent and thus it is possible to stable the amorphous as well as t-ZrO2 up to moderate temperature.

Further, zirconia (ZrO2) having porous nature with adequate pore diameter, pore volume, and surface area have opened new possibilities for different applications in structural, dielectric, catalyst, adsorption, optical and biomedical. Various synthesis methods have been optimized to fabricate porous nature of zirconium oxide using hard or soft templates as well as self-assembly process with the help of specific surfactants. In addition, zirconium oxide in the form of tetragonal (t) at room temperature was considered to be an important ceramic material and the above synthesis methods have been employed for the development of porous structure. However, poor structural stability and difficult in removal of template led to affect the potential application of porous zirconium oxide. In order to develop a porous structure in nanoparticles, it is necessary to form loose or agglomeration free powders in the as-synthesized condition. But it is too difficult to prepare agglomeration free powders through the wet-chemical route without adding any surfactants or templates, because the highly energetic nuclei particles in solution coagulate with each other to form large agglomerate particles by decreasing their surface energy. So, in this context, a new progress has been made in tailoring the synthesis of agglomerated free or loose particles in the as-synthesized condition through gas-bubbles template mechanism. This method help to control the agglomeration of nuclei and thus control the size of the particle in the as-synthesized condition. The agglomeration free smaller size particle may form stable tetragonal zirconium oxide up to moderate temperature. In the same time, the as-synthesized loose particles may develop moderate temperature stable porous particle during calcination process due to coarsening of particles as well as coalescence of voids. Also, the loose nature of in-born nanopowders may lead to develop porous zirconium oxide with various pore morphologies during calcination process. Additionally, it may also help to develop and stable both the amorphous as well as tetragonal zirconium oxide up to moderate temperatures, so that it may be suitable for various possible applications. Based on the critical literature review, the major challenge is to develop agglomeration-free ultra-fine nanopowders via wet-chemical synthesis method without

adding any surfactants or additives. However, different types of surfactants or additives have been utilized during wet-chemical synthesis method using different precipitating agents in order to achieve agglomeration-free nanopowders. The first concept of this research work is to avoid the addition of surfactants or additives during synthesis and to find out a suitable precipitating agent, which may evolve in-situ gas-bubbles during synthesis. These gas-bubbles create numerous gas-liquid interfaces and may help to develop agglomeration-free or loose nanopowders.

Without any addition of additives such as Y2O3/MgO/CaO in to zirconium-salt precursor to develop stable tetragonal/cubic phase of zirconium oxide (considered to be an important ceramic material for different applications) up to moderate temperature (600 °C- 800 °C) is a big challenge. The formation of t/c phase of zirconia without any additives also depends on the initial nuclei size and also depends on the particle growth during calcination process. The second concept of this research work is to develop as-synthesized loose zirconium hydroxide nanopowders using gas-bubbles evolving precipitating reagent. During synthesis, these gas-bubbles act as nucleation center, which led to form loose nanoclusters. These loose structures may help to suppress the coarsening of particles during calcination process and may form a stable t/c phase zirconia up to moderate temperature. This novel synthesis process will be not only stable the t/c phase zirconia up to moderate temperature but also form porous nanostructures within the zirconia nanoparticle due to the presence of loose structure, which was develop due to the use of gas-bubbles evolving reagent. The third concept of this research work is to study the temperature mediated phase transformation, powder morphology and pore geometry. The fourth concept of this research work is to apply this porous material for two suitable applications. One of the applications of porous zirconia is the adsorption of heavy metal ions in order to purify the industrial waste water for environmental applications. Another application of porous zirconia is in the luminescence area. It was observed that smaller particle size of porous structure may show luminescence properties in un-doped zirconia. If the rare earth ions are incorporated in porous structure, then it may lead to stable the t- and / or c-phase of zirconia up to moderate temperature. In addition, colour tuning is possible at various excitation wavelengths starting from near vacuum ultra-violet (near VUV i.e. 205 nm-210 nm) to ultra-violet (UV) range (i.e. 220 nm to 350 nm). White light in zirconia is also possible by incorporating combination of suitable rare earth ions.

2.9: Objectives

In this research work, borohydride synthesis method was adopted to develop nanostructured transition metal oxide such as zirconium oxide. In this context, hydrogen

(H2) gas-bubbles evolving aqueous sodium borohydride (NaBH4) have been used as an effective precipitating agent to develop loose nanoclusters of zirconium hydroxide in the as- synthesized condition and porous zirconia at different calcination temperatures. Based on the statement of problems and critical literature review, this research work consists of the following objectives.

 To justify the importance of borohydride synthesis using aqueous NaBH4 over the

common precipitating agent such as NH4OH for the development of zirconium oxide nanopowders.

 To study the new concept of reaction mechanism between aqueous ZrOCl2·8H2O

and aqueous NaBH4 for developing agglomeration free or loose nanopowders of zirconium hydroxide.  To enhance the activation energy of crystallization of zirconia through an efficient

way of synthesis using NaBH4.  To develop a thermally stable porous tetragonal zirconium oxide nanomaterials as well as analyze the kinetics and regeneration study of Pb (II) toxic ions removal using porous zirconia for environmental applications.  To analyze the temperature mediated phase transformation, pore geometry and pore hysteresis transformation of borohydride derived porous zirconium hydroxide nanopowders.  To study the phase transformation, powder morphology and pore morphology as well as distribution of rare earth ions in borohydride derived rare earth based porous zirconia.  To study the photoluminescence at various excitation wavelengths starting from near vacuum ultra-violet (near VUV i.e. 205 nm-210 nm) to ultra-violet (UV) range as well as powder colour under UV light source for the rare earth based porous zirconium oxide powders.

Chapter 3 Experimental work This chapter deals with a discussion of the synthesis procedures developed for preparing the zirconia nanopowders using sodium borohydride. Thermal, Phase, Particle morphology, elemental distribution, pore morphology, adsorption removal of toxic ions, photoluminescence behaviour of un-doped and doped zirconia were studied using different instrumental techniques. The detail characterization techniques were also discussed in this chapter. 3.1 Raw Materials Two basic raw materials (purity 99.5%) were used in this research work and these are zirconium oxychloride (ZrOCl2·8H2O), and sodium borohydride (NaBH4). Europium oxide (Eu2O3), and terbium oxide (Tb4O7) were used as one of the starting material for luminescence study. Nitric acid (HNO3) was also used for mixing the Eu2O3 and Tb4O7 to form the nitrate solution. 3.2 Powder synthesis Conventional gelation-precipitation route was followed for preparing zirconium hydroxide powders in the as-synthesized condition, but only using ZrOCl2·8H2O as a precursor and NaBH4 as a precipitating reagent. Other than gelation-precipitation, a constant pH method was also adopted to prepare zirconium hydroxide powders in the as-synthesized condition. For synthesis, one particular composition 1 M ZrOCl2·8H2O and 0.5 M NaBH4 was chosen (based on our previous work). If concentration of ZrO2.8H2O and NaBH4 were varied, then there will be chance in the variation of the following factors: (i) evolution of hydrogen gas-bubbles, (ii) the exothermic peak temperature (from DSC), (iii) particle size and (iv) phase transformation behavior of t-zirconia. So, to understand the above said factors, one particular composition was chosen by varying the way of synthesis condition such as gelation, precipitation and constant pH method. The detail experimental procedure for powder synthesis for un-doped zirconia and rare-earth based zirconia are described in experimental section of individual chapter of the results and discussion. 3.3 General characterization

3.3.1 Thermal The thermal behaviour of the as-synthesized powders were analyzed by Differential Scanning Calorimetry-Thermogravimetry (DSC-TG) (NEFTZSCH STA 409C) in nitrogen

(N2) atmosphere at a heating rate of 10°C/min using -alumina as reference material. This gives information about the crystallization and weight loss/gain behaviour of the as- synthesized powders. 3.3.2 Phase analysis To understand the phases present in the as-synthesized as well as calcined powders, the room temperature X-ray diffractometer (Rigaku Ultima-IV, Japan) was performed using cupper target (Cu-Kα radiation i.e. 1.5405 Å). The volume percentage of m and t-phase of

퐼푚(1̅11)+퐼푚(111) zirconia was determined using the formula 푉푚 = and 푉푡 = 1 − 푉푚 퐼푚(1̅11)+퐼푚(111)+퐼푡(111) where I is the integrated intensity in each peak and m and t indicates monoclinic and tetragonal phase [171]. 3.3.3 Powder morphology Generally, powder morphology was studied using Scanning Electron microscopy or Transmission Electron Microscopy (TEM). We performed SEM on some of our samples and it shows agglomerated in nature. In order to visualize the internal porous structure (based on many literatures), we have used TEM and HAADF-STEM. The particle morphology was studied in Transmission Electron Microscopy (TEM). For preparation of TEM sample, the powders are dispersed in isopropyl alcohol and sonicated for half an hour. One drop of the well-dispersed sample solution is deposited on to a carbon coated copper grid (400 mesh). The dried grid was used for microscopy. In order to study the pore in the particle, high angle annular dark field scanning transmission electron microscopy (HAAD- STEM) image was analyzed. 3.3.4 Particle and crystallite size The particle size is usually obtained from TEM micrograph. However, The size corresponds to the mean value of the crystalline domain size of the particles is determined from the X-ray line broadening using Scherrer’s formula [201] i.e. crystallite size (nm) = 0.9

/(βCosθB), where  is the wavelength of Cu K (0.154 nm), β is the Full width at half 2 2 1/2 maximum (FWHM) and θB is the Bragg’s angle. β is expressed as (βs – βstd ) , where βs is the FWHM of sample and βstd is the FWHM value (0.0511) for standard Si-wafer sample. 3.3.5 Distribution of rare earth ions To understand the rare earth distribution in porous zirconia, elemental mapping and EDX line profile scan as well as EDX line scan were performed using TEM instrument. 3.3.6 Fourier transformation infrared spectroscopy (FTIR) Fourier Transform infrared (FTIR) spectrum represents a fingerprint of a sample with transmission peaks which correspond to the frequencies of vibrations between the

bonds of the atoms making up the material. It provides information about the presence or absence of a functional group along with the type of vibration associated with the particular functional group. Prior to analysis the sample preparation was done by properly mixing about 0.1 g of the powder sample with KBr. This mixture was finely grounded and then it was made a pellet. The pellet was used for the analysis. The FTIR spectrums are recorded in the spectral range of 4000-400 cm-1 with a Perkin Elmer FTIR spectrometer. 3.3.7 BET-isotherm and surface area The shape and size of various types of pores present in porous material are analyzed by Brunauer–Emmett–Teller (BET) method by using nitrogen adsorption–desorption hysteresis isotherms. The nature of the hysteresis loop can be divided into four different categories (H1, H2, H3 and H4) based on the IUPAC guideline [131]. According to IUPAC classification the H1 hysteresis are often associated with porous materials consisting of well-defined cylindrical-like pore with uniform sizes and shapes. Isotherms revealing type H2 hysteresis corresponds to channels with a pore mouth smaller than the pore body (this is the case of ink-bottle-shaped pores). The desorption branch for type H3 hysteresis contains a very wide distribution of pore size having slit like pores. Similarly, type H4 loops corresponds to limited amounts of mesopores limited by micropores. Surface area of zirconium oxide powders was determined from five points of adsorption curve. Pore size distribution was analyzed using Barrett-Joyner-Halenda (BJH) method by considering BET- desorption behavior. There are two ways to plot the pore size distribution from BET- isotherm curve: one is dV/dD versus D and dV/logD versus D. In this thesis, dV/logD i.e pore volume (unit is cc/g) as a function of pore diameter (nm) was plotted. 3.3.8 Adsorption of toxic ions and regeneration study To find out the removal efficiency (in %) of Cr(VI) or Pb(II) at different interval of time, two different solutions of potassium chromate (10 ppm) and lead nitrate (10 ppm) were prepared separately using distilled water at normal pH (~7). Calcined (600 °C) porous zirconium oxide of 10 mg was added in 10 ml of the above chromium and lead solutions. These solutions were stirred continuously at room temperature. At different interval of time (2, 3, 6, 9, 12, 15, 30, 60 and 120 minutes), the samples were collected via filtration. The filtered solutions were analyzed using atomic absorption spectroscopy. The removal efficiency (in percentage) of Cr(VI) or Pb(II) at different interval of time was calculated using the formula: [{(Co-Ce)/Co}100], where Co (ppm) and Ce (ppm) are the initial and equilibrium concentration of adsorbate in the solution.

In order to check the regenerative capacity of adsorbent, desorption study was carried out using the filtered Pb (II) loaded zirconium oxide sample, collected after 30 minutes of adsorption time (because the adsorption process was more than 98 % complete within this time). Desorption was carried out by agitating the Pb(II) loaded zirconium oxide with 5 ml of desorbing agent HNO3 (0.5 M) solution. After agitating for 1 hour, it was filtered and the filtrates were analyzed using atomic absorption spectroscopy for determination of recovery percentage of Pb(II) during desorption process. The filtered samples were dried and then again suspended in Pb(II) containing solution for next adsorption run. Five cycles of adsorption–desorption were carried out to examine the capability of the zirconium oxide sample to retain Pb(II) removal capability for regeneration. 3.3.9 Photoluminescence Photoluminescence (PL) studies were performed using Fluoromax-4 spectrofluorometer. The measurements were done at room temperature. The data were collected in the 200 nm to 700 nm range through a computer interfaced to the spectrophotometer at various excitation wavelengths. 3.3.10 Colour purity (R/O ratio) Colour purity is one of the important characteristics of a phosphor materials. Generally, the colour purity is denoted as Intensity ratio of red to orange [denoted as R/O 5 7 5 7 ratio i.e. I( D0/ F2)/I( D0/ F1)] and is calculated by considering the sum of integral intensity 5 7 of red emission peaks observed at 614 and 630 nm for contribution of D0 / F2 transition. 3.3.11 CIE-chromaticity coordinate graph To confirm the colour of rare-earth doped zirconia powders, the commission International de I’Eclairage (CIE) 1931 XY chromaticity coordinates were determined using CIE-software. The CIE 1931 graph contains triangle of R (red), G (green) and B (blue) and the position of a colour in the diagram is called chromaticity point of the colour. In general, the CIE coordinates were calculated by taking into account the visible range (400–700 nm) of the emission spectrum. 3.3.12 Powder colour under UV light source The powder sample was excited at a wavelength of 254 nm and 365 nm using UV source. Then, the photograph of the sample was taken.

Results and discussion Chapter 4 Motivation for selecting borohydride synthesis route to prepare zirconium oxide nanopowders

In this chapter, zirconium oxide nanopowders was prepared through conventional gelation-precipitation method using two different precipitating reagents; one is aqueous

NaBH4 and the other is NH4OH, a common precipitating reagent. Importance of selecting aqueous NaBH4 for preparing zirconium oxide nanopowders have been justified and discussed based on thermal, structure and powder morphology. 4.1 Experimental

In one of the experimental, zirconia nanopowders were prepared using 1.0 M Zr-salt

(ZrOCl28H2O) and 0.5 M NaBH4 through borohydride route. In this experiment, aqueous

NaBH4 was added drop wise to a beaker containing Zr-salt solution and the reaction was conducted at room temperature. While adding the aqueous NaBH4, gel network was formed at pH ~ 2.8 and this gel was converted to precipitate form at a pH ~ 10.5, by further addition of aqueous NaBH4. The precipitates were thoroughly washed with hot distilled water several times and then dried in oven to form as-synthesized borohydride derived powders. In another experiment, zirconia nanopowders were prepared by drop-wise addition of NH4OH to a beaker containing1.0 M Zr-salt (ZrOCl28H2O) solution. In this case, the gel was formed at pH ~ 4 and it was converted to precipitates at pH 10.5. In a similar way, the precipitates were thoroughly washed with hot distilled water several times and then dried to form as- synthesized ammonium hydroxide derived powders. Thermal study of these as-synthesized powders were performed using DSC-TG. Further, the dried powders prepared using NaBH4 and NH4OH were heat-treated at 600°C for 1h. Phase analysis, and powder morphology were studied using XRD and TEM, respectively.

4.2 Results and discussion

4.2.1 Thermal

Fig. 4.1 (a) and (b) show DSC-TG curve of the as-synthesized zirconia powders prepared using NaBH4 and NH4OH, respectively.

622 (a) (b) 3 100 3 442 100

2 95 95 2

90 1 90 1

85 85 0

Mass loss (g)

DSC/(mW/mg)

DSC/(mW/mg) 0 80 80 -1 75 -1 75 -2 0 200 400 600 800 1000 0 200 400 600 800 1000 Temperature (°C) Temperature (°C)

Fig. 4.1: DSC-TG curve of as-prepared zirconia synthesized using (a) NaBH4 (b) NH4OH.

A total weight loss of around 25 % was observed from room temperature to 1000 °C for these two samples, as confirmed from the TG curves. In both samples, the endothermic peak at ~ 125 °C was observed and this was due to desorption of physically adsorbed water. The DSC-TG behaviour of these two samples look similar, however, there is a difference in the position of the exothermic peak temperature that corresponds to the crystalline growth of zirconium oxide from the amorphous zirconium hydroxide. The exothermic peak at 622°C and 442°C correspond to crystalline growth of zirconium oxide, synthesized using

NaBH4 and NH4OH as a precipitating agent, respectively. The exothermic peak temperature or the crystallization temperature values of zirconia derived from NaBH4 was much higher than that of the crystallization temperature of zirconia prepared using a commonly precipitating agent, NH4OH. The reason for shifting of peak is discussed in chapter 6. The shifting nature of exothermic peak of crystalline zirconia has more advantageous for sustaining the amorphous nature as well as stabilizing the tetragonal phase of zirconia up to moderate temperature and the borohydride derived zirconia powders may find some potential application in different fields.

4.2.2 Structure

Phase analysis of un-washed precipitated powders collected after borohydride and ammonium hydroxide reaction were performed using XRD and shown in Fig. 4.2. The un-

washed precipitated powder prepared by ammonium hydroxide was showing amorphous in nature. However, the borohydride process, the powders have contained impurity phase of NaCl. So, the borohydride powders after reaction were thoroughly washed with hot distilled water and further XRD was performed. The powders after washing is showing amorphous in nature without any peak of NaCl.

un-washed (NH OH) 4

washed (NaBH ) 4 #

Intensity (a. Intensity u) # NaCl # un-washed (NaBH ) # 4 # # #

20 30 40 50 60 70 80 2 (degree)

Fig. 4.2: XRD patterns of un-washed as-synthesized powders, derived via NaBH4 and NH4OH. XRD pattern of washed sample prepared using NaBH4 shows amorphous in nature.

It was observed that the as-synthesized powders derived either by using NaBH4 or

NH4OH are purely amorphous in nature. Further, based on thermal (DSC-TG) analysis, the as-synthesized powders are calcined at 450 °C and 600 °C for 1h. The phase analysis of calcined powders, prepared using both precipitating reagents (NaBH4 and NH4OH) were further analysed using XRD patterns. Fig. 4.3 shows XRD patterns of borohydride derived zirconia powders calcined at 450 °C and 600 °C. 450 °C t 600 °C

t t t t t

(a.u.)Intensity

Intensity (a.u.)Intensity

20 30 40 50 60 70 80 20 30 40 50 60 70 80 2 (in degree) 2 (in degree)

Fig. 4.3: XRD patterns of calcined zirconia powders, derived via NaBH4. [‘t' stands for t-phase of zirconia]

Similarly, Fig. 4.4 shows XRD patterns of NH4OH derived zirconia powders, calcined at 450 °C and 600 °C.

t 450 °C m 600 °C

t t m m m m t t t m m m m m

Intensity (a.u.)Intensity t (a.u.)Intensity m m m

30 40 50 60 70 80 30 40 50 60 70 80

2 (in degree) 2 (in degree)

Fig. 4.4: XRD patterns of calcined zirconia powders derived from NH4OH. [‘t' and ‘m’ stands for t- and m-phase of zirconia, respectively]

The borohydride derived sample calcined at 450 °C showed purely amorphous in nature. However, the NH4OH derived sample calcined at same temperature showed a mixture phase of t-zirconia (major) [as per the JCPDS file number: 79-1768.] and m- zirconia (minor) [as per the JCPDS file number: 83-0943.]. For this sample, the percentage of t-ZrO2 and m-ZrO2 was determined and found to be 71 vol % and 29 vol %, respectively. It indicates that amorphous phase of zirconium hydroxide or oxide, derived via borohydride route was much more stable up to 450 °C. When the powders were calcined at 600 °C, pure phase of t-zirconium oxide was formed when it was derived via borohydride route. All the peaks are assigned to pure t-ZrO2 in a single crystalline phase. At this temperature, an average crystallite size of borohydride derived t-ZrO2 is found to be 17 nm determined using

Scherrer’s relation. However, the growth of the NH4OH derived particles led to increase the monoclinic phase and decrease the t-ZrO2 at 600 °C. For this sample, the volume percentage of t-ZrO2 and m-ZrO2 at 600 °C was found to be 19 % and 81 %, respectively. From the phase analysis, it was confirmed that borohydride process is a novel process to develop t-

ZrO2 with smaller particle size up to moderate temperature. From the above results, it is quite possible that the borohydride reaction may extend the temperature stability of the amorphous and t-phase may be related to the incorporation sodium ion in the zirconium oxide lattice. There are a number of research papers, which claim that sodium ion stabilizes zirconia [202-204]. But, in these papers, the concentration of Na ion was around 8 mole % or 3 wt %. So, based on the literatures, it was found that Na

ion can stabilizes the cubic or tetragonal nature of zirconia, but depends on the concentration of Na. Moreover, various synthesis methods were adopted using NaOH as a precipitating agent to stabilize the t-phase of zirconia. However, no one had claimed the role of Na in their systems [205-207]. So, in the borohydride method, NaCl was washed away and forming amorphous nature of zirconium hydroxide in the as-synthesized condition. In this system, Na may be present, but in negligible amount. However, presence and quantification of Na in the current work has not been performed due to unavailability of instrument like XPS. But, EDX analysis have performed (see Chapter 8) and found that there was no evidence of

Na (K of Na is 1.04 eV) in this current system. So, Na ion may not be stabilize the t- zirconia up to moderate temperature in this work.

4.2.3 Powder morphology

Powder morphology of as-synthesized sample prepared using NaBH4 and NH4OH was studied using TEM. Fig. 4.5 (a) and (b) shows TEM micrograph of as-synthesized powders prepared using NaBH4 and NH4OH, respectively. The borohydride derived zirconia powders was found to be much more agglomeration free or loose in nature than the sample prepared using NH4OH.

Fig. 4.5: TEM micrograph of as-synthesized powder derived from (a) NaBH4 and (b) NH4OH.

Further, the powder morphology of the calcined powder samples were performed using TEM. Fig. 4.6 (a) and (b) shows TEM micrograph of calcined (600 °C) zirconia powders prepared using NaBH4 and NH4OH, respectively. The presence of pores are clearly

visible for the borohydride derived calcined zirconia samples. However, there is no indication of pores present in the zirconia samples derived using NH4OH.

Fig. 4.6: TEM micrograph of calcined (600°C) powders, derived via (a) NaBH4 and (b) NH4OH. 4.3 Remarks

The as-synthesized zirconia powders prepared using NaBH4 or NH4OH were found to be amorphous in nature. The powders are loose in nature, while preparing through borohydride route, whereas the powders are agglomerated in nature, when it was synthesized using NH4OH. The crystallization temperature from amorphous to crystalline zirconia was found to be higher (~622 °C) in case of sample prepared through borohydride route and thus the powders calcined at 450 °C was found to be amorphous in nature and the powders calcined at 600 °C was found to be crystalline (t-phase of ZrO2) in nature. However, both t-phase and m-phase of zirconium oxide were present in the sample prepared through NH4OH, when calcined at both 450 °C and 600 °C. Interestingly, fine pores are present in the calcined zirconia sample prepared through borohydride process. The enhancement of the exothermic peak (i.e. crystallization temperature from amorphous nature to t-phase of zirconia) as well as formation of loose nature of as-synthesized zirconia and development of porous nanoparticles at moderate temperature led to motivate for choosing NaBH4 as a precipitating agent. As the gelation-precipitation mechanism between zirconium salt with NaBH4 was not disclosed anywhere, so, gelation-precipitation mechanism of borohydride synthesis was discussed in the following chapter.

Chapter 5 Gelation-precipitation reaction mechanism of borohydride synthesis to prepare zirconium oxide nanopowders

Sodium borohydride is widely used as a reducing agent. However, in this work, zirconium oxide powders were prepared using sodium borohydride via gelation- precipitation route. Further, it is necessary to understand the reaction mechanism between aqueous ZrOCl28H2O and NaBH4. So, in this chapter, gelation-precipitation reaction mechanism of borohydride synthesis method was discussed based on the FTIR spectra of un-washed and washed samples during the synthesis process. 5.1 Experimental Borohydride synthesis via gelation-precipitation method was adopted to prepare as- synthesized zirconium hydroxide powders. For preparation of zirconium hydroxide powders, two different aqueous solutions of 1 M ZrOCl2·8H2O and 0.5 M NaBH4 were prepared separately. Borohydride reaction was conducted at room temperature with drop wise addition of aqueous NaBH4 to a beaker containing aqueous ZrOCl2·8H2O, with constant stirring using a magnetic stirrer. During addition of aqueous NaBH4, gelation took place at pH ~ 2.8 and further the pH of the precursor was increased to ~ 10.5, with addition of aqueous NaBH4 via precipitation process. FTIR analysis was performed on different samples mentioned in the results and discussion section. 5.2 Results and discussion

5.2.1 Gelation-precipitation mechanism between Zr-salt and NaBH4

Powder morphology of the borohydride derived as-synthesized zirconium hydroxide was found to be loose in nature, as observed from Fig. 4.5 (a). Also, the as-synthesized powders were found to be an amorphous in nature and this amorphous phase was stable up to 450 °C, as confirmed from the XRD pattern of Fig. 4.3. So, it seems that the borohydride

based gelation-precipitation reaction is favourable for the formation of loose porous nature (a) of zirconium hydroxide in the as-synthesized condition. Thus, it is justify looking into the

insights of the borohydride reaction mechanism between aqueous ZrOCl2·8H2O and

NaBH4. The following reaction mechanism was discussed in detail. It is well-known that

in aqueous solution, hydrolysis of NaBH4 proceeds to form two species such as tetrahydral

[58] boron [B(OH)4¯] and hydrogen (H2) gas bubbles as per equation (1)

NaBH4 + 4H2O NaB(OH)4 + 4H2 (1)

When aqueous NaBH4 (initial pH ~ 11) associated with two active species such as

B(OH)4 ¯ and H2 gas-bubbles was added to the aqueous solution of ZrOCl2·8H2O (initial

pH ~ 0.3), zirconium hydroxide [Zr(OH)4] nuclei are start to grow via hydroxide ion

exchange reactions along with the formation of trigonal boron [B(OH)3] and H2 gas-bubbles

as per equation (2). The B(OH)3 so formed in aqueous medium converts to tetrahedral boron, as per equation (3).

4 B(OH) ¯+ Zr4+ + H Zr(OH) + 4B(OH) + H (2) 4 2 4 3 2 ¯ + B(OH)3 + 2 H2O  B(OH)4 + H3O (3)

The percentage of B(OH)3 and B(OH)4 ¯ in solution depends on the pH of the precursor

solution. At lower pH value, B(OH)3 is more dominate than B(OH)4 ¯, whereas B(OH)4 ¯ is more dominate at higher pH [208]. During initial stage of synthesis, the highly energetic

Zr(OH)4 species are formed in presence of trigonal boron and H2 gas-bubbles. As the

reaction proceeds with continuous addition of aqueous NaBH4, pH of the precursor solution increases. At pH ~2.8, a viscos gel-network polymeric chain was observed. Generally, the

gelation pH of Zr(OH)4 is ~4, while using the common precipitating agent such as NH4OH

[47] . The sharp decrease of gelation to a pH ~ 2.8, while using aqueous NaBH4 indicates that the boron species strongly participates during gelation. To support this gelation mechanism, Fourier Transformation Infra-red spectroscopy (FTIR) of the gel sample (un-washed) was performed and shown in Fig. 5.1. The FTIR peak around 1420 cm-1 and 1195 cm-1 indicates the stretching and bending vibration of B-OH bond in trigonal boron respectively [209]. Both the stretching and bending vibration of trigonal boron are shifted to a higher frequency in

compared to the stretching and bending vibration of B-OH of pure B(OH)3 boric acid solution (stretching at 1410 cm-1 and bending at 1148 cm-1). This shifting is likely due to

[210] the strengthening of B-OH bond during gelation with Zr(OH)4 . In addition, the broad

band in the range of 1000 cm-1 to 850 cm-1 centered at 950 cm-1 was assigned to stretching vibration of tetrahedral boron [211]. This broadening nature is due to the formation of

[212] intermolecular hydrogen bonding of Zr(OH)4 with B(OH)4 during gelation process . Thus, FTIR spectra of the gel sample indicate the participation of both boron species during gelation process. The nature of peak in the range of 3000 cm-1 to 3500 cm- is also a strong indicator for the participation of boron species. The narrow sharp band at 3220 cm-1 indicates the polymerization via intermolecular hydrogen bonding between Zr(OH)4 with boron species. A lower intense peak at 3400 cm-1 indicates a minute polymerization among [213] −1 Zr(OH)4 units via inter molecular hydrogen bonding . The additional band at 800 cm is due to bending vibrations of B–O–B [214]. Furthermore, the band at 660 cm-1 and 450 cm- 1 are assigned to vibrational modes of Zr-O [47]. The peak at 1630 cm-1 corresponds to O-H of water. Also during gelation process, the H2 gas-bubbles were considered to be trapped within the gel-network [215].

(a)

1630

850

1000 800

660

1195 450

1420

Transmittance (a.u.) Transmittance

3400 3220 4000 3500 3000 2500 2000 1500 1000 500 -1 Wave number (cm )

Fig. 5.1: FTIR spectrum of un-washed gel sample

From the FTIR analysis, it was confirmed that the gel-network polymeric chain was due

[210] to the polymeric nature of both Zr(OH)4 and boron species . This boron species are able to form inter molecular crosslinking through hydroxyl group with Zr(OH)4 nuclei as shown in the schematic diagram of Fig. 5.2.

(b) H H O H H H O B O O B O H H H

O O H H H H O O H H H Zr O Zr O O O H H H O O H H H H O O H H H O B O O B O H H H O

H H

Fig. 5.2: Schematic representation of three-dimensional network of Zr(OH)4 with boron species

Again precipitation process was followed with vigorous stirring and further addition of aqueous NaBH4. In this process, the pH of the precursor solution gradually increases along with dissociation of gel-network as well as the trapped H2 gas-bubbles became active and mobile. The process of precipitation was continued till the pH reaches at 10. To understand the precipitate mechanism, FTIR of the dried un-washed precipitate powders (collected at pH 10.5) was performed and is shown in Fig. 5.3.

1640

Transmittance (a.u.)

660

900

3200

1100

3400 450 1407

4000 3500 3000 2500 2000 1500 1000 500 Wave number (cm-1) Fig. 5.3: FTIR spectrum of un-washed precipitate sample.

The band at 660 cm-1 and 450 cm-1 are assigned to vibrational modes of Zr-O. The reduction -1 -1 in FTIR intensity of 3200 cm along with increase in FTIR intensity of 3400 cm indicates the detachment of boron species from zirconium hydroxide species during gelation to precipitation process. The detached boron species is basically an inter-coordinated boron complex and it was possible through intermolecular hydrogen bonding among boron species only when pH of the solution increases beyond 9 [216]. The detachment of boron complex was confirmed from the decrease of stretching and bending vibration of B-OH from 1420 cm-1 to 1407 cm-1 during gelation to precipitation process. The broad peak starting from 1100 cm-1 to 900 cm-1 was also a strong indicator for the formation of boron complex [217]. However, the formation of a solid piece of boron complex from the precipitation solution is a slow process. A solid piece of boron complex was phase separated out from the solution when the precipitate solution was kept for two to three days. The solid pieces of boron complex in the precipitated solution were shown in the inset of Fig. 5.4. To understand the band position, FTIR of the solid piece was performed and is shown in Fig. 5.5.

Fig. 5.4: Solid pieces of boron complex were phase separated from precipitation solution.

3450

1630 (a.u.) Transmittance

630

540

1500

900

1350 1100 4000 3500 3000 2500 2000 1500 1000 500 Wave number (cm-1)

Fig. 5.5: FTIR spectrum of solid borate sample

The broad peak at 3450 cm−1 was assigned due to stretching mode of O–H band of boron complex. The two major broad peak ranges from 1100 cm-1 to 900 cm-1 and 1500 cm-1 to 1350 cm-1 corresponds to vibrational modes of boron complex. These two broad peaks are due to the polymerization of the boron species [212,214]. The peaks at about 630 cm- 1 , and 540 cm-1 are the characteristic frequencies of symmetric vibration of boron complex [217]. The peak at 1630 cm-1 was assigned as the O—H vibrational mode of water. From above FTIR spectra, it was confirmed that boron species are strongly participated during gelation and these boron species in the form of boron complex were phase separated out during precipitation process. So, the aqueous borohydride process is not only act a reducing agent but also act as a precipitating agent (analogous with NH4OH) to produce zirconium hydroxide precipitate powders. Further to understand the nature of vibrational modes of washed sample, the un-washed precipitate powders were washed several times and dried. The washed and dried as-prepared sample was calcined at 800 °C. The FTIR spectra of both washed as-prepared and calcined (800°C) zirconium oxide powders are shown in Fig. 5.6.

(II)

1630

(i) 660

935 1630

1350

1560

450

660

3400

Transmittance (a.u.) Transmittance

(i) Washed precipitate sample

(ii) Calcined at 800 °C 450

4000 3500 3000 2500 2000 1500 1000 500

Wave number (cm-1)

Fig. 5.6: FTIR spectra of washed precipitate and calcined (800 °C) sample

The complete absence of vibrational modes of boron species was observed in the as- prepared as well as calcined samples. The band position at 3400 cm-1, 1630 cm-1, 1560 cm- 1 -1 and 1350 cm of as-prepared sample are due to O-H vibration modes of H2O. The peak position at 450, 660 and 935 cm-1 are due to the Zr-O vibration modes [47]. Thus, it was confirmed that aqueous NaBH4 participated in gelation-precipitation to form borate free as- synthesized zirconium hydroxide powders. 5.2.2 Thermal Further DSC analysis was performed on the washed samples collected after gelation and precipitation, in order to understand the exothermic peak temperature i.e. crystallization behaviour of zirconium hydroxide powder. Also for comparison, DSC curve of the gelation and precipitated powder, prepared via NH4OH were also plotted. DSC curves of four

samples are represented in Fig. 5.7. The exothermic peak was observed at 600 °C and 622 °C, when sample was collected after gelation and precipitation, which is much higher than the sample prepared using NH4OH (peak temperature is 435 °C for gelation and 422 °C for precipitation). However, the exothermic peak temperature values of zirconia derived from gelation is found to be lower than that of exothermic peak temperature of precipitated powder. (d) 622°C

442°C 435°C (a) (b) (c)

600°C

flow (a.u.) Heat

200 400 600 800 Temperature (°C) Fig. 5.7: DSC of borohydride derived as-prepared zirconia powders. [a and b stand for gelation and precipitation derived sample prepared using NH4OH; c and d stand for gelation and precipitation derived sample prepared using NaBH4]

5.3 Remarks

For the first time, we are reporting the advantages of aqueous NaBH4 for facilitating in gelation and precipitation with aqueous ZrOCl2·8H2O and producing agglomeration free zirconium hydroxide powders in the as-synthesized condition. The boron species such as trigonal and tetrahedral boron are strongly participate in gel-network polymeric chain with

Zr(OH)4 during gelation process. However, the boron complex in the form of solid pieces was phase separated from the precipitated solution after completion of the precipitation reaction and thus forming boron free zirconium hydroxide [Zr(OH)4] in the as-synthesized condition. In addition, the exothermic peak temperature was found to be enhanced with compared to the sample prepared using the common precipitating agent. So the exothermic peak temperature can be further enhanced via different ways of borohydride synthesis. So, an efficient way was developed to enhance the exothermic peak temperature and was discussed in next chapter. In addition, the role of H2 gas-bubbles to develop porous zirconia and adsorption kinetic of removal of lead ion as well as regeneration study was also discussed for suitability of this porous zirconia for environmental application.

Chapter 6 Enhanced activation energy of crystallization, development of porous zirconia and application of porous zirconia for removal of Pb (II) toxic ions

In this chapter, an efficient way such as gelation, precipitation and constant pH method via borohydride synthesis route was performed to enhance the activation energy of crystallization of pure zirconia in order to stable the amorphous as well as t-zirconia up to moderate temperature. Formation of porous stable t-zirconia while using sodium borohydride as a precipitating agent was also discussed in this chapter. One of the potential application of this porous zirconia is adsorption of Pb(II) toxic ions. So, adsorption kinetic mechanism as well as regeneration of Pb (II) loaded zirconium oxide sample was also discussed in this chapter. 6.1 Experimental

Pure t-ZrO2 powders were prepared using 1.0M Zr-salt (ZrOCl2·8H2O) and 0.5M

NaBH4 through borohydride route via three different reaction conditions such as gelation, precipitation, and constant pH. In gelation route, aqueous NaBH4 was added drop wise from a burette to a beaker containing Zr-salt solution at room temperature. The gel was formed at pH ~2.8. In precipitation route, the gel was converted into precipitate form at a pH ~10.5, by further addition of aqueous NaBH4 with the help of stirring. In case of constant pH route, an aqueous Zr-salt solution was added to a beaker containing NaBH4 and during the reaction a constant pH ~10.5 was maintained by addition of extra amount of aqueous NaBH4 from another burrette. The gelation, precipitation, and constant pH products were thoroughly washed, dried, and calcined at 600°C and 800°C for 1 h. Thermal behavior of zirconia powders were studied using differential scanning calorimetric (DSC) at a heating rate of 10°C/min in argon atmosphere. Phase analysis was performed by X-ray diffraction (XRD) using Cu-K radiation. The average crystallite size of zirconia powders was determined from the X-ray line broadening using Debye–Scherrer formula with correction factor. Particle size as well as morphology was studied by transmission electron microscopy (TEM).

6.2 Results and discussion

6.2.1: Thermal DSC curves of as-synthesized zirconia powders prepared via gelation, precipitation and constant pH routes are compared in Figure 6.1 .

717°C

Constant pH

622°C

Precipitation

600°C Gelation

Heat Flow (a.u)

0 100 200 300 400 500 600 700 800 900 Temperature (°C) Fig. 6.1: DSC curves of as-prepared zirconia synthesized through three different routes.

The endothermic peak at ~ 125 °C was due to desorption of physically adsorbed water. The exothermic peak at 600°C, 622°C and 717°C correspond to crystalline growth of amorphous zirconia synthesized via gelation, precipitation and constant pH route, respectively. These exothermic peak temperature values are much higher than that of the crystallization temperature of zirconia prepared using a commonly precipitating agents like [47] NH4OH .

The enhanced crystallization temperature of t-ZrO2 was generally observed either in [218] composite systems like Al2O3-ZrO2 or hydrous zirconia digested in solutions containing + + + [219] [46] Na , K , and NH4 cations or surfactants assisted synthesis of ZrO2 or by varying the concentration (0.06 M to 0.0025 M) of Zr-salt [220]. To the best of our knowledge, for the first time, we are reporting higher exothermic peak or crystallization temperature of pure t-

ZrO2 via constant pH route using 1.0 M Zr-salt and 0.5 M NaBH4. To understand higher crystallization temperature of t-ZrO2, activation energy of crystallization (Ea) was calculated from DSC curves by using the Redhead equation: Ea = RTp[ln(νTp/β) – 3.64],where ν is 13 -1 taken as 10 s ,β is the heating rate, Tp the exothermic peak temperature, and R is the ideal

gas constant [221]. The zirconia sample prepared through constant pH exhibits higher activation energy (Ea ~ 260 kJ/mol), when compared with the Ea of gelation (Ea~ 230 kJ/mol) or precipitation (Ea ~ 236 kJ/mol). These values are much higher than that of the Ea [219] (~ 180 kJ/mol) of zirconia prepared using NH4OH . Similar Ea values in the range from 178 to 213 k J/mol were reported by Chuah et. al in the crystallization of hydrous zirconia + + + [219] digested in solutions containing Na , K , and NH4 cations . Higher activation energy leads to stabilize pure t-ZrO2 with smaller crystallites up to moderate temperature, which was further confirmed from XRD patterns. 6.2.2: Structure Figure 6.2 shows XRD patterns of zirconia powders calcined at 600 °C, synthesized in three different ways. The as-prepared zirconia remains amorphous up to 600°C in constant pH route, whereas, pure t-ZrO2 with a crystallite size of ~ 4 nm are developed in the samples synthesized via gelation or precipitation route.

* t-ZrO 2 Constant pH

(101) *

(002)

(112)

(110)

(200) * * (103) (211) Precipitation

* * * (202) * *

Intensity (a.u.) * ** * * Gelation *

20 30 40 50 60 70 80 2 (in degree)

Fig. 6.2: XRD patterns of zirconia powder calcined at 600°C.

* t-ZrO , # m-ZrO 2 2 *(101)

Constant pH

(112)

* (211)

(110)

(200)

(002)

(103) * * * (202) * * * *

* * * Precipitation * * * * Intensity (a.u.) Intensity

(110)

# (111) # (002)

(110)

(101)

(002)

(020) # (220) Gelation # * # # #

20 30 40 50 60 70 80 2(in degree) Fig. 6.3: XRD patterns of zirconia powder calcined at 800 °C.

XRD patterns of ZrO2 powders calcined at 800 °C, synthesized via three different routes are compared in Figure 6.3. Mixture of 15 vol % monoclinic (m)-ZrO2 along with 85 vol % t-ZrO2 were obtained via gelation route, whereas, only t-ZrO2 was observed for both the samples prepared through precipitation and constant pH route. The crystallite size of t-

ZrO2 through precipitation and constant pH routes are found to be 23 nm and 20 nm, respectively. 6.2.3: Powder morphology

Fig. 6.4 (a), (b) and (c) represent TEM micrographs of ZrO2 (calcined at 800 °C) synthesized through gelation, precipitation and constant pH route, respectively. The particle size of zirconia was larger via gelation route, whereas, smaller size particles were obtained in both precipitation and constant pH route. But, the particle size of t-ZrO2 via constant pH was much smaller (~ 20 nm) than that of precipitation route (particle size ~ 30 nm), as seen from TEM micrographs, which was also consistent with crystallite size of t-ZrO2 obtained from XRD.

(a)

Fig. 6.4: TEM micrographs of calcined (800°C) ZrO2 powders, synthesized through (a) gelation, (b) precipitation and (c) constant pH.

Based on our results, it was found that the way of synthesis using NaBH4 is strongly effect the stability to sustain pure t-ZrO2 up to moderate temperature and also it was confirmed that controlled growth of smaller crystallites lead to enhance the activation energy of crystallization, which further effect the stability of pure t-ZrO2. The formation of smaller crystallites and its controlled growth is again strongly depending on the way of reaction using NaBH4. When NaBH4 react with water, it forms tetra hydroxy borate ion - [55] [B(OH)4 ] with the evolution of hydrogen (H2) gas . In all three ways (gelation, - precipitation and constant pH) of synthesis, the formation of B(OH)4 ions help to produce zirconium hydroxide nuclei with high surface energy and simultaneously, the released H2 gas bubbles create numerous gas–liquid interfaces inside the solution. Depending on the amount of released H2 gas bubbles, the gas-liquid interface creates as many numbers of nucleation or agglomeration centers to further attract zirconium hydroxide nuclei in order to minimize the interfacial energy [110, 222]. To become thermodynamically stable, high surface energy unstable zirconium hydroxide tried to combine either with another zirconium hydroxide or attach with the gas-liquid interface. So, if the number of zirconium hydroxide nuclei is more than that of released H2 gas bubbles during synthesis, then less number of gas-liquid interfaces is not sufficient to reduce the surface energy of all nuclei. Thus, the probability for formation of gel network structure among zirconium hydroxides is higher.

From borohydride synthesis route, it is well-understood that the amount of released H2 gas bubbles in the aqueous solution is higher in constant pH and decreases from precipitate to gelation. Thus, in constant pH, the evolution of huge amount of H2 gas bubbles inhibits the formation of gel network structure among the zirconium hydroxide nuclei. Further, with evolution of huge amount H2 gas during constant pH, the agglomeration of zirconium hydroxide nuclei occurred around the H2 gas bubbles. So, the presence of huge amount of

H2 gas in the form of smaller dimension may enable this agglomeration process to proceed in a controllable way and finally loose nanoclusters are formed [222]. These loose smaller crystallites grow in a controlled way during calcination and thus enhance the activation energy of crystallization of t-ZrO2. Thus, constant pH is more efficient way of synthesis to form amorphous zirconia up to 600°C and stable pure t-ZrO2 at 800°C using borohydride route. A detail study of the phase transformation, powder and pore morphology of the borohydride derived constant pH method was presented in the next chapter. However, the precipitated sample shows porous in nature as per TEM 6.4 (b). So, powder morphology of the precipitated sample was further studied at different calcination temperatures. Also, in

order to find out the existence of loose nature in zirconium oxide, the as-prepared amorphous zirconium hydroxide powders were thermally heat-treated at different calcination temperatures in the range of 400 °C to 800 °C. Powder morphology were analyzed using TEM. Figure 6.5 (a), (b) and (c) show TEM micrographs of as-synthesized powders calcined at 400 °C, 600 °C and 800 °C, respectively.

Fig. 6.5: TEM micrographs of as-synthesized powders (constant pH) calcined at (a) 400 °C, (b) 600 °C and (c) 800 °C. Inset of each micrographs show electron diffraction pattern.

Electron diffraction pattern was also performed on these three samples and are shown in the inset of Fig. 6.5. Electron diffraction pattern in the inset of Fig. 6.5 (a) indicates that the amorphous nature of zirconium hydroxide remains the same. Comparing the two different types of electron diffraction pattern such as hazy ring in inset of Fig. 6.5 (b) and sharp ring in inset of Fig. 6.5 (c), it was further confirmed that minute amount of amorphous

nature still remain along with crystalline nature at 600 °C, but forms purely crystalline at 800 °C. The sharp ring pattern at 800 °C was indexed with the tetragonal form of zirconium oxide. The TEM micrograph of Fig. 6.5 (a) indicates that the loose nature of zirconium hydroxide remains unchanged up to 400 °C. At this temperature, the particles are also well separated with each other by voids with analogues with TEM micrograph (Fig. 4.5 a) of as- synthesized powders. A trapped pore based porous nature of zirconium oxide was observed at 600 °C, as confirmed from Fig. 6.5 (b). The porous nature was also found to be present in the zirconium oxide sample, calcined at 800 °C, as shown in Fig. 6.5 (c). At this temperature, trapped pore and occasional large voids are also observed. However, the particles are agglomerated in nature at this temperature and to visualize one single particle and its porous nature, a higher magnified TEM micrograph of calcined (800 °C) zirconium oxide was represented in Fig. 6.6 (a). It was found that nanoscale pores (~3/4 nm) were well preserved in each individual zirconium oxide nanoparticles, having particle size of ~30 nm. Further, the pore size distribution of calcined (800 °C) zirconium oxide was studied using BJH curve and shown in Fig. 6.6 (b). The BJH curve also indicates that the porous zirconium oxide exhibits wide pore size distribution in the range of 3.6 nm to 15.8 nm, with an average pore diameter of 4.6 nm, which was well correlated with TEM result.

0.14 (b) 0.12

0.10

0.08

0.06

0.04 (cc/g) Volume 0.02 0.00 0 2 4 6 8 10 12 14 16 18 20 Pore diameter (nm)

Fig. 6.6: Higher magnified TEM micrograph (a) and BJH curve (b) of porous zirconium oxide, calcined at 800 °C.

6.2.4: Pore evolution mechanism Further, it was necessary to understand the pore evolution mechanism of the precipitated powder. So, it was discussed based on the following steps. The first step is the involvement

of evolved H2 gas-bubbles during the nucleation of Zr(OH)4 species to form loose zirconium hydroxide in the as-synthesized condition. During borohydride synthesis, the H2 gases generated in the precursor solution were released as gas-bubbles, which act as free-templates [110, 111, 113, 223]. These gas-bubbles create numerous gas–liquid interface aggregation centres throughout the precursor solution during synthesis process [223]. These gas-liquid interface centres help to reduce the interfacial energy of highly energetic Zr(OH)4 species through [224] their surface attachment . The surface attached Zr(OH)4 nuclei were well separated by gas-bubbles and thus it help to prevent continuous agglomeration of Zr(OH)4 species among [113] themselves . The size of Zr(OH)4 may be controlled depending on the quantity of gas- bubbles created during synthesis. However, the quantity of evolved gas-bubbles is difficult to quantify. But, in our previous observation [46], it was confirmed that the quantity of gas- bubbles in the aqueous solution is higher in constant pH method and decreases from precipitate to gelation process. It was also found that the crystallite or particle size of zirconium oxide strongly depends on the way of synthesis and it was due to the participation of different amount of gas-bubbles. In the same time, the participation of different quantity

(based on way of synthesis) of gas-bubbles affects the agglomeration of the Zr(OH)4 during synthesis. The higher amount of the gas-bubbles, more is the nucleation centres and thus supresses the agglomeration of Zr(OH)4. Whereas lower the amount of gas-bubbles during synthesis led to agglomerate the Zr(OH)4 nuclei and finally the cluster size of Zr(OH)4 became higher. However, in this present study, the precipitation derived as-synthesized zirconium hydroxide powder prepared via borohydride route was found to be loose in nature and thus it was assumed that sufficient amount of gas-bubbles are participating during precipitation process and help to suppress the agglomeration process of Zr(OH)4 nuclei during synthesis. In the solution state, the H2 gas-bubbles are surrounded by Zr(OH)4 nuclei, but these gas-bubbles create interconnected voids due to escaping of gas-bubbles in dry state. Further, the particles of as-synthesized powder are well separated with each other by voids as seen from Fig. 1 (a), and thus assume that there is no presence of trapped hydrogen gas in the dry state. The second step involves the formation of trapped pore (within one particle) and or large voids (within some agglomerated porous particles) when the loose nature of as- synthesized zirconium hydroxide undergoes calcination process. During initial stage of calcination process i.e. up to 400 °C, it is basically the decomposition of attached crystalline water in Zr(OH)4. The uniformity of the as-synthesized amorphous zirconium hydroxide powders help to maintain the original pore structure i.e. loose nature of zirconium hydroxide

up to 400 °C. This indicates that the presence of intra-particle voids may inhibit the mass transfer between loose nanoparticles and also help to restrict the coarsening of particles, during calcination up to 400 °C [59]. With increase in calcination temperature up to 600 °C, the seed crystals start to develop within the loose amorphous matrix. These growing nanocrystals impinge in the matrix and thus trapping inter-particulate voids as trapped pore. Again, while heating the sample from up to 800 °C, the coarsening of fine particles along with coalescence of existing pores take place. Occasional large voids can be observed which are formed as a result of excessive necking between particles and coalescence of the adjoining pores. 6.2.5: Adsorption kinetics mechanism The formation of moderate temperature stable porous zirconium oxide obtained from borohydride synthesis is found to be attractive and may find potential applications for the adsorption of heavy metal ions for industrial waste water treatment. In this context, the removal efficiency (in %) of Cr(VI) or Pb(II) at different interval of time was performed and is shown in Fig. 6.7. The adsorption process by the porous zirconium oxide is time independent and within 15 minutes, the removal efficiency of Cr(VI) and Pb (II) was found to be ~ 10 % and ~ 99%, respectively. The removal percentage of toxic metal ions by the porous zirconium oxide is selective in nature. So, the borohydride derived porous zirconium oxide may be a strong candidate for almost complete removal of Pb (II) toxic ions from water solution.

100 12 90 10 80 8

70 6 60 4 Pb (II) 50 Cr (VI) 2

Pb (II) removal efficiency (%) 40 0 Cr (VI) Removal efficiency (%) 0 15 30 45 60 75 90 105 120 Time (min.) Fig. 6.7: Removal percentage of Cr(VI) and Pb(II) with different interval of time.

Further, adsorption kinetic mechanism as well as regeneration of Pb (II) loaded zirconium oxide sample was studied. In order to examine the controlling mechanism of the adsorption process, pseudo-first order, pseudo-second order, Elovich, intra-particle

diffusion (Weber and Morris’ equation) and Bangham’s (pore diffusion) kinetic models were used. The experimental data were treated with the above five kinetic models and the equations of the different kinetic models are given in Table 2.1. The kinetic parameters determined for five kinetic models are given in Table 6.1. The graphs for the pseudo-first order, pseudo-second order, Elovich, intra-particle diffusion (Weber and Morris’ equation) and Bangham’s model (pore diffusion) are shown in Fig. 6.8 (a)-(e), respectively.

0.8 14 Calculated data 0.7 Calculated data (a) (b) ) 12 Fitted data Fitted data ) -1 0.6 -1 10 0.5

mg g 8

(

0.4

)

0.3 6

-q

min g mg

(

q 0.2 4

(

t 0.1 t/q 2 log 0.0 0 -0.1 2 4 6 8 10 12 0 20 40 60 80 100 120 t (min) t (min)

11 (c) Calculated data (d) Stage 2 10 Fitted data 10

)

9 -1

)

-1 8 8

mg g

( Stage 1 Calculated data 7

mg g

(

q Fitted data t 6

q 6

4 4 0.5 1.0 1.5 2.0 2.5 3.0 2 4 6 8 10 12 ln (t) t (min)1/2

(e) Calculated data

] 0.0 Fitted data

) t

-C i -0.2

(

[ -0.4

log log -0.6 0.2 0.4 0.6 0.8 1.0 1.2 log (t)

Fig. 6.8: Pseudo-first order (a) and pseudo-second order (b), Elovich (c), Intra-particle diffusion (d) and Bangham (pore diffusion) (e) kinetic plot for adsorption of Pb (II) by porous zirconium oxide powders.

The pseudo-first order and pseudo-second order are the most widely used rate equations to describe the adsorption of adsorbate from the liquid phase [225]. The pseudo-first order graph [Fig. 6.8 (a)] was found to be linear with a correlation coefficients of R2 = 0.9747, indicating the possible applicability of pseudo first-order model in the present study. But, the pseudo-second order graph [see Fig. 6.8 (b)] was also found to be linear with a correlation co-efficient value [R2 = 0.9995] higher than pseudo first-order model. Moreover, the qe cal (mg/g) value determined from pseudo-first order and pseudo-second order equation was found to be 6.605 and 10.07, respectively. The qe cal (mg/g) value obtained from pseudo- second order was very close to the experimental qe exp (mg/g) value (9.88). So, the experimental data to pseudo-first order rate equation suggested the non-applicability of pseudo-first order kinetic in predicting the mechanism of Pb (II) adsorption process. However, higher value of correlation co-efficient (R2) and nearly close matching value of calculated and experimental qe value indicates the applicability of pseudo-second order kinetics in this present study. The superior fit of the pseudo-second-order model along with matching of calculated and experimental qe (mg/g) data implies that the adsorption process was surface reaction controlled with chemisorption involving valence forces through sharing or exchange of electrons between adsorbent and adsorbate [225, 226]. Further, the quick adsorption of Pb (II) by porous zirconium oxide powder was observed after a short time period, which led to examine the experimental data with the Elovich kinetic model and the graph is shown in Fig. 6.8 (c). The Elovich equation assumes the presence of active sites on the adsorbent surface, which led to chemisorption [227, 228]. Additionally, the correlation coefficient (R2) value for Elovich model was lower than the pseudo-second order, but still rather high. So, Elovich model suggests that chemisorption was the main adsorption controlling mechanism. However, the initial rapid Pb(II) adsorption within 15 minutes also suggests that more than one mechanism may also involve in the process [229]. The pseudo- first order, pseudo-second order and Elovich model can not identify the influence of diffusion on adsorption. So, Weber and Morris’ equation (intra-particle diffusion) and Bangham’s model (pore diffusion) were further analyzed. Rate of adsorption is frequently used to analyze nature of the ‘rate-controlling step’ and the use of the intra-particle diffusion model has been greatly explored. It was evident from the graph of Fig. 6.8 (d) that the first linear portion (Stage I) was attributed to the immediate utilization of the most readily available adsorbing sites on the adsorbent surfaces. The second plateau path (Stage II) indicates very slow diffusion of adsorbate from surface site into the inner pores [230]. Thus, the rapid initial portion of Pb (II) adsorption may be governed by the surface diffusion

process and later part is controlled by pore diffusion. The correlation coefficient (R2) obtained from Weber and Morris’ equation was found to be 0.97, which is still higher value and also the intercept of the line fails to pass through the origin due to difference in the rate of mass transfer in the initial and final stages of adsorption and indicates some degree of boundary layer control which implies that intra-particle diffusion is not only the rate controlling step [230, 231]. Further, the experimental data were used to confirm the pore diffusion as one of the rate-controlling steps using Bangham’s equation. The graph obtained [Fig. 6.8 (e)] using Bangham’s equation was found to be linear with a quite good correlation coefficient R2 = 0.958 indicating that the contribution of pore diffusion to the overall mechanism of Pb(II) adsorption could not be neglected and may play a role in controlling the rate of adsorption. The adsorption kinetics was pore diffusion controlled and the diffusion into the pores of the adsorbent was not the sole rate-determining process. Table 6.1: Kinetic parameters of Pb (II) using porous zirconium oxide as adsorbent

Kinetic parameters

qe, exp (mg/g) 9.888

Pseudo-first order model qe, cal (mg/g) 6.605 -1 K1 (min ) 0.1523 R2 0.9747 Rate equation y = -0.06617 x + 0.81993

Pseudo-second order model qe, cal (mg/g) 10.07 -1 -1 K2 (g mg min ) 0.0608 R2 0.9995 Rate equation y = 0.0993 x + 0.16197

-1 -1 Elovich model ae (mg g min ) 9.319 -1 be (g mg ) 0.4308 R2 0.963 Rate Equation y = 2.32126 x + 3.22667

-1 -0.5 Inter-particle diffusion model Ki (mg g min ) 1.91154 C 2.39608 R2 0.97003 Rate equation y = 1.91154 x + 2.39608

-1 -1 Bangham’s model Kb (mL g L ) 122.26  0.62723 R2 0.958 Rate equation y = 0.62723 x -0.7250

6.2.6: Regeneration of the adsorbent When an adsorbent is applied for adsorption of toxic ions, the possibility of regeneration of the adsorbent is of great importance from application point of view. Efficient removal of loaded metal from the adsorbent was necessary to ensure their long term use for repeated adsorption-desorption cycles. The percentage removal efficiency during adsorption and recovery percentage of Pb(II) during desorption for each cycle was determined and shown in Fig. 6.9. This result indicates that the porous zirconium oxide could be employed on several times for adsorption process without significant losses of its initial capacity of adsorption. It was observed that the removal percentage of Pb(II) was slightly decreases from ~ 99 % to ~93 % from the first to the fifth cycle. This may be due to the non-leaching of previously adsorbed Pb(II) ions that resisted to the desorption process. So, the efficient reuse of the Pb(II) loaded zirconium oxide sample was found to be possible and can be applied to the removal of toxic metals from wastewater efficiently.

Adsorption Desorption 100

(%) efficiency 60

Recovery 20

Removal/ 0 1 2 3 4 5 Number of cycles

Fig. 6.9: Removal (during adsorption) and recovery (during desorption) percentage of Pb(II) as a function of number of cycles.

6.3 Remarks Zirconia powders prepared through constant pH route show highest activation energy of crystallization (Ea = 260 kJ/mol) or higher exothermic peak temperature (717 °C), when compared with gelation or precipitation route due to its controlled growth of smaller crystallites. The released huge amount of H2 gas bubbles during borohydride synthesis via

constant pH route play a major role for formation of loose smaller crystallites and thus enhances the activation energy of crystallization of pure zirconia. So, the as-prepared zirconia powders prepared through constant pH route remain amorphous up to 600°C and pure t-ZrO2 (~20 nm) was stable up to 800°C. Additionally, the evolution of in-situ H2 gas- bubbles creates numerous gas–liquid interface aggregation centres during borohydride synthesis and plays an important role to develop loose zirconium hydroxide powders in the as-synthesized condition. The trapped pore and or large voids are observed, when the loose nature of as-synthesized zirconium hydroxide undergoes calcination process and the porous structure of zirconium oxide was stable up to 800 °C. Without any surface modification the porous zirconium oxide powders were able to adsorb almost all Pb(II) ions within few minutes at normal pH condition. The superior fitting of pseudo-second order as well as quite good fitting of Elovich model indicates that the main adsorption controlling mechanism was chemisorption. Based on intra-particle diffusion and Bangham’s model, it was further suggests that the quick adsorption was attributed to the immediate utilization of the most readily available adsorbing sites on the adsorbent surfaces and the adsorption kinetics was limited by pore diffusion. Further, quite high efficiency of five cycles of regeneration suggest the importance of porous zirconium oxide for the removal of toxic ions for environmental application. Further, phase formation, pore geometry and pore hysteresis were studied at different calcination temperatures for constant pH derived zirconium hydroxide samples, prepared via borohydride method in the following chapter.

Chapter 7 Temperature-mediated phase transformation, pore geometry and pore hysteresis transformation of borohydride derived in-born porous zirconium hydroxide nanopowders

In this chapter, a constant pH method using NaBH4 was employed for the development of loose and porous zirconium hydroxide powders in the as-synthesized condition. A novel information on the temperature-mediated phase transformation, pore geometry as well as pore hysteresis transformation of in-born porous zirconium hydroxide nanopowders with the help of X-ray diffraction (XRD), Brunauer–Emmett–Teller (BET) isotherm and Transmission Electron Microscopy (TEM) images are presented in this chapter.

7.1: Experimental

Borohydride route via constant pH method was performed using aqueous Zr-salt

(ZrOCl2·8H2O) and aqueous NaBH4. First, aqueous solutions of 0.5 M ZrOCl2·8H2O (pH

~0.1) and 0.5 M NaBH4 (pH~11) were prepared separately. While adding the acidic Zr-salt solution to a beaker containing of basic aqueous NaBH4, the pH of the solution decreases, so to maintain a constant pH environment of around 10.5, additional aqueous NaBH4 was added simultaneously along with Zr-salt solution. White precipitates were formed in the aqueous NaBH4 solution during the addition of both solution of Zr-salt and NaBH4 to the aqueous NaBH4. After completion of the reaction, the precipitate powders were filtered and thoroughly washed, dried, and calcined at different temperatures (500, 600°C, 800 °C, 900 °C and 1000°C) for 1 h. Phase analysis was performed using X-ray diffraction (XRD). Particle morphology was studied using Transmission Electron Microscopy (TEM). Pore geometry and pore hysteresis was studied using BET-isotherm.

7.2: Results and discussion 7.2.1: Structure Figure 7.1 shows X-ray diffraction patterns of borohydride derived as-syntesized and calcined powders. The amorphous nature of borohydride derived as-synthesized powders remain amorphous up to 500°C and further it slowly converts to crystalline in nature at 600 °C. All the peaks of powder calcined at 600 °C are identified and indexed with tetragonal form of zirconium oxide (t-zirconium oxide), as per JCPDS file no: 79-1768. Further, the tetragonal nature of zirconium oxide was found to be stable up to 800 °C. A mixture phases of t-zirconium oxide (78 vol %) and monoclinic (m) phase [peaks are indexed as per JCPDS file no: 83-0943] of zirconium oxide (22 vol%) were present in the sample, calcined at 900 °C. At further higher calcination temperature of 1000 °C, major phase of m-zirconium oxide (98 vol %) along with minute amount of t-zirconium oxide (2 vol %) were observed. The growth of crystallite size of zirconium oxide was strongly affected by the thermal treatment and also helps for stabilizing t-zirconium oxide up to 800 °C. Further, the crystallite size and volume percentage of t-zirconium oxide as a function of calcination temperature was represented in Fig. 7.1 (b). It was found that the crystallite size increases with calcination temperature, but at a slower rate and thus it may help to stable the t-zirconium oxide up to 800 °C. Further, the computed values of the crystallite size of t- zirconium oxide powders calcined at 600°C to 1000°C enabled the calculation of the activation energy of particle growth (Q, kJ/mol) of t-zirconium oxide phase using Arrhenius (-Q/RT) equation: Dt = Ae , where Dt denote the crystallite size (nm) calcined at temperature T, A is the frequency factor of Arrhenius equation, and Q is the activation energy of particle growth, R is the gas constant (J/mol.K) and T is the calcination temperature (K) [171]. To determine the activation energy (Q) of particle growth, Figure 7.1 (b) was re-plotted to Figure 7.1 (c), assuming that crystallite growth in nanopowders of zirconium oxide, being a thermally activated process and is dependent on the calcination temperature. The activation energy of particle growth of borohydride synthesized zirconium oxide was found to be ~ 36.8 kJ/mol, which is much lower than that of yttria-stabilized zirconium oxide, but similar to that for pure nano zirconium oxide [232]. The slow growth of the particles during calcination process led to stable the amorphous nature up to 500 °C and also stables the t- zirconium oxide up to 800 °C. Additionally, the nature of powder morphology of as- synthesized sample may also play a role for stabilizing the amorphous phase up to 500 °C and t-zirconium oxide up to 800 °C. So, powder morphology was performed using TEM.

m (a) m 1000 °C

m mm mm t m m mm mm m m m m m

m t m 900 °C mtm m tm m

m m mm m mm m m m m t 800 °C t

t t t t t t t t t t 600 °C

(a.u.) Intensity 500 °C

as-synthesized

20 30 40 50 60 70 80 2 (in degree)

35 (b) 100 3.6 Q/R= 4427 (c) 30 80 3.3 25 3.0 20 60

)

D 2.7 15 (

40 ln 2.4 10

Crystallite size (nm)

20 Tetragonal phase (%) 5 2.1

0 0 1.8 600 700 800 900 1000 0.0008 0.0009 0.0010 0.0011 0.0012 Calcination temperature (°C) 1/T (K-1) Fig. 7.1: X-ray diffraction pattern of the as-synthesized as well as calcined samples (a) [Note: ‘t’ stands for tetragonal and ‘m’ stands for monoclinic zirconium oxide] and crystallite size and volume percentage of tetragonal phase as a function of calcination temperature (b). To determine the activation energy (Q) of particle growth, Figure 7.1 (b) was re-plotted to Figure 7.1 (c).

7.2.2: Powder morphology and BET-isotherm The as-synthesized powders are found to be amorphous in nature, as it shows hazy rings of electron diffraction pattern [inset of Fig. 7.2(a)]. The amorphous powders were found to be loose nature, as observed from TEM micrograph of Fig. 7.2(a) and it indicates that the individual fine particles are well-separated by voids/pores. It was also observed that the nature of pores in the as-synthesized condition seems to be interconnected. To further justify the pore geometry, BET adsorption-desorption isotherm was performed on the as- synthesized powders and was shown in Fig. 7.2 (b). The existence of wide broad hysteresis

loop in BET-isotherm indicates that the in-born as-synthesized powders are porous in nature and was well correlated with TEM micrograph [Fig. 7.2 (a)]. Further, the type of pore was analyzed from the BET hysteresis loop. The wide hysteresis loop also indicate a delay in both condensation and evaporation process [233].

120 (b)

Volume adsorbed (cc/g) adsorbed Volume

0 0.0 0.2 0.4 0.6 0.8 1.0 Relative pressure (P/P ) 0.24 0 (c)

3.6 4.6 0.24 3.6 4.6

0.16 0.16

0.08 6.1

Volume (cc/g) Volume 0.08 15.7 6.1 (cc/g) Volume 0.00 0 5 10 15 20 25 30 Pore diameter (nm) 15.7 0.00 0 25 50 75 100 125 150 175 200

Pore diameter (nm)

Fig. 7.2: TEM micrograph (a) of as-synthesized porous zirconium hydroxide and the corresponding electron diffraction pattern in the inset of (a) indicates that the powder was amorphous in nature. BET-isotherm (b) and pore size distribution (c) of the as-synthesized porous zirconium hydroxide. Inset of (c) is the enlarge view of the pore size distribution.

In this BET-isotherm, the adsorption follows a slow increase in adsorbed volume with increases in partial pressure without any saturation point, but with a slight slope change at P/P0 at 0.6. However, desorption curve follows a different path forming a wide hysteresis loop of type H2 (according to IUPAC classification) with a slight slope change at similar

P/P0 of 0.6 and closes at the starting point of adsorption. This H2 type hysteresis loop indicates the existence of ink bottle-neck type pores in the gas-bubbles derived as- synthesized zirconium hydroxide powders [234, 235]. Further, due to non-saturating behavior of adsorption curve and delay in desorption curve along with the closure point at P/P0 of 0.1 indicates that these ink bottle-neck pores are well interconnected in nature [235-237]. Analyzing the hysteresis loop by considering the ink bottle-neck type pores, it was further suggests that the first half of the BET adsorption isotherm (up to P/P0 at 0.6) represents the condensation of surface necks and second half of this adsorption isotherm associated with continuous condensation of the inner interconnected bottles through narrow necks without [236] a saturation point . During desorption process, first half (up to P/P0 at 0.6) represents the evaporation from surface necks of all bottles and second half of desorption closes at P/P0 ~ 0.1 is thought to depend not only on size of the bottles, but also on the connectivity of the pore network [237]. The interconnected network nature of ink bottle-neck pores gives rise to pore-blocking effects, which also delays both adsorption and desorption mechanism [237]. The pore size distribution of as-synthesized porous zirconium oxide was also calculated from BJH desorption curve and was shown in Fig. 7.2 (c). The highest pore volume was found to be ~0.21 cc/g for the pore diameter of 3.6 nm to 4.6 nm. However, the pore diameter of as-synthesized zirconium hydroxide varies from 3.6 nm to 15.7 nm (see enlarge view of pore size distribution in the inset of Fig. 7.2 (c)]. So, it seems that the pore size of as- synthesized zirconium hydroxide was fall in the meso-range (2 to 50 nm), but may also contain some amount of micro pores (< 2 nm), as the volume content of minimum pore size (3.6 nm) was same as that of 4.6 nm [238]. In addition, the surface area of as-synthesized zirconium hydroxide powders was found to be 182 m2/g. The pore geometry of the in-born interconnected loose porous nanoparticles may strongly modify the pore morphology of powders during calcination process. So, the loose porous as-synthesized powders were further calcined at different temperatures. Fig. 7.3 (a) shows TEM micrograph of porous as-synthesized powders calcined at 500 °C. Up to this temperature, the amorphous nature still remains same, as confirmed from the hazy rings of electron diffraction pattern, indicated in the inset of Fig. 7.3 (a). TEM micrograph of Fig. 7.3 (a) clearly indicates that the powders calcined at 500 °C are still porous in nature. The

presence of intra-particle voids may inhibit the mass transfer between loose nanoparticles and also help to restrict the coarsening of particles, during calcination up to 500 °C.

120 (b)

(cc/g) adsorbed volume 0 0.0 0.2 0.4 0.6 0.8 1.0 Relative pressure (P/P0)

0.4 (c) 4.6 0.4 4.6 0.3 0.3 3.6 0.2

3.6 0.2

(cc/g) Volume 0.1 6.1 8.8

(cc/g) Volume 0.1 0.0 6.1 0 5 10 15 20 25 30 Pore diameter (nm)

8.8 0.0 0 25 50 75 100 125 150 175 200 Pore diameter (nm) Fig. 7.3: TEM micrograph (a) of the as-synthesized powders calcined at 500 °C and the corresponding electron diffraction pattern in the inset of (a) indicates that the powder was amorphous in nature. BET-isotherm (b) and pore size distribution (c) of the calcined (500°C) porous powders. Inset of (c) is the enlarge view of the pore size distribution.

So, formation of smaller sized particles with interconnected voids along with lower activation energy of particle growth during calcination process led to sustain the amorphous phase of up to 500 °C. BET adsorption-desorption isotherm was performed on this calcined powders and was shown in Fig. 7.3 (b). The nature of BET-isotherm was found to be same of H2 type as that of as-synthesized powders. In this isotherm, the high pressure unsaturated adsorption and a comparable change in the slope at P/P0 = 0.6 in desorption curve with a slightly higher closure point at P/P0 = 0.2 (as compared to desorption behavior of as- synthesized sample) reflects that the pores are well interconnected, but the inter-connection are blocked due to neck closing of a very few bottle-neck pores during calcination process and thus leads to easy evaporation of condensed gas during desorption process [239]. The pore size distribution was evaluated from BJH desorption curve and was shown in Fig. 7.3 (c). The pore size distribution of this sample was becoming narrow [see enlarge view of pore size distribution in the inset of Fig. 7.3 (c)] as compared to pore size distribution of as- synthesized sample, which further suggests that this calcined powders was purely mesoporous in nature, as the pore size varies from 3.6 nm to 8.8 nm with as maxima of 4.6 nm. In addition, the pore volume was found to ~ 0.36 g/cc, which is slightly higher than the pore volume of as-synthesized sample of same particular pore size of 4.6 nm. In the case of as-synthesized sample, the lower volume containing pores having larger diameter in the range between 6.1 nm to 15.7 nm indicates that some of the pores may connect with each other, but without a neck similar to dumb-bell shape. Heating at 500 °C, the size of lower volume containing dumb-bell shaped pores varies in between 6.1 nm to 8.8 nm. The decrease of dumb-bell shaped pores size from 15.7 nm to 8.8 nm at 500 °C was mainly due to the formation of additional bottle-neck pore from dumb-bell shaped pores via coarsening during calcination process. So, the increase in pore volume for particular pore diameter of 4.6 nm was due to conversion of interconnected dumb-bell shaped pores to bottle-neck pores. Additionally, the surface area of the powder calcined at 500 °C was decreased to 160 m2/g. Fig. 7.4 (a) shows TEM micrograph of the powders calcined at 600 °C and the electron diffraction pattern was indicated in the inset of Fig. 7.4 (a). Hazy dotted ring pattern of electron diffraction indicates that the amorphous nature was not fully converted to crystalline nature of t-zirconium oxide at this temperature. While increase in calcination temperature, the seed crystals start to develop within the loose amorphous matrix and these nanocrystals grow as well as traps the inter-particulate voids and form a trapped pore within the particle. So, pores are entrapped within individual zirconium oxide particles and

developed porous structure. The pore morphology of this crystalline t-zirconium oxide was analyzed using BET-isotherm and was shown in Fig. 7.4 (b).

40 (b)

(cc/g) adsorbed volume 0 0.0 0.2 0.4 0.6 0.8 1.0

Relative pressure (P/P0)

4.6 (c) 0.20 4.6 0.20 0.15 0.15

0.10

0.10 Volume (cc/g) 0.05 3.6 6.1 0.00 (cc/g) Volume 0 5 10 15 20 25 30 0.05 Pore diameter (nm)

3.6

6.1 0.00 0 25 50 75 100 125 150 175 200 Pore diameter (nm) Fig. 7.4: TEM micrograph (a) of the as-synthesized powders calcined at 600 °C and the corresponding electron diffraction pattern in the inset of (a) indicates that the powder was in crystalline in nature. BET-isotherm (b) and pore size distribution (c) of the calcined (600°C) porous zirconium oxide. Inset of (c) is the enlarge view of the pore size distribution.

In this BET-isotherm, adsorption curve increases with increase in partial pressure with non-saturating behavior and forming a hysteresis loop with the help of desorption curve. This desorption curve suddenly drops at P/P0 = 0.6 and matches with the adsorption curve at P/P0 = 0.5. The curvature of BET hysteresis loop is typically H2 type indicating ink bottle-neck type pores and some pores are interconnected due to non-saturating behavior of adsorption curve. The pore size distribution was evaluated from BJH desorption curve and was shown in Fig. 7.4 (c). The pore size distribution was becoming still narrow [see enlarge view of pore size distribution in the inset of Fig. 7.4 (c)] as compared to pore size distribution of 500 °C heated sample, but the pore size varies from 3.6 nm to 6.1 nm with as maxima of 4.6 nm. It also further indicates that with increase in calcination temperature from 500 °C to 600 °C, the mesoporosity remain same in porous t-zirconium oxide. But, the pore volume of the particular diameter of 4.6 nm decreases to 0.20 g/cc at 600 °C and this may be due to coarsening of nanoporous particles. The decrease in pore volume with narrow size distribution (as compared to BET-isotherm of 500 °C heated sample) leads easy evaporation of condensed gas and thus shifts the closure point of desorption branch at P/P0 = 0.5. The surface area of crystalline t-zirconium oxide further decreases to 55 m2/g at 600 °C. Further, TEM was performed on 800 °C heated sample in order to study the particle morphology and pore geometry of t-zirconium oxide. Fig. 7.5 (a) shows TEM micrograph of calcined (800 °C) t-zirconium oxide and the electron diffraction pattern was indicated in the inset of Fig. 7.5 (a). Sharp dotted ring pattern of electron diffraction indicates that the zirconium oxide was purely crystalline in nature. At this temperature, particle shape is nearly spherical to polyhedral in shape. The sizes of these particles are in the range between ~20 to ~40 nm. In each particle, the fine pores are well distributed as well as the surface of the particle seems to be in different morphology, as observed from TEM micrograph in Fig. 7.5 (a). In addition, the coarsening of fine particles along with coalescence of existing pores take place at higher temperature and develop a stable phase of porous t-zirconium oxide at 800 °C. To justify the pore as well as surface morphology of t-zirconium oxide, a higher magnified TEM was analyzed and shown in Fig. 7.5 (b). It was confirmed that the zirconium oxide particle consists of two different types of pore geometry such as nearly spherical and lamellar type [marked as arrow in Fig. 7.5 (b)]. The size of spherical pores varies from ~ 3 nm to ~ 12 nm and the thickness of lamellar type pores was ~ 3 nm. It was also observed that a thin disordered layer was covering each particles and this layer was visualized in a clear way, as shown in the inset of Fig. 7.5 (b). This thin layer consists of ultra-fine loose particles [marked as arrow in the inset of Fig. 7.5 (b)] with a thickness of around 2 to 4 nm.

The pore morphology of calcined (800 °C) t-zirconium oxide was also analyzed using BET- isotherm and was shown in Fig. 7.5 (c). The presence of hysteresis loop indicates that the t- zirconium oxide particles retain their porous nature up to 800 °C. However, the behavior of hysteresis loop was quite different from the hysteresis behavior of lower temperature calcined samples. This hysteresis loop is typically consists of mixture of ink bottle-neck and slit type pores (H2 + H3 type).

40 (c) 0.14 (d) 4.6 0.14 0.12 4.6 30 0.12 0.10 0.10 0.08

20 0.08 0.06

0.04

Volume (cc/g) Volume 15.5 0.06 0.02

3.6 0.00 10 0.04 0 5 10 15 20 25 30 (cc/g) Volume 15.5 Pore diameter (nm) 0.02

volume adsorbed (cc/g) adsorbed volume 3.6 150 0 0.00 0.0 0.2 0.4 0.6 0.8 1.0 0 25 50 75 100 125 150 175 200

Relative pressure (P/P0) Pore diameter (nm)

Fig. 7.5: TEM micrograph (a) of the as-synthesized powders calcined at 800 °C and the corresponding electron diffraction pattern in the inset of (a) indicates that the powder was in crystalline in nature. Higher magnified TEM image (b) indicates the presence of two types (spherical as well as lamellar) of pore. Inset of (b) indicates the presence of ultra-fine loose particles on the surface of the particle. BET-isotherm (c) and pore size distribution (d) of calcined (800°C) zirconium oxide. Inset of (d) is the enlarge view of the pore size distribution.

At this temperature, the pore hysteresis transformed from H2 type to mixed H2 and H3 type and was also well-correlated with TEM micrographs. It was well understood that the coarsening of porous particles having lower diameter (~ 3.6 nm) bottle-neck pores is faster than the higher diameter bottle-neck pores at 800 °C. At the same time, the entrapped air in pores expands at this temperature and migrates towards the surface with high pressure without closing the interconnected paths and thus forming a thin layer of ultra-fine loose particles on the surface of t-zirconium oxide. As the particles are ultra-fine in nature at the surface of the zirconia particles (see Fig. 7.5 b), there may be some amorphous zirconia is still present at the surface of the zirconia particles. However, the degree of crystallinity of the materials at 800 °C is too difficult to quantify with the existing facilities. Further, the lower diameter pore diffuses and matches with the interconnected paths and thus creating lamellar type, whereas the higher diameter pore remains as bottle-neck type pore. The pore size distribution was evaluated from BJH desorption curve and was shown in Fig. 7.5 (d). From this figures, it was found that the pore size varies from 3.6 nm to 150 nm with a maximum of 4.6 nm with wide pore size distribution [see enlarge view of pore size distribution in the inset of Fig. 7.5 (d)]. In this t-zirconium oxide sample, the pores are mostly mesoporous in nature along with some lower amount of macro porous in the range between 50 nm to 150 nm. Additionally, the wider distribution was mainly due to the additional contribution of lamellar type pores. High temperature coarsening of particles leads to decrease the bottle-neck pore volume of the particular diameter of 4.6 nm to 0.129 g/cc. In addition, the surface area of porous zirconium oxide, calcined at 800 °C was found to be ~ 29 m2/g. At still higher temperature of 900 °C, the particle morphology of zirconium oxide was studied using TEM micrograph and was shown in Fig. 7.6 (a). The powder morphology was seem to be porous in nature, however, from inset of Fig. 7.6 (a), it was observed that the calcined (900 °C) zirconium oxide contain lamellar type pores. BET-isotherm in Fig. 7.6 (b) also confirms that the pore morphology was typically lamellar of H3 type. So, at 900 °C, remaining bottle-neck pores diffuse with the interconnected paths and forming lamellar type. The average pore diameter decreases from 4.6 nm to 3.6 nm at 900 °C and pore volume also decrease to 0.005 cc/g. The surface area of zirconium oxide calcined at 900 °C was found to be 6 m2/g. However, at higher temperature of 1000 °C, all the pores were diffused completely and forming a fully non-porous spherical and polyhedral particles having size ranges from 150 nm to 400 nm, as confirmed from TEM micrograph of Fig. 7.6 (c).

5.0 0.006 (b) 0.005 4.5 0.004

4.0 0.003 0.002

Volume (cc/g) Volume 3.5 0.001 0.000 0 20 40 60 80 100 120 140 160

3.0 Pore diameter (nm) 2.5 2.0

Volume adsorbed (cc/g) adsorbed Volume 1.5

0.0 0.2 0.4 0.6 0.8 1.0

Relative pressure (P/P0)

Fig. 7.6: TEM micrograph (a) of the as-synthesized powders calcined at 900 °C. BET- isotherm (b) and inset of (b) shows the pore size distribution. TEM micrograph (c) indicates the non-porous zirconium oxide, calcined at 1000 °C.

Based on the XRD, BET-isotherm and TEM image analysis, the H2 gas-bubbles assisted borohydride route was found to be a potential synthesis method for development of thermally stable porous nanopowders of zirconium hydroxide or oxide having adequate pore size, pore volume and surface area within a temperature range from 500 °C to 800 °C. More importantly, these porous nanopowders can be used efficiently in different areas of applications. So, in this context, it is justified to compare the textural properties of borohydride derived as-synthesized as well as calcined porous nanopowders with the available reported literatures. Further, Table 7.1 summarizes the textural properties of the borohydride derived as-synthesized as well as calcined porous nanopowders. Moreover, the phase stability, pore size, pore size distribution and surface area are the important parameters to find out the potential use of these zirconium hydroxide or oxide nanopowders in various fields such as adsorption of heavy metal ions, catalytic reaction, oxygen sensors, storage of gases and luminescent applications [1, 240-243]. In view of the phase stability of t- zirconium oxide up to moderate temperature, many researchers have been successfully synthesized t-zirconium oxide via various synthesis methods using surfactants or additives [244-246]. However, in most of the cases, the retention of porous nature in t-zirconium oxide was limited in the range between 300°C to 600 °C and which further led to prevent the use of these porous powders for specific applications [135, 138]. In this perspective, without the use of any additives or surfactants, the borohydride route is found to be much more advantageous because of the development of in-born porous nature of amorphous zirconium hydroxide in the as-synthesized condition, stable the amorphous as well as porous structure nature up to 500 °C as well as stable t-zirconium oxide nanopowders with porous structure up to 800 °C. Further, the surface area, mean pore diameter and pore volume of borohydride derived powders calcined at 500 °C was comparable with the calcined (450 °C) amorphous zirconium oxide powders prepared by Cui et al.[247], which was having a surface area of 161.8 m2/g, mean pore diameter ~9 nm and pore volume of 0.43 cc/g. It was further suggested that this type of large surface area, high pore volume and loose porous structure may be a strong candidate for heavy metal adsorption for environmental application [158]. Similarly, Kuai el. al. [95] have developed disordered pores of various transition metal oxides by surfactant assisted aerosol spray method and it was observed that the amorphous nature of zirconium oxide sample heated at 400 °C having a surface area of 116 m2/g and a pore volume of 0.16 cc/g. This type of large surface area and pore volume can be widely used in various fields of applications [241]. In addition, the obtained surface area of porous amorphous powders derived via borohydride route was found to be quite comparable with the literatures based on the synthesis of highly ordered porous powders developed with addition of additives or surfactants [134, 138]. Also, the pore volume and pore diameter of the

borohydride derived porous t-zirconium oxide (calcined at 600 °C) can be well comparable with the pore volume and pore diameter of ordered porous t-zirconium oxide prepared using surfactants [248]. Further, the surface area, mean pore diameter and pore volume of 600 °C calcined t-zirconium oxide can also be well comparable with the disordered structure of porous zirconium oxide prepared using surfactants [133]. The calcined (500 °C) t-zirconium oxide prepared from the thermal decomposition of metal– organic frameworks by Yan et al. [98] shows a pore diameter of 5-8 nm with pore volume of 0.208 cc/g. Similarly, Chen et al. [249] have developed t-zirconium oxide at 500 °C which have a lower surface area of 95 m2/g and total pore volume as high as 0.19 m2/g. In addition, Mokhtar et al. [171] have prepared t- zirconium oxide (calcined at 600 °C) with using surfactant and achieve a lower surface area of 41 m2/g and pore volume of 0.3315 cc/g. Further, the retention of porous nature of t- zirconium oxide up to 800 °C was found to be difficult, while synthesized using surfactant- assisted synthesis routes [47, 133 250]. Based on the above discussion, the borohydride derived porous nature of amorphous as well as crystalline t-zirconium oxide nanopowders can efficiently be used for different applications.

Table 7.1: Textural properties of borohydride derived porous nanopowders.

Sample Phase Crystall Particle Pore morphology Pore Surface Range of Maximum condition ite size* size hysteresi area pore size Pore volume (nm) range s (m2/g) /Mean at mean (XRD) (nm) pore size pore size (TEM) (cc/g) (in nm)

As- A - - Interconnected ink H2 type 182 3.6 to 15.7 0.21 synthesiz bottle-neck (major) + / ed dumb-bell shaped pores 3.6-4.6 (minor) Calcined A - - Interconnected ink H2 type 160 3.6 to 8.8 0.36 at 500 °C bottle-neck + neck / closing of a very few 4.6 bottle-neck pores Calcined t 7.5 10-20 typically ink bottle-neck Typically 55 3.6 to 6.1 0.20 at 600 °C type pores H2 type / 4.6 Calcined t 18.8 20-40 mixture of ink bottle- H2 + H3 29 3.6 to 150 0.129 at 800 °C neck and slit type pores type / 4.6 Calcined t + m 28.9 30-50 typically lamellar or slit typical 6 3.6 to 150 0.005 at 900 °C type pores H3 type / 3.6 Calcined t + m 33.8 150------at 1000 400 °C

Note: A: amorphous; t: tetragonal phase of zirconium oxide; m: monoclinic phase of zirconium oxide * Crystallite size of t- zirconium oxide

7.3: Remarks Interconnected porous structure with loose amorphous nature of zirconium hydroxide was successfully prepared through constant pH method using sodium borohydride. The presence of voids or pores in the as synthesized powders as well as the slower growth of particles during calcination process led to sustain its amorphous nature as well as loose porous structure up to 500 °C. Temperature mediated phase transformation from amorphous to crystalline nature of t-zirconium oxide took place at 600 °C and the t- zirconium oxide with porous structure was found to be stable up to 800 °C. In addition, the pore morphology of calcined (600 °C) t-zirconium oxide was found to be typically ink bottle-neck (H2-type) pores and it transformed to mixed ink bottle-neck and slit (H3 type) pores at 800 °C. Also, the mean pore diameter of 4.6 nm remains unchanged within a temperature range from 500 °C to 800 °C. Further, the t-phase of zirconium oxide transformed to mixed t- and m- phase with only slit type pores at 900 °C. The adequate surface area and pore volume along with nearly constant mean pore diameter of porous amorphous as well as crystalline t-zirconium oxide, developed through borohydride route may be suitable for different applications including adsorption of heavy metal ions for environmental application, shape-selective heterogeneous catalysis, ion exchange and proton conduction as well as gas-sensing application. Further, phase formation, phase stability, powder morphology and pore size was studied for the rare earth based porous zirconia samples, prepared using borohydride synthesis. As the incorporation of rare earth in porous zirconia may lead to develop a phosphor material. So, photoluminescence behaviour and powder colour under UV light source was also studied in detail.

Chapter 8 Phase formation, pore morphology and photoluminescence behaviour of rare earth (single-doped of Eu3+ or Tb3+ and co-doped of Eu3+ and Tb3+) based porous zirconia nanopowders

The photoluminescence behaviour of un-doped porous zirconia nanopowders was discussed in the first section of this chapter. Phase formation, phase stability, powder or pore morphology, elemental distribution of rare earth ions, photoluminescence and sample colour under UV light for the rare earth based [single-doped (Eu3+ or Tb3+) and co-doped (Eu3+ and Tb3+)] porous zirconia nanopowders were studied and discussed in the second section of this chapter. 8.1 Experimental

Borohydride route via constant pH method was performed to prepare both un-doped [the procedure was followed as discussed in experimental section of the chapter 7] and rare earth doped porous zirconia. An aqueous solutions of ZrOCl2·8H2O, europium chloride/terbium chloride [5 mol % europium oxide / 5 mol % terbium oxide mixed with an appropriate amount of hydrochloric acid (HCl) to form europium/terbum chloride] and 3+ 3+ NaBH4 were prepared separately. However, for developing doped (Eu or Tb ) and co- doped (2 mol % Eu3+ and 5 mol % Tb3+) porous zirconia nanopowders, the mixture solution of Zr-salt and Eu and/or Tb-salt was added to a beaker containing of basic aqueous NaBH4. As the above solution was added, then the pH of the solution decreases, so to maintain a constant pH environment of around 10.5, additional 0.5 M NaBH4 aqueous solution was added simultaneously along with mixture solution of Zr-salt and Eu and/or Tb -salt. White precipitates were formed during reaction. After completion of the reaction, the precipitate powders were filtered and thoroughly washed, dried, and calcined at different temperatures for 1 h. Further, phase stability, powder or pore morphology, EDX line profile, photoluminescence and sample colour under UV light was studied in detail.

8.2: Results and discussion 8.2.1: Photoluminescence behaviour of un-doped porous zirconia Based on the literature study, it was found that photoluminescence properties of un- doped zirconia was due to several factors including smaller size, different morphologies as well as surface defects [168-171]. In this research work, the borohydride derived un-doped zirconia powders was found to be smaller in size with porous nature as well as a thin layer of ultra-fine loose particles were observed on the surface of the zirconia particles. So, the porous un-doped zirconia may show luminescence, when excited under UV. In addition, the photoluminescence properties of un-doped zirconia may arise due to the movement and diffusion of atoms and a change of the electronic structure, especially band gap [172-175]. So, in the present work, the band gap of as-synthesized and calcined (600°C and 800°C) t-ZrO2 was determined from UV-vis diffuse reflectance data by transforming it into a function of reflectance (R) as proposed by Kubelka-Munk [251]. The Kubelka-Munk plot for determining band gap energy was shown is Fig. 8.1 for the borohydride derived zirconia. KM plot was plotted with the nth power of product of function of reflectance F(R) [= (1-R)2/(2R)] and photonic energy (Eg = hυ) against the photonic energy.

45 as-prepared, Band gap = 5.55 40 600 °C, Band gap = 5.75 800 °C, Band gap = 5.65

] 35 1/2

[ 30



 25

F(R)*h

[ 15

10 5.0 5.2 5.4 5.6 5.8 6.0 Energy (eV)

Fig. 8.1: Kubelka-Munk plot of borohydride derived zirconia.

Considering as a direct band gap of zirconia, the value of n is taken as 2 for allowed transitions. The energy band gap was found out by extrapolating the linear portion of the graph to the X-axis. The band gap of as-prepared zirconia was found to be 5.55 eV ( ~223

nm), which was well matched in the range within the experimental band gap of amorphous [177-179] zirconia . However, t-ZrO2 calcined at 600 °C and 800 °C was found to be 5.75 eV

( ~ 215 nm) and 5.65 eV ( ~219 nm), respectively. The decrease of band gap of t-ZrO2 from 600 °C to 800 °C is due to increase in particle size [252]. However, the obtained results are well matched in the range within the experimental band gap of t-ZrO2. As studied from the literature, it was well accepted that the photoluminescence and the band gap energy are well associated with each other [180]. The photoluminescence band position was strongly depends on the dopants (mostly rare earth (RE) ions elements), native defects, excitonic wavelengths [182, 183]. In this context, photoluminescence property of borohydride synthesized amorphous and calcined t-zirconia nanopowders was studied by fluorescence spectroscopy. Further, the excitation spectrum was performed on the as- synthesized as well as calcined t-ZrO2 powders. The typical excitation spectrum, performed at an emission wavelength of 420 nm was shown in Fig. 8.2. Based on this excitation spectra, the photoluminescence emission spectra of borohydride derived zirconia was analysed at an excitation wavelength of near vacuum ultra-violet (VUV is characterized with wavelength shorter than 200 nm) region of 205 nm and 210 nm as well as ultra-violet region of 220 nm, 240 nm and 300 nm.

em = 420 nm

(a.u.)Intensity

200 225 250 275 300 Wavelength (nm) Fig. 8.2: Typical excitation spectra versus wavelength of borohydride synthesized un- doped zirconia.

The typical PL spectra of borohydride synthesized amorphous and crystalline zirconia nanopowders, recorded using excitation wavelengths of 205 nm, 210 nm, 220 nm, 240 nm and 300 nm are shown in Fig. 8.3.

8 1.4x108 1.4x10  = 210 nm as-prepared ex = 205 nm as-prepared ex 8 8 1.2x10 600 °C 1.2x10 600 °C 800 °C 800 °C 8 1.0x10 8 1.0x10

7 8.0x10 7 8.0x10

7 6.0x10 7 6.0x10

7 4.0x107 4.0x10

Intensity (a.u.)Intensity (a.u.)Intensity 7 2.0x10 7 2.0x10 0.0 0.0 200 250 300 350 400 450 500 550 600 650 700 200 300 400 500 600 700 Wavelength (nm) Wavelength (nm)

8x106 1.2x10 8 ex = 220 nm as-prepared ex = 240 nm 6 as-prepared 600 °C 7x10 800 °C 600 °C 8 1.0x10 6 6x10 800 °C

7 8.0x10 5x106

6 7 4x10 6.0x10

3x106

7 (a.u.)Intensity Intensity (a.u.)Intensity 4.0x10 2x106 2.0x107 6 1x10 0.0 200 300 400 500 600 700 0 200 300 400 500 600 700 Wavelength (nm) Wavelength (nm)

6 1.8x10  = 300 nm 6 ex as-prepared 1.6x10 600 °C 6 800 °C 1.4x10

1.2x106

6 1.0x10

8.0x105

Intensity (a.u.)Intensity 6.0x105

4.0x105 350 400 450 500 550 600 650 700 Wavelength (nm) Fig. 8.3: Photoluminescence spectra of borohydride derived un-doped zirconia powders.

The PL signal of both amorphous and crystalline t-zirconia was broad starting from 250 nm to 700 nm with a centred peak at ~ 430 nm, when excited at a wavelength of 205 nm. When the excitation wavelength increases to 210 nm or 220 nm, there was no adequate change of PL spectra, but, the PL intensity slightly decreases. However, the PL intensity of zirconia decreases two order magnitudes, when the excitation wavelength is used as 240

nm. Also, the PL intensity marginally decreases with increase in calcination temperature for any particular excitation wavelength. In addition, the PL intensity of zirconia suddenly decreases without affecting the peak position, when the excitation wavelength was used as 300 nm. At this excitation wavelength, the amorphous and calcined t-zirconia show a broad PL signal, centered peak at same 430 nm. However, the PL signal of amorphous zirconia was prominent, as compared with calcined t-ZrO2 samples. It is well understood that if the excitation energy h > 5.55 eV (or excitation wavelength < 223 nm), h > 5.75 eV (or excitation wavelength < 215 nm), and h > 5.65 eV (or excitation wavelength < 219 nm), then there could be a direct band to band transition in case for as-synthesized amorphous zirconia, calcined t-ZrO2 at 600 °C and 800 °C, respectively. In addition, it could also be speculated that the defects are more concentrated at the surface than in the interior in case ultra-fine nano-sized materials. In the same time, defects can induce the formation of different energy levels in the bandgap for nano-sized materials. Also, it is well-reported that the near VUV wavelength could only penetrate part of the particle surface [253]. So, the higher intensity of emission of zirconia samples at an excitation wavelength of 205 nm (~ 6.04 eV) or 210 nm (~5.90 eV) is due to the total contribution of both surface defect and induced different energy levels in the band gap. When the excitation wavelength of 205 nm or 210 nm is used then there is an unchanged nature of emission position and shape, indicating that the luminescence involves the same initial and final states. This can also be explained by fast relaxation from the final state reached by photo-excitation to those states from which the luminescence originates [254]. Additionally, a minute decrease of luminescence intensity with increase in calcination temperature was observed at excitation of either 205 nm or 210 nm. This could be explained on the basis of the decrease of surface defects with increase in calcination temperature. Similar behaviour of PL spectra of zirconia samples were observed with slightly lower emission intensity (compared to PL spectra at ex = 205 nm or 210 nm) when excitation wavelength of 220 nm (~ 5.63 eV) was used and it was due to decrease of excitation energy. However, a drastic decrease of PL intensity with unchanged peak position of 430 nm was observed, when excited at a wavelength of 240 nm (i.e. 5.16 eV) and 300 nm (i.e. 4.13 eV) indicates that the emission may be due to the presence of different energy levels in the band gap, as the energy of excitation is not enough for direct band to band transition. Also decrease of PL intensity with calcination temperature either at 240 nm or 300 nm was due to decrease of defects within the band gap.

It was also well-understood that the broad luminescence band excited at a wavelength starting from 205 nm to 300 nm seems to be mostly caused by the nano-sized particles leading to an inhomogeneous broadening from a distribution of surface or defect states [254]. Also, according to photoluminescence band position, the observed PL spectra fall in the range of 2.0−3.5 eV, which further confirmed that the luminescence is due to defects. Additional emission band peaks at 328 nm, 408 nm, 450 nm, 465 nm, 495 nm and 533 nm in the PL spectra are mainly due to involvement of mid-gap trap states [185, 255]. So, defects can induce the formation of new energy levels in the bandgap and thus different excited wavelengths induce PL emission results from the radiative recombination of photo- generated hole with an electron occupying the defects [172]. Based on the band gap and PL analysis, the smaller sized particle along with surface defects (thin layer of ultra-fine porous particles on surface of zirconia) induces luminescence in borohydride derived un-doped as- synthesized and calcined zirconia samples. The PL signal centered at ~430 nm for un-doped zirconia powders could be assigned as a violet-blue emission [256-258]. To further confirm the colour of zirconia powders, the commission International de I’Eclairage (CIE) 1931 XY chromaticity coordinate graph (XY diagram) of the borohydride synthesized un-doped zirconia upon different excitation wavelengths are compared in Fig. 8.4.

Fig. 8.4: Excitation wavelength variation chromaticity diagram for borohydride derived un-doped zirconia (a) and fluorescence microscope image (b) indicate that the sample is nearly violet-blue in colour.

The chromaticity coordinates of all samples are added in Table 8.1. With increase in excitation wavelength, the emitting colour shifts from violet-blue to nearly light violet-blue.

The intensity variation in the PL spectra in as-synthesized and calcined zirconia samples manifests itself in the form of a change in the chromaticity coordinate positions in CIE diagram, which shows its position spanning from violet-blue to nearly light violet-blue region with change in excitation wavelength. Further, when the UV light of wavelength of 254 nm was excited on the powder sample (as-synthesized, calcined at 600°C and 800°C), violet-blue colour was observed, as shown in Fig. 8.4 (b). So, it was confirmed that the un- doped zirconia sample is nearly violet-blue in colour, which is well-correlated with the chromaticity coordinates. Table 8.1: Chromaticity co-ordinates of borohydride synthesized un-doped zirconia at different excitation wavelengths

Excited at 205nm x-coordinates y-coordinates As-prepared zirconia 0.234 0.235 Zirconia, 600°C 0.233 0.234 Zirconia, 800°C 0.228 0.229

Excited at 210nm x-coordinates y-coordinates As-prepared zirconia 0.235 0.236 Zirconia, 600°C 0.233 0.235 Zirconia, 800°C 0.228 0.229

Excited at 220nm x-coordinates y-coordinates As-prepared zirconia 0.235 0.236 Zirconia, 600°C 0.235 0.236 Zirconia, 800°C 0.230 0.230

Excited at 240nm x-coordinates y-coordinates As-prepared zirconia 0.235 0.241 Zirconia, 600°C 0.236 0.246 Zirconia, 800°C 0.232 0.243 Excited at 300nm x-coordinates y-coordinates As-prepared zirconia 0.281 0.283 Zirconia, 600°C 0.285 0.288

Zirconia, 800°C 0.287 0.289

8.2.2: Remarks The luminescence behaviour of un-doped zirconia was mainly due to smaller size particle with porous structure as well as surface defects. In addition, the colour of the un- doped porous zirconia was found to be violet-blue. So, by incorporation of Eu and / or Tb ion in porous zirconia may led to develop a phosphor materials for lighting applications. In the same time, incorporation of rare earth ion may stable the t- or c-phase of zirconia. So, next section deals with the phase stability, powder or pore morphology, rare earth distribution, luminescence and powder colour of the rare earth based porous zirconia.

8.2.3: Phase analysis of doped (Eu3+/Tb3+) and co-doped ((Eu3+-Tb3+)) porous zirconia nanopowders Phase analysis of single doped [5 mol % Eu3+ or5 mol % Tb3+] and co-doped [2 mol % Eu3+-5 mol % Tb3+] zirconia samples were analyzed using XRD. Fig. 8.5 (a), (b) and (c) shows XRD patterns of calcined (600 °C and 800 °C) Eu3+-doped, Tb3+-doped and Eu3+- Tb3+-co-doped zirconia, respectively.

(a) (b)

c-zirconia phase t-zirconia phase 800 °C 800 °C

Intensity (a.u.)Intensity

Intensity (a.u.)Intensity t-zirconia phase 600 °C c-zirconia phase 600 °C

66 68 70 72 74 76 78 66 68 70 72 74 76 78

2 (in degree) 800 °C 2 (in degree)

Intensity (a.u.)Intensity 800 °C

Intensity (a.u.)Intensity 600 °C 600 °C 20 30 40 50 60 70 80 20 30 40 50 60 70 80 2 (in degree) 2 (in degree)

(c)

c-zirconia phase 800 °C

c-zirconia phase 600 °C (a.u.)Intensity 800 °C

66 68 70 72 74 76 78 2 (in degree)

Intensity (a.u.)Intensity 600 °C

20 30 40 50 60 70 80 2 (in Adegree)

Fig. 8.5: XRD patterns of calcined (600 °C and 800 °C) samples of Eu3+-doped (a), Tb3+- doped (b) and Eu3+-Tb3+-co-doped zirconia (c).

All samples calcined either at 600 °C and 800 °C are found to be crystalline in nature except the Tb3+-doped zirconia sample calcined at 600 °C. The Tb3+-doped zirconia sample calcined at 600 °C is found to be partly crystalline and partly amorphous in nature. In order to justify the t- and or c-phase of zirconia, a enlarge version of XRD patterns in the range from 66 degree to 78 degree were inserted in each XRD patterns of Fig. 8.5. After analysis of the inset XRD patterns, it was found that all the peaks of the samples calcined at 600 °C

and 800 °C are indexed with the cubic structure of zirconia, as per JCPDS file number 27- 0997, except the Eu3+-doped sample calcined either at 600 °C or 800 °C. This particular sample was indexed with tetragonal structure of zirconia as per JCPDS file number: 79- 1768. The crystallite size of all samples are determined with the help of Scherrer’s formula and is given in Table 8.2. Phases obtained at 600 °C and 800 °C are also presented in Table 8.2. It was observed that the crystallite size was found to be increased slowly with increases in calcination temperature from 600 °C to 800 °C. However, the crystallite size of Eu-doped zirconia sample was found to be larger than the Tb-doped and co-doped (Eu and Tb) zirconia samples. The higher crystallite size may led to form t-phase of zirconia in case of Eu-doped zirconia, whereas, lower crystallite size stable the c-phase of zirconia even up to 800 °C. It was also well known that Y and Tb are trivalent and the stabilization mechanism of Tb is similar to Y [259]. The decrease of particle size as well as similar stabilization mechanism with Y-doped zirconia may be the possible reason for stabilizing the c-phase of zirconia in case of Tb-doped zirconia sample.

Table 8.2: Phase and crystallite size of doped and co-doped zirconia sample Sample Phase Crystallite size (nm) 600 °C 800 °C 600 °C 800 °C Eu3+-doped zirconia T T 18 26 Tb3+-doped zirconia C C 14 21 Eu3+-Tb3+-co-doped C C 13 14 zirconia Note: ‘C’ corresponds to cubic structure, ‘T’ corresponds to tetragonal structure. 8.2.4: Powder morphology of doped (Eu3+/Tb3+) and co-doped (Eu3+-Tb3+) porous zirconia nanopowders In order to examine the morphology of doped zirconia, TEM was performed on these samples. Fig. 8.6 (a-c) shows TEM micrographs of Eu3+-doped, Tb3+-doped and Eu3+-Tb3+- co-doped zirconia, calcined at 800 °C, respectively. These doped samples show nearly spherical in nature. The particle size of Eu3+-doped zirconia was in the range between 20 to 30 nm. However, the Tb3+-doped zirconia was seems to be smaller in size than other two samples. The particle size of co-doped zirconia was nearly uniform in nature and was found to be around 20 nm, as confirmed from inset of Fig. 8.6 (c). The particle size obtained from TEM micrographs was inconsistent with the crystallite size obtained from Scherrer’s formula. All samples are porous in nature. The pores are present within the particles as confirmed from TEM micrographs. One such enlarge view of TEM micrograph indicating

the porous nature of the particles was shown in the inset of Fig. 8.6 (c). Two different types of pores are present in the samples such as the pores are entrapped within the particles and larger pores are formed due to coagulation of the agglomerated particles. The size of pores within the particle was in the range of 5-10 nm which comes under mesopores type. The size of larger pores due to formation of agglomerated nature was found to be in the range between 30 nm to 150 nm. In addition, the pores are ultrafine and uniform in nature in case of Tb3+-doped zirconia sample, whereas, the pores are large in size as well as dual modal (within the particles and around the particles) pores are present in Eu3+ as well as co-doped sample. Further to analyse the pore in the particles, high angle annular dark field scanning transmission electron microscopy (HAADF-STEM) images were performed on these samples and was shown in Fig. 8.7. From HAADF-STEM images, it was clear that the pores are present within the particles (intra-particular) as well as around the agglomerated particles (inter-particular). The pore size of around 10 nm was observed for Eu3+-doped zirconia sample. The pores are finer in nature in case of Tb3+-doped sample and pores are found to be larger (as compared to Eu3+ and Tb3+ sample) in size in case of co-doped zirconia sample.

Fig. 8.6: TEM micrographs of Eu3+-doped (a), Tb3+-doped (b) and Eu3+-Tb3+-co-doped (c) zirconia, calcined at 800 °C.

Fig. 8.7 HAADF-STEM images of (a) Eu3+-doped, (b) Tb3+-doped and (c) co-doped (Eu3+- Tb3+) zirconia.

8.2.5: Distribution of rare earth ions in porous zirconia particles Further, rare earth elemental distribution in porous zirconia was analysed using HAADF-STEM image. From energy dispersive X-ray spectroscopy (EDX) line profile scan, rare earth signals were detected. Fig. 8.8 indicate the EDX line profile of Eu3+-doped porous zirconia particles. The similar trend of EDX line profile was also observed for the Pd/Pt core-shell nanoparticle synthesized via area-selective atomic layer deposition (ALD)

by Cao et al. [260]. Based on this literature as well as compositional line profile of Fig. 8.8, it was found that the Eu signals was not observed on the surface (within few nm) of the zirconia particle, when compared with the Zr signal [see the dotted line]. The Eu signals are detected only in the inner core of the zirconia particles.

Zr (K)

Eu (L)

(a.u) Intensity

0 20 40 60 80 100 Point number

Fig. 8.8: EDX line profile of Eu3+-doped porous zirconia particles. Inset represents the HAADF-STEM image of Eu3+-doped porous zirconia. Line scan along the direction denoted by the green line. Dotted line represents that the Eu signal was not present on the surface of the zirconia particle.

Similarly, EDX line profile scan of Tb3+-doped zirconia was performed. Fig. 8.9 shows the EDX line profile of Tb3+-doped porous zirconia of two particles. In this case, it was also observed that Tb3+ ions was also not prominently present on the surface of the two particles, as indicated by dotted line.

(a.u)

)



(

Intensity Tb

0 20 40 60 80 100

Point number

Fig. 8.9: EDX line profile of Tb3+-doped porous zirconia particles. Inset represents the HAADF-STEM image of Tb3+-doped porous zirconia. Line scan along the direction denoted by the green line. Dotted line represents that the Tb signal was not present on the surface of the zirconia particle.

Similarly, EDX line profile scan of co-doped (Eu3+-Tb3+) porous zirconia was performed. Fig. 8.10 indicates the EDX line profile of co-doped porous zirconia of single particle. Based on the dotted line, it was also observed that Eu as well as Tb signals were also not prominently present on the surface of the particles. In all the three samples, it was confirmed that the rare earth ions are present within the core of the particle, but not prominently present near the surface of the zirconia particles. Further in order to understand the distribution of rare earth ions in the zirconia particles, elemental mapping from HAADF-STEM image was performed on co-doped porous zirconia sample.

(a.u.)

)



(

Intensity Eu

0 20 40 60 80 100

Point number

(a.u.)

)



L (

Intensity Tb

0 20 40 60 80 100 Point number

Fig. 8.10: EDX line profile of co-doped (Eu3+-Tb3+) porous zirconia particles. Inset represents the HAADF-STEM image of co-doped porous zirconia. Line scan along the direction denoted by the green line. Dotted line represents that the Eu and Tb signal was not present on the surface of the zirconia particle.

Fig. 8.11 (a), (b) and (c) show HAADF-STEM image and elemental mapping of Zr, Eu and Tb, and EDX line scan of Zr, Eu and Tb, respectively. The elemental mapping (Fig. 8.11 b) was obtained from the HAADF-STEM image (see box region) and it indicates that the rare earth ions of Eu and Tb were well distributed within the zirconia particles. However, Eu and Tb ions were not prominently present on the surface of zirconia particle. The above elemental line profile as well as elemental mapping analysis confirmed that the rare earth

ions were more prominent inside the zirconia particle, as these elements were not prominent on the individual surface of the particle. The EDX line scan (see Fig. 8.11 c) indicates the presence of Eu, Tb and Zr in the particles. It was found that there was no other impurity phases were present. Further, the Eu and Tb concentration was calculated by the Eu or Tb atomic concentration divided by the sum of Zr, Eu and Tb atomic concentration [253]. The Eu and Tb concentration was found to be 1.65 % and 4.41 %, respectively.

(a) (b)

cps/eV Eu-LATb-LA Zr-KA 4.5 4.0 (c) 3.5 Zr Zr 3.0 Tb 2.5 Tb Tb

2.0 Eu Eu Zr Eu Tb

1.5 1.0 0.5 0.0 5 10 15 20 25 30 35 40 keV Fig. 8.11: EDX elemental mapping (b) and line scan (c) of co-doped (Eu3+-Tb3+) porous zirconia particles from HAADF image (a).

8.2.6: Pore morphology of rare earth doped and co-doped porous zirconia Pore morphology of calcined doped and co-doped porous zirconia was analysed using BET-isotherm. Fig. 8.12 shows BET isotherm and pore size distribution (inset of Fig.

8.12) of Eu3+, Tb3+ and co-doped (Eu3+-Tb3+) porous zirconia sample calcined at 800 °C. All three samples, the BET isotherm looks similar behaviour. In these BET-isotherm, adsorption curve slightly increases up to P/P0 = 0.6 and then suddenly increases with increase in partial pressure with non-saturating behaviour and forming a hysteresis loop with the help of desorption curve. In general the curvature of BET hysteresis loop of these three sample was typically H3 type indicating lamellar or slit type pores, however, bottle neck pores may present in Tb and co-doped sample due to change in slope of desorption curve at

P/P0 = 0.9. The pore size distribution was evaluated from BJH desorption curve and was shown in the inset of each BET-isotherm of Fig. 8.12. From the inset figures, it was found that the pore size varies from 4.2 nm to 130 nm with a maximum of 9.6 nm with wide pore size distribution for Eu-doped zirconia sample. Similarly, the pore size varies from 5.2 nm to 63 nm with a maximum of 18.6 nm with wider pore size distribution for Tb-doped zirconia sample. However, a broad pore size distribution in the range of 4.2 nm to 160 nm with a maximum of 6.5 nm and 31.2 nm was observed for co-doped zirconia sample. The wider distribution was mainly due to the lamellar type pores. In addition, the surface area of Eu, Tb and co-doped porous zirconium oxide, calcined at 800 °C was found to be ~ 36 m2/g, 29 m2/g and 32 m2/g, respectively.

250 210 (b) (a) 180 0.30 18.6

0.40 9.6 200 0.25 0.32 150 0.20

150 0.24 0.15

0.16 120 0.10 Volume (cc/g) Volume

Volume (cc/g) Volume 0.08 0.05 100 90 0.00 0.00 0 50 100 150 200 250 300 0 20 40 60 80 100 120 140 60 Pore diameter (nm) 50 Pore diameter (nm) 30

volume adsorbed (cc/g) adsorbed volume 0 (cc/g) adsorbed volume 0 0.0 0.2 0.4 0.6 0.8 1.0 0.0 0.2 0.4 0.6 0.8 1.0 Relative pressure (P/P ) Relative pressure (P/P0) 0 300 (c) 0.5 6.5 250 0.4 0.3

200 0.2

Volume (cc/g) Volume 150 0.1 0.0

0 20 40 60 80 100 120 140 160 100 Pore diameter (nm)

volume adsorbed (cc/g) adsorbed volume 0 0.0 0.2 0.4 0.6 0.8 1.0 Relative pressure (P/P0) Fig. 8.12: BET-isotherm and pore size distribution (see inset) of (a) Eu3+, (b) Tb3+ and (c) co-doped (Eu3+-Tb3+) porous zirconia sample, calcined at 800 °C.

8.2.7: Photoluminescence behaviour of Eu-doped porous zirconia As observed from EDX line profile analysis, it was found that the distribution of rare earth ions is more prominent in the centre of the zirconia particle. So, the luminescence behaviour may strongly affect for the rare earth doped and co-doped zirconia sample. In order to study the photoluminescence behaviour, first, the excitation spectra was performed on Eu3+-doped, Tb3+-doped and co-doped (Eu3+-Tb3+) zirconia samples. The behaviour of excitation spectra for all sample was found to be similar. The typical excitation spectra of Eu-doped sample calcined at 600 °C, performed at emission wavelenght of 614 nm, was shown in Fig. 8.13. It was observed that the intensity was maximum in the range from 200 nm to 230 nm. However, in this present study, the emission spectra of Eu3+-doped, Tb3+- doped and co-doped (Eu3+-Tb3+) zirconia samples, calcined at 600 °C and 800 °C were analyzed at different excitation wavelengths in the range between 205 nm to 350 nm.

7x108 em = 614 nm 6x108

8 5x10 8 4x10 3x108

2x108

Intensity (a.u.) Intensity 8 1x10

200 250 300 350 400 Wavelength ( nm) Fig. 8.13: Excitation spectra of Eu-doped sample calcined at 600 °C.

Figure 8.14 show the photoluminescence spectra of calcined (600 °C) Eu3+-doped zirconia powders, excited at different wavelengths. The PL signal of calcined (600 °C) Eu3+- doped zirconia was broad starting from 250 nm to 700 nm with a centered peak at ~ 430 nm, when excited at a wavelength of 205 nm. When the excitation wavelength increases to 210 nm or 220 nm, there was no adequate change of PL spectra. Moreover, the PL intensity of zirconia slightly decreases, when the excitation wavelength was used as 230 nm, whereas PL intensity decreases one order of magnitude when excited at 240 nm. Also, there was no indication of emission spectra due to the Eu3+ ions when the sample was excited in the range between 205nm to 220 nm. When the sample was excited at 230 nm, a broad peak ranging from 250 nm to 700 nm along with sharp peak at 585 nm and 610 nm were observed. The 5 7 emission peak at 585 nm corresponds to D0 → F1 transition and the emission peak at 610

5 7 5 7 nm corresponds to D0 → F2 transition. The transition of D0 → F2 (red emission) is due to 5 7 the typical electric dipole transition, and the transition of D0 → F1 (orange emission) is attributed to the magnetic dipole transition.

7 6x10 1.8x107  = 230 nm  = 205 nm ex ex 7  = 240 nm 7  = 210 nm 1.6x10 ex 5x10 ex 7  = 250 nm  = 220 nm 1.4x10 ex ex 7 7 4x10 1.2x10

7 7 1.0x10

3x10

8.0x106

Intensity 7 Intensity 2x10 6.0x106 4.0x106 1x107 2.0x106 0 0.0 250 300 350 400 450 500 550 600 650 700 250 300 350 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm) 7x106  = 260 nm 6  = 290 nm ex 4x10 ex 6 6x10  = 270 nm  = 300 nm ex ex 6  = 280 nm  = 310 nm 5x10 ex 3x106 ex  = 350 nm ex 4x106 6

2x10 3x106

Intensity

6 2x10 1x106 1x106 0 0 300 350 400 450 500 550 600 650 700 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm)

Fig. 8.14: Photoluminescence spectra of calcined (600 °C) Eu3+-doped zirconia powders.

In addition, the intensity of the emission peak observed at 585 nm and 610 nm increases with the increase in excitation wavelengths when it was changes from 230 nm to 250 nm. Whereas, the intensity of the broad peak (i.e. observed in the range between 250 nm to 500 nm with a maximum at around 430 nm) decreases with the increase in excitation wavelengths i.e. from 230 nm to 250 nm. Further, when the excitation wavelength used as 260 nm, 270 nm and 280 nm, there was slightly decrease of intensity of emission spectra 5 7 5 7 position at 585 nm (i.e. D0 → F1 transition) and 614 nm (i.e. D0 → F2 transition). In the same time, the intensity of the broad peak region from 300 nm to 550 nm also slightly decreases. If further increase in excitation wavelengths (i.e. 290 nm, 300nm, 310 nm, and 350 nm), there was decrease of intensity of the emission peak position at 585 nm and 614 nm. In a similar way, photoluminescence spectra of calcined (800 °C) Eu3+-doped zirconia powders was performed at different excitation wavelengths and the emission spectra was shown in Fig. 8.15. Similar trend of PL spectra was observed for the Eu-doped zirconia

sample calcined at 800 °C, as observed for the sample calcined at 600 °C. However, the 5 7 5 7 major difference is the emission peak width of D0 → F1 and D0 → F2, which was sharper in the sample calcined at 800 °C than the sample calcined at 600 °C. Also, the peak intensity at each excitation wavelength decreases if the calcination temperature increases from 600 °C to 800 °C. The sharp nature of the emission peak for the sample calcined at 800 °C was due to the particle size effect as the particle size increases from 600 °C to 800 °C. Also due 5 7 5 7 to increase in particle size, the peak intensity of D0 → F1 and D0 → F2 decreases with increase in calcination temperature for same excitation wavelength.

3.0x107 1.0x107  = 230 nm  = 205 nm ex ex  = 240 nm 7  = 210 nm ex 2.5x10 ex 6 8.0x10  = 250 nm  = 220 nm ex ex 2.0x107 6.0x106 7

1.5x10

Intensity

Intensity 4.0x10 1.0x107

2.0x106 5.0x106

0.0 0.0 250 300 350 400 450 500 550 600 650 700 250 300 350 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm)

7x105 6  = 260 nm  = 290 nm 2.5x10 ex ex 6x105  = 270 nm  = 300 nm ex ex 6 2.0x10  = 280 nm 5  = 310 nm ex 5x10 ex  = 350 nm ex 5 1.5x106 4x10

5 6 3x10

Intensity

1.0x10 Intensity 2x105 5.0x105 1x105

0.0 0 300 350 400 450 500 550 600 650 700 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm) Fig. 8.15: Photoluminescence spectra of calcined (800 °C) Eu3+-doped zirconia powders.

For lighting applications, poor colour purity hampers the use of the phosphor, which corresponds to the fact that the intensity of red emission of Eu3+-doped sample is lower than 5 7 5 7 that of the orange emission. Colour purity is denoted as R/O ratio [I( D0 → F2)/I( D0 → F1)] and was calculated by considering the sum of integral intensity of red emission peaks 5 7 3+ observed at 614 and 630 nm for contribution of D0 → F2 transition. The R/O ratio of Eu - doped zirconia sample calcined at 600 °C and 800 °C was determined at different excitation wavelengths and was shown in Fig.8.16. The trend of R/O ratio for both the sample was found to be nearly similar. It was found that the R/O ratio of 600 °C calcined sample shows

higher value than 800 °C calcined sample at above the excitation wavelength of 250 nm. The R/O ratio of 600 °C calcined sample was in range in between 1.86 to 2.53 and the R/O ratio of 800 °C calcined sample was in range in between 1.75 to 2.20. The higher R/O ratio for Eu3+-doped zirconia sample (calcined at 600 °C) may be due to smaller particle size. However, there is no direct correlation between size and colour purity [91]. But, in fact, the distribution of Eu3+ in zirconia matrix associated with smaller particle size was responsible for the change in R/O values. As observed from elemental profile, the europium concentration on the interior and the surface of nano-sized zirconia was different. The 5 7 contribution of D0 → F2 transition was more prominent in the interior of particle than the surface, and thus leads to higher R/O ratio in UV range (above 230 nm).

3.0 600 °C 2.7 800 °C 2.4

2.1

1.8

Ratio R/O 1.5

1.2 220 240 260 280 300 320 340 360 Wavelength (nm) Fig.8.16: Excitation wavelength dependent R/O ratio of Eu-doped zirconia.

To further confirm the colour of Eu-doped zirconia powders, the CIE-1931 XY chromaticity coordinate graph (XY diagram) of the borohydride synthesized zirconia upon different excitation wavelengths are compared in Fig. 8.17. The chromaticity coordinates of these samples are given in Table 8.3. It was observed from Fig. 8.14 and Fig. 8.15 that the emission spectra consists of a broad peak, centered at 420 nm and sharp peaks at 595 nm and 614 nm. The broad nature of peak was mainly due to the contribution of the host, whereas, the sharp peaks are a result of the contribution of activator i.e. europium ions. The concentration of europium ions in zirconia matrix at the surface and interior was different and thus this intensity difference manifests itself in the form of a change in the chromaticity coordinate positions in CIE diagram. The CIE coordinates position spanning from bluish to reddish-orange region, when excitation wavelength changes from 205 nm to 280 nm,

whereas the colour changes from reddish-orange to nearly pinkish-white, when the excitation wavelength used in the range from 290 nm to 350 nm for the Eu-doped zirconia sample calcined at 600 °C. However, for the Eu-doped zirconia sample, calcined at 800 °C show a colour spanning from bluish to slight reddish-orange, when the excitation wavelength changes from 205 nm to 250 nm and then the colour changes from slight reddish-orange to nearly pinkish-white, when the excitation wavelength used from 260 nm to 350 nm.

Table 8.3: Chromaticity co-ordinates of Eu-doped zirconia obtained at different excitation wavelengths

Excited at Sample calcined at 600 °C Sample calcined at 800 °C wavelength x-coordinates y-coordinates x-coordinates y-coordinates (nm) 205 0.226 0.222 0.226 0.222 210 0.226 0.222 0.226 0.223 220 0.229 0.224 0.229 0.224 230 0.247 0.233 0.254 0.238 240 0.356 0.277 0.356 0.282 250 0.455 0.322 0.411 0.314 260 0.488 0.334 0.406 0.325 270 0.496 0.335 0.380 0.317 280 0.491 0.332 0.364 0.309 290 0.486 0.335 0.365 0.307 300 0.484 0.335 0.363 0.307 310 0.428 0.315 0.353 0.301 350 0.317 0.289 0.364 0.308

Further, photograph of the powder sample was taken under UV lamp source at wavelength of 254 nm and 365 nm. Fig. 8.18 shows the powder colour of Eu-doped samples under UV lamp. The both calcined (600 °C and 800 °C) zirconia sample’s colour was observed as reddish-orange when excited at 254 nm. When excited at 365 nm, the colour of both sample was found to be pinkish-white in colour. It was found that the colour of these samples nearly matches with the colour mentioned in CIE-diagram.

(600 °C) (800 °C)

Fig. 8.17: Excitation wavelength variation chromaticity diagram for Eu-doped zirconia.

Fig. 8.18: Powder colour image of Eu-doped zirconia powder under UV lamp.

8.2.8: Photoluminescence behaviour of Tb-doped porous zirconia Photoluminescence spectra of calcined (600 °C) Tb-doped zirconia powders, excited at different wavelengths (ranges from 205 nm to 350 nm) were shown in Fig. 8.19.

8x106 6  = 230 nm  = 205 nm ex 6 ex 2.0x10 7x10  = 210 nm  = 240 nm ex ex 6  = 220 nm  = 250 nm 6x10 ex ex 1.5x106 6 5x10

4x10 1.0x106 6

Intensity 3x10 Intensity 6 2x10 5 5.0x10 1x106

0 0.0 250 300 350 400 450 500 550 600 650 700 250 300 350 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm) 2.5x105  = 260 nm  = 290 nm 4x105 ex ex  = 270 nm  = 300 nm ex 2.0x105 ex  = 280 nm  = 310 nm ex ex 3x10 5 1.5x105

2x105 1.0x105

Intensity

5 1x10 5.0x104

0 0.0 350 400 450 500 550 600 650 700 300 350 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm)

5 1.2x10  = 320 nm ex 5  = 330 nm 1.0x10 ex  = 340 nm ex 4 8.0x10  = 350 nm ex 6.0x104

Intensity 4 4.0x10 2.0x104 0.0 400 450 500 550 600 650 700 Wavelength (nm) Fig. 8.19: Photoluminescence spectra of calcined (600 °C) Tb-doped zirconia powders.

The PL signal of calcined (600 °C) Tb-doped zirconia was also similar type of Eu- doped zirconia sample. A broad peak starting from 250 nm to 700 nm with a centered peak at ~ 430 nm was observed when excited at a wavelength of 205 nm. When the excitation

100

wavelength increases to 210 nm or 220 nm, there was no adequate change of PL spectra. However, the PL intensity slightly decreases, when the excitation wavelength was used as 230 nm, whereas PL intensity decreases one order of magnitude when excited at 240 nm or 250 nm. It was also found that there was no indication of emission spectra due to the Tb3+ ions when the sample was excited in the range between 205nm to 220 nm. When the sample was excited at 230 nm, there was broad peak ranging from 250 nm to 700 nm along with sharp peak at 540 nm. When the sample was excited at 240 and 250 nm, the characteristic emission of a Tb3+ ion was prominently observed. Generally, the emission spectra display 3+ 5 the characteristic emission of a Tb ion composed of four bands associated with the D4 →

7 5 7 5 7 5 7 Fj (j = 3-6) transition, at 488 nm ( D4  F6), 543 nm ( D4  F5), 584 nm ( D4  F4), and

5 7 622 nm ( D4  F3) nm, respectively. Also, the broad peak intensity decreases with the increase in excitation wavelengths from 230 nm to 250 nm. When the excitation wavelength used as 260 nm, 270 nm and 280 nm, there was slight decrease in intensity of emission spectra. In the same time, the intensity of the broad peak region also decreases. Further increase in excitation wavelengths (290 nm, 300nm, 310 nm, 320 nm, 330 nm, 340 nm and 350 nm), the four emission peaks due to Tb3+ ion was prominent and the intensity slightly decreases with in excitation wavelengths. In a similar way, photoluminescence spectra of the calcined (800 °C) Tb-doped zirconia powders were performed at different excitation wavelengths and was shown in Fig. 8.20. Similar trend of PL spectra was observed for the sample calcined at 800 °C, as compared with the PL of the sample calcined at 600 °C, when the sample was excited at wavelengths of 205nm, 210nm and 220 nm. However, when the excitation wavelength increases to 220 nm, 230 nm and 250 nm, the emission spectra due to Tb ion was not prominent. In addition, the intensity of the broad peak decreases with increase in excitation wavelengths. The emission peak due to Tb ion was found to be prominent when the excitation wavelength was used as 260 nm. The four characteristic emission of a Tb3+ ion was also observed when the excitation wavelength increases from 270 nm to 350 nm. However, the intensity of these peaks were decreases with increase in excitation wavelengths. Comparing PL spectra of calcined (600 °C) Tb-doped sample with 800 °C calcined Tb-dopes sample, it was found that the intensity of characteristics emission peak of Tb ion was decreased for 800 °C calcined sample and it was due to increase of particle size when calcined from 600 °C to 800 °C. In addition, a low intensity of the characteristics emission peaks was observed for 800 °C calcined Tb-doped zirconia sample when it was excited in the range of 340 nm to 350 nm.

101

6 6x10 7 1.6x10  = 205 nm  = 230 nm ex ex 6 7  = 210 nm  = 240 nm 1.4x10 ex 5x10 ex 7  = 220 nm  = 250 nm 1.2x10 ex ex 4x106 1.0x107 6 6

3x10 8.0x10

Intensity 6

6.0x10 Intensity 2x10 4.0x106 6 2.0x106 1x10

0.0 0 250 300 350 400 450 500 550 600 650 700 250 300 350 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm)

5 4x10 5 5x10  = 260 nm ex  = 290 nm ex  = 270 nm ex  = 300 nm 5 4x105 ex 3x10  = 280 nm  = 310 nm ex ex

3x105 2x105

5 Intensity 2x10

Intensity 1x105 1x105

0 350 400 450 500 550 600 650 700 300 350 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm)

1.5x105  = 320 nm ex  = 330 nm 1.2x105 ex  = 340 nm ex  = 350 nm ex 9.0x104

Intensity 6.0x10

3.0x104

0.0 400 450 500 550 600 650 700 Wavelength (nm) Fig. 8.20: Photoluminescence spectra of calcined (800 °C) Tb-doped zirconia powders.

To further confirm the colour of Tb-doped zirconia powders, CIE-1931 XY chromaticity coordinate graph (XY diagram) of the Tb-doped zirconia upon different excitation wavelengths are compared in Fig. 8.21. The chromaticity coordinates of Tb- doped zirconia samples are given in Table 8.4. It was observed from Fig. 8.19 and Fig. 8.20 that the emission spectra consists of a broad peak centred at 420 nm and a sharp peaks at 488 nm, 543 nm, 584 nm, and 622 nm. The broad nature of peak was mainly due to the contribution of the host, whereas, the sharp peaks are a result of the contribution of activator

102

i.e. terbium ions. The concentration of terbium ions in zirconia matrix at the surface and interior was different and thus this intensity difference manifests itself in the form of a change in the chromaticity coordinate positions in CIE diagram. The CIE coordinates position spanning from sky blue to light green region, when excitation wavelength changes from 205 nm to 280 nm and then the colour changes from light blue to nearly bluish-white, when the excitation wavelength changes from 290 nm to 350 nm for the Tb-doped zirconia sample calcined at 600 °C. However, for the Tb-doped zirconia sample, calcined at 800 °C show a colour spanning from sky blue to light green, when the excitation wavelength changes from 205 nm to 300 nm and then the colour changes from greenish white to nearly bluish-white, when the excitation wavelength used from 310 nm to 350 nm.

Table 8.4: Chromaticity co-ordinates of Tb-doped zirconia obtained at different excitation wavelengths. Excited at Sample calcined at 600 °C Sample calcined at 800 °C wavelength x-coordinates y-coordinates x-coordinates y-coordinates (nm) 205 0.231 0.231 0.239 0.253 210 0.232 0.232 0.241 0.254 220 0.233 0.232 0.241 0.253 230 0.236 0.241 0.242 0.256 240 0.238 0.266 0.241 0.259 250 0.245 0.321 0.245 0.281 260 0.259 0.365 0.254 0.309 270 0.266 0.382 0.268 0.327 280 0.274 0.384 0.281 0.341 290 0.281 0.366 0.291 0.357 300 0.284 0.353 0.294 0.351 310 0.283 0.324 0.291 0.329 320 0.279 0.323 0.281 0.311 330 0.276 0.319 0.284 0.312 340 0.273 0.311 0.286 0.306 350 0.269 0.312 0.283 0.299

Further, photograph of the powder sample was taken under UV lamp source at wavelength of 254 nm and 365 nm. Fig. 8.22 shows the powder colour of Tb-doped samples under UV lamp. The 600 °C and 800 °C calcined zirconia sample’s colour was observed as light green and greyish green, respectively, when excited at 254 nm. When excited at 365 nm, the colour of 600 °C and 800 °C calcined zirconia sample’s was found to be bluish- white and light greyish white in colour. It was found that the colour of these samples nearly matches with the colour mentioned in CIE-diagram.

103

(600 °C) (800 °C)

Fig. 8.21: Excitation wavelength variation chromaticity diagram for Tb-doped zirconia.

Fig. 8.22: Powder colour image of Eu-doped zirconia powder under UV lamp.

104

8.2.9: Photoluminescence behaviour of co-doped (Eu-Tb) porous zirconia Photoluminescence spectra of calcined (600 °C) co-doped (Eu-Tb) zirconia powders was anlyzed at different excitation wavelengths and was shown in Fig. 8.23.

7 6x10 7  = 230 nm  = 205 nm ex ex 1.8x10  = 240 nm 7  = 210 nm ex 5x10 ex 7  = 220 nm 1.5x10  = 250 nm ex ex 4x107 1.2x107 3x107

9.0x106

Intensity Intensity 2x107 6 6.0x10 1x107 6 3.0x10 0 250 300 350 400 450 500 550 600 650 700 0.0 Wavelength (nm) 250 300 350 400 450 500 550 600 650 700 Wavelength (nm)

6 3.5x10  = 260 nm 6 ex 3.5x10  = 290 nm ex  = 270 nm 6 ex 6  = 300 nm 3.0x10 3.0x10 ex  = 280 nm ex  = 310 nm 6 6 ex 2.5x10 2.5x10 6 2.0x106 2.0x10

6 6 Intensity 1.5x10 1.5x10

Intensity 6 6 1.0x10 1.0x10 5.0x105 5.0x105 0.0 0.0 300 350 400 450 500 550 600 650 700 350 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm)

6  = 320 nm 1.2x10 ex  = 330 nm ex 6  = 340 nm 1.0x10 ex  = 350 nm ex 8.0x105

5 6.0x10 Intensity 4.0x105

2.0x105

0.0 400 450 500 550 600 650 700 Wavelength (nm) Fig. 8.23: Photoluminescence spectra of calcined (600 °C) co-doped (Eu-Tb) zirconia.

A broad peak in the range between 250 nm to 700 nm with a maxium at 440 nm was observed similar to PL spetra of Eu-doped or Tb-doped sample when it was was excited at 205 nm, 210 nm and 220 nm. In addition, there was no indication of change in PL spetra in these excitation wavelenghts. When the co-doped sample was excited at 230 nm, 240 nm and 250 nm, the broad peak decreases along with development of peaks related to Eu- transition (588nm, 604 nm, 630 nm and 650 nm) and Tb-transition (488 nm, 540 nm). The

105

combination of Eu-transition and Tb-transition peaks are more prominent when the sample was excited at 260 nm, 270 nm and 280 nm. The intensity of these peaks was found to be slightly decreases with increase in excitation wavelenghts. When the sample was excited at higher excitation wavelengths i.e. from 290 nm to 350 nm, the characteristic peaks related to Eu- and Tb- transition was also found to be prominent, but the intensity decreases with excitation wavelenghts. Similarly, PL spectra of calcined (800 °C) co-doped (Eu-Tb) zirconia powders was anlyzed at different excitation wavelengths and was shown in Fig.

8.24.

7  = 205 nm 7  = 230 nm 3.0x10 ex 1.0x10 ex  = 210 nm  = 240 nm 7 ex ex 2.5x10  = 220 nm 6  = 250 nm ex 8.0x10 ex 2.0x10 7 6.0x106 7

1.5x10

Intensity 6 Intensity 4.0x10 1.0x10 7

6 5.0x10 6 2.0x10

0.0 0.0 250 300 350 400 450 500 550 600 650 700 250 300 350 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm)

 = 260 nm 6 ex 2.0x10 6  = 270 nm 1.5x10 ex  = 290 nm ex  = 280 nm  = 300 nm ex 1.5x106 ex  = 310 nm ex 1.0x106 6 1.0x10

Intensity Intensity 5 5.0x10 5 5.0x10

0.0 300 350 400 450 500 550 600 650 700 350 400 450 500 550 600 650 700 Wavelength (nm) Wavelength (nm)

6.0x105  = 320 nm ex  = 330 nm ex  = 340 nm ex 5 4.0x10  = 350 nm ex

Intensity 5 2.0x10

0.0 400 450 500 550 600 650 700 Wavelength (nm) Fig. 8.24: Photoluminescence spectra of calcined (800 °C) co-doped (Eu-Tb) zirconia.

106

Similar trend of PL spetra was obtained for the sample calcined at 800 °C (as compared to 600 °C), when it was excitaed at different excitation wavelengths. There was no indication of emission peaks of Eu and Tb transtion were observed when the sample was excited in the range of 205 nm to 220 nm. However, the emission peaks related to Eu and Tb was found to developed when the sample was excited in the range of 230 nm to 250 nm. In adiditon, the borad peak was found to decreases with increaes in excitation wavelenghts. The emission peaks of Eu and Tb was also prominent when the sample was excited in the range between 260 nm to 350 nm. However, the peak intensity slightly decreases when the excitation wavlenth was increased from 260 nm 280 nm, but the intensity decrease with increase in excitation wavelenghts starting from 290 nm to 350 nm. Further, colour purity was determined for both calcined co-doped zirconia sample and the R/O ratio as a function of wavelength was shown in Fig. 8.25. The colour purity behaviour with excitation wavelength for the both calcined co-doped zirconia sample was found to be nearly similar. The colour purity of calcined (600 °C) was found to be in the range between 1.92 to 2.42 whereas the colour purity of calcined (800 °C) was found to be in the range between 1.97 to 2.62. The both calcined sample shows similar colour purity and the colour purity nearly constant from 240 nm to 280 nm and then increases maximum up to around 2.6 at excitation wavelength of 330 nm.

3.0 600 °C 2.7 800 °C

2.4

2.1

1.8

Ratio R/O 1.5

1.2

220 240 260 280 300 320 340 360

Wavelength (nm)

Fig.8.25: Excitation wavelength dependent R/O ratio of co-doped (Eu-Tb) zirconia.

To further confirm the colour of co-doped (Eu-Tb) zirconia powders, CIE-1931 XY chromaticity coordinate graph (XY diagram) of the co-doped zirconia upon different excitation wavelengths are compared in Fig. 8.26. The chromaticity coordinates of co-doped zirconia samples are given in Table 8.4. It was observed from Fig. 8.23 and Fig. 8.24 that

107

the emission spectra consists of a broad peak centred at 420 nm and a sharp peaks corresponds to Eu and Tb transition. The broad nature of peak was mainly due to the contribution of the host, whereas, the sharp peaks are a result of the contribution of activators i.e. europium and terbium ions. The concentration of europium and terbium ions in zirconia matrix at the surface and interior was different and thus this intensity difference manifests itself in the form of a change in the chromaticity coordinate positions in CIE diagram. The CIE coordinates position spanning from bluish to reddish-yellowish region, when excitation wavelength changes from 205 nm to 300 nm and then the colour changes from reddish-yellowish to nearly pinkish-white, when the excitation wavelength used from 310 nm to 350 nm for the co-doped zirconia sample calcined at 600 °C. In addition, for the co-doped zirconia sample, calcined at 800 °C show a colour spanning in a similar way to that of co-doped sample calcined at 600 °C.

Table 8.5: Chromaticity co-ordinates of co-doped (Eu-Tb) zirconia obtained at different excitation wavelengths. Excited at Sample calcined at 600 °C Sample calcined at 800 °C wavelength x-coordinates y-coordinates x-coordinates y-coordinates (nm) 205 0.231 0.238 0.236 0.251 210 0.231 0.238 0.236 0.251 220 0.232 0.239 0.238 0.251 230 0.237 0.243 0.242 0.254 240 0.264 0.256 0.261 0.262 250 0.321 0.288 0.308 0.286 260 0.366 0.321 0.351 0.318 270 0.408 0.334 0.395 0.333 280 0.442 0.342 0.433 0.343 290 0.457 0.342 0.454 0.345 300 0.463 0.343 0.458 0.344 310 0.433 0.331 0.421 0.331 320 0.391 0.321 0.376 0.316 330 0.345 0.316 0.336 0.315 340 0.331 0.314 0.322 0.311 350 0.323 0.311 0.314 0.306

Further, powder image of these sample was taken under UV lamp at wavelengths of 254 nm and 365 nm and was shown in Fig. 8.27. The 600 °C and 800 °C calcined sample’s colour was observed as reddish-yellowish, when excited at 254. However, the powder image of 600 °C calcined sample was found to be pinkish white when excited at 365 nm, but the powder image 800 °C calcined sample was found to be nearly white in colour, when excited

108

at 365 nm. It was also found that the colour of the powder taken under UV lamp was found to be matches with the colour mentioned in CIE-diagram.

(600 °C) (800 °C)

Fig. 8.26: Excitation wavelength variation chromaticity diagram for co-doped zirconia.

Fig. 8.27: Powder colour image of co-doped zirconia powder under UV lamp.

109

8.2.10: Proposed mechanism for showing different colours at different excitation wavelength Based on the structural analysis, it was found that the incorporation of rare earth ions such as Eu or Tb or mixture of Eu and Tb led to stable the t and / or c-phase of zirconia at 600 °C and / or 800 °C. In addition, the rare-earth ions based zirconia sample shows porous in nature as observed from TEM as well as STEM-HAADF images. From energy dispersive X-ray spectroscopy (EDX) line profile scan, it was confirmed that the rare earth signal was not prominently present on the surface of the particle, but rather distributed within the core of the porous zirconia particles. The existence of rare earth ions in the core of porous zirconia particles may be originated during borohydride synthesis. During synthesis, aqueous zirconium salt and rare earth salts (chloride form) were added drop wise to a beaker containing aqueous sodium borohydride and to maintain a constant pH, an additional amount of aqueous sodium borohydride was added separately along with aqueous zirconium salt and rare earth salts. It is assumed that the rare earth ions get started to nucleate first and then entrapped in the zirconium hydroxide matrix. In addition, the evolved gas-bubbles act as a nucleation centre and further help to coagulate the nuclei of rare earth ions at a faster rate than the nuclei of zirconium hydroxide. So, after calcination process, the concentration of rare earth ions was found to be higher at core of the zirconia particle than the surface of the zirconia particle. Further, the rare earth based zirconia samples were excited at different excitation wavelengths starting from 205 nm (i.e. near vacuum ultra-violet) to 350 nm (ultra-violet). It was observed that the rare earth based porous zirconia sample shows different colours at different excitation wavelengths. It is known that UV light has a greater penetration depth for phosphors (several micro-meters) [253]. But, it is generally considered that the penetration depth of the near VUV for phosphors is in nanoscales [253]. It is also considered that the near VUV light could only irradiate on the part of the phosphor particle where it is near the surface [253]. On the basis of these view, the activator (i.e. rare earth ions) in zirconia particles is not a homogeneous distribution. So, the rare earth concentration on the zirconia surface and in the interior is different, which causes different change trends in their PL spectra, while excited using near VUV and UV wavelengths. In this work, the existence of rare earth ions are found to be more prominent in the interior of the particle than the surface of the particle. So, the phosphor particles are assume to be core-shell type, where core is activators (rare earth ions) and shell is the matrix (zirconia). Further, the excitation of near VUV wavelength (205 and 210 nm) light could

110

penetrate part of the particle and the penetration depth is up to a certain depth from the surface [253]. The penetration depth of near VUV and UV light on the porous zirconia particle is schematically represented in Fig. 8.28 (a). As, there is no presence of rare earth ions on the surface of the particles (observed from EDX line profile scan), the PL spectra showed only a broad peak starting from 250 nm to 700 nm with a centred peak at 430 nm without any characteristics peak of the rare earth ions. In addition, there was no variation in PL spectra (without any characteristics peaks of rare earth ions), when the sample was excited using the wavelengths of 205 nm to 220 nm. As the excitation wavelength changes from near VUV to UV, the penetration depth of UV light increases. So, with increase in excitation wavelength, the characteristics peaks of rare earth ions appears, when it was excited at a wavelength of around 230 nm. With further increase in excitation wavelength, the penetration depth of UV light increases and the characteristics peak intensity increases with increase in excitation wavelength along with decrease of the broad peak (i.e. ranges from 250 nm to 700 nm). The appearance of the characteristic peaks of rare earth ions at 230 nm was due to the interaction of UV light and the activators. As the penetration depth increases, the concentration of activators increases and thus the characteristic peak intensity increases with increase in excitation wavelength. Further, at higher excitation wavelengths, the characteristic peak intensity decreases due to less interaction of UV light with the activator ions. So, the change in the intensity of the characteristic peak in PL spectra of the sample led to show different colours, when excited using the wavelengths of 230 nm to 350 nm. Different penetration depth of UV light led to show colour tuning in rare earth based porous zirconia particle and it was schematic represented in Fig 8.28 (b).

(a) (b)

Fig. 8.28: Schematic diagram indicating the penetration depth of near VUV and UV light on the porous zirconia particle (a) and different penetration depth of UV light led to show colour tuning in rare earth based porous zirconia particle (b). [Black small circles indicate the rare earth ions within the zirconia particle]

111

8.3 Remarks

Borohydride synthesis via precipitation route using NaBH4 was effectively utilized for synthesizing un-doped and rare-earth doped porous zirconia phosphor nanomaterials.

Both the as-synthesized and t-ZrO2 show a broad PL signal, centered peak at ~430 nm, when excited at a wavelength ranging from near VUV (205 nm and 210 nm) to UV (220 nm, 240 nm to 300 nm). The presence of smaller sized porous particles along with surface defects in un-doped sample led to show violet-blue in colour, as confirmed from CIE diagram as well as the colour obtained under UV lamp source. In addition, the rare earth based porous zirconia sample shows t- and / or c-phase of zirconia. The powders of rare earth based zirconia is porous in nature and the rare earth ions are distributed non-uniformly within the core of the particles. However, there was no indication of rare earth ions on the surface of the particles. So, this lead to develop different colours based phosphor material at different excitation wavelengths. As per the CIE diagram, Eu-based porous zirconia sample shows a colour spanning from bluish to reddish-orange via pinkish region, whereas Tb-based porous zirconia sample shows a colour spanning from bluish to light green via sky blue. However, mixture of Eu and Tb based zirconia sample shows a colour spanning from bluish to reddish- yellowish via nearly white region. This indicates that rare-earth based porous zirconia sample may suitable for lighting applications.

112

CONCLUSIONS

This Ph.D. research work unveiled the importance of borohydride method of synthesis using aqueous sodium borohydride (NaBH4) for the development of porous zirconium oxide nanopowders for suitable applications in different areas of research fields. Based on the results and discussion of this research work, the following major points are focused in the conclusion section.

 Aqueous sodium borohydride (NaBH4) is well known for its reducing property and well- established for the development of metal or metal boride nanoparticles through reduction

method. In contrary, this research work discloses the importance of aqueous NaBH4 as a precipitating agent towards development of porous metal oxide such as zirconium oxide.

 Comparing the powders prepared using aqueous NaBH4 and NH4OH (a common precipitating agent), it was found that the powders are loose in nature for borohydride

route, whereas, agglomerated powders are formed while using NH4OH. In addition, fine pores are present in the calcined zirconia powder sample prepared through borohydride process due to involvement of in-situ gas-bubbles evolved during reaction process.

 The crystallization peak temperature increases from 442 °C (while using NH4OH) to 622

°C while using NaBH4. This shifting nature was further enhanced to still higher temperature by creating numerous gas bubble nucleation centres using constant pH method and it was found that crystallization peak temperature shifted to 717 °C. So, the

as-prepared zirconia remains amorphous up to 500°C and phase pure t-ZrO2 was stable at 800°C.

 It was also proven that aqueous NaBH4 took part in gelation and precipitation for developing boron free zirconium hydroxide powders in the as-synthesized condition.  The gelation and precipitation reaction mechanism indicates that the active boron species help to nucleate zirconium hydroxide nuclei and then participate in gelation process with zirconium hydroxide nuclei through intermolecular hydrogen bonding. In precipitation process, the borate species were phase separated out by forming pieces of borate complex and leaving behind borate free as prepared zirconium hydroxide powders.

 The evolved in-situ hydrogen (H2) gas-bubbles also play an important role to develop as- synthesized loose zirconium hydroxide. The presence of intra-particle voids in the loose zirconium hydroxide help to develop porous zirconium oxide during calcination process.

 Evolution of H2 gas bubbles and the way of borohydride synthesis (gelation, precipitation and constant pH method) strongly affect the crystallite size of zirconia and thus different

polymorphs of ZrO2 are formed at different temperatures. However, the porous nature of zirconia in the form of tetragonal phase was found to be stable up to 800 °C.  The developed thermally stable porous zirconia powder through gas bubble template method was employed for adsorption of toxic ions from industrial waste water. It was

113

found that borohydride derived porous zirconia adsorbed toxic lead ions from aqueous medium almost completely with faster rate (within 15 minutes) at normal pH condition. The adsorption kinetics of borohydride derived porous zirconia powders was analyzed by using five kinetic models (pseudo-first order, pseudo-second order, Elovich, intra- particle diffusion and Bangham’s model). The good fitting with pseudo-second order kinetics indicates that the nature of adsorption was chemisorption type. The good fitting with Elovich model indicates that, the chemisorption type adsorption is due to presence of active sites on the exposed surface of porous zirconia powders. The rapid adsorption behaviour of lead ions was studied by using intra-particle diffusion and Bangham kinetics. the quick adsorption was attributed to the immediate utilization of the most readily available adsorbing sites on the adsorbent surfaces and the adsorption kinetics was limited by pore diffusion. Further, quite high efficiency of five cycles of regeneration suggest the importance of porous zirconium oxide for the removal of toxic ions for environmental application.  The evolution of hydrogen gas bubbles at constant pH method helps to develop agglomerate free loose zirconia powders and along with inborn interconnected pores having mixture of ink bottle-neck (major) and dumb-bell shaped pores (minor) with surface area of 182 m2/g. In addition, the pore morphology of calcined (600 °C) t- zirconium oxide was found to be typically ink bottle-neck (H2-type) pores and it transformed to mixed ink bottle-neck and slit (H3 type) pores at 800 °C. Also, the mean pore diameter of 4.6 nm remains unchanged within a temperature range from 500 °C to 800 °C. Further, the t-phase of zirconium oxide transformed to mixed t- and m- phase with only slit type pores at 900 °C.  Phase stability of t- and / or c- phase of zirconia was found to be stable up to 800 °C, when the rare earth ions [such as single doped (Eu or Tb) and co-doped (Eu and Tb)] are

incorporated in porous zirconia prepared through borohydride method using NaBH4. Further, it was also found that un-doped as well as rare earth based porous zirconia is a strong candidate as a phosphor material for lighting application.  The presence of porous nature, smaller particle size and surface defects led to show a broad PL signal with a centred peak at ~ 430 nm for the un-doped sample of zirconia. From Chromaticity coordinate positions in CIE diagram as well as powder color under UV lamp, it was found that the un-doped porous zirconia nanopowders is violet-blue in color.  The rare earth based zirconia was found to be porous in nature as confirmed from TEM as well as STEM-HAADF image. Also, the distribution of rare earth ions in center of the porous zirconia is prominent than the surface of the particles.

114

 Considering the concept of depth of penetration power of near VUV and UV light, the emission of different colours at different excitation wavelength (205 nm to 350 nm) was observed for rare-earth based porous zirconia sample.  Based on the CIE diagram, Eu-based porous zirconia sample emits a color from bluish to reddish-orange via pinkish region, whereas Tb-based porous zirconia emits a colour from bluish to light green via sky blue with increases in excitation wavelengths. However, mixture of Eu and Tb based zirconia emits a colour from bluish to reddish- yellowish via nearly white region. It was also found that the colour of the powder taken under UV lamp (at 254 nm and 365 nm) was found to be matches with the colour mentioned in CIE-diagram. Finally, the borohydride synthesis method is novel and the developed porous zirconia nanopowders have found potential applications in the field of removal of toxic ions for environmental applications as well as rare earth based porous zirconia nanopowders may be a potential candidate as a phosphor materials for lighting applications. Scope for future work

Based on the results obtained in this Ph.D. work, some of the following points may be taken care and can carry forward as a future work.  Synthesis of other metal oxide systems (such as zinc oxide, aluminium oxide and yttrium boron oxide) using aqueous sodium borohydride and study its gelation and precipitation mechanism. The chloride forms of these metal ions may develop smaller particle size, while using sodium borohydride which can be used in different potential applications in the field of structural, photocatalysis and optical applications.  It may be possible to quantify the hydrogen gas-bubbles released during borohydride synthesis of zirconium oxide. So, study the particle or crystallite size as well as stability of tetragonal zirconium oxide with a function of the released gas-bubbles.  Detail adsorption study of heavy metal ions including effect of pH, time, adsorbent dose and adsorbate dose as well as thermodynamic study.  Modification of pore size, pore volume and surface area of metal oxide systems using various surfactants through borohydride method. The major disadvantage is the stability of pore network at high temperatures. However, based on the literatures, different surfactants such as Pluronic P-123, polyvinyl alcohol (PVA), PEG–PPG– PEG, Hydroxypropyl cellulose (HPC) and agarose can be used to prepare mesoporous and or ordered porous frameworks of various materials.  Study of photoluminescence behavior of porous metal oxide systems with the addition of different concentration of rare earth ions.

115

References 1. T. Yamaguchi, Catal. Today, 20, 199 (1994). 2. R.H.J. Hannink, P.M. Kelly, and B.C. Muddle, J. Am. Ceram. Soc., 83, 461 (2000). 3. M.Z-C. Hu, R.D. Hunt, E.A. Pazant, and C.R. Hubbard, J. Am. Ceram. Soc., 82, 2313 (1999). 4. R.C. Garvie, R.H. Hannink, and R.T. Pascoe, Nature, 258,703 (1975). 5. M.L. Balmer, F.F. Lange, and C.G. Levi, J. Am. Ceram. Soc., 77, 2069 (1994). 6. S. Moreau, M. Gervais, and A. Douy, Solid State Ionics, 101,625 (1997). 7. A. Mondal and S.Ram, Mater Lett., 57,1696 (2003). 8. E.H. Kisi, and C.J. Howard, Key Eng. Mater., 153, 1 (1998). 9. G.D. Yadav, and J.J. Nair, Micropor. Mesopor. Mater., 33, 1 (1999). 10. X-M. Liu, G.Q. Lu,and Z-F. Yan, Appl. Catal. A: General., 279, 241 (2005). 11. S.A. Steiner III, T.F. Baumann, B.C.Bayer, R. Blume, M.A. Worsley, W.J. MoberlyChan, E.L. Shaw, R.Schlogl, A.J. Hart, S. Hofmann,and B. L. Wardle, J. Am. Chem. Soc., 131,12144 (2009). 12. F. Tietz, H.-P. Buchkremer, and D. Stover, Solid State Ionics, 152, 373 (2002). 13. S.H. Kim, G.Y. Jin, M. Kim, and Y.S. Yang, J. Korean Phys. Soc., 61, 980 (2012). 14. C. Bernard, B. Laurence, and V. Franceline, US patent 20020031675. 15. M. Sivakumar, and I. Manjubala, Mater. Lett., 50,199 (2001). 16. R.J. David, and D. Ramgopal, US patent, 20020172838. 17. I.-W. Chen, and L.A. XueJ. Am. Ceram. Soc., 73, 2585 (1990). 18. S.A. Ghom, C. Zamani, S. Nazarpour, T. Andreu, and J.R. Morante, Sens. Actuators, B, 140, 216 (2009). 19. E. Karapetrova, R. Platzer, J.A. Gardner, E. Torne, J.A. Sommers, and W.E. Evenson, J. Am. Ceram. Soc., 84, 65 (2001). 20. J. Guan, R. Doshi, G. Lear, K. Montgomery, E. Ong, and N. Minh, J. Am. Ceram. Soc., 85, 2651 (2002). 21. W.L. Gong, W. Lutze, and R.C. Ewing, J. Nucl. Mater., 277, 239 (2000). 22. M. Wang, S.X. Wang, and R.C. Ewing, Phil. Mag. Lett., 80,341 (2000). 23. C. Renger, P. Kuschel, A. Kristoffersson, B. Clauss, W. Oppermann, and W. Sigmund, J. Ceram. Process. Res., 7, 106 (2006). 24. I. Nukatsuka, and H. Yamane, Adsorpt. Sci. Technol., 29, 1025 (2011). 25. Y-M. Zheng, L. Yu and J. P. Chen, J. Colloid Interface Sci., 367, 362 (2012). 26. Y-M. Zheng, L. Yu, D. Wu and J. P. Chen, Chem. Eng. J., 188, 15 (2012). 27. Y.-S. Ko, and Y.-U. Kwon, ACS Appl. Mater. Interfaces, 5, 3599 (2013). 28. A. Zelcer, and G.J.A.A. Soler-Illia, J. Mater. Chem. C, 1,1359 (2013). 29. A. Mondal, and S. Ram, J. Am. Ceram. Soc., 91, 329 (2008). 30. J. Torralvo, M.A. Alario, and J. Soria, J. Catal., 86,473 (1984). 31. I. Osendi, I.S. Moya, C.I. Serna, and I. Soria, J. Am. Ceram. Soc., 68, 135 (1985). 32. A. Benedetti, G. Fagherazzi, and F.Pinna, J. Am. Ceram. Soc., 72, 467 (1989). 33. J. S. Lee, T. Matsubara, T. Sei, and T. Tsuchiya, J. Mater. Sci. Lett., 32, 5249 (1997). 34. W. Z. Zhu and M. Yan, J. Mater. Sci. Lett., 16, 1540 (1997). 35. R.M. Dickerson, M.V. Swain, and A.H. Heuer, J. Am. Ceram. Soc., 70, 214 (1987). 36. H. Tsubakino, K. Sonoda, and R. Nozato, J. Mater. Sci. Lett., 12, 196 (1993). 37. D.L. Porter, and A.H. Heuer, J. Am. Ceram. Soc., 60, 183 (1977). 38. M.S. Khan, M.S. Islam and D.R. Bates, J. Mater. Chem., 8, 2299 (1998). 39. F. del Monte, W. Larsen, and J.D. Mackenzie, J. Am. Ceram. Soc., 83, 628 (2000). 40. Li.I.-W. Chen, and J.E.P.-Hahn, J. Am. Ceram. Soc., 77, 118 (1994). 41. S.-M. Ho, Mat. Sc. & Eng., 54, 23 (1982).

116

42. R.C. Garvie, J. Phys. Chem., 69, 1238 (1965). 43. E. Ryshkewitch, Oxide Ceramics: Physical Chemistry and Technology; Academic Press, New York, 1960, p. 350. 44. V.S. Nagarajan, and K.J. Rao, J. Mater. Sci., 24, 2140 (1989). 45. P. Nayak, B.B. Nayak, and A. Mondal, Mat. Chem. Phys., 127, 12 (2011). 46. B.B. Nayak, S.K. Mohanty, Md Q.B. Takmeel, D. Pradhan, and A. Mondal, Mater. Lett., 64, 1909 (2010). 47. A. Mondal, and S. Ram, J. Am. Ceram. Soc., 87, 2187 (2004). 48. B. Xia, L. Duan, and Y. Xie, J. Am. Ceram. Soc., 83, 1077 (2000). 49. B. H. Davis, J. Am. Ceram. Soc., 67, C-168 (1984). 50. K. Matsui and M. Ohgai, J. Am. Ceram. Soc., 80, 1949 (1997). 51. G. Štefanić, S. Popović, and S, Musić, Thermochim. Acta, 303, 31(1997). 52. G. N. Glavee, K. J. Klabunde, C. M. Sorensen, and G. C. Hadjipanayis, Langmuir, 9, 162 (1993). 53. B. B. Nayak, S. Vitta, A. K. Nigam, and D. Bahadur, Thin Solid Films, 505, 109 (2006). 54. B. B. Nayak, S. Vitta, and D. Bahadur, J. Mater. Res., 22, 1520 (2007). 55. L. Yu and M. A. Matthews, Int. J. Hydrogen Energy, 36, 7416 (2011). 56. Y. T. Zhu and A. Manthiram, J. Solid State Chem. 110, 187 (1994). 57. A. Manthiram, A. Dhananjay, and Y. T. Zhu, Chem. Mater. 6, 1601 (1994). 58. J. Zhang, T. S. Fisher, J. P. Gore, D. Hazra and P. V. Ramachandran, Int. J. Hydrogen Energ. 31, 2292 (2006). 59. J. R. Groza, (Ed.Koch, C. C.) Ch. 5, 173–217 (William Andrew Inc., 2007). 60. L. B. Kong, L. Y. Zhang and X. Yao, J. Mater. Sci. Lett., 16, 824 (1997). 61. J.H. Ha, J.H. Kim, and D.K. KIM, J. Ceram. Soc. Jpn., 112, 1084(2004). 62. S.H. Jo, P. Muralidharan, and D.K. Kim, Solid State Ionics, 178, 1990 (2008). 63. P. Riachy, F. Roig, , M.J. García‐Celma, M.J. Stébé, A. Pasc,, J. Esquena,, C. Solans and J.L. Blin, Eur. J. Inorg. Chem. 2016, 1989 (2016). 64. S Singh, KC Barick, and D Bahadur, Int. J. Nanosci. 10, 1001 (2011). 65. W. J. Kim, B. K. Min, D. Pradhan, Y. Sohn, Cryst Eng Comm, 17, 1189 (2015). 66. B.B. Srivastava,, S. Jana,, N.S. Karan,, S. Paria,, N.R. Jana,, D.D. Sarma, and N. Pradhan, J. Phys. Chem. Lett., 1,1454 (2010). 67. M Pradhan, P Bhargava, J. Am. Ceram. Soc., 88, 216 (2005). 68. M. Tahmasebpour, A.A. Babaluo, M.K. Razavi Aghjeh, J. Eur. Ceram. Soc. 28, 773 (2008). 69. M. M. Rashad, H. M. Baioumy, J. Mater. Process. Techno., 195, 178 (2008). 70. R. Srinivasan, L. Rice and B. H. Davis, J. Am. Ceram. Soc., 73, 3528 (1990) 71. F. J. Berry, S. J. Skinner, L. N. Bell, R. J. H. Clark and C. B. Ponton, J. Sol. State. Chem., 145, 394 (1999). 72. S. Shukla, S. Seal, R. Vij, S. Bandyopadhyay and Z. Rahman, Nano Lett., 2,989 (2002). 73. M Bhagwat, V Ramaswamy, Mater. Res. Bull., 39, 1627 (2004). 74. M. H. Lee, H. Y. Lin and J. L. Thomas, J. Am. Ceram. Soc., 89, 3624 (2006). 75. J. Liang, X. Jiang, G. Liu, Z. Deng, J. Zhuang, F. Li, Y. Li, Mater. Res. Bull., 38, 161 (2003). 76. N. L. Wu and T. F. Wu, J. Am. Ceram. Soc., 83, 3225 (2000). 77. M. Rezaei, S. M. Alavi, S. Sahebdelfar, Zi-Feng Yan, J Porous Mater, 15,171(2008) 78. H .Zhu, D. Yang, L. Zhu, J. Amer. Ceram. Soc., 90, 1334 (2007). 79. E. A. Sterling, J Stolk, L. Hafford, M. Gross, Metall. Mater. Trans. A, 40, 1701 (2009). 80. Z. Wang, Z. Yan, M. Wang and J. Zhao, Int. J. Electrochem. Sci., 10, 1953 (2015).

117

81. Y.W. Zhao, R. K. Zheng, X. X. Zhang, and John Q. Xiao, IEEE Transaction on Magnetics, 39, 2764 (2003). 82. G. N. Glavee, K. J. Klabunde, C. M. Sorensen and G. C. Hadjipanayis, Inorganic chemistry, 32, 474 (1993). 83. G. N. Glavee, K. J. Klabunde, C. M. Sorensen and G. C. Hadjipanayis, Langmuir, 10, 4726 (1994). 84. J. Legrand, A. Taleb, S. Gota, M.J. Guittet and C. Petit, Langmuir, 18, 4131 (2002). 85. R. Caputo, F. Guzzetta, A. Angerhofer, Inorg. Chem., 49, 8756 (2010). 86. C. Wu, F. Wu, Y. Bai, B. Yi and H. Zhang, Mater. Lett., 59, 1748 (2005). 87. S.W. Heinzmanand B. Ganem,J. Am. Chem. Soc., 104, 6801 (1982). 88. G.N. Glavee, K.J. Klabunde, C.M. Sorensen, and G.C. Hadjapanayis,, Langmuir, 8, 771(1992) 89. W. Ye, H. Zhang, D. Xu, L. Ma and B. Yi, J. Power Sources, 164, 544 (2007). 90. B.B. Nayak, S.Vitta, A.K. Nigam, and D. Bahadur, IEEE Transaction on Magnetics, 41, 3298-3300. (2005), 91. S. Srivastava, A. Mondal, N. K. Sahu, S. K. Behera and B. B. Nayak, RSC Adv., 5, 11009 (2015). 92. S. Srivastava, S. K. Behera and B. B. Nayak, Dalton Trans., 44, 7765 (2015). 93. M. E Davis, Nature, 417, 813 (2002). 94. S.J. Smith, B. Huang, S. Liu, Q. Liu, R.E. Olsen, J. Boerio-Goates and B.F. Woodfield, Nanoscale 7, 144 (2015). 95. L. Kuai, J. Wang, T. Ming, C. Fang, Z. Sun, B. Geng and J. Wang, Sci. Rep. 5, 9923 (2015). 96. R. Liu, Y. Li, C.A Wang and S. Tie, Mater. Des. 63, 1 (2014). 97. W.G. Fahrenholtz, D.M. Smith, and D.W. Hua, J. Non-Cryst. Solids, 144, 45 (1992). 98. X. Yan, N. Lu, B. Fan, J. Bao, D. Pan, M. Wang and R Li, Cryst. Eng. Commun. 17, 6426 (2015). 99. M. P. Albano, L. A. Genova, L. B. Garrido and K. Plucknett, Ceram. Int. 34, 1983 (2008). 100. X. Guo,, G. Hao,, Y. Xie,, W. Cai, and H. Yang, J. Sol-Gel Sci. Technol. 76, 651 (2015); 101. M.A. Alves Rosa, E.P. Santos, C.V. Santilli, S.H. Pulcinelli, J. Non-Cryst. Solids, 354, 4786 (2008). 102. F. Prete, A. Rizzuti, L Esposito, A. Tucci and C. Leonelli, J. Am. Ceram. Soc. 94, 3587 (2011). 103. X-Y.Yang, A. Léonard, A. Lemaire, G. Tian and B.L. Su, Chem. Commun. 47, 2763 (2011). 104. H.L. Jiang and Q. Xu, Chem. Commun. 47, 3351 (2011). 105. J. Yin, X. Qian, J. Yin, M. Shi, J. Zhang and G. Zhou, Inorg. Chem. Commun. 6, 942(2003). 106. X. Liu, G. Lu and Z. Yan, J. Phys. Chem. B 108, 15523 (2004). 107. G-R. Duan, X-J. Yang, G-H. Huang, L-D. Lu and X. Wang, Mater. Lett. 60, 1582 (2006). 108. A. Benedetti, G. Fagherazzi, F. Pinna and S. Polizzi J. Mater. Sci. 25, 1473(1990). 109. V. Bajpai, P. He and L. Dai, Adv. Funct. Mater. 14, 145(2004). 110. M. Mozafari, F. Moztarzadeh, A. M. Seifalian and L. Tayebi,. J. Lumin. 133, 188 (2013). 111. Y. Yan, L. Chen, X. Li, Z. Chen and X. Liu, Polym. Bull. 69, 675 (2012). 112. H. G. Yang, and H. C. Zeng, Angew. Chem. 116, 6056 (2004). 113. H. Wang, H. Yang, L. Lu, Y. Zhou and Y. Wang, Dalton Trans. 42, 8781 (2013).

118

114. Y. Li, W-Z. Jia, Y-Y. Song and X-H. Xia, Chem. Mater. 19, 5758 (2007). 115. Wan, Y. & Zhao, D. Chem. Rev. 107, 2821 (2007). 116. L. Guo, F. Liang, X. Wen, S. Yang, L. He, W. Zheng, C. Chen, and Q. Zhong, Adv. Funct. Mater. 17, 425 (2007). 117. Q. Peng, Y. Dong and Y. Li, Angew. Chem. Int. Ed. 42, 3027 (2003). 118. F. Dong, W. Guo and C.S. Ha, J Nanopart Res.,14, 1303 (2012). 119. F. Gu, C. Zhong Li, S. F. Wang, and M. K. Lu, Langmuir, 22, 1329 (2006). 120. Y. Chen, H. Xia, D. Zhang, Z. Yan, F. Ouyang, X. Xiong, and X. Huang, RSC Adv.,4,8039 (2014). 121. J. Xu, Y. Wang, and Y. Zhu, Langmuir,29, 10566 (2013). 122. J. Suk, D. Y. Kim, D. W. Kim, and Y. Kang, J. Mater. Chem. A 2, 2478 (2014). 123. W. Huang, M. Wang, J. Zheng and Z. Li, J. Phys. Chem. C, 113, 1800 (2009). 124. W.Huang, L. Fu, Y. Yang, S. Hu, C. Li and Z. Li , Electrochem. Solid-State Lett. , 13, K46 (2010). 125. X. Chen, K. Sun, E. Zhang and N. Zhang, RSC Advances 3, 432 (2013). 126. X. Guo, X. Li, Y. Zheng, C. Lai, W. Li, B. Luo and D. Zhang, J. Nanomaterials, 2014, (2014) DOI 10.1155/2014/358312 127. D. Liu, B. B. Garcia, Q. Zhang, Q. Guo, Y. Zhang, S. Sepehri, and G. Cao, Adv. Funct. Mater., 19, 1015 (2009). 128. S. Naumov, Hysteresis Phenomena in Mesoporous Materials, Universität Leipzig, Dissertation, 2009. 129. S.S. Chang, B. Clair, J. Ruelle, J. Beauchêne, F. Di Renzo, F. Quignard, G.J. Zhao, H. Yamamoto, and J. Gril, J. Exp. Bot., 60, 3023 (2009). 130. Weilun Wang, Peng Liu, Ming Zhang, Jiashan Hu, Feng Xing, Open J. Compos. Mater. 2, 104 (2012). 131. K. Sing, D. Everett, R. Haul, L. Moscou, R. Pierotti, J. Rou-querol, and T. Siemieniewska, Pure Appl. Chem. 57, 603 (1985). 132. V.I. Pârvulescu, H. Bonnemannb, V. Pârvulescu, U. Endruschat, A. Rufinska, Ch. W. Lehmann, B. Tesche, G. Poncelet, Appl. Catal. A, 214, 273 (2001). 133. Q Chang, JE Zhou, Y Wang, G Meng, Adv. Powder Technol., 20, 371 (2009). 134. G. Pacheco, E. Zhao, A. Garcia, A. Sklyarov and J. J. Fripiat, Chem. Commun., 103, 491 (1997). 135. U. Ciesla, M. Fröba, G. Stucky, F. Schüth, Chem. Mater. 11, 227 (1999). 136. X.Z. Lin and Z.Y. Yuan, RSC Adv., 4, 32443 (2014). 137. G.S. Armatas, G. Bilis and M. Louloudi, J. Mater. Chem., 21, 2997 (2011). 138. Q. Yuan, L.L. Li, S.L. Lu, H.H. Duan, Z.X. Li, Y.X. Zhu and C.H. Yan, J. Phys. Chem. C, 113, 4117 (2009). 139. H.R. Chen, J.L Shi, Z.L. Hua, M.L. Ruan, D.S. Yan, Mater. Lett., 51, 187 (2001). 140. J. He, J. Chen, L. Ren, Y. Wang, C. Teng, M. Hong, J. Zhao and B. Jiang, ACS Appl. Mater. Interfaces, 6, 2718 (2014). 141. X. Guo, J. Song, Y. Lvlin, K. Nakanishi, K. Kanamori and H. Yang, Sci. Technol. Adv. Mater. 23,867 (2016). 142. X. Yang, X. Song, Y. Wei, W. Wei, L. Hou and Fan, X., J. Nanosci. Nanotechnol. 11, 4056 (2011). 143. D.D. Maksin, A.B. Nastasović, A.D. Milutinović-Nikolić, L.T. Suručić, Z.P. Sandić, R.V. Hercigonja and A.E. Onjia, J. Hazard. Mater., 209,99 (2012). 144. S. Lagergren, Kungliga Svenska Vetenskapsakademies Handlingar, 24, 1, (1898). 145. Y-S. Ho, J. Hazard. Mater. 136, 681 (2006). 146. D. Sparks, Academic Press. New York, 119 (1989). 147. W. J. Weber and J. C. Morris, J. Sanit. Eng. Div. 89, 31 (1963).

119

148. E. Boyd , A. W. Adamson , L. S. Myers. J. Am. Chem. Soc. 69, 2836 (1947). 149. E Tütem, R. Apak,and Ç.F. Ünal.. Water Res. 32, 2315 (1998). 150. M. Sui and L. She Front. Environ. Sci. Eng., 7, 795 (2013). 151. A. M. Herrera, A. A. Martins de Oliveira, A. P. Novaes de Oliveira, and D. Hotza, J. Ceramics, 2013, 1 (2013). 152. Yu D Z, Deng Z S, Zheng M M, Cai R X., Environ. Sci. Technol., 27, 75 (2004). 153. Y. Chang, C. Wang, T. Liang, C. Zhao, X. Luo, T. Guo, J. Gong and H. Wu, RSC Adv., 5, 104629 (2015). 154. Q. Zhang, Q. Du, M. Hua, T. Jiao, F. Gao and B. Pan, Environ. Sci. Technol. 47, 6536 (2013). 155. D. Chen, L. Cao, T. L. Hanley and R. A. Caruso, Adv. Funct. Mater. 22, 1966 (2012), 1966–1971. 156. N. Rahman and U. Haseen, Ind. Eng. Chem. Res. 53, 8198 (2014). 157. Y. Jia, Y. Zhang, R. Wang, J. Yi, X. Feng, and Q. Xu, Ind. Eng. Chem. Res., 51, 12266 (2012). 158. Y. Zhu, T. Shimizu, T. Kitajima, K. Morisato, N. Moitra, N. Brun, T. Kiyomura, K. Kanamori, K. Takeda, H. Kurata, M. Tafu and K. Nakanishi, New J.Chem., 39, 2444 (2015). 159. J. Luo, X. Luo, J. Crittenden, J. Qu, Y. Bai, Y. Peng and J. Li, Environ. Sci. Technol. 49, 11115 (2015). 160. L. M. Sharygin, M. L. Kalyagina, and O. L. Borovkova, Russ. J. Appl. Chem., 82, 815 (2009). 161. L. A. Rodrigues, L. J. Maschio, R. E. da Silva, M. L. C. P. da Silva, J. Hazard. Mater., 173, 630 (2010). 162. T. M. Suzuki, J. O. Bomani, H. Matsunaga, T. Yokoyama, React. Funct. Polym., 43, 165 (2000). 163. A. Bortun, M. Bortun, J. Pardini, S. A. Khainakov, J. R. Garcı´a, Mater. Res. Bull., 45 (2010). 164. Y. Su, H. Cui, Q. Li, S. Gao and J. K. Shang, Water Res., 47, 5018 (2013). 165. D. Manoharan, A. Loganathan, V. Kurapati and V.J., Ultrason. Sonochem., 23 174 (2015). 166. S. Kumar and A. K. Ojha , J. Alloys Compd., 644, 654 (2015). 167. A. Emeline, G.V. Kataeva, A.S. Litke, A.V. Rudakova, V.K. Ryabchuk, and N. Serpone, Langmuir, 14, 5011 (1998). 168. X. Xu and X. Wang, Nano Res., 2, 891 (2009). 169. L. Kumari, G. Du, W. Li, R. S. Vennila, S. Saxena and D. Wang, Ceram. Int., 35, 2401 (2009). 170. A. Mondal, A. Zachariah, P. Nayak and B. B. Nayak, J. Am. Ceram. Soc., 93, 387 (2010). 171. M. Mokhtar, S. N. Basahel and T. T. Ali, J. Mater. Sci., 48 2705 (2013). 172. Y. Cong, B. Li, S. Yue, D. Fan and X.-j. Wang, J. Phys. Chem. C, 113, 13974 (2009). 173. F. Trivinho-Strixino, F. E. GuimarÃ£es and E. C. Pereira, Chem. Phys. Lett., 461, 82 (2008). 174. F. Davar, A. Hassankhani and M. R. Loghman-Estarki, Ceram. Int., 39, 2933 (2013). 175. R Espinoza-González, E Mosquera, Í Moglia, R Villarroel, V. M. Fuenzalida, Ceram. Int., 40, 15577 (2014). 176. H. Jiang, R. I. Gomez-Abal, P. Rinke and M. Scheffler, Phys. Rev. B, 81, 085119 (2010). 177. R. French, S. Glass, F. Ohuchi, Y.-N. Xu and W. Ching, Phys. Rev. B, 49, 5133 (1994).

120

178. D. W. McComb, Phys. Rev. B, 54, 7094, (1996) 179. S. Miyazaki, M. Narasaki, M. Ogasawara and M. Hirose, Microelectron. Eng, 59, 373 (2001). 180. K. D. Sattler, Handbook of Nanophysics: Nanoparticles and Quantum Dots, CRC Press, 2010. 181. T. H. Gfroerer, Encyclopedia of Analytical Chemistry, 2000. 182. Y. Cong, B. Li, B. Lei and W. Li, J. Lumin., 126, 822 (2007). 183. J. Joo, T. Yu, Y. W. Kim, H. M. Park, F. Wu, J. Z. Zhang and T. Hyeon, J. Am. Chem. Soc.,125, 6553 (2003). 184. A. K. Singh, and Umesh T. Nakate. Sci. World J., 2014 (2014). 185. J. Liang, Z. Deng, X. Jiang, F. Li and Y. Li, Inorg. Chem., 41, 3602 (2002). 186. L. Kumari, W. Z. Li, J. M. Xu, R. M. Leblanc, D. Z. Wang, Yi Li, H. Guo, and J. Zhang, Cryst. Growth Des., 9, 3874 (2009). 187. C. Lin, C. Zhang and J. Lin, J. Phys. Chem. C, 111, 3300 (2007). 188. H. Che, S. Han, W. Hou, A. Liu, S. Wang, Y. Sun, X. Cui, J. Porous Mater., 18, 57 (2011). 189. L.N. Nagy,, J. Mihály,, A. Polyák,, B. Debreczeni,, B. Császár,, I.C. Szigyártó,, A. Wacha,, Z. Czégény,, E. Jakab,, S. Klébert, and E. Drotár,,. J. Mater. Chem. B, 3, 7529 (2015). 190. H.Q. Cao, X.Q. Qiu, B. Luo, Y. Liang, Y.H. Zhang, R.Q. Tan, M.J. Zhao and Q.M. Zhu, Adv. Funct. Mater. 14, 243 (2004). 191. G.K. Sidhu, A.K. Kaushik, S. Rana, S. Bhansali and R. Kumar, Appl. Surf. Sci., 334, 216 (2015). 192. C. Zhang, C. Li, J. Yang, Z. Cheng, Z. Hou, Y. Fan and J. Lin, Zhang, Langmuir, 25, 7078 (2009): 7078-7083. 193. G. Chu, J. Feng, Y. Wang, X. Zhang, Y. Xu and H. Zhang, Dalton Trans., 43, 15321 (2014). 194. C. L. Hardin, Y. Kodera, S. A. Basun, D. R. Evans, and J. E. Garay, Opt. Mater Express, 3, 893 (2013). 195. X. Bai, A. Pucci, V.T Freitas, R. A Ferreira, N. Pinna, Adv. Funct. Mater., 22, 4275 (2012). 196. S. López-Romero, M. García-Hipólito, and A. Aguilar-Castillo. World J. Cond. Matt. Phys., 3, 173 (2013). 197. S. Mukherjee, D. P. Dutta, N. Manoj and A. K. Tyagi, J. Nanopart. Res., 14, 1 (2012). 198. S. D. Meetei and S.D. Singh, J. Alloys Compd., 587, 143 (2014). 199. S. D. Meetei, S.D. Singh and V. Sudarsan, J. Alloys Compd., 514, 174 (2012). 200. B. K. Moon, I.M. Kwon, J.H. Jeong, C.S. Kim, S.S. Yi, P. S. Kim, H. Choi, J. H. Kim, J. Lumin., 122, 855 (2007). 201. B.D. Cullity, Elements of X-ray diffraction, addition-Wesley publishing company, INC. reading massachusetts, United States of America (1956). 202. H. Näfe, N. Karpukhina, J. Am. Ceram. Soc., 90, 1597 (2007). 203. P. Canton, G. Fagherazzi, R. Frattini, P. Riello, J. Appl. Crystallogr., 32 475 (1999). 204. G. Fagherazzi, P. Canton, A. Benedetti, F. Pinna, G. Mariotto and E. Zanghellini,, J. Mater. Res., 12 318-321 (1997). 205. A. Esmaeilifar, S. Rowshanzamir, A. Behbahani, Iran. J. Hydrog. Fuel Cell, 3, 163 (2014). 206. A. Rizzuti, A. Corradi, C. Leonelli, R. Rosa, R. Pielaszek and W. Lojkowski, J. Nanopart. Res., 12, 327 (2010). 207. S.N. Bramhe, Y.P. Lee, T.D. Nguyen and T.N. Kim J. Mater. Res., 23, 441 (2013). 208. C. G. Salentine, Inorg. Chem., 22, 3920 (1983).

121

209. J. Schott, J. Kretzschmar, M. Acker, S. Eidner, M.U. Kumke, B. Drobot, A. Barkleit, S. Taut, V. Brendler and T. Stumpf, Dalton Trans., 43, 11516 (2014). 210. C. Su and D. L. Suarez, Environ. Sci. Technol., 29, 302-311 (1995). 211. D. Peak, G. W. Luther and D. L. Sparks, Geochim. Cosmochim. Acta., 67, 2551 (2003). 212. D. L. Griscom, Ed., Pye, L. D., Fréchette, V. D. & Kreidl, N. J. 11-138 (Plenum Press, 1978). 213. U. Manna and S. Patil, J. Phys. Chem. B 113, 9137 (2009). 214. C. Gautam, A. K. Yadav and A. K. Singh, ISRN Ceram., 2012, 1 (2012). 215. W. Senior and W. Thompson, Nature, 205, 170 (1965). 216. C. R. Cloutier, A. Alfantazi and E. Gyenge, J. Fuel Cell Sci. Technol., 4, 88 (2007). 217. Y. Jia, S. Gao,Y. Jing, Y. Zhou, and S. Xia, Chem. Pap., 55, 162 (2001). 218. J. R. Sohn and E. S. Cho, Appl. Catal. A, 282,147 (2005). 219. G. K. Chuah, Catal. Today, 49, 131(1999). 220. W. Stichert and F. Schuth, Chem. Mater., 10 2020 (1998). 221. J. L. Falconer and J. A. Schwarz, Catal. Rev., 25, 141(1983). 222. P. Hu, L. Yu, A. Zuo, C. Guo, and F. Yuan, J. Phys. Chem. C, 113, 900 (2009). 223. R. Wang, Y. Ma, H. Wang, J. Key and S. Ji, Chem. Commun. 50, 12877 (2014). 224. X. Wang, Q. Peng and Y. Li, Acc. Chem. Res., 40, 635 (2007). 225. H. Qiu, et al. Critical review in adsorption kinetic models. J. Zhejiang Univ. Sci. A, 10, 716–724 (2009). 226. Y.-S. Ho, J. Hazard. Mater., 136, 681 (2006). 227. H. Tahermansouri, Z. Dehghan and F. Kiani, RSC Adv., 5, 44263 (2015). 228. R.S. Juang and M.L. Chen, Ind. Engg. Chem. Res., 36, 813 (1997). 229. M. S. Rahma, and K. V. Sathasivam, Res. Int. 2015, 1 (2015). 230. C. Chakrapani, C. Babu, K. Vani and K. S. Rao, J. Chem., 7, 419 (2010). 231. A Daifullah, S. Yakout and S. Elreefy, J. Hazard. Mater., 147, 633 (2007). 232. S. Shukla, S. Seal, R. Vij and S. Bandyopadhyay, Nano Lett., 3, 397 (2003). 233. M. Thommes, B. Smarsly, M. Groenewolt, P. I. Ravikovitch and A. V. Neimark, Langmuir, 22, 756 (2006). 234. A. Grosman, and C. Ortega, Langmuir 24, 3977 (2008). 235. P. T. Nguyen, C. Fan, D. Do and D. Nicholson, J. Phys. Chem. C, 117, 5475 (2013). 236. B. Coasne, A. Galarneau, R. J. Pellenq and F Di Renzo, Chem. Soc. Rev., 42, 4141 (2014). 237. A. Bumajdad, M. I. Zaki, J. Eastoe and L. Pasupulety, J. Colloid Interface Sci., 302, 501(2006). 238. Y Ren, Z Ma, RE Morris, Z Liu, F Jiao, S Dai and P.G. Bruce, Nat. Commun., 4, 2015 (2013). 239. D. Wallacher, N. Künzner, D. Kovalev, N. Knorr and K. Knorr, Phys. Rev. Lett., 92, 195704 (2004). 240. M. Hasan, M.C Nguyen, H. Kim, S.W. You, Y.S. Jeon, D.T. Tong, D.H. Lee, J.K. Jeong and R. Choi, Thin Solid Films, 589, 90(2015). 241. Q.X. Gao, X.-F. Wang, X. C. Wu, Y.R. Tao and J.-J. Zhu, Micropor. Mesopor. Mater., 143, 333 (2011). 242. V. Valtchev, and L. Tosheva, Chem. Rev.,113, 6734 (2013). 243. V. Hornebecq, C. Knöfel, P. B. Kuchta, and P. L. Llewellyn, J. Phys. Chem. C, 115, 10097 (2011). 244. R. Dwivedi, A. Maurya, A. Verma, R. Prasad and K. Bartwal, J. Alloys Compd., 509, 6848 (2011).

122

245. F Kazemi, A Saberi, S Malek-Ahmadi, S Sohrabi, H.R. Rezaie, and M., Tahriri,. Ceram.-Silik. 55, 26 (2011). 246. H. Eltejaei, J. Towfighi, H. R. Bozorgzadeh, M. R. Omidkhah, and A. Zamaniyan, Mater. Lett., 65, 2913 (2011). 247. H. Cui, Q. Li, S. Gao, and J. K. Shang, J. Ind. Eng. Chem., 18, 1418 (2012). 248. X. Ma, L. Klosterman, Y.Y. Hu, X. Liu and K. Schmidt‐Rohr, J. Am. Ceram. Soc., 95, 3455 (2012). 249. H. R. Chen, J. L. Shi, J. H. Gao, L. Li and D. S. Yan, Solid State Phenom., 121, 5 (2007). 250. S. Shukla, and S. Seal, Int. Mater. Rev. 50, 45 (2005). 251. J. H. Nobbs, Rev. Prog. Color. Relat. Top., 15, 66 (1985). 252. M. Mikhailov and A. Verevkin, Russ Phys J., 47, 600 (2004). 253. Q. Dong, Y. Wang, Z. Wang, X. Yu and B. Liu, J. Phys. Chem. C, 114, 9245 (2010). 254. M. N. Tahir, L. Gorgishvili, J. Li, T. Gorelik, U. Kolb, L. Nasdala and W. Tremel, Solid State Sci., 9, 1105 (2007). 255. N. C. S. Selvam, A. Manikandan, L. J. Kennedy and J. J. Vijaya, J. Colloid Interface Sci., 389 (2013). 256. S.G. Chen, Y.S. Yin, D. P. Wang and J. Li, J. Cryst. Growth, 267, 100 (2004). 257. Z. Wang, B. Yang, Z. Fu, W. Dong, Y. Yang and W. Liu, Appl. Phys. A, 81 691 (2005). 258. H. R. Chen, J. L. Shi, Y. Yang, Y. S. Li, D. S. Yan and C.-S. Shi, Appl. Phys. Lett., 81, 2761 (2002). 259. C. L. Hardin, Y. Kodera, S. A. Basun, D. R. Evans, J. E. Garay, Opti. Mater. Expres. 3, 893 (2013). 260. K. Cao, Q. Zhu, B. shen and R. Chen, Scientific Reports, 5, 8470 ( 2015)

123

CURRICULUM VITAE Nadiya Bihary Nayak Research Scholar (Ph.D.) Department of Ceramic Engineering National Institute of Technology Rourkela – 769008, Odisha, INDIA E-mail: [email protected] Mobile No.: +91 8763366093 Research Area: Porous, Nanopowder synthesis, Adsorption, Luminescent materials Educational Qualification:

Education Institute/University/Council/Board Year of Division/Grade Passing Ph.D. Department of Ceramic Engineering, (Submitted, - National Institute of Technology Rourkela 2016) M. Tech (Research) Department of Metallurgy and Material 2010 First Sciences, NIT Rourkela M. Sc. (Chemistry) Department of Chemistry, National 2003 First Institute of technology Rourkela B. Sc. (Chemistry Utkal Universiy, Bhubaneswar Odisha 2001 First Hons) 12th Council of Higher Secondary Education, 1998 Second Bhubaneswar, Odisha 10th Board of Seconday Education, Cuttack, 1996 First Odisha Publications based on Ph.D. work: 1. Nadiya B. Nayak, Bibhuti B. Nayak, A. Mondal, “Enhanced activation energy of crystallization of pure zirconia nanopowders prepared by an efficient way of synthesis using NaBH4”, Journal of the American Ceramic Society, 96, 3366 (2013). 2. Nadiya B. Nayak, Bibhuti B. Nayak, Aqueous sodium borohydride induced thermally stable porous zirconium oxide for quick removal of lead ions, Scientific Reports 6, 23175 (2016) [Nature publishing Group]. 3. Nadiya B. Nayak, Bibhuti B. Nayak, Temperature-mediated phase transformation, pore geometry and pore hysteresis transformation of borohydride derived in-born porous zirconium hydroxide nanopowders, Scientific Reports 6, 26404 (2016) [Nature publishing Group]. 4. Nadiya B. Nayak, Bibhuti B. Nayak, “Violet-blue luminescence in borohydride derived un-doped porous zirconia” (to be submitted). 5. Nadiya B. Nayak, Bibhuti B. Nayak, “Multi-colour luminescence in borohydride derived rare earth ions doped porous zirconia” (to be submitted). Other publications: 1. S. C. Mishra, Nadiya B. Nayak, “An investigation of dielectric properties of chicken feather reinforced epoxy matrix composite”, Journal of Reinforced Plastics and Composites 29, 2691 (2010). 2. S. C. Mishra, Nadiya B. Nayak, A. Satapathy, “Investigation on bio-waste reinforced epoxy composites”, Journal of Reinforced Plastics and Composites 29, 3016 (2010). 3. S. C. Mishra, Nadiya B. Nayak, “Wear response prediction of TiO2-polyester composites using neural networks”, International Journal of Plastics Technology, 14, S24 (2010). 4. Bibhuti B. Nayak, Nadiya B. Nayak, Rahul K. Mallik, Aparna Mondal, “Synthesis and magnetic properties of cobalt ferrite with different morphologies”, International Journal of Modern Physics: Conference Series, 22, 164 (2013). Personal Information: Father’s Name: Mr. Mahesh Prasad Nayak Mother’s Name: Mrs. Sushama Nayak Date of Birth: 07th July 1981 Sex: Male Marital Status: Married Religion: Hindu Nationality: Indian Languages Known: Odia, Hindi and English Permanent Address: Vill/PO: Gadapokhari, Dist: Cuttack, Odisha

124

125