Material Properties of RE- Doped Ln (Ln= Y, La) 
Oxides and Oxysulfides Phosphors For Red-
Emitting Devices 
By 
Abdub Guyo Ali 
(MSc) 
A thesis submitted in fulfilment of the requirements f or the degree 
Doctor of philosophy 
in the 
Faculty of Natural and Agricultural Sciences 
Department of Physics 
University of the Free State 
Republic of South Africa 
Promoter: Prof. F.B. Dejene 
Co-Promoter: Prof. H.C. Swart 
November 2015 
Dedication 
This thesis is dedicated to my Son Guyo 
Acknowledgements 
First and foremost, l express my heartfelt gratitude to T he Almighty God, for granting me 
the opportuni ty to pursue thi s study. I thank him also for enabl ing me to complete my studies 
successfully. I bow before Almighty God for giving me the strength and courage to pursue 
this study. 
l would like to express my heartfelt gratitude to my two advisors Prof. Francis Birhanu 
Dejene and Prof. Hendrik C. Swart, for giving me the opportunity to work in their research 
group, and for their guidance, support and encouragement. The two of them have been very 
valuable in the development of my PhD investigation and they made this interdiscipl inary 
project an exciting adventure. l would like to thank Prof. F.B. Dejene for making me feel 
from the beginning part of his group and also for introducing me to the fi eld of Material 
Sciences. I would particularly like to express my indebtedness to my co-supervisor Prof. H. 
C. Swart for his guidance and encouragement during the entire course of my studies. He has 
taught me a lot on the writing skills of which he is very good. I would like to acknowledge 
the moral support from Prof. JJ . Terblans, Prof. W.D. Roos, Prof. P.J. Meintjes, Prof. M.J.H 
Hoffman, Prof. T. Kroon and Mr. D.P. van Jaarsveldt. 
I would like to thank members of staff of the Department of Physics, University of the Free 
State for the positive interactions and support: Dr. Koao L.F, Dr. Tshabalala K.G, Dr. 
Motloug S.V, Mr. Ocaya R.O, Mr. Motloung S.J, Mrs Cronje, K.The late Dr. Doto, J.J. Ms 
Lebeko, K.M. Pro f. Mothudi, B.M, to mention but a few. Fellow researchers: Mr Wako, A.H, 
Dr Coetsee, E. Pro f. Dlamini , M.S. Dr. Duvenhage, M.M, Miss Foka, K. E, Miss Lephoto, 
M.A, Mr Ungula J, Mrs Jattani (Sharon), Miss Tebele A, Miss Mulwa W.M, Dr Roro, R. 
(DST), Mr Malevu T.D, Miss Tshabalala, M.A. among others. Prof Van Wyk, P.W.J. and 
Janecke, B. of the Centre of Microscopy for their support and advice during SEM 
measurements. My stay at University of the Free State and in particular Qwaqwa was an 
enjoyable journey due to many of my fiiends. I am greatly indebted to the Afiican Laser 
Center (ALC), National Research Foundation (NRF) and University of the Free State for their 
financial support. I would like to extend my special gratitude to the staff of the National 
Laser Center (NLC) for their valuable support during my visits to the N LC for my 
experimental work. 
i i 
Abstract 
Structural and optical properties of Eu3+-doped Ln (Ln=Y, La) oxide and oxysulfide nano 
crystals synthesized by sol-combustion method were analysed as a function of host to fuel 
ratio. Structural characterization shows crystallite nanosized particles and the hexagonal 
phase as the dominant structure. The red emission of Eu3+ doped Y 20 2S, La20 2S and Y 20 3 
nanocrystals appearing near 624 nm was assigned to the 5Do-7F2 transition of Eu3+. Due to 
insufficient quantities of thiourea at the higher Ln/S mole ratio, the bright red emission has 
been quenched. Fourier-transform infrared spectrometry analysis showed that there was a 
negligible difference in the absorbed impurities with various molar ratios. The Ln/S 
concentration also affects the decay time of the red emission of the Eu3+ ions from 140 µs for 
Ln/S= I to 76 µs for the higher concentrations. Structural and optical properties of 
La20 2S:Eu
3+micro crystals synthesized by sol-combustion method were analyzed as a 
function of La/S concentration. Structural characterization shows a crystallite size o f about 
178 nm and the hexagonal phase as the dominant crystalline structure. The red emission of 
Eu3+ doped La20 2S microcrystals appearing near 624 nm was assigned to the 
5D 70- F2 
transition of Eu3+. Due to insufficient quantities of thiourea at the higher La/S mole ratio, the 
bright red emission has been quenched. Fourier-transform infrared spectrometry analysis 
showed that there was a negligible difference in the absorbed impurities with various molar 
ratios. The La/S concentration also affects the decay time of the red emission of the Eu3+ ions 
from 140 µs for La/S= I to 7 6 µs for the higher concentrations. 
To investigate the effect of co-doping a series of red-emitting phosphors Y 20 3:Eu
3+:Ho3+ 
were prepared by the solution combustion method. X-ray diffraction (XRD) patterns indicate 
that the Eu3+ and Ho3+ doping do not show obvious effect on the cubic Y 20 3 crystal. Their 
crystall ite size estimated by x-ray diffractometry and scanning electron microscopy was 
about 8 nm. Under UV 325 nm excitation, emission wavelengths at 626 nm was quenched at 
higher mole percent of Ho3+ and energy was transferred from Eu3 ~ to Ho3+. Y20 3:Eu3+: Ho
3+ 
phosphor shows a red-emitting afterglow phenomenon, and the Eu3' ions are the luminescent 
center during the decay process. The bright red emission near 626 nm has been noticeable 
due to the 5D 70- F2 transition of Eu
3 
... . The intensity of the luminescence has decreased with an 
increase of concentration of Ho3+. In sufficient quantities of Eu3+ to Ho3+, the bright red 
emission near 626 nm has been predominant due to 5D 70- F2 transition of Eu
3+. The decay 
iii 
characteristic of Y 20 :Eu
3 Ho33 +: + phosphor is according with the double exponential 
equation. 
The as-prepared powder Y 320 2S:Eu + was deposited on Si ( 100) substrates by using a pulsed 
laser deposition technique. The thin films grown under different oxygen deposition pressure 
conditions have been characterized using structural and luminescent measurements. The X-
ray diffraction patterns showed mixed phases of cubic and hexagonal crystal structures. As 
the oxygen partial pressure increased, the crystallinity of the films improved. Further increase 
of the 0 2 pressure to 140 mtorr reduced the crystallinity of the fi lm. Similarly, both scanning 
electron microscopy and atomic force microscopy confirmed that an increase in 0 2 pressure 
affected the morphology of the films. The average band gap of the films calculated from 
diffuse reflectance spectra using the Kubeika-Munk function was about 4.75 eV. The 
photoluminescence measurements indicated red emission of Y20 2S:Eu
3
+ thin fi lms with the 
most intense peak appearing at 619 nm, which is assigned to the 5Do-7 F2 transition of Eu3+ . 
This most intense peak was totally quenched at higher 0 2 pressures. X-ray photoelectron 
(XPS) indicated that Y20 3 thin films are formed on the surfaces of the Y20 2S: Eu
3
+ thin films 
during prolonged electron bombardment. The films grown in a lower 0 2 ambient consist of 
smaller but more densely packet particles relative to the films grown at a higher 0 2 ambient. 
In order to study the effect of annealing temperature on the films, four samples were annealed 
at various temperatures whi le one was kept unannealed. X-ray diffraction measurements 
show that the un-annealed thin film was amorphous, while those annealed were crystalline. 
At lower annealing temperature of 600 °c to 700 °c cubic bixbyite Y20 3:Eu3+ was formed . 
As the annealing temperatures were increased to 800 °c, hexagonal phase emerged. The 
average crystall ite size of the fi lm was 64 nm. Photoluminescence (PL) measurement 
indicates intense red emission around 612 nm due to the 500~ 7F2 transition. Scanning 
electron microscopy (SEM) indicated that agglomerates of non-crystalline particles with 
spherical shapes were present for the un-annealed films. After annealing at high temperature, 
finer morphology was revealed. Atomic fo rce microscopy (AFM) further confi rmed the 
formation of new morphology at the higher annealing temperatures. UV-vis measurement 
indicated a band gap in the range of 4.6 to 4.8 eY. It was concluded that the annealing 
temperature played an important role in the luminescence intensity and crystallinity of these 
films. 
To investigate the effect of different species of gases Y20 S:Eu
3
2 + thin films have been grown 
on Si ( I 00) substrates by using a pulsed laser deposition technique. The thin films grown 
iv 
under different species of gases have been characterized using structural and luminescent 
measurements. The X-ray diffraction patterns showed mixed phases of cubic and hexagonal 
crystal structures. The crystallinity of the film deposited in vacuum is poor, but improved 
significantly in argon and oxygen atmosphere. Similarly, both scanning electron microscopy 
and atomic force microscopy confirmed that different species of gases affected the 
morphology of the fi lms. The average band gap of the films calculated from diffuse 
reflectance spectra using the Kubeika-Munk function was about 4.69 eV. The 
photoluminescence measurements indicated red emission of Y 0 S:Eu32 2 + thin films with the 
most intense peak appearing at 6 12 nm, which is assigned to the 5D 70- F2 transition of Eu
3
+. 
The intensities of thi s most intense peak greatly depend on the species of gas with argon 
having the highest peak. This phosphor has applications in the flat panel displays. 
v 
Key words 
Y20 2S: Eu3+ 
Y20 3: Eu3+ 
La20 2S: Eu3 3+, La20 3: Eu + 
Y202S: Eu3+:Ho3+ 
Solution - Combustion Method 
Morphology 
Excitation 
Band gap 
Luminescence 
Rare Earth Ions 
PLO 
Laser ablation 
decay time 
red-emitting 
vi 
Acronyms and symbols 
• PL- Photoluminescence 
• XPS -X-ray photoelectron spectroscopy 
• XRD -X-ray diffraction 
• HRTEM - High resolution Transmission electron microscopy 
• SEM- Scanning electron microscopy 
• EDS -Energy dispersive spectroscopy 
• PLO -Pulsed laser deposition 
• AFM- Atomic force microscopy 
• FTIR - Fourier-Transform infrared 
• He-Cd- Helium Cadmium 
• RE- Rare earth 
• KrF-Krypton fluoride 
• Y- Ylttrium 
• La- Lanthanum 
• Al- Aluminium 
• 02- Oxygen molecule 
• 0- Oxygen atom 
• VB- Valence band 
• CB- Conduction band 
• VO- Oxygen vacancy 
• CRTs- Cathode Ray Tubes 
• LPP-Long Persistent Phosphors, 
• TEM- Transition Electron Microscopy 
• TL-Thermoluminescence 
vii 
List of figures 
Figure. 2.1. Atomic representations of La20 2S along the a) < 11 O> and c) <00 I> 
directions. Six possible anion vacancies are noted (i.e. A 1, A2, BI , 82, 
CI, C2, with uppercase letter indicating the corresponding anion layer). 
Atomic representations of La20 3 along the b) <11 O> and d) <00 I> 
directions---------------------------------------------------------------------------9 
Figure. 2.2. a) Schematic illustration for self-assembled Na-doped La20 2S 
nanoplates with OA as capping agents; the orange box highlighted in a) is 
enlarged in b), which shows the thickness of one nanoplate, indicating the 
three layers of primitive cells along c-axis with La3+ as ending ions on 
both sides of the nanoplates.----------------------------------------------------10 
Figure 2.3. Nanomaterials with a variety of morphologies-------------------------------1 2 
Figure 2.4. Schematic illustration of the preparative methods of nanoparticles 
--------------------------------------------------------------------------------------1 2 
Figure 2.5. Model showing persistant luminescence mechanism ------------------- 16 
Figure 3.1: Bruker 0 8 Advance model x-ray diffractometer ----------------------------29 
Figure 3.2: Schematic representation of a S EM ------------- ------------------------------3 0 
Figure 3.3: Shimadzu Superscan SSX-550 model Scanning Electron Microscope 
---------------------------------------------------------------------------------------3 l 
Figure 3.4: The cavity structure of He-Cd Laser ------------------------------------------32 
Figure 3.5: Cary Eel i pse Florescence S pectrophotometer--------------------------------3 3 
Figure 3.6(a, b): Schematic illustra tion of common recombinatio n processes -------------34 
Figure 3.8: Schematic diagram of a pulsed laser deposition chamber setup-----------36 
Figure 3.9: Examples of picture of plume developed during PLO [ 45] ----------------37 
Figure 3.10: 248 nm KrF Lambda Physic excimer laser with PLD setup----------------39 
vii i 
Figure 3.11 : Schematic diagram of the XPS process in copper----------------------------40 
Figure 3.12: PHI 5400 Versaprobe scanning x-ray photoelectron spectrometer--------40 
Figure 3.13 (a): Schematic diagram of a transmission electron microscope [ 18] ----------42 
Figure 3.13 (b): JEOL JEM-2 100 model transmission electron microscopy-----------------43 
Figure 3.14: CIE chromaticity chart-----------------------------------------------------------44 
Figure 3.15: Visible light spectrum and corresponding wavelengths---------------------45 
Figure 4.1: Representative XRD pattern of one of the sample with S!Y= 1.8 molar 
ratios obtained by Sol- Combustion method----------------------------------51 
Figure 4.2 : The SEM images of the Y20 3: Eu3+ with (a) l.9 (b) 2.0 (c) 2.5 (d) 4.0 S!Y 
molar ratios. 0.5 run field of view----------------------------------------------53 
Figure 4.3. (a): Emission spectrum of the different S!Y molar ratio Y20 3: Eu
3
+ phosphor 
excited at 260nm obtained by the Sol-Combustion method. (b) CIE 
coordinate diagram of the different emissions as indicated-----------------54 
Figure 4.4: The decay curve of Y 20 3:Eu
3
+ phosphor--------------------------------------56 
Figure 4.5: Effect of S!Y molar ratios on the intensity of the broad PL peaks and 
corresponding emission wavelength-------------------------------------------5 7 
Figure 5.1. X-ray diffraction patterns of La20 2S with different La/S ratios as well as 
the standard XRD pattern--------------------------------------------------------63 
Figure 5.2. X-ray di ffraction powder patterns at (*) plane for different La/S mole 
ratios------------------------------------------------------------------------------ 64 
Figure 5.3. X-ray diffraction powder patterns at (101) plane for different La/S mole 
ratios-------------------------------------------------------------------------------64 
Figure 5.4. The effect of fuel on the formation of La20 2S and La20 3 prepared by the 
sol- combustion process---------------------------------------------------------65 
Figure 5.5. Fourier-transform infra-red spectroscopy spectra of the as-prepared 
La 320 2 S: Eu +p owders for various La/S mole ratios--------------------------67 
Figure 5.6 (a). XPS survey spectrum of the La20 2S microcrystals prepared with a La/S 
ratio of 2.5-------------------------------------------------------------------------68 
ix 
Figure 5.6 (b). XPS survey spectrum of the La20 2S microcrystals prepared with a 
La/S ratio of 1. 0-------------------------------------------------------------------69 
Figure 5.6 (c). La 3d XPS peakfitted with peaks for the La20 2S and the La203 
--------------------------------------------------------------------------------------70 
.Figure 5.6 (d ). XPS spectra of La (3d5/i) and 0 (I s) of the as-prepared Eu
3+- doped 
La20 2S microcrystals.La 3d region for La202S and La20 3 with the peak 
fitting components for the 4ps12 peak------------------------------------------70 
Figure 5.6 (e). XPS S2p peak for La20 2S with the peak fi tting components for the S2ps12 
peak---------------------------------------------------------------------------------71 
Figure 5.7. SEM micrographs of the as-prepared La20 2S: Eu
3+powders with La/S 
molar ratios of (a) 1.0, (b) l.8, (c)2.0, (d)2.5 with 5000 nm field of view 
--------------------------------------------------------------------------------------72 
Figure 5.8 (a). Excitation spectra of La20 2S with different La/S molar ratios-------------73 
Figure 5.8 (b). PL emission spectra of La20 2S with La/S= l.8, l.9, 2.0, 2.5 and 3.0 molar 
ratios--------------------------------------------------------------------------------73 
Figure 5.8 (c). CIE coordinate of emission of La20 2S phosphor----------------------------75 
Figure 5.9. Graph of maximum peak intensity versus La/S molar ratios---------------76 
Figure 5.10. Afterglow characteristics of La20 2Swith different La/ S molar ratios.Inset: 
A graph of In log of intensity versus decay time showing a double 
exponential function--------------------------------------------------------------77 
Figure 5.11. Thermoluminescence plots of the La 320 2S: Eu + phosphor------------------79 
Figure 6.1. X-ray diffraction patterns of films deposited in vacuum and various 0 2 
partial pressures and the standards JCPDS card Nos: 24- 1424 and 22-
0993--------------------------------------------------------------------------------86 
Figure 6.2. Crystallite sizes and axial ratio as functions of oxygen partial pressure 
--------------------------------------------------------------------------------------87 
Figure 6.3: SEM images of the thin films ablated in a) vacuum, b) 20 mtorr, c) 60 
mtorr and c)l40 mtorr 0 2 ambient at 300 °C with a fluence of 0.767 ± 0. 1 
Jcm-2 (5 kV beam energy, magnification of x 20 000 and a scale of 1 µ m 
(FOY: 2 x Iµ m). As insets: 30 Height AFM images done in contact 
x 
mode for the thin fi !ms ablated in a) Vacuum, b) 20 mtorr and c) l 40 
m torr oxygen am bi ent------------------------------------------------------------8 8 
Figure 6.4. Excitation spectra for fi lms deposited in vacuum and at different oxygen 
partial pressure. The inset show excitation spectrum of Y20 2S:Eu
3
+ 
powder phosphor-----------------------------------------------------------------9 l 
Figure 6.5. Emission spectra for films deposited in vacuum and at different oxygen 
partial pressure. The inset show emission spectrum of Y 320 2S:Eu + powder 
phosphor---------------------------------------------------------------------------92 
Figure 6.6. The plot of maximum peak intensity versus oxygen partial pressure-----93 
Figure 6.7: Decay curves for the thin films deposited in vacuum atmosphere and at 
different oxygen am bi ent-------------------------------------------------------94 
Figure 6.8. UV-vis diffuse reflectance spectra of nanocrystalline Y 320 2S:Eu + thin fi lm 
deposited in vacuum and different oxygen pressure-------------------------96 
Figure 6.9. (a) Graph of F[(R)*hv]2 as a function of band gap energy, (b) 
Dependance of band gap energy on partial oxygen pressure---------------96 
Figure 7.1 . X-ray diffract ion pattern of Y 20 3: Eu3+: Ho3+ phosphor------------------- 103 
Figure 7.2. Crystallite sizes and Lattice constant as funct ions of oxygen partial 
pressure--------------------------------------------------------------------------- 105 
Figure 7.3. SEM micrographs of Y 320 3: Eu +: Ho
3
+ samples with (a) O. l (b) 0.2 (c) 0.3 
( d) 0.4 ( e) 0.5 % of Ho3+ ions. 4. 77 µm field of view-----------------------106 
Figure 7.4 Excitation spectra of Ho3+ co-doped Y 20 3: Eu
3
+ phosphor when Ho3+ ion 
concentration was vari ed from O. l to 0.5%---------------------------------- 107 
Figure 7.5. Photoluminescence emission spectra of Ho3+ co-doped Y20 3 : Eu
3
+ 
phosphor when Ho3 1 ion concentration was varied from 0.1 to 0.5% 
-------------------------------------------------------------------------------------108 
Figure 7.6. Chromaticity colour coordinates of the Y 3 320 3:Eu ... :Ho + powder under 325 
nm UV excit ation--------------------------------------------------------------- 109 
Figure 7.7. Decay curves of Ho3+ co-doped Y 20 33: Eu + phosphor when Ho
3+ ion 
concentration was varied from 0. 1 to 0.5%---------------------------------- 1 10 
xi 
Figure 7.8. Concentration of Ho3+ ions vs. maximum peak intensity graph of 
Y 20 3:Eu3 3+ :Ho + phosphor------------------------------------------------------111 
Figure 7.9. Uv-vis absorbance spectra of Y 3 320 3: Eu +: Ho + red- emitting phosphor 
with% mole concentration of Ho3+ from O. l to 0.5%---------------------113 
Figure 7.10. Graph of F [( R)*hv] 2 as a function of band gap energy------------------114 
Figure 7 .11. Band gap energy a~ a function of Ho3+ mole concentration--------------- I 15 
Figure 8.1. X-ray diffraction pattern of un-annealed and annealed Y20 3: Eu3+ thin 
film deposited on a (100) Si substrate after firing at temperatures between 
600 and 900 °c in air for 2 hours-------------------------------------------122 
Figure 8.2. The crystallite sizes and lattice parameters as a function of temperature 
-------------------------------------------------------------------------------------123 
Figure 8.3. SEM micrographs of (a) un-annealed and annealed samples (b) 600 (c) 
800 and (d ) 900 °c-------------------------------------------------------------- l 25 
Figure 8.4. 30 Height AFM images done in contact mode for the thin films which are 
(a) Un-anneal ed----------------------------------------------------------------126 
(b) annealed at 600°C---------------------------------------------------------- l 27 
( c) annealed at 9000C---------------------------------------------------------128 
Figure 8.5. Diffuse reflectance measurements for un-annealed and those annealed at 
different temperatures for Y 20 3: Eu
3
+t hin films-----------------------------129 
Figure 8.6. Graph of F[(R)*hv] 2 as a function of band gap----------------------------1 30 
Figure 8.7. The excitation spectrum of Y20
3
3: Eu + thin films for un-annealed and 
those annealed at 600, 700, 800 and 900 °c---------------------------------132 
Figure 8.8. The emission spectrum of Y2 0 3: Eu3+ thin films for un-annealed and those 
annealed at 600, 700, 800 and 900 °C---------------------------------------- l 33 
Figure 8.9. The CIE co-ordinates for samples that were un-annealed and those 
annealed at various temperatures---------------------------------------------134 
xii 
Figure 8.10. Showing decay characteristics of Y203: Eu3+ phosphor thin fi lms for un-
annealed and films annealed at 600, 700, 800 and 900 oC ----------------135 
Figure 9.1. The XRD spectra of the Y 20 2S: Eu
3
+ thin fi lms deposited in vacuum and 
different gas atmospheres------------------------------------------------------144 
Figure 9.2. SEM micrographs for thin fi lms deposited in (a) vacuum, (b) argon and 
(c ) oxygen atmosphere---------------------------------------------------------1 53 
Figure 9.3. AFM images of thin films deposited in (a) vacuum------------------------ 145 
AFM images of thin films deposited in (b) argon--------------------------146 
AFM images of thin films deposited in ( c) oxygen------------------------ 14 7 
Figure 9.4(a). Excitation spectra for Y 20 2S: Eu
3+ thin fi lms deposited in vacuum, argon 
and oxygen atmosphere. Inset: Excitation spectrum for sample deposited 
in oxygen atmosphere---------------------------------------------------------- 148 
Figure 9.4(b). Deconvoluted excitation spectra for Y 20 2S: Eu
3
+ thin films deposited in 
vacuum, argon and oxygen atmosphere-------------------------------------- 149 
Figure 9.4(c). Excitation spectra of Y 0 S:Eu32 2 + thin film deposited in vacuum recorded 
at 590, 6 12, 625 and 655 nm emission wavelengths-----------------------150 
Figure 9.4(d). Excitation spectra of Y 20 2S:Eu
3
+ thin film deposited in argon recorded 
at 590, 6 12, 625 and 655 nm emission wavelengths----------------------- 151 
Figure 9.4(e). Excitation spectra of Y 0 S: Eu32 2 + thin film deposited in oxygen recorded 
at 590 6 12 625 and 655 nm emission wavelengths-----------------------151 
' ' 
Figure 9.5 (a). Emission spectra for Y2 0 S: Eu
3
2 + thin films deposited in vacuum, argon 
and oxygen atmosphere. Inset. Emission spectrum for thin film deposited 
in oxygen------------------------------------------------------------------------- I 5 3 
Figure 9.S(b). Emission spectra fo r Y 0 S: Eu32 2 + thin films deposited in vacuum, argon 
and oxygen atmosphere. Inset. Emission spectrum for thin film deposited 
in argon---------------------------------------------------------------------------1 54 
xi ii 
Figure 9.S(c). Emission spectra of Y 20
3
2S:Eu + thin fi lm deposited in vacuum 
atmosphere recorded at 237, 245, 260 and 290 nm excitation wavelengths 
-------------------------------------------------------------------------------------1 5 5 
Figure 9.S(d). Emission spectra of Y 20 2S:Eu
3
+ thin film deposited in oxygen 
atmosphere recorded at 237, 245, 260 and 290 nm excitation wavelengths 
------------------------------------------------------------------------------------- 156 
Figure 9.6. The decay curves for PLO Y20 2S: Eu
3
+ thin films deposited in vacuum 
and different gas atmospheres-------------------------------------------------158 
Figure 9.7. UV-vis diffuse reflectance spectra of nanocrystalline Y2 0 2S: Eu
3
+ thin film 
deposited in vacuum, argon and oxygen atmospheres---------------------159 
Figure 9.8. Graph of F[ (R)*hv] 112 as a function of band gap energy------------------1 60 
xiv 
List of tables 
Table 4.1 The average grain size as a function of SN molar ratio--------------------52 
Table 4.2 Decay constants for the fitted decay curves of the phosphor powders with 
different SN mo Jar ratios-------------------------------------------------------56 
Table 5.1. The concentration and calculated crystall ine size of Eu3+ ion doped 
La20 2S  mi crocrys ta!s ------------------------------------------------------------66 
Table 5.2. Results for the fitted decay curves of the phosphor--------------------------67 
Table 5.3. Trap energy levels for different concentration of La20 2S-------------------80 
Table 6.1: Showing how oxygen partial pressures affect lattice parameters and 
particle size of the films---------------------------------------------------------87 
Table 6.2 : Decay constants for the fitted decay curves of the thin fi lms ablated in 
vacuum and various oxygen partial ambient----------------------------------94 
Table 7. 1. The crystallite sizes as a function of % concentration of H03+ ions-----! 04 
Table 7.2 : Decay constants for the fitted decay curves of the Y 20 3 : Eu
3+:Ho3+ 
powder with various mole concentration------------------------------------ 11 9 
Table 8.1. Showing lattice parameters and crystalline sizes of Y 20 3: Eu
3+ thin fi lms 
-------------------------------------------------------------------------------------1 24 
Table 8.2. Results fo r the fi tted decay curves of the un-annealed and films annealed 
at different temperatures-------------------------------------------------------1 36 
Table 9.1 Decay constants for fi tted decay curves of the films ablated in vacuum, 
argon and oxygen atmospheres----------------------------------------------157 
xv 
.=....---------------
Table of Contents 
Dedication ..... .... ..... ......................................... ..... ............. ..... .... .......................... ... ................ ................. i 
Acknowledgements ................ ..... ................................................. .... .......... .. .... .............. .. ....... ..... ........... ii 
Key words .. ... .... ........ ............. ...... ........ ................... .. ... ... .......... .... ..... .. .......... .. ..... ... ....... ........ ............... vi 
Acronyms and symbol s .......... ...................... ...................................................... .... .. .... ..... .. .... .............. vii 
L isr of figures .... .. ................... .. ... .... ....... .. ......... .... .... ....................................................... .................... viii 
L ist o f tables ....... .. ............. ..... ......... ............ ... ..... ... .. ... ..... .. ...... ..... ......................................... ... .... ... ..... xv 
Chapter I .............................................. ................. ......... .. ... .. ............... .. ........ .. ....... .... ....................... .. ... 1 
Introduction ........................................................... .... .. .... ... .. .... ..... .. ....... ...... ... .... ........... ....... ............. ..... l 
1. 1. Background .......... ........ .. ................................................................... ... ... .... ...... ...... ... ... ... ........ .... 1 
1.2 An overview of past Phosphor research ....... ... .......... ........ ........... ......... .... ................ .. ......... ......... 2 
1.3 Statement of the problem ........ ............... ........ ........................................... ................. ... ...... .. .. .. .... 3 
1.4 Research objectives .............. ............................................................................... ..... .. ................... 4 
1.5 Thesis layout .. .............................................................. ................................................................. 4 
References ....................................................................................................................................... .... 6 
Chapter 2 ............... ........ ................... ...... ....... .... ................................................... .... .. ... .................. .. ...... 7 
Theory ..................................................................................................................................................... 7 
2.1 An overview of phosphors ......... ... .. ....... ........................................................ ....... ... ......... ............ 7 
2.2 Fluorescence ........................ .................. .. .......... ........ ........ ......... .... .......... ............... ............ ..... ... . 8 
2.3 Phosphorescence ..... ......................... .................... ....................... ..... ...... ....................................... 8 
2.4 Properties and Applications of Nanomateria ls ..... .. .... ................. ....... ...... ... .............. .. .. ................ 8 
2.4.1 Some Properries o f Nanomaterials ......... .. ........... ................. .......... .......... .... ... .... .... ... ...... ..... . 8 
2.4.2 Nanomaterial - symhesis and proct!ssing ............. ................................................... ............. 12 
2AA Applicat ions of Nanometer-sized Y20 1: Eu.1 .. ..... ............. .. ................... ....... ............... ... .... 13 I 
2.4.5. Mechanism of the Persistent Luminescence .. .. .... ............. .. ... ..... ...... .......... ..... .......... .... ..... 15 I 
2.4.6. The Luminescent Centre ............................. .............. ..... .. ...... ............. ...... ....... .. ...... ........ ... 16 I 
2.4.7. Phase Transfonnation .... ...................... .... .... ..... ..... ............... ....... .... .. ...... ........................... 17 I 
2.4.8 . Effect of Lattice Det~c t s on Persistent Luminescence .................................. ...... ... ............. 17 I 
2.4.9. Energy Transport and torage in Luminescent solids .... ..................... .... .. .... ...... ................ 18 
I 
2.5.1: Emission and Excitation 1'v1echanis111s of Phospliors ... ....... .. ... .............. ... ..... .. .... .... ........ ... 18 
2.5.2: General ( \rnsiclerntion , - Fluorescent La111p-. .... ... ..... ....... .. ........ ... ............... .. .. ... ............. 19 I 
xvi 
--------·-......-. 
2.5.3 Genera l Considerations - Cathode Ray Tub..:s ...... .... ........ ................ ............ ...... .... .. ...... .. ... 19 
2.5.4. Emission and Excitation Nkchanisrns o f Phosphors ........... .. .. .... .......... ...... .... ................ ... 20 
2.5.5. Luminescence Mechanisms ...... ....................... ... ....... ...... ....... .............. ... ...... ..................... 20 
2.5.6. Center Luminescence ... ... ............ ....... ... ........ ............... ..... .................................................. 21 
2.5.7. Charge Transfer L uminesce1m.: .. .. .............. .... ... .. ....... ...... .................................... ............... 21 
~ . 5.8 . Donor Acceptor PGir Lum in..: ~ c..:11..:..: .................................... .. .... .. ............ .. ...... ................... 21 
2.6.1. Mechanisms Underlying Energy Transfer ........ .. .................... .. .... ........ .. .. .... .......... .. ...... .... 22 
2.6.3. Exchange interaction bet ween sensiti zer and activator ion .... ............... ..... .. .. ................ .. ... 22 
............. ... .. .. .... .......................... .... ... ........ ..... ...... ..... .... .............. ....... .... .... ............. .... ........ ... ....... . 22 
2.6.4 The Energy does not reach the Luminescent ion ..................... ... .. .......................... ............. 23 
References .................. ..... ........ ... .... ......... ....... ..... .. ... ... ... ... ........ ............. .. ...... .. ... .... .... .. .. .... ............ .. 23 
Chapter 3 ........ ..... ..... ... ....................... ...................... ........ ....... ................................................. ...... ....... 26 
3. Experimental Techniques ... .. .. ................ ......................................... ..... ........................ .... .......... ... .... 26 
3.1. lntroduction ..... .. ............. ...... ... .. ...... ...................... .... ......... .... ... .. ....................................... ... ..... 26 
3.2. Synthesis and deposition technique ............................. ........ ... ................................................... 26 
3.2. l. Sol- combustion method ........ ........ .............. .............. .. ............... .. .. ......... .. ...... .. ........... ...... 27 
3.3 .2 . Scanning electron microscopy (SEM) .. ......... .......... .. ................ ........... ........ ................. .. .. . 28 
3.3.3. Photoluminescence spectroscopy ( Helium cadmium laser) ............. .. ................................. 30 
3.3.4. Radiative recombination mechanisms observed in PL .................... ....... ........ .. .................. 32 
3.3.4. Rad iative recombination mechanisms observed in PL .... .. ................................................. 33 
3.3.5 Pu lsed laser deposit ion (PLD) ............................. ............................................. .. ............ .. .... 3S 
3.3.6 X-Ray Photoe lectron Spectroscopy (XPS) ............................................. .......... .. ......... ....... 39 
3.3.7 Atomic Force Microscopy (AFM) .......................... .. ........................ .. ....... ....... ....... .... ........... 41 
3.3.8 Transmission electron microscopy (TEM) ................... ....... ... ..................... ..... ... .. ............... 41 
3.4 Evaluation of Phosphor ................ ...... ............... ...... .... ....... ....... .. ...... ...... .. ............................... ... 43 
3.4.1 Chro1nat ici ty ...... ... .......... ... ...... ....... ...... ......... ... ...... ... ............... ...... ............. ...................... .. 43 
3.4.2 Spectral Distribution .. ....... ... .. ............ ....... .. ......... ............... .. .. .. ........................................... 4S 
Retcrences .. .......... .... ........................ .. ................. .. ..... .. ... ...... ... .... ............ .... ... ................ ... ...... ... .... ..... . 46 
4.1 lntroduct ion ................... ..... ............... .. ..... ........... ............ .. .... ............ .. ... ... .. ..... ... ......... ... ................ 49 
4.2. Experimental detai ls. ......... .................... ... .................. ..... .... ..... ........ .. .... ......... ..... ... ..... .... ..... ..... SO 
-1.2.1 yn tht:sis proc<:dure ... .................... .. ............ ......... ....... ... .... .. .......... .. ........... ...... .... ...... .... .... SO 
4.2.2 Characterization .... ...... ....... ............... ... .. .... ...... ..... ..... ..... .. .... .... ....................................... ... . SO 
4.3. Results and di scussion ............. ... ....... ...... ...... .. ........ .................................... .. .... .. .. ..... ... ...... ...... SO 
4.J.I Crystal structure ........... ....... ............. ....... ..... ... ....... ........ .... ............. ... ...... ... .. ... .................... 50 
xvi i 
I 
4.3.2 tvlorphology ..................................................................................................... ..................... 52 
4.3.3. Photoluminescence ............................................................................................................ .. 53 
4.3.4 . Afterglow decay curves of the red phosphors .............................................. ....... ... .... ......... 55 
4.4 Conclusion .......................................................... ............ .......... ........ .............................................. 57 
References ................................................................................................... .. ...... ... ........... ... ...... ...... ... .. 58 
Chapt..:r 5 .................. ................................ ... ............... ... ... .... ... ... .... .. .. .... ...... ........... .. .... .... .. ...... .... .. ... ... 60 
Characterization of Eu3Tactivated lanthanum oxysu l tide syn thesized by sol-combustion method ....... 60 
5. I. Int rod uction ..................................................... ... ....................................................................... 60 
5.2 Experimental ............................. ...................................... ............................. ................. .............. 61 
5 .2.1 Characteri zation ............................. .. ..................... ... ........ ... ................. .... .......... .... .. ............ 61 
5.3 Results and D iscussion ........... .................. .. .................. ....................................... .. ..... ................ 62 
5.3. I Crystal structure ............................................. ...................................................................... 62 
5.3.2. Fourier transforms infrared spectroscopy .............. ............................................................. 66 
5.3.3. X-ray photoelectron spectroscopy ...... .......... ... .. .......... ...... ................. ................................ 67 
5.3.4 Morphology ................................. ......... ............. ............................................ .......... ............. 71 
5.3.5 Photoluminescence ..... .. ........................................ .. ....... ...... ........... ................ .......... .......... .. 72 
5.3.6 Thermoluminescence ......................................... ......................................... ......................... 78 
5.4 Conclusions .... ....... ........................................................... ....... ....... ................ ........... ... ............... 80 
References ................................................................. ... .. ..... .............. ..... ............................. .. ... ......... 80 
6.1 lnt roduction .......................... ....................... ............. .......................... ....................................... .. 83 
6.2 Experimental procedure .......... ............. ..... ... ............... .. .. ....................... .. .... ................... .. ... ... .... 84 
6.2. I Powder synihesis ........................................ .... ... .. ..... .. .. .................. ............... .. ... .. ................ 84 
6.2.2 Pulsed Laser Deposition (PLO) .............. .. .. ..... .. .................................................................. 84 
6.3 Resul ts and discussion .... .... ................................ .. ... .. ................................................................. 85 
6.3. I X-ray diffraction ............................................ .... ......... .. ........ ...... ............ ..... ........................ 85 
6.3 .2 Morphology ..... .. ...................................... .......... ... .. ........ ....... .... ............ .. ... ....... .... .. ... ... ...... . 88 
6.3.3 Atomic ForceM icroscopy(A FM) ..... ............... ...... ....................... .... .............................. .... 89 
6.3.4. Photoluminescence spectra .............. ......... ... .... .. ....... ... ........................ .. ................. .. ...... .... 89 
6.3.6 Optical properties ............... ............................... .. ................ .......................................... ...... . 94 
6.4 Conc lusion ..... ...................................................... .. ..... ... ....... ......... ........................................... 97 
References. .. ................... .. ................................................ .. .. ................................................................. 97 
Chapter 7 .. ......... ............................................. ................................ ..................................................... 100 
Energy transfer and material properties of Y 203: Eu:i' :H o3 • rwnophosphors synthesized by so l-
combust ion method ................................................ ........... .. ........... .... .. ... .................................... ........ 100 
xvi ii 
7. 1 lntroduction ............................................ .......... .......... ..... ..... .... ... .................................... .. ............ 100 
7 .2 Experimental .... ..... ................. .. ... ..... .. ................... .. ... ....... .... .. .... .............................................. 101 
7.2. I Nanocrystal synthesis ....... .. .. .... ... ...... ... .... ...... ....... .. .... ...... ....... ..... ... ..... ... ........ .. ..... ... ........ 101 
7.2.2 Characterization ................ ........ .. ... .......... ....... .. ..... ..... ..................... ..... ............. ................ 101 
7.3 Results and discussions ................ ......... ... ................ .................. .. ... ..... ............... ....... .... ...... .. .. . 102 
7.3 . I X -ray d iffraction study ......... ..... ........ ... .......................... .... .... .. ..... .. ...... .. ....... .. ....... .. ... ... .. . 102 
7.3.2 Scann ing elect ron microscopy ........ .. ........... .. ... .... .... .. .. .. .... ... ........... ....... .... ..... ... .... .. .... .. .. 105 
7.3 Photoluminescence ............................................................ ............................................... ..... ... 107 
7.3. 1 Exci tation .... .. .............. ...... ..... ................. ..... ........ ........ ... ............................................ .. ..... 107 
7.3.2 Emission .................................................................... ........... .......... ........... ... ........... ........... 107 
7.3.3 Decay charac teristics ................................ .... ......... ..... ... ... ... ...... ... .. ........... .... ...... .. .. .. ....... .. 109 
7.3.4 Optical properties ............. ......... ............ ... ...... .. ...... ...... .. .... ... ..... ........ ....... .......... ....... .... ... . 111 
7.4.2 Determination of Eg. from reflectance spectra .................. .... .. .... ....... ...... .... .... .. .. .. ....... ... 112 
7.4 Concl usion .................. ..... ...................................... ............... ............... ... ........ ......................... ..... 115 
Referenccs ...... ...................................... ..... .... ........... .............. ... .... ...................... .. ... ....... ....... .. .. ......... 116 
C hapter 8 ........ ..... .. .. ............. .............. ......... ........ .. ...................................... ........................................ 118 
Temperature dependence of structural and luminescence properties ofEu3 ' -doped Yi OJ red- emitting 
phosphor thin ti I ms by Pu lsed Laser Deposi tion ... ......... ............ .. .. ......... .... ..... ..... .. .. .. ....... ........ ... ..... 118 
8.1 lntroduction .... .. .. ...... .............................................................. ........ ..... ... .. ... ......... ... ................ ...... 118 
8.2 Experimental decails ............ .... .... ..... ..... ................... .... ........................ .. ...... ............ ................. 119 
8.2.1 Powder synthesis ............................................ .... ....... ..... ... .... .. ............ ... .................. .......... 119 
8.2.2 Pulsed Laser Deposition (PLO) .. ......... .. ....... .. ..... ............. ... ..................... ... ............ .......... 120 
8.2.3 Characterization ........................... ...... .... .. ... .......... .... .. ........ .... ... ... ... .... .............. ................ 120 
8.3 Results and di scussions ............... ... ...... ... ... ... ...... ......... ... .... .... ... ... .. ..... .. .... .. ....... ...... .... ..... ....... 120 
8.3. I Structural and morphological ana lysis ........ .. .. ...... ... .... ..... ... .... ........ ... .. .. .. .. .................. .... . 120 
8.3 .2 Optical properti es ... ........... .......... ....... .. .... .. .... .. .. .. ...... .... .... ............. ...... ... ... ................. ...... 128 
8.3.4.3 Decay curve ....... ............... ...... ...... .... ............ .... .... ... ...... ...... .................. .. .......... ...... ..... .. 134 
8.4 Conclusion ..... ......... .. .................... .............................. ......... .. ...... ... ...... .............................. ...... 136 
References ...... ............ .. ..... ... .............................. .. ............ ............................ ....... .......... .... .................. 136 
Chapter 9 ......... ......... ....... ....... ... .. ........................................ ..... .. ... ............. .. ....................................... 139 
The influence of different species of gases on the luminescent and structural propert ies of pulsed laser 
ab lated Y20 2S:Eu3T thin films ............................ .. ................. .. .... .... .... .............. .. .................... ...... ...... 139 
9.1 lntroduction ... .. ....... ... ............. ...... ....... .... .... .. ..... ..... ..... ... .......... .... .......... ....... ....... ....... ................. 139 
9.2 Experimental procedures ....... ..... ...... .... ......... .. ............ ..... .. .... ............. ...... ..... .......................... .... 140 
9.2.1 Powder preparati on ................. ........ .... .... ............... .............. ...... ...... ... ...... ... ....... .... ....... .... 140 
xix 
9.2.2 Pu lst:d lasi.:: r dt:position (PL 0 ) ................ ...... .. .......... .. .. .... .. .. ........... .. ..... .. ...... .. .................. 141 
9.2.3 Characteri zation ........ ............ .... .... ...... ......... .... .. ...... .. ......... .. ... .. .... .. ... .......... ........ ............. 141 
9.3 Results and discussion .. ........ ... ............ ..................... ...... .................. ........................................ 142 
9.3. 1 X -ray diffraction analys is ................................. .. ....... ...... .... ...... ... .. ... ............... ... ..... .. ........ 142 
9.3.2 Scanning electron microscopy (SEM) .............. .. ......... ... ..... .. .. ... ... ...... ................... .. ......... 143 
9.3.3 Atomic force microscopy (AF1'vlJ ..................... .. ................................................... ....... ..... 144 
9.3.4 Photoluminescence resul ts ....... .. ... .................. ................. ........................ ..... ............. .. ...... 148 
9.3 .5 Decay curves ...... .. .. ... .............................. ..... ...... .. ........ ............................ .......... ...... .. .. ...... 156 
9.3 .6 Optical properties ... ...... ....... .. ... .. ...................... ........................ ................ ......................... 158 
9.4 Conclusion ................ ............. ............. ..................... .. ................... ... ......... ..................... ... ........ 161 
Reference ... ......... .............. ..... .. ........ ..... ... .... .. .. .... ...... .. ... ... .... ... ........... ....... ... .... .. ..... ... ...... .... .. ..... ...... 161 
Chapter I 0 ...................... ...... .. .. .......................................... ..................... ...... ........................... ... ..... ... 165 
Summary and suggestions for future work ....... .. ... ..... .. .... .................................................................. 165 
I 0. 1 Thesis summary .... ............... .. ................ ................... .............. .. .... ... .. .................... ............... .. 165 
I 0.2 Suggestions for future work .......................... .. ........ .. .................. .. .......................................... 167 
Pub licati ons .. .. ........................ ... ............ ... .......................... .. ....................................... .... ...... .............. 167 
Papers presented at con ferences .. .. ........... .......................... .. ........... ............. ............. ....... .... ....... ...... .. 169 
A ppendix: ... ...... ................... .... ................. ...... .................... ... .. ....... .... .......... ........... .. ......... .... .... .... ..... 170 
xx 
Chapter 1 
Introduction 
1.1. Background 
Optical materials have a broad range of applications in a variety of aspects of human life. 
Among those are medicine, military, communications, computing, manufacturing, and 
various industrial applications. Rapid progress of nanotechnology opens new opportunities in 
designing optical materials with improved optical properties. Current research in 
nanotechnology is focused on new materials, novel phenomena, new characterization 
technique and fabrication of nano devices. Y 20 2S:Eu
3
+ and La20 2S:Eu3-.- are excellent 
materials of current interest [ 1-3] owing to their interesting optical and opto-electronic 
properties. The crystal structure of M20 2S (M = Y, La and including all lanthanides) are 
discussed in detail [3-5). The crystal symmetry of the above two systems is hexagonal, with 
the space group P3m 1 (D33d), as determined by X-ray diffraction. These systems are 
grouped under wide band gap (4.6 - 4.8 eV) semiconductors. Y20 2S:Eu3+, Y20 3:Eu
3
+ and 
La20 2S:Eu
3
.,. as red- emitting phosphors, with its sharp emission line for good calorimetric 
definition and high luminescence efficiency, is extensively used in the phosphor screen of 
display devices, fluorescent lamps used for lighting purposes, television sets used fo r 
entertainment and information gathering, X-ray imaging instruments used in hospitals and 
laser instruments used for experimental purposes and, many other electrical and opto-
electronic equipment. They employ luminescent materials for [6, 7) electronic portal imaging 
devices (EPID), radioisotope distribution and so on [8). Due to the large size and weight of 
CRTs, developments of flat-panel displays (FPOs) are of great interest. Among several FPO 
technologies, liquid-crystal displays (LCDs) dominate the FPO market and plasma di splay 
panels (PDPs) are now commercially available in the large area TV market [9-11). New and 
enhanced properties are expected due to size confinement in nanoscale dimensions that can 
revolutioni ze the display devices market in future. Commercially available bulk oxysulfides 
are quite expensive and are not easily available. So, for the time being, Y 320 2S:Eu + and 
1 
La 0 S: Eu32 2 + nanostructures are relatively a good choice whi le compared with the bulk 
systems. However, fo r an extensive use in the commercial applications, Y20 2S: Eu3+, Y20 3: 
Eu3+ and La20 2S:Eu
3+ nanocrystals must be prepared at lower temperatures. Therefore, it is 
necessary to develop a low-temperature synthesis technology for the growth of both oxide 
and oxysulfide nanophosphors. In this background, this chapter has been devoted to the 
nanophosphors development using these two systems. The realm of novel devices from this 
wonderful material is yet to be accomplished in ful l. To give a quantitative report on the state 
of art o f Y20 2S:Eu
3+ Y20
3
3: Eu + and La20 2S:Eu
3
+ is quite diffi cult and an attempt has been 
made to give an account of the synthesis of the nanophosphors in this thesis. 
1.2 An overview of past Phosphor research 
The scientific research on phosphors has a long history going back more than I 00 years. A 
prototype of the ZnS-type phosphors, an important class of phosphors for television tubes, 
was first prepared by Theodore Sidot, a young French chemist, in 1866 rather accidentally. It 
seems that this marked the beginning of scientific research and synthesis of phosphors [ 12]. 
From the late 19th century to the earl y 20th century, Phi lip E.A. Lenard and coworkers in 
Germany performed active and extensive research on phosphors, and achieved impressive 
results. They prepared various kinds of phosphors based on alkali ne earth chalcogenides 
(sulfides and selenides) and zinc sulfide, and investigated their luminescence properties. They 
established the principle that phosphors of these compounds are synthesized by introducing 
metallic impurities into the materials by firing. Lenard and co-workers tested not only heavy 
metal ions but various rare-earth ions as potenti al activators. P. W. Pohl and co-workers in 
Germany investigated Tl+-activated alkali halide phosphors in detail in the late 1920s and 
1930s. They grew single-crystal phosphors and performed extensive spectroscopic studies. In 
co-operation with F. Seitz in the U.S. they introduced the configurational co-ordinate model 
of luminescence centres and established the basis of present-day luminescence physics. 
Humbolt Leverenz and co-workers at Radio Corporation of America (U.S.) also investi gated 
many practical phosphors with the purpose o f obtai ning materials with desirable 
characteristics to be used in television tubes. Detailed studies were perfo rmed on ZnS type 
phosphors. Since the end of World War II , research on phosphors and solid-state 
luminescence has evolved dramatically. This has been supported by progress in solidstate 
physics, especially semiconductor and latti ce defect physics. Advances in the understanding 
2 
of the optical spectroscopy of solids, especially that of transition metal ions in general and 
rare-earth ions in particular, have also helped in these developments. The concept of the 
configurational coordinate model of luminescence centres was established theoretically. 
Spectral shapes of luminescence bands were explained on the basis of this model. The theory 
of excitation energy transfer successfully interpreted the phenomenon of sensitized 
luminescence. Opti cal spectroscopy of transi ti on metal ions in crystals clarified their energy 
levels and luminescence transition on the basis of crystal field theory. In the case of trivalent 
rare-earth ions in crystals, precise optical spectroscopy measurements made possible the 
assignment of complicated energy levels and various luminescence transitions. 
1.3 Statement of the problem 
A lot of research has been devoted to luminescence of nano and microsized sulphide 
phosphors since they are used in many display applications including cathode ray tube and 
field emission display. A mechanism that shows the relationship between their luminescence 
and surface chemical reactions has been established. Since these phosphors are not very 
efficient at low voltages required for field emission displays, micro-sized and nanoparticle 
oxide phosphors are being investigated to replace them [ 13]. Several effective red-emitters 
phosphors, for example La20 S:Eu
3
2 +, Y2~S:Eu3+ and Gd 0 S: Tb32 2 ~, have been investigated 
for their luminescent properties of both powders and thin films, because of their better 
emission properties as compared to their counterparts, vanadate. A proper way to evaluate 
these phosphors for application displays would be to study their luminescent properties. It is 
important to determine the mechanism that shows the correlation between their PL intensities 
and changes on the surface chemical composition during electron beam exposure. It is well 
known that the reduction in particle size of crystalline systems in the nanometer regime gives 
rise to some important modifications of their properties with respect to their bulk 
counterparts. Two main reasons for the change of electronic properties of the nanosized 
particles can be identified as: (I) the ' quantum confinement ' effect due to the confinement of 
dclocalized electrons in a small s ized particles, which results in an increased electronic band 
gap and (2) the increase of the surface/volume ratio in nanostructures, which enhances 
'surface' and 'interface' effects over the vo lume effects. In case of rare-earth ions, the 
electronic f-f transitions involve localized electrons in the atomic orbital of the ions. 
Therefore, no size dependent quantum confinement effect is found in the electronic 
transitions of the rare-earth doped nanosized particles. However, the 'surface effect ' plays a 
3 
vital role in the photoluminescence properties of these ions. Although there has been an 
explosive growth in the synthesis of nanosized materials, it is still a challenge for material 
chemists to design a process for the fabrication of highly luminescent nanosized materials 
with high degree of crystallinity. Somewhat more recently, the focus of interest has shifted to 
nanosized luminescent materials with tunable morphologies such as nanorods, nanowires, 
nano tubes etc [ 14]. 
1.4 Research objectives 
The specific objectives of this study were; 
I . Synthesis and characterization of the Y20 2S:Eu
3
+ and La20 2S:Eu
3 
... phosphor powder. 
2. Deposition of the Y20 2S:Eu
3
+ and Y 320 3: Eu + phosphor thin films onto Si (100) substrates 
with the use of a KrF excimer laser in pulsed laser deposition. 
3. Characterisation of both powder and thin films using Scanning Electron Microscopy 
(SEM), X-Ray Diffraction (XRD), Photoluminescence Spectroscopy (PL), X-ray 
Photoelectron spectroscopy (XPS), Uv- vis measurements (UV), Fourier Transform Infrared 
(FTIR) and Atomic Force Microscopy (AFM). 
4. Monitor changes in the material properties, d ue to ablation of Y 20 2S: Eu
3
+ thin films in 
vacuum, oxygen and argon gas ambient, using Pulsed Laser Deposition. 
5. Monitor changes in material properties of Y20 2S:Eu
3
+ after annealing at several different 
temperatures. 
1.5 Thesis layout 
The thesis is organized into ten chapters; 
Chapter 1 
ln this chapter the background information, overview of research contributions on classical 
phosphors, rationale and aims o f the research project are given. The issues, perspectives and 
general advantages of nanostructured phosphor are briefly discussed. Finally, a summary of the 
subjects treated in the succeeding chapters of this thesis, is presented. 
Chapter 2 
This chapter provides background infonnation on the fundamentals of phosphors and types of 
4 
luminescence in so lids. Detailed information about the phosphorescence mechanism of long 
persistent phosphors as well as the electronic transition of rare earth ions (Eu3+) is provided. 
The structural properties of the lanthanides (Ln20 2S:Eu
3+, Ho3+) are briefly discussed. 
Chapter 3 
A brief description of the experimental techniques used to synthesize and characterize 
lanthanides phosphors is provided in this chapter. The sol-gel, sol-combustion and solid state 
reaction methods used to synthesize the phosphors are discussed in detail. Detailed 
information on the principle and operation of the experimental techniques used to investigate 
the luminescence and the structure of the phosphors are presented. 
Chapter 4 
Luminescent and structural properties of Y 20 3: Eu
3
+ phosphors prepared by a so l-combustion 
reaction method are discussed in this chapter. 
Chapter 5 
In this chapter, structural and optical properties of La20 2S:Eu
3
+m icro crystals synthesized by 
sol-combustion method were analyzed as a function of La/S molar ratios. 
Chapter 6 
The influence of oxygen partial pressure on materi al properties of Eu3+- doped Y 20 2S thin 
films phosphor deposited by Pulsed Laser Deposition was studied and analyzed in the 
chapter. 
Chapter 7 
ln this chapter energy transfer and material properties of Y 3 320 3: Eu .,.:Ho "'" nanophosphors 
synthesized by so l- combustion method were discussed in detail. 
Chapter 8 
The chapter presents the temperature dependence of structural and luminescence properties of 
Eu3+ -doped Y 20 3 red- emitting phosphor thin films by Pulsed Laser Deposition. 
Chapter 9. 
In this chapter, the infl uence of different species of gases on the luminescent and structural 
properties of pulsed laser ablated Y20 2S:Eu
3
+ thin films were discussed in detail. 
5 
Chapter JO 
This chapter gives general concluding remarks on the overall study and suggestions for 
possible future studies. 
References 
[l] G. Blasse, B.C. Grabmaier, Luminescence Material, Springer-Verlag, New York, (1994) 
[2] S. Shionoya, W.M. Yen, Phosphor Handbook, CRC Press, ( l 998) 
[3] K.F. Braun, Ann. Phys. Chem. , 60 (l 987) 552 
[ 4] H. Nalwa and L.S. Rohwer, Eds. Handbook of Luminescence, Display, ( 1999) 
[5] Materials and Devices, Vols. 1-3, American Scientific Pub. , Stevenson Ranch, CA, 
(2003) 
[6] T. Peng, H Yang, X. Pu, B. Hu, Z. Jiang, C. Yan, Materials Letters 58 (2004) 352 -356 
[7] D. Wang, Y. Li, Y. Xiong, Q. Yin, Journal of the Electrochemical Society, 152 (l) 
(2005) Hl-Hl4 
[8] X. Luo, W. Cao, Z. Xiao, Journal of Alloys and Compounds 416 (2006) 250-255 
[9] S.K Tokuno, S. Komuro, H. Aizawa, T. Katsumata, T. Morikawa, C IS E - ICASE, 
International Joint Conference, Bexco, Busan, Korea, 18 - 2 l Oct, 2006 
[ 1O J C. Chang, D. Mao, Thin Solid Films 460 (2004) 48-52 
[11) H. Chander, Mat. Sci. and Eng., R49: l 13, (2005). 
[ 12) A.M, Srivastava. C.R, Ronda, Luminescence from Theory to Applications. Ronda, C. 
(Ed.), Wiley-VCH, Germany, Chap. 4, (2008) . 
[13) U, Vater. G, Kunzler. W, Tews. J. Fluores .. 4: 79, (1994). 
[14) C.H, Seager. D.R, Tallant. J. Appl. Physics., 87: 4264, (2000). 
6 
Chapter 2 
Theory 
2.1 An overview of phosphors 
Luminescence can be defined as a process by which chemical substances/materials emit 
photons during an electron transition from the excited to the ground state. The materials can 
be excited by irradiating them with high energy electrons or photons. Accordingly, the 
luminescence resulting from excitation by high energy electrons 1s called 
cathodoluminescence and that from the excitation by high energy photons is called 
photoluminescence. The class of materials which emit characteristic luminescence are called 
phosphors. Phosphors consist of a host material which constitutes the bulk and intentional 
impurities introduced to the host. The characteristic luminescence properties are obtained 
either directly from the host or activators/dopants introduced intentionally to the host 
material. An activator is an impurity ion which when incorporated into the host lattice gives 
rise to a centre which can be excited to luminesce. If more than one activator is used, they are 
called co-activators or co-dopants. One activator (sensitizer) tends to absorb energy from the 
primary excitation and transfer to the other activator to enhance its luminescent intensity [I]. 
Luminescence in sol ids, i.e. inorganic insulators and semiconductors, is classi fied in terms of 
the nature of the electronic transitions producing it. It can either be intrinsic or extrinsic. In 
the intrinsic process, the luminescence results from the inherent defects present in the crystal 
structure [2]. This type of luminescence does not involve impurity atoms. Extrinsic 
photoluminescence on the other hand, results from the intentionally incorporated impurities 
in the crystal structure [3]. This type can be divided into two categories; namely localized and 
delocalized luminescence. In the localized luminescence excitation and emission processes 
are confined to a localized luminescence center, whereas in the delocalized luminescence the 
electrons and holes participate in the luminescence process (free electron in the conduction 
band and free holes in the valence band) [ 4] . Luminescence processes can be di vided into two 
main categories, namely fluorescence and phosphorescence based on the time the excited 
electrons takes to return to their ground states after the excitation has been stopped. 
7 
I 
_j 
2.2 Fluorescence 
Fluorescence is the process in which emission of photons stops immediately when excitation 
is cut off. It is the process in which the excited electrons return to the ground state in a time 
not greater than 10-6 sec, the resulting emissions is described as fluorescence [5]. In 
fluorescence there are no traps but many luminescent centres. 
2.3 Phosphorescence 
Phosphorescence occurs when the recombination of the photo-generated electrons and holes 
is significantly delayed in a phosphor. If one of the excited states of a luminescent center is a 
quasistable state (i.e., an excited state with very long life time) a percentage of the centers 
will be stabili zed in that state during excitation. Excited electrons and holes in the conduction 
and valence bands of a phosphor can often be captured by impurity centers or crystal defects 
before they reach emitting centers. When the probability for the electron (hole) captured by 
an impurity or defect center to recombine with a hole (electron) or to be reactivated into the 
conduction band (valence band) is negligibly small , the center or defect is called a trap [6]. 
The decay time of phosphorescence due to traps can be as long as several hours and is often 
accompanied by the photoconductive phenomena (6). 
2.4 Properties and Applications of Nanomaterials 
2.4.1 Some Properties of Nanomaterials 
Nanomaterials are materials with particle sizes less than one micrometer, usually less than 
I 00 nm. These small particle sizes impart different physical and chemical properties 
compared to the bulk forms. Different phases are also found in some nanocrystall ine 
materials. For example, bulk Er20 3 exists in two hexagonal phases, but its nanocrystalline 
Er20 3 exhibits two phases (fee cubic and monoclinic) that are not found in the bulk. 
8 
a b 
c 
Fig. 2.1. Atomic representations of La20 2S along the a) < 11 0> and c) <001 > directions. 
Six possible anion vacancies arc noted (i.e. Al , A2, B 1, B2, C I , C2, with uppercase letter 
indicating the corresponding anion layer). Atomic representations of La20 _, a long the b) 
< 110> and d) <00 I> direc tions. 
A we ll -known prope rty of nanomaterials is that their surface areas are tremendously 
increased. The ir surface-to-vo lume rati os are very high, so that most of the mo lecules/atoms 
are on the surface or at the grain boundaries. Since surface molecules/atoms don't have any 
force above the particle surface to balance the attractive force from ins ide the partic le, they 
are in high energy states. In addition , molecules/atoms at the grai n boundaries arc in highly 
disto rted lattice s tructu res, and forces exerted on a molecule/atom from surrounding species 
are not balanced, so molecules/atoms at the grain boundaries are also in high energy states. 
Therefore, the surface energy of a nanomaterial is very high . The large surface area and 
number of grain boundaries of nanomate ria ls provides a high concentration of low-energy 
diffusion paths. The refore , nanomaterials have higher self-d iffusivity and solute diffusivity 
than the bulk forms. Nanoparticles have electrical and optical properties that are not observed 
in the bulk . These "quantum-size" effects appear when particle sizes are comparable with o r 
smaller than some charac teristic lengths, such as a phonon wavelength , an e lectron de Brog lie 
wavelength , o r an effecti ve Bohr radius around impurity centres. T he energy states of doped 
9 
impurity atoms arc strongly modulated in nanocrystall ites, whose sizes are smaller than the 
Bohr radius of the impurity atoms. This phenomenon is called quantum confinement. 
Quantum confinement effect changes overl aps of the wave functions of the impurity atoms 
with those of host atoms, leading to more efficient interact ions between impurity atoms and 
the host atoms. For example, luminescent propertie · of activators in nanocrystalline 
phosphors are enhanced. In nanocrystalline Y20 3: Tb, the luminescent efficiency increase · 
proportionally with square of decreasing partic le size ( 10 nm to 4 nm), which is predicted 
accurately by a quantum confinement model. Decrease in particle sizes causes localization of 
exciton wave-functions near the impurities (acti vators), which results in higher overlaps of 
the exciton wave-functions wi th those of the impurities. so energy transfer rate from the 
cxcitons to the impurities is higher. Therefore, non-radioacti ve decay rate is relatively 
reduced, and luminescent efficiency increases. However, if there are many extinguishing 
defects at the grain boundaries, the luminescent efficiency of a nanocry~talline phosphor 
decreases. Controlling grain-boundary defects is an important facto r to further improve 
efficiencies of nanocrystalline phosphors. 
...  ................. ..................  .... .............. ...............  a b ...  .. . .   
..: .   ...
:    
.·· ·······························.· ·······-······················· 
0 s Na La 
Fig. 2.2. Schematic illustration fo r se lf-assembled Na-doped La202S nanoplates with OA as 
capping agents; the orange box highlighted in a) is enlarged in b), which shows the thickness 
of one nanoplate. indicating the three layers of primitive ce lls along c-ax is with La3+ as 
ending ions on both sides of the nanoplates. 
Nanocrystalline monocl inic Y 20:i: Eu I+ prepared by laser ablation method has a longer 
5D 70- F2 transition lifetime than that of the bulk. In addition, line widths in the excitat ion 
spectra increase with decreasing particle sizes. This phenomenon has been attributed to 
10 
inhomogeneous broadening from lattice distortion. A blue shift in the emission spectra of 
nanocrystalline Y2 0 3: Eu
3
+ was also observed. Phonons of wavelengths greater than the 
particle sizes cannot propagate in nanocrystalline materials, so phonon distributions (density-
of-states) in nanocrystall ine materials change greatly from the bulk materials (the phonon-
confinement effect). Due to this effect, as particle sizes become smaller, nanocrystalline Si 
becomes more emissive. lt is suggested that the emission centers in porous Si are actually 
nanocrystallites of Si. Because of their novel properties due to the quantum-size effects, 
nanocrystallites of semiconductors are often called quantum dots. The phonon-confinement 
effect is also observed in Raman spectra of nanocrystalline Y 20 3 and Ti02. As the particle 
size decreases from 40 nm to 7 nm, the characteristic Raman lines of nanocrystalline Y 20 3 
shift to lower frequencies, accompanied by significant broadening. The term super-plasticity 
is used to describe the ability of a material to exhibit high tensile ductility (elongation) 
without significant necking. If treated at high homologous testing temperatures, 
conventionally brittle polycrystalline ceramic materials of average grain sizes smaller than I 0 
nm, such as Y20 3-stabilized Zr0 2, exhibit super-plasticity (elongation > 100%). Decreasing 
particle sizes further into the nanometer range will not only increase the overall ductility of a 
materi al prior to failure, but also decrease the super-plasticity-appearance temperature of the 
material. Room-temperature super-plasticity is observed in nanocrystalline Ti02 (rutile). The 
origin of super-plasticity is grain-boundary sliding with some true sliding contribution 
accommodated by matter transportation, grain-boundary migration, grain rotation, and 
diffusion or dislocation motion. Hardness and fracture toughness of a bulk ceramic material 
increase with increasing sintering temperature. However, same hard ness and fracture 
toughness can be achieved by the nanocrystalline form, such as nanocrystalline T i02, sintered 
at much lower temperatures. This observation indicates that nanocrystall ine compacts densify 
much more rapidly than polycrystalline compacts. 
11 
Au nanoparticle Buckminsterfull e rene FePt nanosphere 
Titanium nanoflower S il ver nanocubes S n02 nano fl ower 
Fig. 2.3. Nanomateria ls w ith a varie ty of morphologies 
2.4.2 Nanomaterial - synthesis and processing 
Nanomateri a ls deal with very fine structures: a nanometer is a billionth of a meter. This 
indeed allows us to think in both the 'bo ttom up ' or the ' top down· approaches (Fig.2.4) to 
synthesize nanomateri als, i.e. e ither to assemble atoms together or to dis-assemble (break, o r 
dissoc iate) bulk solids into finer pieces until they are constituted of onl y a few atoms. This 
domain is a pure example of interdi sciplinary work encompas'.'. ing physics , che mistry, and 
engi nee ring upto medic ine. 
12 
Bulk 
fupi.Jll \D 
/ • 
' 11n op11111dn I 
/ 
• \ .•.•.  
Fig. 2.4. Schematic illustration of the preparative methods of nanoparticles 
2.4.3 Applications of Nanomaterials 
Because of the novel properties of nanomaterials compared to their bulk fo rms, they are 
promising candidates fo r many advanced technical applications. Nanomaterial inherentl y 
have a very high surface-to-volume ratio. Therefore, nanometer-sized catalyst supports, or 
nanometer-sized catalysts have greatl y improved efficiencies. Nanocrystallites of opticall y 
acti ve materials (such as Cr: Mg2Si04, Cr: CaMgSi20 6), whose single crystal are di fficult to 
be grown and are sensiti ve to their environment, can be embedded in tran parent host 
material (typically a polymer) to form optical composi tes. The optical composites have the 
propertie of nanocrystallite and the processability of the polymer hosts. Nanoparticle · of 
magnetic materi als exhibit greatly improved magnetic properties and much smaller particle 
sizes, which find many potential applications in magnetic recording, magnetic refrigerat ion, 
and ferroflu ids. Nanometer-sized semiconductor clusters are promising materi als to prepare 
devices fo r efficient conversion of light into electric ity (for example, rutheniu m polypyridyl 
sensiti zers anchored to porous colloidal Ti02 films), or e lectrici ty into light (for example, 
nanocrystalline S i). Nanocrystall ites of emiconductor material s are considered as quantum 
dot due to quantum confinement effects , and doped quantum dots are candidates for 
advanced displays (High Definition TV, Field Emiss ion Display, Pia ma Di play and 
Electroluminescent Display), ultra-fast en ors, and lasers. Super-plastici ty of nanometer-
sized ceramic materials creates a new processing technology for ceramics, the super-plastic 
forming techno logy. Superior hardness and fracture toughness of some nanomaterial s make 
the m ideal materi als for cutting tools. The mechanical properties of nanocrystalline ceramics 
13 
lead them to be called "ceramic steel". Commercial realization of ceramic engmes also 
depends on the development of such nanocrystalline ceramics. 
2.4.4 Applications of Nanometer-sized Y203: Eu3+ 
Y 20 3 is one of the most important host materials for phosphors, scintillators, lasers, and 
fiber-optic communications. Eu-doped Y20 3 is an important red-emitting phosphor, and 
Tbdoped Y20 3 is a green-emitting phosphor. Because of the quantum-size effects, 
luminescent properties of nanocrystalline phosphors are different from their bulk forms, 
which may greatly improve their performance and extend their applications. Due to a 
maximum field gradient before charge leaking, the maximum voltage (approximately 1 kV) 
that can be applied on flat panel display devices (FED, EL, and PD) is much smaller than that 
on normal di splay devices (approximately 5 kV). As nanocrystalline phosphors generally 
have higher efficiencies, lower voltages are adequate to achieve a same efficiency. Therefore, 
nanocrystalline phosphors are ideal candidates fo r flat panel di splays . HDTV requires 
phosphors of very small particle size, narrow size distribution, uni form shape, and high 
intensity without light saturation. Nanocrystalline phosphors have very fast luminescent 
recombination rate, so the saturation can be eliminated, while their nanometer-scale sizes 
fulfill the other requirements of HDTV. Si02 is the gate oxide/dielectric material in metal 
oxide semiconductors (MOS) in verylarge- scale-integrated (VLSI) circuits. By decreasing 
the thickness of the Si02 layer, increases in charge-storage capacity and trans-conductance 
are achieved. However, the smallest practical thickness of the Si02-gated dielectric layer is 
being approached in modem silicon devices. Further improvement needs materials of higher 
dielectric constants as the gate material. Y 20 3 thin fi lms are excellent substitute gate 
materials. They can be made very thin (25 nm), and the dielectric constant is approximately 
four times higher than that of Si02. In addition, they have lower leakage current for a given 
gate voltage, and higher breakdown strength. Y 20 3 is an important additive in many 
structural and functional ceramics. Zr02 is valued for strength and toughness in industrial 
ceramic applications. Y 20 3-stabilized Zr02 avoids the destructive phase transformations 
(volume change during phase transformation causes material cracking) from monoclinic to 
tetragonal and further to cubic at elevated temperatures. Y2 0 3- stabilized Zr02 is used in 
high-temperature refractory, heating cells in oxidation atmospheres, ceramic engines. Y 20 3-
stabilized Zr02 also has high oxygen-ion conductivity at elevated temperatures, which makes 
it suitable for use in oxygen sensors and oxygen pumps at elevated temperatures. Similarly, 
14 
Y20 3 is also used to stabilize Hf02, which is a promising ultrarefractory ceramic material for 
nuclear applications (control rods and neutron shielding). Y 20 3- doped Th02 is also used in 
oxygen sensors. Addition of Y 20 3 nanoparticles influences density and elastic moduli of 
Si3N4 ceramics, and improves their sinterabi lity. Nanocrystalline Y20 3 additives also help to 
prepare AlN ceramics of higher density and thermal conductivity. Nanocrystalline Y 20 3 is 
also used in high-density magnetic recording. To achieve highdensity magnetic recording, the 
recording medium must have high coercivity (>3000 Oe) with thin or no overcoat. Particle 
sizes of the medium must be very small (<  I 0 nm), but magnetically iso lated to minimize the 
transition noise. Y 20 3-doping effectively reduces particle sizes in thin films of 
nanocrystalline BaFe120 19 magnetic material from several hundreds of nanometer to 
approximately 50 nm, while keeping the high coercivity of the material. Y 20 3, Ah03, MgO, 
and Zr02 are novel transparent ceramic materials that can be used in severe environments 
instead of traditional glasses. Y20 3 has a higher melting point and better chemical stabi lity, 
which makes it suitable for heat-resistant transparent windows and walls for high-pressure 
sodium electric discharge light bulbs. There are two methods to prepare transparent Y 20 3 
material. One is the traditional sintering method, and the other is the hotpressing of Y 20 3 
nanoparticles in vacuum. Due to the improved sinterability of nanocrystalline Y20 3, the hot-
pressing method has the advantage of much lower operation temperature ( 1300 °c) and much 
lower operation pressure (44 MP a) than those of the sintering method (2300 °c and 980 MPa , 
respectively). 
2.4.5. Mechanism of the Persistent Luminescence 
Although the overall mechanism of the persistent luminescence of CA}i0 24:Eu + is now quite 
well agreed on [7, 8-1 OJ, the details involved are largely unknown. Long persistent 
luminescence of CA120
2
4: Eu + is thought to have originated from alkaline earth vacancies 
[ 11]. The formation of both electron and hole trapping and subsequent slow thermal 
excitation of the traps followed by emission from Eu2+ ions (Figure 2.5) are being taken to be 
the root causes of the persistent luminescence[ l 2- 14, 15 , 16] . According to this model the 
trapped electrons and holes act as pairs and luminescence can take place as a result of indirect 
centre to centre transitions. Ln other similar systems (e.g. photo- stimulated materi als [ 17] 
the main charge carriers were observed to be electrons and ions but the effect of holes has 
gained more importance in the persistent luminescence materials. However, with the addition 
15 
of some trivalent RE3+ 10ns the persistent luminescence lifetime and intensity can be 
improved further [ 18] 
Conduction Band Localiztd Transitions .\lodel 
Coduction Band Model 
E 
Activt Actin 
Traps Traps 
Deep 
Traps 
Recombinution 
Recombinati on 
Center Center 
Valence Band 
Figure 2.5: Model showing Pers istent Luminescence Mechanism. 
The knowledge of the underlying mechanism of long persistence is very necessary and would 
significantl y assist in the search for persistent luminescence materials. In the present study, a 
detailed investigation was carried out on the Eu2+ doped alkaline earth aluminates 
{CAh04:Eu
2+). Especially, the role of co-doping with different trivalent rare earth [RE 3+] 
ions (D/~ and Nd3+) in the enhancement of the afterglow of CAh04:Eu
2+, RE3+ was studied 
by several spectroscopic methods viz Photoluminescence (PL) and Thermoluminescence 
(TL). 
2.4.6. The Luminescent Center 
Despite the fact that considerable amount of study on the aspects of luminescence could be 
carried out by taking into account a simple model for the centre it is quite hard to find out 
what is exactly going on inside the centre. Several theories or approaches have to be put to 
trial depending on the complexity of the centre. One such famous approach is the 
configurational coordinate model. This approach assumes that the luminescent centre has 
some equilibrium position in the crystal lattice and that a change in energy occurs due to 
some displacement trom this position. The interaction of the centre with the crystal lattice in 
terms of its electronic state and the vibrations of the lattice can be seen to be as a function of 
the position of the nuclei and the model can be employed to easily explain such effects as the 
Stokes shift between absorption and emission. But owing to the fact that electronic transitions 
16 
of the centre are coupled to the movements of the lattice around the centre the simple model 
is not generally acceptable. 
To differentiate the electronic state from vibrational state of the luminescent centre the Bom-
Oppenheimer approximation is used but it has been shown by Fowler and Dexter ( 1962) that 
the potential energy curves in the configurational coordinate diagram are also a function of 
the electronic state. 
In condensed systems the Einstein relations are not valid as such and the complex relaxations 
which occur after an excitation do not simpli fy the scenario because the electronic states in 
emission are likely not to be the same with those in absorption. Furthermore, because of the 
Jahn-Teller effect, which tends to remove degeneracy of an excited state by creating 
asymmetry in the centre, there may be a separation in the excited state. There are also 
transition probabilities for the absorption and the emission. One of the most important points 
is that the matrix element for the absorption transition may be different from that for the 
emission transition. A lot of interest in luminescence now needs to be taken in quantitative 
studies o f phonon-photon interactions (preferably at very low temperatures) [ 19]. 
2.4.7. Phase Transformation 
In spite of a great deal of research work on Y 20 2S: Eu
3
+ phosphors, the phase trans formations 
of Eu3+ doped lanthanides compounds and their effects on luminescent properties have been 
rarely reported until now. 
2.4.8. Effect of Lattice Defects on Persistent Luminescence 
Generally, when the mean particle size of phosphors is smaller than 1-2 µm , there is a drop in 
their luminescence efficiency. This is due to the fact that surface defects become more 
important with decreasing particle size and an increase in the surface area. This can often lead 
to the reduction of the emission intensity [20-24]. The presence of any kind of lattice defect 
in the host lattice in most cases has been found to greatly reduce the efficiency of 
luminescence. It also brings about the long afterglow observed in some potentiall y efficient 
luminescent materials. These defects are usuall y considered to be disadvantageous as far as 
the properties of a phosphor are concerned when the practical applications are considered 
[25]. Consequently, the luminescence applications based on phosphors with lattice defects are 
rare. 
17 
2.4.9. Energy Transport and Storage in Luminescent solids 
The dominant role played in luminescence mechanisms by the transport of energy was 
pointed out by Broser in an invited paper [26] . ln condensed systems the interatomic distance 
is considerably smaller than in gases and the probability of interaction between a luminescent 
centre and distant atoms are much greater. Energy transfer may take place by free charge 
carriers, excitons, quantum mechanical resonance, photons or phonons, and may be studied 
by direct measurements of such properties as velocity, lifetime and carrier range, or by 
indirect measurements. Two important parameters are lifetime and diffusion coefficient. 
Great advances have been made in the study of energy transport by direct mechanisms in 
phosphors during the last ten years. New experiments have been devised (particularly for 
excitons) to measure transport parameters and older experiments have been perfected. 
Nevertheless in the field of energy transport in phosphors there are ample problems 
remaining to be so lved in the next decade. 
Storage of energy in phosphors is still being studied extensively by thermal ejection 
measurements on trapped charge carriers [27-30]. Interpretation of glow curves in terms of 
traps or metastable states is obviously more difficult in organic compounds than inorganic 
compounds. Even in inorganic compounds it is not likely that the method gives the true or 
complete distribution of trapping states in a phosphor. 
The conventional method of filling the traps at low temperature is by illuminating the 
specimen, but if the traps are filled by space-charge injection of charge using a high field 
across the specimen it is possible to remove any ambiguity as to the sign of the charge carrier 
responsible for the peak [31]. lf a blocking electrode is used as the cathode during the 
heating-up process it is also possible to distinguish between surface and volume states. The 
use of high fi elds causes the carrier transit time to be reduced, the probability of re-trapping 
to be reduced, and the kinetics to be more like those of the monomolecular theoretical model. 
For the investigation of trap spectra with a continuous energy distribution the fractionated 
glow technique is proving to be of value [32]. 
2.5.1: Emission and Excitation Mechanisms of Phosphors 
Basic concepts involved in luminescence will be discussed. We will take a closer look at a 
number of excitation mechanisms which are involved in generating luminescence and 
processes which lead to luminescence, tak ing illustrative examples from luminescent 
18 
materials applied in fluorescent lamps and cathode ray tubes. With respect to fluorescent 
lamps, we will restrict ourselves to discharge lamps based on the low-pressure Hg discharge. 
2.5.2: General Considerations - Fluorescent Lamps 
On passing a current through an Hg discharge, UV light is generated as a consequence of 
electronic transitions on the Hg atoms. In low-pressure Hg discharge, the main emission line 
is located at 254 run. This light is invisible and harmful; therefore it has to be converted into 
visible light. This is done by the application of luminescent materials. These materials have to 
show a strong absorption at 254 run and have to convert this into visible light very efficiently. 
In most of the fluorescent lamp phosphors, the optical Luminescence processes leading to 
luminescence do not involve host lattice states, implying that the energy gap is at least 4.9 
eV, this being the energy of a photon with wavelength 254 nm. Therefore, the luminescent 
materials applied in fluorescent lamps are insulators. The conversion efficiency of 
luminescent materials is very high: about 90% of the UV photons are absorbed, and also 
about 90% of the absorbed photons are converted into visible light. This implies that such 
materials cannot be improved any further in terms of conversion efficiency unless materials 
can be found that generate more than one visible photon after absorption of a UV photon. 
Compact fluorescent lamps have lower light generation efficiency (only 15 %). As the 
luminescent materials applied are the same or very similar, this must due to the lower 
discharge efficiency in these devices, which, in tum, is due to the smaller diameter of the 
lamp envelope and therefore to the increased wall losses: excited Hg atoms reach the ground 
state on interacting with the lamp wall without generating UV light: energy and momentum 
can be conserved by interaction of excited species with the wall without generation of light. 
2.5.3 General Considerations - Cathode Ray Tubes 
Though the importance of cathode ray tubes is rapidly decreasing, we will treat the 
luminescence mechanism in these materials in view of its historical importance. In addition, 
the excitation mechanism that comprises excitation with high-energy particles (electrons, X-
ray photons, or g-rays) is also operative in phosphors used in scintillators for, e.g. , medical 
applications. Luminescent materials appl ied in cathode ray tubes in general differ from those 
applied in fluorescent lamps. Excitation proceeds via the band gap. To achieve high 
19 
efficiencies, small values for the band gap are needed, as will be elucidated below. For this 
reason, quite a few luminescent materials applied in cathode ray tubes are semiconductors. 
2.5.4. Emission and Excitation Mechanisms of Phosphors 
The luminescence mechanism operating in the blue and green emitting phosphors applied in 
cathode ray tubes is a beautiful example of luminescence involving defect states in 
semiconductors. We wi ll therefore also discuss this mechanism in some detail. The maximum 
energy efficiency of the cathode ray phosphors is relatively low, at most about 25 %, as wi ll 
be outlined below. Also for these phosphors, the maximum efficiencies have been reached. 
2.5.5. Luminescence Mechanisms 
Luminescent materials, also called phosphors, are mostl y solid inorganic materials consisting 
of a host lattice, usually intentionally doped with impurities. The impurity concentrations 
generally are low in view of the fact that at higher concentrations the efficiency of the 
luminescence process usually decreases (concentration quenching, see below). In addition, 
most of the phosphors have a white body color. Especially for fluorescent lamps, this is an 
essential feature to prevent absorption of visible light by the phosphors used. The absorption 
of energy, which is used to excite the luminescence, takes place by either the host lattice or 
by intentionally doped impurities. In most cases, the emission takes place on the impurity 
ions, which, when they also generate the desired emission, are called acti vator ions. When the 
activator ions show too weak an absorption, a second kind of impurities can be added 
(sensitizers), which absorb the energy and subsequently transfer the energy to the acti vators. 
This process involves transport of energy through the lumi nescent materials. Quite 
frequently, the emission color can be adjusted by choosing the proper impurity ion, without 
changing the host lattice in which the impurity ions are incorporated. On the other hand, quite 
a few acti vator ions show emission spectra with emission at spectral positions which are 
hardly influenced by their chemical environment. This is especially true fo r many of the rare-
earth ions. Luminescent material containing activator ions (ions showing the desired 
emission) and sensitizing ions (on which, e.g., UV excitation can take place). 
20 
2.5.6. Center Luminescence 
In the case of center luminescence, the em1ss10n 1s generated on an optical center, in 
contradiction to, e.g., emission, which results from optical transitions between host lattice 
band states or from a transition between two centers. Such an optical center can be an ion or a 
molecular ion complex. One speaks of characteristic luminescence when, in principle, the 
emission could also occur on the ion in a vacuum, i.e. when the optical transition involves 
electronic states of the ion only. Characteristic luminescence can consist of relatively sharp 
emission bands (spectral width typically a few nm), but also of broad bands, which can have 
widths exceeding 50 nm in the visible part of the spectrum. Broad emission bands are 
observed when the character of the chemical bonding in the ground and excited state differs 
considerably. This goes hand in hand with a change in equilibrium distance between the 
emitting ion and its immediate chemical environment. 
2.5.7. Charge Transfer Luminescence 
In the case of charge transfer, the optical transition takes place between different kinds of 
orbitals or between electronic states of different ions. Such an excitation very strongly 
changes the charge distribution on the optical center, and consequently the chemical bonding 
also changes considerably. In these cases, therefore, very broad emission spectra are 
expected. A very well-known example is CaW04, used for decades for the detection of X-
rays, which shows luminescence originating from the (W04) 2 groups. A similar compound, 
also showing blue emission, was used in early generations of fluorescent lamps: MgW04. 
The transition involves charge transfer from oxygen ions to empty d-levels of the tungsten 
ion. In this material no intentional dopant is introduced, and for this reason it is also called 
self-activated. 
2.5.8. Donor Acceptor Pair Luminescence 
This luminescence mechani m is found in some semi-conducting materia ls doped with both 
donors and acceptors. Emission and Excitation Mechanisms of Phosphors results in 
luminescence. Electrons that are excited into the conduction band are captured by ionized 
donors, and the resulting holes in the valence band are captured by ionized acceptors. The 
emission invo lves electron transfer between neutral donors and neutral acceptors. The final 
state (with ionized donors and acceptors) is Coulomb stabilized. Therefore, the spectral 
21 
position of the emission generated on a donor-acceptor pair depends on the distance between 
the donor and the acceptor in a pair: the smaller the distance, the higher the energy of the 
photon generated. 
The energies involved in these processes are: 
1. The absorption of energy with the band gap energy. 
2. Neutralization of the ionized donor. The Coulomb term originates from the electrostatic 
interaction between ionized donor and acceptor. 
3. Neutralization of the ionized acceptor. 
4. The luminescence process. 
2.6.l. Mechanisms Underlying Energy Transfer 
For energy transfer, the sensitizer ion and the activator ion have to show physical interaction. 
This energy transfer can find its origin in electrostatic and exchange interaction. In addition, 
the emission spectrum of the sensitizer ion and the absorption spectrum of the activator ion 
have to show spectral overlap, for energy conservation reasons. 
2.6.2. Cross-relaxation and Energy Transfer 
A phenomenon not discussed until now is cross-relaxation. In such a process, which can also 
be looked upon as energy transfer, the excited ion transfers only part of its energy to another 
ion. In this case, the energy difference between the 50 3 and 50 4 excited states matches 
approximately the energy difference between the 7F6 ground state and higher 
7FJ states . As in 
the energy transfer processes discussed above, the process of cross-relaxation has a low rate. 
In many host lattices, therefore, at low Eu3+ concentration, emission from both the 50 0 and 
50 2 excited states is observed (unless the gap between these two states is bridged by phonon 
emission, for which relatively high-energy phonons are required). 
2.6.3. Exchange interaction between sensitizer and activator ion 
This mechanism does not require allowed optical transitions. This is the mechanism which is 
operative in the one-component white fluorescent lamp phosphor Ca5(P04) 3(F,Cl):Sb,Mn, as 
deduced from an analysis of the decay curve for some Mn2+ concentrations [ 45]. The same 
study did not reveal ev idence for energy transfer between antimony ions, indicating the 
22 
necessity of nearest neighbor Sb-Mn interaction, which is a prerequisite for energy transfer 
via exchange interaction. Please note, in addition, that in view of the large Stokes shift 
between absorption and emission on the Sb3+ ion in this lattice, no energy transfer between 
the antimony ions is expected. Both for electric dipole - electric quadrupole and exchange 
interaction, the distance between sensitizer ion and activator ion has to be rather small , not 
larger than about 1 nm. This requires high activator and/or sensitizer ion concentrations, 
which is a disadvantage, considering the high costs of these materials. 
2.6.4 The Energy does not reach the Luminescent ion 
When there is more than one origin of optical absorption at the wavelength at which the 
excitation takes place, the quantum efficiency can be less than unity, even if the ion showing 
luminescence has a quantum efficiency of one. This is, e.g., the case if both the luminescent 
ion and the host lattice show optical absorption at the excitation wavelength, or the energy 
transfer probability of the host lattice to the luminescent ions is smaller than unity. 
Comparing the absorption or refl ection spectra with the excitation spectra can disentangle the 
different contributions to the absorption. Degradation of luminescent materials can be due to 
creation of additional absorption centers in the spectral range where the activators or 
sensitizers also absorb. 
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[ 17] C. B, Murray,; C. R, Kagan. ; M. G, Bawendi,.: Science 270, 1335 ( 1995). 
[1 8) C, Bruch,; J. K, Kruger. ; H. G, Unruh,; W, Krauss,; 8 , Zimmermeier,; C, Beck,; R, 
Hempelmann,: Ber. Bunsenges. Phys. Chem. I 01 , 176 1 ( 1997). 
[19] D, Bersani,; P.P, Lottici, P.; X. Z, Ding,: Appl. Phys. Lett. 72, 73 (1998). 
[20] T. G, Langdon,: Metal!. Trans. I 3A, 689 (1982). 
[2 1) T. G, Langdon,: Key Eng. Mater. 97-98, 109 (1 994). 
[22] J , Karch,; R, Birringer,; H, Gleiter,: Nature 330, 556 ( 1987). 
[23) F, Wakai,; N, Kondo,; H, Ogawa,; T, Nagano,; S, Tsurekawa,: Mater. Character. 37, 33 1 
( 1996). 
[24] R.S, Averback,; H, Hahn,; H.J, Hofler,; J. L,Logas,; T. C, Shen,: Mater. Res. Soc. Symp. 
153, 3 (1989). 
[25] T, Masui,; K, Fujiwara,; K, Machida,; G, Adachi,; T, Sakata,; H, Mori ,: Chem. Mater. 
9, 2197 (1997). 
[26] D. B, Barber,; C.R, Pollock,; L.L, Beecroft,; C.K, Ober,.: Opt. Lett. 22, 1247 (1997) . 
[27] R, Dagani,: Chem. & Eng. News 70, 18 ( 1992). 
[28] R, Argazzi,; C.A, Bignozzi,: J. Am. Chem. Soc. 117, 11 8 15 ( 1995). 
[29] R.P, Chin,; Y.R, Shen,; Y, Petrova-Koch,: Science 270, 776 ( 1995). 
[30] R.N, Bhargava,: J. Lumin. 70, 85 (1996). 
24 
[31] J, Wittenauer,: In Plastic Deformation of Ceramics, R. C. Bradt, C. A. Brookes, and J. L. 
Routbort (eds.), Plenum, New York, 321 (1995). 
[32] G, Skandan,; B.H, Kear,: Mater. Sci. Forum 243-245, 217 ( 1997). 
[33] K. H Jack, : High Technology Ceramics, Past, Present, and Future; Ceramics and 
Civilization 3, W. D. Kingery (ed.), 259 (1986). 
[34] R.W, Cahn,: Nature 332, 761 ( 1988). 
[35] E.T, Goldburt,; B, Kulkarni ,; R. N, Bhargava,; J, Taylor, ; M, Libera,: Mater. Res. Symp. 
Proc. 424, 441 ( 1997). 
[36] J.M.P.J, Yerstegen,; J. L, Sommerdijk,; J .G. Yerriet,; (1973) J. Lumin. , 6, 425. 
[37] T.F, Soules,; R.L, Bateman,; R. L, Hewes,; E. R, Kreidler: ( 1973) Phys. Rev. 
87, 1657. 
[38] D.J , Robbins, (1980) J. Electrochem. Soc., 127, 2694. 
[39] R. H, Bartram,; A. Lempicki, (1996) J . Lumin., 69, 225. 
[40] A, Meijerink,; J, Nuyten,; G, Blasse. (1989) J. Lum., 44, 19. 
[41] G.S, Ofelt. (1962) J. Chem. Phys., 37, 511. 
[42] B.R, Judd. (1962) Phys. Rev. , 127, 750. 
[43] K.J.B.M, Nieuwesteeg. (1989) Philips J. Res., 44, 383. 
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Singapore, Hong Kong, 377-408. 
25 
Chapter 3 
3. Experimental Techniques 
3.1. Introduction 
This chapter gives a brief account of the experimental techniques used to synthesize and 
characterize the phosphor. X-ray diffraction (XRD), Scanning electron microscopy (SEM) 
and X-ray photoelectron spectroscopy (XPS) were used to investigate the crystalline 
structure, particle morphology and elemental composition of the phosphor respectively. 
Atomic force microscopy (AFM) was used to analysis the surface roughness of the 
phosphors. Fourier transform infrared spectroscopy (FTIR) was used to identify and/or verify 
compounds. A 325 nm He-Cd laser fitted with a SPEX 1870 0.5m spectrometer and a 
photomultiplier detector was used to collect photoluminescence data in air at room 
temperature. A Cary Eclipse fluorescence spectrometer fitted with a monochromatic xenon 
lamp was also used to investigate the photoluminescence (excitation and emission) properties 
and decay characteristics of the phosphors. The powders investigated in this research were 
prepared by sol-combustion method, whi lst the thin fi lms were deposited by pulsed laser 
deposition. Theory on the PLD technique, XRD, SEM, PL, AFM, FTIR, TL and XPS can be 
found in this chapter as well as the experimental procedure. 
3.2. Synthesis and deposition technique 
The growth of nanoscience and nanotechnology in the last decade has been made possible by 
the success in the synthesis of nanomaterials. The synthesis of the nanomaterials includes 
control of size, shape and structure of the materials. Over the past few years, nanoparticles 
(powders) of ceramic materials have been produced on large scales by employing physical 
and chemical methods. The considerable scientific progress in preparation of nanomaterials 
such as ceramics and semiconductors has been made possible by synthesis techniques such as 
co-precipitation, sol-gel, combustion method, sol-combustion etc [ I]. In th is study, sol-
combustion method was used to synthesize the lanthanides oxides and oxysulfides phosphors 
doped with rare earth ions. 
26 
depending on the desired measurements. The X-ray diffractometer fall broadly into two 
classes: single crystal and powder. The powder diffractometer is routinely used for phase 
identification and quantitative phase analysis [7] . X-ray diffractometer consist of three basic 
elements: an X-ray tube, a sample holder, and an X-ray detector. The X-rays are produced in 
a cathode ray tube by heating a filament to produce electrons. When the voltage is applied, 
the electrons will accelerate towards the target material. When electrons have sufficient 
energy to dislodge the inner shell electrons of the target material, characteristic X-ray spectra 
will be produced. These X- ray spectra consist of several components and the most common 
are Ka and K~. The target materials that are usually used are Cu, Fe, Mo and Cr. Each of 
these has specific characteristic wavelengths [8]. In order to generate the required 
monochromatic X-rays needed for diffraction, a filter such as a foil or crystal 
monochrometers is usually used. Copper is the most commonly used target material for 
single-crystal diffraction, with Cu Ka radiation = 1.54 l 8A. The resulting X-rays are 
collimated and directed onto the sample. As the sample and detector are rotated, the intensity 
of the reflected X-rays is recorded. When the geometry of the incident X-rays impinging on 
the sample satisfies the Bragg Equation, constructive interference occurs and characteristic 
diffraction peaks of the sample will be observed [9]. The X-ray diffractometer Bruker XRD 
D8 Advance shown on Figure 3.1, from University of the Free State, Department of Physics, 
Bloemfontein Campus was used to analyze the samples in this study. 
3.3.2. Scanning electron microscopy (SEM) 
Scanning electron microscopy is a technique in which a beam of finely focused electrons is 
used to examine materials on a nanometer to micrometer scale. It is often used as the 
analytical instrument of choice when the light microscope no longer provides adequate 
resolution. The SEM consists of an electron optical column mounted on a vacuum chamber 
with electron gun placed on top of the column, as illustrated in Figure 3. l. The electron gun, 
which consists of a tungsten or LaB6 filament or a field emission electron gun is used to 
generate electrons, when the applied current causes resistance heating which generates the 
electrons [ 1O ]. When a high energy electron beam impinges on the sample, a variety of 
electrons and/or x-rays will be generated. Depending on the nature of the sample, these can 
include secondary electrons (electrons from the sample itself), backscattered electrons (beam 
electrons from the filament that bounce off nuclei of atoms in the sample) and X-rays. 
28 
Figure 3.1: Bruker D8 Advance mode l x-ray diffractometer 
This technique can be used to investigate the topography, morphology, and elemental 
composition of materials (if coupled with an EDS) on the micrometer to nanometer scale. 
The secondary electrons can be used to investigate the image and the topographic features of 
so lid samples. The SEM coupled with EDS, is used during the analysis of the characteristic 
x-rays emitted as a result of electronic transitions between the inner core levels to provide a 
quantitati ve and qualitative elemental composition of the sample [ 11 ]. 
29 
~ E.LECTRO. .. GUN 
r. ilJ MAG NIF!'ICATION COffTROl 
.... 
CL 
0 -. ;i--+~--0-BJ_E_CTI_V-E~D-1_ _ _ 
~ 
.0.. . 1• • f-- LEl'tS ":; • .. 
._ r-1- - --- SCAN INC COfLS OENEAAlOR 
~I ~--
~ S YNCHROUOUSL 'V 
VACUUM SPECIMEN I \ 
SCANNED CRf 
SYSTE M 
Figure 3.2: Schematic representation of a SEM 
The analysis of elemental compositions can be perfo rmed by measuring the energy and 
intensity distribution of the x-ray signals generated by a focused electron beam. Due to a 
well-defined nature of the various atomic energy levels, it is clear that the energies and 
associated wavelengths of the set of x-rays will have characteristic values for each of the 
atomic species present in a sample [ 12]. The morphologies and elemental compositions of the 
phosphor powders were obtained by using a Shimadzu Superscan SSX-550 SEM coupled 
with an EDAX EDS from the Center of Microscopy at the University of the Free State as 
shown in Figure 3.3. 
3.3.3. Photoluminescence spect roscopy (Helium cad mium laser) 
Photoluminescence (PL) is a powerful and a relatively simple method, extensively used as 
characterization technique of semiconductor physics fo r a number ofreasons [ 13, 14]. 
• It is non-destructive because it is based on pure optical processes. 
• No sample preparation is required. 
30 
• Highl y sensitive. 
• Detailed information on the electronic structure tn the semiconductor can be 
deduced from the experiments. 
Figure 3.3: Shimadzu Superscan SSX-550 model Scanning Electron Microscope 
Information that could be deduced from a PL study includes the s ize of the band gap, 
impurity levels, interface, and surface properties as well as density of states and exocitonic 
states. Basically in PL measurements, a semiconductor sample is optically excited by an 
excitation source such as a laser which produces photons havi ng energies larger than the band 
gap of the semiconductor. The incident photons are absorbed under creation of electron-hole 
pai rs in the sample. After a short time the electrons eventually recombine with the holes, to 
emit photons, and light or luminescence wi ll emerge from the sample . The energy of the 
emitted photons reflects the energy carrier in the sample. The emitted luminescence is 
collected, and intensity is recorded as a function of the emitted photon energy, to produce a 
PL spectrum. In a PL measurement, the excitati on energy is kept fixed , while the detection 
31 
energy is scanned. The energy of emitted photon is characteristic for radiative recombination 
process. 
~ 
\ 
' 
Titt plate 
Figure 3.4: The cavity structure of He-Cd Laser 
PL technique is particularly helpful in the analysis of discrete detect and impurity states. To 
gain more knowledge about the electronic structure, magnetic and electric fi elds can be 
applied in a controlled manner. Moreover external forces can be used in PL investigations, 
e.g. the strain by exposing the material to mechanical pressure. Since PL relies on radiative 
recombination, so it is very difficult for the investigation of non-radiative processes needs 
indirect methods, and the material having poor quality are hard to characterize through PL. 
3.3.4. Radiative recombination mechanisms observed in PL 
In semiconductors, the luminescence can be achieved by several radiati ve transitions between 
the conduction band and valence band, exciton, donor and acceptor levels, as shown in Figure 
3.6. Upon excitation at energy above the bang gap, free electrons are created in the 
conduction band together with the free holes in the valance bond. These carriers wi ll 
energetically relax down the band edge. Due to mutual coulomb interaction, electron-hole 
pair is formed. This electron-hole is usually called a free exci ton (FE). Its energy is s lightly 
smaller than the bang gap energy. This energy difference is the binding energy of the free 
exciton. A neutral donor (acceptor) will give rise to an attractive potential, a free exciton 
might be captured at the acceptor (donor). A bound exciton (BE) is formed. 
32 
--------- -------------------- -------·· - ----------------
... --
.1 
Figure 3.5: Cary Eclipse Florescence Spectrophotometer 
3.3.4. Radiative recombination mechanisms observed in PL 
ln semiconductors, the luminescence can be achieved by seve ral radiative transitions between 
the conduction band and valence band , exciton, donor and acceptor levels, as shown in Figure 
3.6. Upon excitati on at energy above the bang gap, free electrons are created in the 
conduction band together with the free holes in the valance bond. These carriers will 
energetically relax down the band edge. Due to mutual coulomb interaction, e lectron-hole 
pair is fo rmed. This electron-hole is usually called a free exciton (FE). Its energy is slightly 
smaller than the bang gap energy. This energy difference is the binding energy of the free 
exc iton. A neutral donor (acceptor) wi ll give rise to an attractive potential, a free exc iton 
might be captured at the acceptor (donor). A bound exc iton (BE) is formed. 
33 
in l c rb a nd p h o l o luminc s ccn cc 
k 0 k k 0 
d i rcc l - gap malcrt als 1ndi r c c l - gap m a l c ria ls 
Figure 3.6(a): Schematic illustration of common recombination proces es 
Conduction Band 
• • • 
~ ~ 
~ ~ =o I • • 
Valence Band 
Band 10 Free Bound f ret to Donor-acceptor Defee! 
band exci1on nciton bound pair transition iniernal 
transition transition trnnsirino 
Figure 3.6(b): Sche matic ill ustration of common recombination processes 
34 
An electron bound to a donor can recombine directly with a free hole from a valence band. 
This kind of recombination is called free-to-bound (FB) transition. The recombination energy 
for such a transition corresponds to the band gap energy reduced with the binding energy of 
donor. Another possibility is that a hole bound to an acceptor recombines with an electron 
bound to a donor in donor-acceptor pair (DAP) transition. Both the donor and the acceptor 
are neutral before the recombination (i.e. the donor positively and the acceptor negatively 
charged). Thus there is a Coulomb interaction between the donor and acceptor after the 
transition and extra Coulomb energy is gained in the final state added to the radiative 
recombination energy. The transition energy E (R) depends on the distance R between the 
donor and acceptor atoms. 
3.3.5 Pulsed laser deposition (PLD) 
PLD is the most recent thin fi lm deposition technique compared to MOVPE, MBE and 
sputtering. Research on the interaction of high power laser beams with solid surfaces have 
been conducted since the 1960s when the first high power ruby laser became available. 
However, it was not until 1987 when a more stable Nd:Y AG laser was used to deposit ternary 
HgCdTe layers that research on PLO began to flourish [ 15]. This method has received 
considerable attention for its ability to deposit the complex ox ides needed to produce 
superconducting thin films [ 16- 19] . It also has the ability to operate in high pressure reactive 
gases, unlike other deposition methods [20]. The advantages of PLO also come from the 
simple hardware and setup. A deposition system usually consists of an excimer laser and 
optical elements to maneuver and focus the laser beam. Some of the optical elements that are 
used in the set up are foc using lens, apertures, mirrors, beam splitters and laser windows. The 
substrate and target are housed in a vacuum chamber, while a high-power laser is mounted 
externally and can be focused via optical lenses to vaporize the target material and deposit 
thin fi lms on the substrate. A schematic of a basic deposition system that uses oxygen as its 
reactive gas is shown in Figure 3.8. 
Excimer lasers with wavelengths between 200 and 400 nm are most often used for pulsed 
laser deposition [21]. Excimer laser below 200 nm are not typically used for PLO due to the 
possibility of absorption by the Schumann- Runge bands of molecular oxygen. As shown in 
Figure 3.10, the laser source is located external to the vacuum chamber. The external energy 
35 
source allows the film growth process to take place in a reactive environment with any type 
and amount of gas. The external source gives the added advantage of the laser being available 
for more than one deposition system. 
/ 
Target Vacuum 
P\lu me 'S u bstrate pump 
Reactive gas 
Figure 3.8: Schematic diagram of a pulsed laser deposition chamber setup 
Once the laser is focused into the chamber, the target absorbs the energy from the laser. The 
ultra- violet (UV) radiation is converted to electronic excitation. This is converted into 
thermal, chemical, and mechanical energy, leading to ablation and evaporation of the target. 
The evaporants form a mixture of energeti c species including atoms, molecules, electrons, 
ions, and micron sized particulates. This mixture is the often referred to as a plume. An 
example what a plume looks like during the fi lm growth process is shown in Figure 2. 10. The 
plume quickly expands in the vacuum from the target to form a "nozzle j et" [22]. As the 
plume reaches the substrate (which may be heated), film nucleation commences. The quality 
of the thin films produced by pulsed laser deposition is dependent on several vari ables. Laser 
power and spot size have a significant effect on particulate size and density. As the laser 
fluence is increased beyond a threshold, the number of particulates that are fonned also 
increases [23, 24]. Laser fluence is defined as the laser energy per unit area and thus, may be 
adjusted by varying the laser power or laser spot size. Background gases may change growth 
36 
parameters such as the deposition rate and the kinetic energy distribution of the depositing 
species [25]. For instance, an oxidizing environment can help oxides to form and stabilize the 
desired crystal phase at the deposition temperature [26]. Substrate temperature has an effect 
on the stoichiometry of the fi lm as well as the fi lm structure [27]. Film structure has also been 
influenced by the deposition rate. Wu et al. found that an increase in the deposition rate led to 
a decrease in the crystallinity of YBa2Cu30 r8 thin films. At the higher deposition rates, the 
arrival rate exceeds the diffusion rate. Equilibrium conditions are not maintained and 
structural defects are formed [28]. 
Figure 3.9: Examples of picture of plume developed during PLD 
In summary, the advantages of this growth method include [29]: 
• Flexibi lity to use energy source with more than deposition chamber 
• Easy process control 
• Ability to use high reactive gas pressures 
37 
• Decreased contamination from outside sources 
• Control of film stoichiometry. 
The disadvantages of PLO are: 
•The generation of particulates during the deposition process, which is not ideal for the 
application field . 
•The non-uniform layer thickness. 
•The ablation plume cross section is generally small and this limits the sample size. 
•The deposition of novel materials usually involves a period of optimization of 
deposition parameters. 
The short laser pulses result in congruent evaporants. Congruent evaporation aids in 
stoichiometry control of the thin films during mass transfer from target to substrate. One 
obvious disadvantage is the presence of micron sized particulates. Also, scale up to large area 
deposition is not easi ly completed. Pulsed laser deposition has been shown as an appropriate 
method for growing phosphor films, specifically oxide phosphors. Yttria and silicate 
phosphors are some of the oxide materials that have been grown by this method [30, 33]. 
While the hardware required in PLO is relatively simple, the laser-target interaction ts 
extremely complex. The mechanism for material ablation depends on properties of the laser 
as well as the optical, topological, and thermodynamic properties of the target. The 
electromagnetic energy absorbed by the solid surface is first converted into electronic 
excitation, and then into thermal, chemical, and mechanical energy that leads to evaporation, 
plasma formation and exfoliation [34]. There are two main drawbacks of PLO. Due to the 
angular distribution of the plume, large area deposits are often not uniform. This is usually 
resolved by rastering the laser or rotating/translating the substrate. The other more intrinsic 
drawback is known as "splashing," where micron-sized molten globules are deposited onto 
the substrate from either subsurface boiling, expulsion of the liquid layer, or exfoliation -
where so lid particulates instead of liquid globules are deposited onto the substrate. In order to 
38 
Figure 3.10: 248 nm KrF Lambda Physic excime r laser with PLO setup 
In order to reduce or e liminate the effect of splashing from liquid layer expulsion and 
exfo li ation , several approaches have been studied . Using a mechanical particle filter, s low-
moving particulates can be screened with a velocity selector placed between the target and 
the substrate. However, the filte rs tend to be bulky and have low transmiss ion rate that lowers 
the deposition rate as wel l. Other attempts include manipulating laser plumes by using two 
synchronized laser beams; using a hot screen to re-evaporate lighter particulates; or placing 
the substrate at an angle from the normal axis (35]. 
3.3.6 X-Ray Photoelectron Spectroscopy (XPS) 
X-ray photoelectron spectroscopy, al so known as electron spectroscopy fo r chemical analysis 
(ESCA), is a widely used surface technique to obtain chemical information at surfaces of 
various materials . The XPS process involves the ejection of an electron (photoe lectron) in 
vacuum from the K level of an atom by an energetic inc ident x-ray photon (36]. 
Photoelectrons are collected and analyzed to produce a spectrum of emiss ion intensity versus 
39 
e lectron binding energy. Jn general, the binding energies of the photoelectrons are 
characteri stic of the element from which they are emitted [37]. The schematic of the XPS 
process is shown in Figure 3.1 1. 
f1's'!o:=C"' l'lf!f C!s&r9!!! 1• t 5 • VJ 
••<~ 01'1)' fl'OM t,.. ".ry !Oji l wri#Ce 
t"TO • f !~A.) o f lM I • "'9 • 
, .....,_ - """'"'>' ...... ~· •"1 
,_..,.,., u1fll•.,,,.Oft •~/'!O " «>-•J 
Figure 3.11 : Schematic diagram of the XPS process in copper 
Figure 3.12: PHI 5400 Yersaprobe scanning x-ray photoe lectron spectrometer 
40 
7 · GullA~I 
COiis 
2 Coooenser 8 Conoens.e• L1o11s # 1 
St1gmat0< 
+ 9 • Ccnaenser Lens 112 3 • Condenser 
Aperture 10 BeamT 11 
Shi" ContrOis 
4 Objective + Speom"n Hol.;ller 
ApertlXe 
1:' Qbiect ve Lens 
13 Ooec1,ve 
5 Selecteo S!ign ator 
/Vea Apen11e + 
14 D.Hract Ol'\i 
6 Intermediate lf:tenned ate L 111 
Shgmator 15 l r.•e!me Leris • 2 
16 Pro,ector L~s # 1 
17 Pro.ector Leos 112 
,.. 
Figure 3.13 (a) : Schematic diagram of a transmission electron microscope 
electron s ignal is greatly magnified by a ene of electromagnetic lenses. The magnified 
transmitted s ignal may be observed in e ither an electron diffraction mode or direct imaging 
mode. Data is accumulated from the beam after it passes through the sample. The electron 
diffraction mode is employed fo r crystalline structure analysis, while the image mode is used 
for investigating the microstructure, e.g. the grain size and lattice defects [40] . A modern 
high-reso lution TEM goes down to a resolution < I 00 pm. While the ability to get atomic-
scale reso lutions from transmission e lectron microscopy is o f great advantage, TEM is 
relatively expensive equipment. It requires exten ive ample preparation, which makes it a 
relative ly time-consuming technique with a low output of samples. 
42 
Figure 3.13 (b): JEOL JEM-2 100 model transmission electron microscope 
3.4 Evaluation of Phosphor 
The re are several experiments that are completed to evaluate the overall performance of 
phosphors . This includes chromaticity, spectral di stribution, and life time of lumjnescence. 
3.4.1 Chromaticity 
A quantitative method has been established that re lates the color produced on d isplay screens 
to a standard value. The 193 1 Commission Internati onale de l'Eclairage (CIE) established a 
standard that defines chromaticity by x, y, and z coordinate . The chromaticity values are 
plotted on a two- dimensional graph, shown in Figure 3. 16, using the x and y coordinates. 
The third value, z, is fo und by knowing x+y+z=l. Thus, when mentioning the chromaticity 
of a phosphor, usuall y the x and y coordinates are onl y given. T he CIE diagram also shows 
43 
how red, green, and blue blend to give all hue needed fo r image production on a display 
screen. 
Jn addition to using chromaticity to define color, for incandescent bodie it is also possible to 
use the color ten1perature to define its color. As seen in Figure 3.1 7, the black body curve is 
in the sequence o f black, red, orange, ye llow, white, and blue-white, which corresponds to the 
increasing temperature of an incandescent object as it radiates thermally. When an object is 
heated to emit light that correspond to the black body curve, its temperature is defined as the 
color te mperature. Therefore, the color temperature of an incande cent light source is the 
temperature of a body on the black body curve that ha the same color, or chromaticity on the 
di agram, as the light source. The correlated color temperature of a li ght source , then, i 
defined as the temperature of the body that is not on the black body curve with a color that is 
closest to the light source. T ypical color temperature of the white region in the diagram 
range between 2500 and !0000 K [41] . 
09 
08 
07 
06 
05 
y 
0 4 
03 
02 
0 1 
00 
00 0 1 02 0 3 0 4 05 06 07 08 
x 
Figure 3.14: CIE chromaticity chart 
44 
3.4.2 Spectral Distribution 
The pectral distri bution gives in fo rmat ion about the characteristic emiss ion of a pho phor. 
The visible light emission of phosphor · i within 400- 700 nm. This range of wavelengths is 
divided into ix major divisions for the fo llowing colors: red, orange, yellow, green, blue, and 
violet. An example of this range with its corresponding colors is shown in Figure 3. 17. The 
color that each phosphor emits corresponds to a wavelength within the visible light spectrum. 
Figure 3.15: Visible light spectrum and corre ponding wavelengths 
45 
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26. M Endo, R.F. Walter, Gas Lasers, CRC Press, New York, 2006 
27. K.R. Nambiar, Lasers: Principles, Types and Applications, New Age International 
Publishers, New Delhi, 2004 
28. 0 Svelto, D.C. Hanna Principles of lasers, Fourth Edition, Kluwer Academic/Plenum 
Publishers, 1998 
29. T Katsumata, S Toyomane, R Sakai, S Komuro, and T Morikawa, J. Am. Ceram. Soc., 
89 (3) (2006) 932- 936 
30. The International Pharmacopoeia, Organization Mondiale de la sante, World Health 
Organization, Fourth Edition, Volume I, Geneva, 2006 
rd 
31. B.D.Cullity and S.R.Stock "Element of x-ray diffraction" 3 edition, Practice-Hall lnc. 
New Jersy. (200 1) 
32. M.R.Sardela Jr. , "The growth, thermal stability, structural and electrical properties of 
doped and undoped Si-based epitaxial structures" PhD thesis Linkoping university 
Press, Dissertation No.341 (2000). 
33. L.T Lin, Y Hu and TH Hsieh, Rev.Adv. Mater.Sci, 5 (2003) 464 
47 
34. T Fukuyo and H Imai, Journal ofC rystal Growth, 241 (2002) 193 
35. K.D Kim, D.N Han and H.T Kim, Chemical Engineering Journal, 104 (2004) 55 
36. M Rong, M Zhang, H Lui and H Zeng Polymer, 40 ( 1999) 6169 
37. G Carotenuto, G.P Pepe and L Nicolais, Eur. Phys.J., 16 (2000) 11 
38.C Luo, Y Zhang, X Zeng, Y Zeng and Y.J Wang, Colloid lntetface Sci., 288 (2005) 444 
39. W Songping and M Shuyuan, Materials Chemistty and Physics, 89 (2005) 423 
40. L Suber, l Sondi E Matijevic and D.V Goia, Journal of Colloidal and Interface Science, 
288 (2005) 489 
41. A Sileikaite, I Prosycevas, J Puiso, A Juraitis and A Guobiene, Materials Science 
(Medziagotyra), 12 (2006) 287 
48 
Chapter 4 
Synthesis and Characterization of Y20 3:Eu
3
+ phosphors Using Sol-
combustion Method 
4.1 Introduction 
Global claim for phosphor materi als as efficient sources of energy that can supply sustained 
competence is growing day by day. The phosphors are facing increased global challenges 
including high production of rare earth materials, environmental and recycling issues, and 
necessity to supply devices very quickly that may be outdated rapidly due to new 
technological developments arisi ng in the industry and market. A number of appl ications 
have emerged in recent years that will change the future of the industry and new technologies 
like innovations and specialty phosphors are garnering increased attention. The primary 
drivers for growth are the expansion of key end-use applications including solid-state lighting 
and fluorescent lighting. Current research in technology is focused on new materials, novel 
phenomena, new characterization techniques and fabrication of devices. The manufacturing 
of the new generation of display equipment in recent years raised a stringent demand on the 
fluorescent materials [ I). Particular interest is focused on the high-brightness phosphors 
having sub-micrometer particle sizes. Conventional solid-state synthesis of phosphors 
fo llowed by milling generally results in a dramatic decline of emission intensities for l -3 µm 
size particles due to the defects introduced by crushing. Therefore, wet-chemical methods 
seem to be an attractive alternati ve to the classical approach. The attempts to prepare fine 
powders of Y 20
3
3 : Eu +, perhaps the simplest ox ide among industrially used phosphors, 
usually lead to broadening of the emission peaks and overall decrease of the luminescence 
perfo rmance.2 3- Trivalent-europium-doped yttrium oxide (Y20 3:Eu
3+) is an important 
phosphor system extensively applied in co lor-television picture tubes owing to higher 
luminescent efficiency and the saturation degree in color. 
Physica B: Physics of Condensed Matter, Volume 439, pp. 181-184. DOI: 
10.1016/j.physb.2013.l 1.051 
49 
Y20 3, with a band gap of 5.8 eV can be considered as a large band gap semiconductor system 
[4 ]. In this paper we investigate how the structure, morphology and the luminescence 
intensities of these phosphors are affected by varying fuel to host ratios. 
4.2. Experimental details 
4.2.1 Synthes is procedure 
All the chemicals used for the preparation of the powders were of analytical grade. It includes 
yttrium nitrate (Y (N0 3)3.6H20), thiourea (NH2CSNH2), europium nitrate ( Eu(N0 3)3.6H20), 
ethanol and disti lled water. During the preparation of the micro-powders, thiourea was used 
as a fuel. Y20 3: Eu
3- micro- powders were prepared as fo llows; 
Y(N0 3)3.6H20 , NH2CSN H2, Eu(N03)3.6H20 , ethanol and distilled water were mixed in 
required stoichiometric ratios and dissolved by stirring using a magnetic sti rrer for 5 - 10 
minutes.The mixture was heated in an air tube furnace to ignition temperature of 400°C. A 
white foamy product was obtained after the combustion reaction. 
4.2.2 Characterization 
The Photoluminescence (PL) spectrum as well as decay curves fo r all the samples were 
investigated by Cary Eclipse fl uorescent spectrophotometer equipped with a 150W xenon 
lamp at an excitation source with the slit of 1.0 nm and scan speed of 240nm min-1• To 
determine the average particle diameter and the phase of the samples, X-ray powder 
diffraction(XRD) spectra were measured with a 0 8 Bruker Advanced AX S GmbH X-ray 
diffractometer using Cu Ka radiation at a wavelength of O. l 54056nm, the size and 
morphology of the as-prepared particles were carried out by using a Scanning electron 
microscope (S EM), SHIMADZU SSX-550 SUPERSCAN. 
4.3. Results and discussion 
4.3.1 Crystal structure 
The XRD pattern of a representati ve as-prepared sample is shown in Fig. I . All peaks can be 
perfectly indexed as the cubic phase with space group D d3 (p31113 ) and the cell parameter 
50 
a0=10.60A (a=b=c) is in good agreement with tandard Y20 3 JCPDS cards (No. 41- 1105). C3 
symmetry operation o f y 3+ ex ists in the cry ·tal, and no new crystal phase arose by doping 
with Eu3+. 
,........., 
C\J 
C\J SIY= 1.9 ....- C\J 
:::::l ,........., 0 
-cu ~ ~ ,........., ......... T""" 
~ 
+-' T""" C
Cf) ....
\..J.. . 
c 
Q) 
+c- ' 
Q) 
> 
+c-u'  JCPDS CARD NO- 41-110 
Q) 
a: 
20 30 40 50 60 70 80 90 
Wavelength (nm) 
Figure 4.1 : Representative XRD pattern of one of the sample w ith S/Y=l.8 molar ratio 
obta ined by Sol- Combustion method. 
The grain sizes ranging between 158 and l 80nm were estimated by using Scherrer ' 
equati on. ln order to investigate marginal decrease, the c rystallographic unit cell volume was 
determined by using the fo llowing equati on; 
1 h 2 + k 2 + [2 
dz az 
Ir is observed that for the Y 20 3 microstructure , there is a marginal decrease (-0.29%) in 
crystallographic unit-cell that tends to contract due to the increase in surface area of the 
grains. This may lead to a decrease in the lattice constant. No peaks attributable to other types 
of oxides are observed in the XRD patterns, indicating the high purity of the phases obtained. 
51 
Table 4.1 
The average grain size of the synthesized Y 20 3: Eu
3
+ nanoparticles as a function of S/Y 
molar ratio. 
Average grain sizes as a function of S/Y molar 
ratios 
S/Y ratio <222> <400> <440> <622> 
1.7 18 1 178 176 179 
1.8 172 172 175 177 
1.9 178 178 179 180 
2 178 178 179 180 
2.5 15 1 151 158 164 
4 150 152 155 154 
4.3.2 Morphology 
The microstructures of Y 20 3: Eu3+ are studied by SEM patterns and presented in Fig.2. It can 
be seen from the images that the samples possess a number of micro particles when the S/Y 
molar ratios are low ( 1.9 and 2.0). With the increase in the S/Y molar ratio to 2.5, the 
particles seem to agglomerates to fo rm nanorods-like structures. Further increase in the S/Y 
molar ratio to 4 reveals agglomeration of the particles. Thus, S/Y molar ratios seem to have 
an effect on the morphology of the phosphors. This current method also allowed obtaining 
agglomerates of contro lled size since synthesis time can be controlled; thi s being an 
important advantage compari ng with other methods [ I, 10, and 15]. An additional advantage 
of spherical agglomerates is its excellent fluid ity ( 18), which is beneficial for the elaboration 
of pressed bodies. 
52 
Figure 4.2: The SEM images of the Y 320 3: Eu + with (a) 1.9 (b) 2.0 (c) 2.5 (d) 4 .0 S/Y molar 
ratios. 0.5nm field of view. 
4.3.3. Photoluminescence 
Upon excitation, Eu3+ doped systems indicate many Stark energy levels in the visible region. 
In particular, luminescence transitions corresponding to the 5D 70---+ F1 manifold in the orange-
red region are of practical significance [ 16]. Fig.3 shows the emission spectrum of the Eu3+ 
ion doped in a Y 20 3 matrix with different S/Y molar ratios. The figure indicates intense red 
emission around 626nm due to the 5D 70---+ F2 transition, and a line e mission around 588 
nm.The peak around 618nm is due to e lectric dipole trans ition of 5 D 70 ---+ F1which is induced 
by the lack of inversion symmetry at the Eu3+ sites and is much stronger than the 5D 70---+ F 1 
transition. 
53 
200 .... 
,._ LL 1 . 8 
180 
- 1 . 9 1 60 ::::i 2. 0 
..c...t.l. .. 140 2. 5 
->-.Cii 120 4 .0 
-c Q.) 1 00 c 
~ 80 
ctl 
Q.) ... 
0.. 60 ,._ LL ,._ LL 
Q.) 0 
0 
> 0 40 0 
~ "' "' ctl 
Q.) 20 
a: 
0 
500 600 700 
Wave length (nm) 
0 .9~~~~~~~~~~~~~~~~~~~~~~~~~~~~---. 
• 1 
• 2 
• 4 
•5 
0 . 6 
>-
w 
(..) 
0 . 3 
0 0 .2 0 . 4 0.6 0.8 
CIE X 
Figure 4.3. (a): Emission spectrum of the different S/Y molar ratio Y20 3: Eu
3
+ phosphor 
excited at 260nm obtained by the Sol-Combustion method. (b) CIE coordinate diagram of the 
different emissions as indicated. 
54 
It is well known that the5Do~ 7F2/5Do~ 7F 1 intensity ratio is a good measure of the site 
symmetry of the rare-earth ions in a doped material. This is because the hypersensitive 
transition 5D0~7F2 tends to be much more intense at a site with no inversion symmetry, while 
the magnetic dipole transition 5D0~7F 1 is constant, regardless of the environment [1 3]. At 
S/Y= 1.8 phosphors have a red emission intensity of transition was the strongest due to the 
dominance of the5Do~7F2 transition. At the S/Y= 4 molar ratio the peak splits due to the 
magnetic dipole transition. Therefore, to have some insights of various luminescence 
transitions of Eu3+ doped in a matrix, this transition can be used as an internal standard [14, 
16, and 17]. Relative intensity of the hypersensitive transition with respect to the magnetic 
dipole transition as internal standard would give an idea on the transition strength of the 
hypersensitive transition. 
rn order to investigate the effect of S/Y molar ratios on the maximum intensity of the 
phosphors, a graph of maximum peak intensity verse various molar ratios was plotted as 
shown in figure 5. The plot indicates that the maximum peak intensities are higher at lower 
S/Y ratios(S/Y= l to 2.5). It drops drastically for higher molar ratio of S/Y= 4.This is in 
agreement with the previous discussion that S/Y = 1.8 molar ratios give the highest intensity 
and longer afterglow as pointed out in the next section. 
4.3.4. Afterglow decay curves of the red phosphors 
The afterglow properti es of samples with different S/Y ratios are compared, as shown in 
Fig.4. It can be seen that afterglow origin of the sample of S/Y= 1.8 and 1.9 has highest 
afterglow and brightness, while the sample with the ratio of S/Y=4.0 has the lowest afterglow 
and brightness. Thus, the ratio S/Y= 1.8 was taken as the optimum. The decay times of the 
phosphor can be estimated by using the fo llowing double exponential equation; 
where 1 is the phosphorescence intensity, Ai, and A2, are constants, t is time, r 1andr2, are 
decay times fo r exponential components, respecti vely. The fit ting results of parameters t 1and 
t2are listed in Table 2 below. 
55 
60 S/Y molar ratio= 4.0 
...- \ - 2 .5 2.0 
:::J 50 ' - - 1 .9 
.c..o_  - - 1 .8 
->. 40 (/) 
-c Q) 30 c 
-Q) 20 > co 1 0 
Q) 
a: --
0 
0 20 40 60 80 100 
Time {ms) 
Figure 4.4: The decay cu rve of Y20 3:Eu
3
+ phosphor. 
S!Y 1.7 1.8 1.9 2 2.5 4 
Components Decay constants(T, s) 
Fast (Tl) 1.3729 1.37 15 1.3656 1.364 7 1.3598 1.3579 
Medium ('t2 ) 1.4622 1.4588 1.457 1 1.4537 1.4499 1.3989 
Table 4.2 
Decay constants for the fitted decay curves of the phosphor powders with different S/Y molar 
ratios. 
56 
- 190 Peak intensity . :J.  180 • 
._ca. . . \ 
170 
Z' 
"c:'  160 
..C..l,)  •• • 
c: 150 
~ 
ca 140 
Cl) 
c. 130 
E 
:J 120 
·E 
c>
-
a<  110 
~ • 
100 
1.5 2.0 2.5 3.0 3.5 4.0 
S/Y molar ratios 
Figure 4.5: Effect of S/Y molar ratios on the intensity of the broad PL peaks and 
corresponding emission wavelength. 
4.4 Conclusion 
A very simple and effi cient chemical route to prepare a promising afterglow red phosphor by 
sol- combustion synthesis is presented. Ethanol has the effect of decreasing the water needed, 
hence simplifying the experimental procedure by dissolving rare earth nitrate and sulfur-
contained organic fuel into an even solution. This prompts the formation of rare earth ox ide 
by igniting first during heating that leads to a combustion decomposition reaction. Y2 0 3 : Eu
3
-
microcrystall ine structures were obtained using thiourea as organic fuel. The increase in the 
57 
ratios of fuel to the host decreased the grain size and the maximum PL intensity of the 
phosphors. 
References 
[ 1] K. Narita. Phosphor Handbook; CRC Press LLC: Boca Raton, FL, 1999. 
[2] M.R. Royce, US Patent no. 341 8, Vol. 246, 1968. 
[3] S.H. Cho, Y.S. Yoo and J.D. Lee, J. Electrochem. Soc. 145 (1998) 10 17. 
[4] M. Mikami and A. Oshiyama, Phys. Rev. 857 (1998) 8939. 
[5] Y. Zhiping, F. Jianwei, L. Xu, Journal of Heibei Uni versity (Natural Science Edition), 
25(3) (2005) 262. 
[6] F. C. M. Van de Pol, Ceram. Bull. 69 (1990) 1959. 
[7] K. Nashimoto, S. Nakamura, H. Mariyama, Japan. J . Appl. Phys. 43 ( 1995) 509 1. 
[8] T. Nagata, T. Shimura, A. Asida, N. Fujimura, T. Ito, J. Cryst. Growth 237 (2002) 537. 
[9] A.J. Freem an, K.R. Paeppelmeier, T.O. Mason, R.P.H, Chang, T.J. Marks. Mater. Res. 
Soc. Bull. 25 (2000 45. 
[ 1O ] L. Spanhel , M.A. Anderson, J .Am.Chem. Soc., 113 ( 199 1) 2826. 
[11 ] P.P. Hoyer, H. Wetter, Chem. Phys. Lett. , 22 1 ( 1994)379. 
[12] P. Hoyer, R. Eichberyer, H. Weller, Ber. Bunsen- Ges. Chem. Phys. Chem., 97 (1993) 
143. 
[ 13] A. Eri c, M. Eulenkamp, J. Phys. Chem B, 102 (2002) 5566. 
[14] D.C. Reyolds, D.C. Look, B. Jogai, H. Mokoc, Solide State Commum lO I (1997) 643. 
[ 15] M. Liu, A.H. Kitai, P. Mascher, J. Lumin. 54 (1992) 35. 
[16] Z.W. Jin, Y.Z. Yoo, T. Sekiguchi, T. Chikyow, H. Ofuchi ,H. Fuj ioka, M. Oshima, H. 
Koinuma, Appl. Phys. Lett. 83 (2003) 39. 
[ 17] X.T. Zhang, Y.C. Liu, J.Y. Zhang, Y.M. Lu, D.Z. Shen, X.W. Fan,X.G. Kong, J. Crystal 
Growth 80 (2003) 254. 
58 
[18] H.J. Lee, S.Y. Jeong, C.R. Cho, C.H. Park, Appl. Phys. Lett. 81(2002)4020. 
59 
Chapter 5 
Characterization of Eu3+ activated lanthanum oxysulfide synthesized by sol-
combustion method 
5. I. Introduction 
Lanthanide (La) oxysulfides with high thennal and chemical stabili ty are known as a wide-
bandgap (4.6 - 4.8 eY) material suitable for doping ion activation. Compared with the 
lanthanide oxides, oxysulfide is a more effi cient phosphor with a broader excitation band. 
Therefore, the lanthanide oxysulfides become a very important family of inorganic materials 
that have high potenti al for applications in various fi elds, such as color television picture 
tubes [ 1, 2], radiographic imaging [3], field emission displays [ 4, 5] and long-lasting 
phosphorescence [6]. Among them, Eu3+ activated lanthanide oxysulfide has been extensively 
investigated due to its excellent efficiency to be used as a red phosphor applied in television 
picture tubes. Currently, studies on the luminescent properties of micro phosphors are 
attracting interest; because of its signifi cance not onl y fo r applications but also for the 
essential understanding of micro crystals, such as the quantum confi nement and surface effect 
[7, 8]. Among them, rare earth doped micro phosphors have attracted particular attention [9, 
1O J. It is expected that in the microsized phosphors, the luminescent quantum yield as well as 
the resolution of display to be considerably improved. Some rare earth ions, such as Eu3+, 
may act as common activators to detect local environments [ I I] due to their supersensitive f:f 
transitions. Until now, a great number of rare earth doped microsized phosphors have been 
prepared and studied [ 12-1 3] especially the ro le of doping concentrations, crystal sizes and 
crystal structures on the luminescence in several oxide hosts have been studied extensively 
[ 14-1 5]. The reports on microsized lanthanide oxysulfides are rather scarce. However, in the 
past few years, there appeared only a few papers regarding the preparation and 
photo luminescence (PL) properties of rare-earth doped lanthanide oxysulfides microcrystals. 
It has been shown that the successful synthesis of lanthanide oxysulfi de microparticles can be 
60 
achieved by a number of processes, including sol-gel [ 16], gas-phase condensation methods 
[ 17] or colloidal chemical methods [ 18-19]. But efforts to make concentrated colloidal 
solution of highly uniform size luminescent micro-oxides in the past have met with technical 
difficulties. To our knowledge, there is no paper reporting the synthesis of La20
3
2S:Eu + by 
sol-combustion method which prompted us to attempt it. 
In this paper, the original sol -combustion approach for the quantitative synthesis of 
lanthanum oxysulfide microparticles is described. Preliminary results of some optical and 
structural investigations as a function of La/S concentrations are reported. 
5.2 Experimental 
Lanthanum (La) oxysulfides phosphor (La 0 S:Eu32 2 +) was prepared by a sol-combustion 
synthesis method. High purity (Aldrich make, 99.99%) raw materials, lanthanum nitrate (La 
(N03)3.6H20), europium nitrate (Eu (N03)3.6H20), ethanol (C2H50H) and thiourea 
(NH2CSNH2) , were used for preparation of the phosphor. The materials were weighed in 
stoichiometric ratios, dissolved in distilled water and mixed by stirring using a magnetic 
stirrer for 5 -10 minutes. The mixture was heated in an air tube furnace to an ignition 
temperature of 400°C. A white foamy product was obtained after the combustion reaction. 
Several samples with different La/S molar ratios were then prepared via a similar route. 
During the preparation of La20 2S, the doping concentration for all the samples was kept 
constant at 0.05 mo!%. 
5.2.1 Characterization 
The PL spectra and decay curve of the powders were obtained with a Cary Eclipse 
fluorescent spectrophotometer equipped with a 150 W xenon lamp as an excitation source 
with the slit of 1.0 nm and scan speed of 240 nm min-1• To determine the average particle 
diameter and the phase of the samples, X-ray powder diffraction (XRD) spectra were 
measured with a 08 Bruker Advanced AXS GmbH X-ray diffractometer using Cu Ka 
radiation at a wavelength of 0. 154056 nm, the size and morphology of the as-prepared 
particles were carried out by using a Scanning electron microscope (SEM), SHIMADZU 
SSX-550 SUPERSCAN. The thermoluminescence (TL) spectra were obtained using a 
Thermoluminescence Reader (Integral- PC Based). The XPS analysis was carried out with a 
61 
PHI 5000 Versa probe-Scanning XPS Microprobe. Fourier Transfonned Infra-red (FTIR) 
spectroscopy was done with a Bruker Tensor 27 FTIR Spectrometer. 
5.3 Results and Discussion 
5.3. l Crystal structure 
Figure 5.1 shows the XRD patterns of the La202S powders synthesized by the sol-combustion 
process at 400 °C. The average crystalline size of the La20 2S particles calculated using the 
most intense reflection at 20= 25.596° are tabulated in table 5.1. Estimated according to the 
Scherrer'sequation, the average crystalline size of the powders is detennined to be 178 
nm.The cell constants a and c were obtained using equation 1 below; 
2 2 2 2
Sm. 2  8 = -A. (-4h +hk+k2 + -1 } 2 ----------- (5.1) 4 3 a c
The most dominant peaks can be indexed as the hexagonal La20 2S phase with space group 
P3m l (164] and cell constants (a) 0.4 128 nm, (c) 0.6985 nm, which are close to the reported 
data (JCPDS Cards File: 27-0263). The other peaks marked * are indexed as La20 3. It must 
also be noted that a s light shift in the XRD peaks occurred, as pointed out in table 5.1, with 
an increase in the La/S ratio. It is well-known that thiourea reacts with water to produce 
gaseous NH3, H1S, and C02: The produced gas NH3 wi ll react with the crystal water. 
Therefore, a reaction may occur between lanthanum ni trate and the crystal water in the 
system at elevated temperature to produce La (OH) 3 : 
3 + 3 la{N03}3~ H20-tla{OH}3(s)+3N02 (g) -02 (g)---------------- (5.2) 2 4 
Since C02 and H2S are both acidic, reactions between La (OH) 3 and them immediately 
occurred; then the following equilibrium were fonned : 
la{OH}3 (s) + C02 (g) -t La{OH}C03 (s) + H2 0(g)------------------ (5.3) 
2la{OH}3(s) + H2S(g) -t La2 0 2S(s) + 4H2 0(g) -------------------- (5 .4) 
The equilibrium given in eq.5.4 is more likely to move to the right than that expressed in 
eq.5.3. Therefore, the final product was crystall ized La20 2S, with small traces of La20 3 phase 
62 
as detected by XRD. The small trace of La20 3 were formed due to the decomposition of 
La20 2S at elevated temperatures according to eq.5.5 below; 
..-....  
::J.  
..c_o.. . * * 
:>t:- 
(/) 
c: 
..(..].),  La/S=--1 .0 
c: --- 1.7 
Card No: 27-0263 --1.8 
(]) 
--- 1.9 
..>...,  ...... 
co N ---2.0 - = (]) O' =.....  ~--2.5 QM 
0:: = ~~ ~ ~ 
10 20 30 .l O 50 60 70 
2theta( degrees) 
Figure 5.1. X-ray diffraction patterns o f La20 2S with different La/S ratios as well as the 
standard XRD pattern 
63 
12000 
11000 ~~ 1.0 
10000 1.7 
1.8 
9000 
- 1.9 8000 :::::J -- 2.0 
~ 7000 .._ 2.5 
>. 6000 
+-" 
(/) 
c 5000 
Q) 
+- 4000 c " 
3000 
2000 
1000 
0 
25.2 25.4 25.6 25.8 26.0 
28 (degrees) 
Figure 5.2. X-ray diffraction powder patterns at(*) plane for different La/S mole ratios. 
12000 
~~ 1.0 1.7 
10000 -- 1.8 
- 1.9 
:::::J 8000 - 2.0 
.~._  2.5 
>. 
+-" 6000 
(/) 
c 
Q) 
+c- " 4000 
2000 
25.5 25 .6 25.7 25.8 25.9 26.0 26.1 26.2 26 .3 26.4 26.5 
28 (degrees) 
Figure 5.3. X-ray diffraction powder patterns at ( 10 I) plane for different La/S mole ratios . 
64 
Figure 5.2 shows a graph of XIX+Y versus La/S molar ratios. X denotes the intensity of the 
most pronounced peak of La20 2S which is in the [ 101] direction, while Y is the intensity of 
the most pronounced peak of La20 3. According to Figure 5.2, the ratio of the two intensities 
drop gradually between La/S= l tol.7 and then drastically from La/S = 1.7 to 2.5. This shows 
that at lower ratio of fuel to oxidizer favors the formation of the La20 2S rather than La20 3 
phosphor. The diffraction peaks shift slightly towards lower angles with an increase in La/S 
mole ratios as shown in Table 5.2, indicating that the lattice parameters are slightly bigger 
than those of lower mole ratios, which is mainly due to larger radius of La3+ 1.18 A than that 
of Eu3+ (0. 95 ° A). This indicates that the dopant ions are well incorporated into the lattice 
sites of La3+ at lower ratios and lead to the decrease in inter-atomic distance. 
0 85 • X= Intensity of [101] Y= Intensity of [0 11] 
0 80 
~ 0.75 
• 
+x \ -.x.....   0 70 \ 
"'. 
0 65 --------. 
oso ......- -...........- -.........., ..... ........., ..... ..................................... ... 
0 8 1 0 1 2 1 4 1 6 1 8 2 0 2 2 2.4 2 6 
La/S molar ratio 
Figure 5.4.The effect of fuel on the formation of La20 2S and La20 3 prepared by the sol-
combustion process. 
65 
Table 5.1.The concentration and calculated crystalline size of Eu3+ 10n doped La20 2S 
microcrystals. 
La/S Grain size (nm) 
--------------------------------------------------------------------------
1.0 25.897 0.3786 178 
1.7 25.827 0.4273 175 
1.8 25.825 0.4519 185 
1.9 25.8 19 0.2411 175 
2.0 25.805 0.2256 182 
2.5 25.798 0.2045 184 
5.3.2. Fourier transforms infrared spectroscopy 
The FTIR measurements have been made in the wave number range 400 cm-1to 4500 cm·' in 
order to identify the presence of La20 2S. The broad absorption band around 3420 cm-
1 can be 
ass igned to 0 - H stretching vibration; the bands around 1500 and 1387 cm- 1 result from C- 0 
asymmetri cal stretching vibration; the peak that appears at 1074 and 1114 cm- 1 can be 
assigned to C- 0 symmetric stretching vibration; the peaks at 856 and 670 cm- 1 correspond to 
C- 0 deformation vibrations (20]. In the spectrum of the samples, the peak around 504 and 
620 cm- 1 associated with the vibration of La-0 and La-S [2 1, 22) is observed, indicating the 
formation of La20 2S. It was observed that an increase in Lal S molar ratios reduce the 
absorption band at 504 cm·' . 
66 
1.0 
• 
0.8 .)"~ (   .JJ 0 . "• q- ,. ,, Cl) in (.) • •' c 0.6 ',  ,',,  ,I  co N \ \ \ . /\ .',  
=m  . •'  •.,,   I \ \ ' .•I .•  I  ' ' I E ' \ '', 
"c'  0.4 
., ' 
• • • 
m 'I \ I I ' La/S 
"-
I- vl t'C '' O 0 N 1.0 0.2 Mc o q-0 
~q- ,...
  ,i.n..  M - - - - 1.9 - . - 2.0 
,0..,.,... ,......  - - 2.5 
0.0 
1000 2000 3000 4000 
Wavenumber (cm-1) 
Figure S.S. Fourier-transform infra-red spectroscopy spectra of the as-prepared La20 2S: 
Eu3+p owders for various La/S mole ratios. 
S.3.3. X-ray photoelectron spectroscopy 
XPS results fo r the powders of La20 2S:Eu
3+, to identify the presence of e lementary states in 
the oxysulfides, are discussed in this section. Figure 5.6 (a) shows the wide-scan spectra of 
the hexagonal La20 2S:Eu
3+microstructures for the sample with La/ S mole ratio of 1.0 in the 
range of 0- 1400 eY. All the expected e lements can be identified from the survey. Figure 5.6 
(b-d) shows the detailed regions of the XPS spectra of the hexagonal 
La 320 2S:Eu +microstructures. The binding energy (BE) o f the phosphors was determined with 
re fe rence to the C 1s  peak at 284 .6 eY, the de-convolution not shown. The high resolution 
XPS spectrum corresponding to the La3d states is characterized by a main doublet composed 
of s ix peaks situated at 833.2, 834.6, 837.9, 850, 85 1.8 and 854.7 eV when fitted by Gaussian 
67 
de-convolution. The peaks at 833.2 and 850 eY are related to 3ds12 and 3d312 bonds in La20 J. 
respective ly. The La 3d peaks show doublets attributed to spin-orbi t coupling splitting of the 
3d sublevels [23]. These are fitted with La 3d512 peaks (837.9, 850 eV) and 3d512 (85 1.8, 854.7 
eY). The fitting also confirms that La3+ ions occupied two diffe rent site . Three distinct 
components of 01 s peak can be consistently fitted by Gauss ian de-convolution , cente red at 
528.9, 530.8 and 532. 1 eY , respecti ve ly. The first BE 528.9 eY is attributed to the La20 3. 
The other two components at the higher BE of 530.8 and 532. l eY can be attributed to the 
La 320 2S:Eu + [24, 25]. The S2p spectrum presents two components peaked at 168.7 and 
169.9 eY. The peaks may be attributed to 2p312 and 2p 112 bonds in La20 2S:Eu
3
+. The S2p 
peaks are slightly weak as compared with that of the othe r major lines . The composi tion 
estimated by XPS using the relative sensiti vity factors o f 0 and S also revealed excess 
oxygen in the samples. The XPS results confirm that some impurit ie are formed (La20 3) 
after combustion synthes is at initial temperature of 400 °C under ambient condition which i 
cons istent with XRD data. 
lSOGJ C1s 01s 
(i) 7B6 (ii) - r 
500 
100 , 
400 r 
u"'  -5 300 
so ~ ' ' 200 · 
100 I '\ '"' 
0- -'--- 0 - l 
288 286 284 536 534 532 530 528 526 
Binding Energy (eV) Binding Energy (eV) 
S2p La3d 
~·~ 
(iii 2500 6 'T (iv) 100 ~ ~ /' 2000 'I 
80 . " ~~ 
1500 ' 
.....I  
u"'  60 • -!!! I <.> 
40 j, 1000[ 
20 l ,'.  '. ! j 500 I • 
0 ' _J j __,_ o ~ -l _J 
175 170 165 160 865 860 855 850 845 840 835 830 
Binding Energy (eV) Binding Energy (eV) 
Figure 5.6 (a). Graph of (i) C 1s  (i i) 0 Is (iii) S2p and (iv) La3d fo r La20 2S microcrystals 
prepared with a La/S ratio of 2.5. 
68 
x 4 
T 
i I ~omil -0 Ols (a) Cls 
La3d5 
S2p 
cl 3 
-01 5 
s 
f -C1s 
'-"'. . ~ 
~· --·-'\ 
... -S2p , .,. f',~ _,.,..,_ · 1 
' \... \ l 
' • 
0 I _....__ J 
140 120 100 8 6 4 2 0 
Bindina Enerav 
Figure 5.6 (b). XPS survey pectrum of the La20 2S microcrystal s prepared with a La/S ratio 
o f 1.0. 
69 
.. --1 
1400 . 
(c) La20 2S 
La20 3 
'"'l 
1000 La20 3 
000 > 
C/s6oo I 
I ..... 400 
200 . 
ol ~ 
-200 t 
I --1 ..l_ I 
865 860 855 850 845 840 835 830 825 
Binding Energy (eY) 
Figure 5.6 (c). La 3d XPS peakfitted with peaks for the La20 2S and the La20 3_ 
1200 
11odf (d) 
La20 2S 
10011 
J 
900 ~ 
800 La20 3 
els 700 ~ 
600 
500 
400 .·.• · ··········..   .... I 
300 ······ 
... -· ·· ... 
•. ..........·· .. . ·.. ... ·· 
j j _L J 
_!_ ---'---
538 536 534 532 530 528 526 
Binding Energy 
Figure 5.6 (d). 3XPS pectra of La (3d5h) and 0 ( Is) o f the a -prepared Eu +- doped 
70 
La20 2S microcrystals.La 3d region for La202S and La20 3 with the peak fitting components 
fo r the 4psn peak. 
36 ~ (e) 
La 20 25 
34 
32 ~ 
els 
30 
28 I 
26 l 
24 ' L j I l J 
178 176 174 172 168 166 164 162 16 158 
Binding Energy (eV) 
Figure 5.6 (e). XPS S2p peak for La20 2S with the peak fitting components fo r the S2p512 
peak. 
5.3.4 Morphology 
Figure 5.7 shows the SEM image o f the La20 2S: Eu
3+microcrystalline powders. It is obvious 
that all the powders yield microparticles and they tend to aggregate together. The images 
show that the morphology consisted o f a foamy agglomeration and a continuou three-
dimensional network at lower fuel to ox idizer ratio. At higher La/S molar ratio the 
morphology appear as regular crystalline latti ces. 
71 
Figure 5.7. SEM micrographs of the as-prepared La 320 2S: Eu +powders with La/S molar 
ratios of (a) 1.0, (b) 1.8, (c)2.0, (d)2.5 with 5000 nm fi eld of view. 
5.3.5 Photoluminescence 
The excitation spectrum of the phosphor in the spectral region 200- 400 nm is shown in 
Figure 5.8 (a) . The spectra mainly consist of two broad peaks, one located near 242 nm and 
the other near 328 nm. The band near 242 nm is attributed to the 0 2 -Eu3- + CTB (charge 
transfer band), while the one near 328 nm is due to the La20 2S host absorption. 
72 
400 
E u s= 2 
c: E 
350 c: u s= 3 E N 
c: "'2' LIS= 1.9 N 
- 300 :::J.  
-('CS 250 >. 
·-c;; 
-c: 
200 
<1> 
c: 150 
<1> 
·->- ('CS 100 
Ci) 
0::: 
50 
0 
200 220 240 260 280 300 320 340 360 380 400 
Wavelength (nm) 
Figure 5.8 (a). Excitation spectra of La20 2S with different La/S molar ratios. 
1 20 
--L 1.8 
--L 1.9 
- 100 --L 2 .0 --L 2 .5 :::J 
-ca 80 --L3 .0 
.~ 
II) "' 
-c 60 Cuf?.  ~ 0 0 G) u. ...L L ~ ... ... c ...L L I 0 ' o 0 0 0 
0 0 ... 
-G) 4 0 0 
"' 
.~ "' "' /j ...L L.., 0 ca 0 
G) 20 "' a: 
\ 
0 
5 50 600 650 700 
Wavelength (nm). 
Figure 5.8 (b). PL emission spectra of La20 2S with La/S=l.8, 1.9, 2.0, 2.5 and 3.0 mo lar 
ratios. 
73 
Figure 5.8 (b) 7shows the emission spectra of Eu3+ which are assigned to the5Do- Fo, 1. 2, 3, 4 
transitions of Eu3+. There are three groups of distinctive emission peaks between 537 and 
624nm, which are rel ated to the 5D 70- FJ (J = 0, I , 2) transitions of Eu
3
+, respectively. The 
strongest emission peak near 615 nm corresponding to forced electron dipole transitions of 
Eu3+ is due to the 5D 70- F2 transition in the C2 symmetry of Eu
3
+ incorporated in La20 2S, 
which is hypersensitive to environmental effects [26, 27]. ln addition, there are other energy 
transitions of Eu3 5+ corresponding to Do- 7Fo (537- 555 nm), 5Do- 7F1- S6 in S6 symmetry 
(579-595 nm), 5D 7 50- F3 (~650 nm) and Do-
7F4 (688- 7 13 nm) in the luminescence spectra 
[28]. In the hexagonal La20 2S there are two crystallographic sites, one with C2 symmetry and 
the other with S6 symmetry. Therefore, it can be concluded that the Eu
3
+ ions replace La3+ 
ions and occupy both the C2 and S6 symmetries because the emissions assigned to both 
symmetries were observed (Figure 5.8 (b)). The La/S molar ratios may be affecting the 
crystallinity of the system and consequently the neighborhood around the Eu3+ ions. 
Compared with the 5Do- 7F2 transition, the intensity of the 5Do- 7F 1 transition corresponding 
to the orange color is much lower, which makes La 0 S:Eu32 2 + a more pure red phosphor. The 
emissions of the samples doped with Eu3+ differ from each other, the peaks for samples 
La/S=2.0, 2.5 and 3.0 are quenched at emission wavelength of 615 nm but slightly enhanced 
at 624 nm wavelength. The 5Do-7F2 trans ition is an electron dipole transition and thus is very 
sensitive to the symmetry in which the Eu3+ is situated. The peak intensity fo r samples 
La/S= l .8 and 1.9 are quenched at a wavelength of 624 nm. Bohus et al. [29] obtained 
emission spectra of core-shell Y 20 3 :( Eu
3+, La3+) samples and compared the spectra to the 
spectra of core-shell Y20 3:Eu
3
+ samples and found new emission peaks that were visible at 
613 and 622 nm. These emission peaks correspond to the 5D 70- F2 transition of the 
La 320 3:Eu + phosphor. In La20 3:Eu
3
+ phosphor Eu3+ ions replace La3+ ions at the C2 symmetry 
sites and that is why new emission peaks appear at 613 and 622 nm in their case. The changes 
in the relative ratio of the 5Do-7F2 to 5Do-7F1 emission and the 6 15 nm to 624 nm emission 
may therefore be explained by the fact that there were presence of both electric and magnetic 
dipoles which favor enhancement of intensities depending on the position of emission 
wavelengths and also the La/S mole ratios. The relative change in the 615 nm to 624 nm 
emission ratios is an indication that for the higher La/S ratios the Eu3 3+ occupied the La + 
position in La20 3. 
74 
520 
CIE Y 
700 
CIE X 
Figure 5.8 (c). CIE coordinate of emis ion of La20 2S phosphor. 
Figure 5.8 (c) represents the chromacity co-ordinates o f the PL pectra for the ample 
La/S= 1.9, which are determined using the C IE (International Commi ion on Illumination) 
Coordinate Calculator software. According to the software the position of the co lor 
coordinates (0.49, 0.45) lie well in the red/orange region. The detailed analysis of the 
phosphor finds its suitability in making the red/orange producing phosphor for di splay 
application · and light emitting diode. 
75 
-. 11 0 ::J. - 10 0 ~ •/ ."" 
• Emi ssion 615 
• ... Em ission 624 
-~ 90 
.."c:' 80 Q..,.)   ~ 
c: 
70 
~ \ 
~ 
Q,) 60 
a.. 
E 5 0 
::J \ • 
E .l 0 
>< 
~ ... ------------ • 3 0 
~ .... 
2 0 
1 8 2 0 2 2 2 .l 2 6 2 8 3 0 
La/S molar ratios 
Figure 5.9. Graph of max imum peak intensity versus La/S molar ratios. 
The graph o f max imum PL intensity of the as-prepared powders as a function of La/S molar 
ratios are shown in Figure 5.9. The emission peak intensity decreased when the La/S molar 
ratio increased, and a max imum value was fo und when La/S= 1.9, thereafter the emission 
intensity quenches graduall y. Persistent luminescence spectra of the phosphor powders are 
shown in Figure 5.8. It can be seen from the spectra that the powders showed di fferences in 
initial intensity and medium persistence when the powders were e ffic ie ntly acti vated by the 
UV lamp. The results indicate that the initial luminescence intensity and the decay time of 
phosphors are enhanced w ith a decrease in the La/S mo lar ratios. The Inset of Figure 5.8 
indicates the plot of In (natural logarithm) of intensities versus the decay times of the 
phosphor. The slopes o f the graphs decrease graduall y and bend to be almost horizontal 
showing a double exponential function. The decay behavior can be analyzed by cu rve fitting, 
re lying on the fo llowing double exponential equation: 
76 
300 La/S= 1 .8 
- 1 .9 :-J- 250 
-al =- 2.0 -" 0 >- 20 0 2.5 
+-' ·2 
(j) 
c 150 3 .0 
Q) 
+-' 10 
c Time (1r6) 100 
Q) 
> 
+-' 50 
al 
Q) 
a: 0 
0 2 4 6 8 10 
Time (ms) 
Figure 5.10. Afte rglow characteristics of La20 2Swith di fferent La/ S molar ratios.Inset: A 
graph of In Jog of intensity versus decay ti me showing a double exponential function. 
Where! is phosphorescence intensity, A 1, and A2 are constants, t is time, t 1 andt 2 are decay 
constants, deciding the decay rate for rapid and slow exponentiall y decay components, 
respecti vely. The fitting results of parameters t 1 and t 2are li sted in Tab le 5.2 below. 
Table 5.2. Results for the fitted decay curves of the phosphor powders with different La/S 
molar ratios. 
La/S 1.0 1.8 2.0 2.5 3.0 
Components( ms) 
Fast (T1) 0.97 0.95 0.9 1 0.88 0 .82 
Slow (T2) 1.08 0.99 0.97 0.95 0.89 
77 
Two components namely fast and slow are responsible for the luminescence properties of the 
as- synthesized phosphor. A trend can be observed (Table 5.2) that the decay constants of the 
phosphors decrease gradually with the increasing La/S molar ratios 
5.3.6 Thermoluminescence 
The peaks between 50 and 11 o0c in the TL plots have been previously associated with 
persistent luminescence [30]. Our TL data for all the samples exhibit peaks below this range, 
which means they have poor afterglow. According to [31 ], defects introduced by doped ions 
play a key role in long-lifetime light emission. The doped Eu3+ ions can replace La3+, creating 
traps of electrons and holes. Upon heating in the TL experiment, the trapped electrons and 
holes escape and recombine, resulting in the emission of light. The characteristic ionization 
energy of a trap can be estimated from the following equation: 
E = 2 k ( Tm) z;T z - Tm ----------------------- ( 5.7 ) 
Here k is the Boltzmann constant, Tm is the TL peak temperature and T2 corresponds to the 
point on the slope to the right of the TL peak where the intensity is half of the peak value. 
The related parameters are calculated using equation (5.7) and are listed in Table 5.3. 
78 
La :S = 1.8 
1200 0 1.9 
2 .0 
10000 2 .5 
3 .0 
- 8000 :::;; 
-~ 
·~v;  6000 
c: 
Q.> 
c: 4000 
r= 
200 0 
0 
20 40 60 80 100 120 140 
Temperature {°)C 
Figure 5. 11. Thermoluminescence plots of the La20 2S: Eu
3
+ phosphor. 
In order to achieve long-lasting phosphorescence, the trapping levels need to be located at 
suitable depths. If the trap level is too shallow, on ly a small amount of electro ns can be 
captured by the traps. Under the action of thermal di sturbances, e lectrons are easily released 
from traps and then recombine with holes in the ground state, which results in a shorter 
afterglow time. On the other hand, if the traps level is too deep, it would be difficult fo r the 
captured e lectrons to gain enough energy to return to the excited state leve ls at room 
temperature, which also results in a poor afterglow property [32, 33]. From Table 3, the 
sample with a La/S= 1.8 andl.9 ratios have the most suitable depth of 0.81 and 0.80eV, 
respectively which also give the highest peak intensities in the PL spectra as shown in Figure 
6(b) [34, 35]. The intensity of the sample with a La/S= 2.0 ratio was low due to a too deep 
trap level ( 1.8 eY), wh ile that of the sample wi th a La/S=2.5 and 3.0 ratio were lower due to 
shallower trap leve ls (0.58eV). 
79 
Table 5.3. Trap energy levels for different concentration of La20 2S. 
La/S Tm(k) T2 (k) E (eV) 
1.8 300.4 331.4 0.8 
1.9 350.0 376.2 0.8 
2.0 350.2 369.1 1.8 
2.5 327.3 378.2 0.6 
3.0 302.8 346.4 0.6 
5.4 Conclusions 
ln conclusion, trivalent-europium doped La20 2S microcrystals were prepared by sol-
combustion method and their material properties were studied at room temperature. The 
phosphors emit bright red phosphorescence originating from the 5D 70- F2 transition of 
Eu3+.TheXRD pattern of the as-synthesized powder revealed the presence of the hexagonal 
La20 2S phase with an average crystalline sizes of 178 nm with a space group P3m 1[  164] and 
cell constants of (a) 0.41 28 nm and (c) 0.6985 nm. The lower ratio of fuel to ox idizer favors 
the fo rmation of the La20 2S rather than La20 3 phosphor. De-hydration of the reactants was 
crucial for the successful synthesis of the oxysulfide phase. SEM images of the as-
synthesized powders showed that the morphology consisted of a foamy agglomeration and a 
continuous three-dimensional network. At higher La/S molar ratio the morphology appear as 
regular crystalline lattices. 
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Korea, (2003) 203. 
80 
[6] V.L. Levshi n, M.A. Konstantinova, E.A. Trapeznikova, on the application of rare-earth 
elements in the chemistry of phosphors, in: Rare-earth Elements, USSR AN Publi shing, 
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[7] T. Hisamune, Technical trend of phosphors for plasma display panels, in: Proceedings of 
the 9th International Display Workshops (2002) p. 685. 
[8] T. Juste!, H. Niko! and C. Ronda. Angew.Chem. Int. 37 ( 1998) 3085. 
[9] S. Shionoya and W.M. Yen, Editors, Phosphor Handbook, CRC Press, Boca Raton, FL 
( 1998). 
( 10] C.F. Bacalski, M.A. Cherry, G.A. Hirata, J. Mckittrick, J . Mourant, J. Soc. Inf. Display 
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Surfaces A: Physicochemical and Engineering Aspects, 405(5) (2012) 6-13. 
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( 1996) 62 l. 
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82 
Chapter 6 
The influence of oxygen partial pressure on material properties of EuJ+-
doped Y2 0 28 thin film deposited by Pulsed Laser Deposition 
6.1 Introduction 
Europium-doped Y20 2S exhibits strong UV and cathode ray-excited luminescence, so it is 
widely used as red phosphors for low-pressure fluorescent lamps, cathode-ray tubes and 
plasma display panels (PDPs) [l]. Also, the hexagonal Y20 2S is a good host material for rare 
earth ions. A lot of attention has been given to the nanoscale Y20 2S:Eu
3
+ for its tremendous 
potential applications in optical display and lighting materials and basic science research on 
special luminescent spectra (2-4]. Nanoscale Y20 2S:Eu
3
+ has remarkably different 
luminescent properties from those of bulk samples: such as emission line broadening, lifetime 
changes and its spectral shift [5]. There are several methods to synthesize nanocrystalline 
Y20 2S:Eu
3
+, such as sol-gel [6], combustion [7], micro emulsion [8], and spray pyrolysis 
method [9], but these methods are limited in the complexity of the preparation methods. Solid 
state reaction at room temperature is a good method to synthesize nanoparticles [I 0-12]. It is 
generally accepted that thin-film phosphors have several advantages over bulk-type powder 
phosphors: better thermal stability, reduced outgassing, better adhesion, and improved 
uniformity over substrate surface. However, the biggest hindrance in the application of thin-
film phosphors is their low brightness and efficiency in comparison to those of powder 
phosphors. Pulsed laser deposition (PLD) technique, which provides a unique process for 
stoichiometric evaporation of target materials and control of film morphology [ 13, 14], has 
been used for the deposition of oxysulfide films ( 15- 17]. 
In this work, we report on a study of the PLD conditions, the consequent crystalline and 
surface morphology structures, and photoluminescence (PL) characteristics of Y 20
3
2S:Eu + 
thin films. 
Physica B: Physics of Condensed Matter (2016), pp. DOI information: 
10.1016/j.physb.2015. 10.005 
83 
Our study showed that the oxygen partial pressure positively affected the crystalline phase, 
the morphology and the PL efficiency of the thin films. The luminescence results show that 
the PL intensity from Y20
3
2S:Eu + films under oxygen partial pressure may be as much as 1.4 
times higher than that from Y20 2S:Eu
3
+ film deposited in vacuum. As far as our knowledge is 
concerned, no one has so far investigated the influence of oxygen partial pressure of materi al 
properties of Y 0 S:Eu32 2 + under these conditions and achieved similar results. 
6.2 Experimental procedure 
6.2. l Powder synthesis 
Y 0 S:Eu32 2 - nanocrystals were synthesized using the sol- combustion route. The method of 
synthesis essentially comprises of mixing the precursors in appropriate stoichemetric ratios, 
fo llowed by firing in an air tube furnace at a temperature of 500 °C. T he white foam y product 
was then grounded and left to dry in an enclosed oven for 24 hours. A pellet with a 2.4 cm 
diameter and 6 mm thickness was prepared by pressing the Y 20
3
2S:Eu + powder for 1h our at a 
pressure of 1.96 x 107 mbar. The pellet was then annealed fo r 3 hours at 600°C temperature 
in an open - air furnace to improve its hardness. 
6.2.2 Pulsed Laser Deposition (PLO) 
The Si ( 100) wafers used as substrate were first cleaned using ethanol, followed by methanol 
each for I 0 minutes in a ultrasonic bath. This was fo llowed by rinsing the substrate in 
disti lled water also for I 0 minutes in a ultrasonic bath. The cleaned substrate was then dried 
in an oven for 2 hours. The deposition chamber was evacuated to a base pressure of 8 x I o·6 
mtorr. The Lambda Physic 248 nm KrF excimer laser was used to ablate the phosphor pellet 
in vacuum and various 0 2 partial pressure atmospheres. A Baratron Direct (Gas Independent) 
PressureVacuum capacitance Manometer (1.33 x 10·2 mtorr) was used for the high pressure 
measurements. The laser energy density, number of pulses and laser frequency were set to 
0.74 J/cm2, 12000 and 10 Hz respectively. The substrate temperature was fixed at 300 °C , and 
the target to substrate distance was 5 cm. The ablated area was I cm2. The Shimadzu 
84 
Superscan SSX-550 system was used to collect the Scanning Electron Microscopy (SEM) 
micrographs. Atomic Force Microscopy (AFM) micrographs were obtained from the 
Shimadzu SPM - 9600 model. X-ray diffraction (XRD) data was collected by using a 
SIEMENS 05000 diffractometer using CuKa radiation of A. = 1.5405 nm. PL excitation and 
emission spectra were recorded using a Cary Eclipse fluorescence spectrophotometer (Model: 
LS 55) with a built-in xenon lamp and a grating to select a suitable wavelength for excitation. 
The excitation wavelength was 230 nm and the slit width was 10 nm. The afterglow curves 
for the films were also obtained with the Cary Eclipse spectrophotometer. 
6.3 Results and discussion 
6.3.1 X-ray diffraction 
Figure 6. 1 shows the XRD patterns of Y 0 S:Eu32 2 + thin films grown at 300 °C in vacuum 
atmosphere as well as in various oxygen partial pressures. The patterns show mixed phases of 
cubic and hexagonal crystal structures. The films grown at low and high oxygen partial 
pressure are predominantly cubic, while that grown at moderate pressure ( 100 mtorr) is 
predominantly hexagonal. The average lattice parameters for hexagonal phase a=3.785nm 
and c=6.589nm, are very close to the standard values provided in the powder diffraction file 
PDF #24-1424. The intensities of the XRD peak from the (100) and ( 11 3) crystal planes were 
found to increase with the oxygen partial pressure in the range from 20 to 140 mtorr 0 2. This 
may be attributed to the enhanced oxidation kinetics and improvement in crystalline nature of 
the fi lms. As the oxygen partial pressure increase, the crystallinity of the films improved. It is 
also clear that the 300 °C thin fi lm consist of nanoparticles as seen from the broad XRD 
peaks. The estimated crystallite size of the fi lms with hexagonal and cubic phase rangi ng 
between 50 and 70 nm and is shown in figure 2. 
It is observed that for the Y20 2S:Eu
3
+ nanostructures, there is a marginal decrease (-0.26%) in 
crystallographic unit-cell that tends to contract due to the increase in surface area of the 
deposited film as compared to that of the substrate itself. This may have been caused by 
stress between the two surfaces and may lead to a decrease in the lattice constant. Eu20 3 
diffraction peaks from XRD patterns were not detected indicating that the Eu3+ was 
incorporated into the Y20 2S host lattice homogeneously [ 18]. 
85 
...-.. 
::s•  
·
."
-
c. Cl.')   
-c 
21w(ctget 
Figure 6.1. X-ray diffraction patterns of films deposited in vacuum 
and various 0 2 parti al pressures and the standards JC PDS card Nos: 
24- 1424 and 22-0993. 
Figure 6.2 shows the role of oxygen partial pressure on crystallite s izes and ax ial ratio. A 
monotonous increase of the ax ial ratio (c/a) and decreasing particle size is established by 
means of X-ray analysis for a series of thin fi lms deposited at vacuum as well as vari ous 
oxygen ambient. It is shown that the effect cannot be explained by impurity or intrinsic 
defects in the thin films. The relations obtained are based on the size dependence of the 
internal stress and the intra-crystalline pressure stipulated by the interaction of the eleme nts 
of the crystal charge latt ice. Calculati ons made fo r the ion charge lattice of quartz agrees with 
the monotonous increase of the lattice parameters with decreasing particle s ize. 
86 
80 
1.8 
• • • 
- 1.6 E 
-c 70 ---
" -Cl')  -- 1.4 ....... .~ ..... .... 
"Cl' ..... ca ) ..... 
~ ...... _ 1.2 0 
-ca .... 60 .... .... >"'  -•- Crystallite sizes .... ..... 1.0 ~
(.) -• -Cla ratio 
0.8 
50 0.6 
20 40 60 80 100 120 140 
Figure 6.2. Crystallite sizes and axial ratio as functions of 
oxygen partial pressure 
Table 6.1: Showing how oxygen partial pressures affect lattice parameters and particle size 
of the films. 
--------------------------------------------------------------------------
0 2 12artial 12ressure 
(mtorr) Lattice 12arameters Particle size (nm) 
_J f~x_a_gg.D~L - _____C l!'2.i~ 
a c a 
Vacuum 
20 3.805 6.650 3.789 70 ± 0.5 
60 3.798 6.642 3.785 67 ± 0.5 
100 3.786 6.633 3.78 1 62 ± 0.5 
14Q 3.145 6.526 3.116 58 ± 0.5 
87 
6.3.2 Morphology 
Figure 6.3 shows the SEM images for Y 20
3
2S:Eu + thin films ablated in a) vacuum, b) 20 
mtorr, c) 100 mtorr and d) 140 mtorr oxygen ambient, at 300 °C and fluence of 1.6 ± 0.1 J 
cm·2. The thin fi lm ablated in vacuum and 20 mtorr 0 2 show rougher layers. A smoother 
urface is however visible from the films ablated in 60 mtorr 0 2 and it appears as a layer 
consisting of spherica ll y shaped nano - particle layer in all the images. The film deposited in 
140 mtorr ambient show some layer on the substrate which is even smoother than that in 
figure c) as shown in fi gure d). The size distribution was broad and the average diameter was 
about 60± 0.5 nm. The ex istence of bigger micron size particles also occur on all the thin 
films. These bigger micron particle · can severely degrade the performance of e lectronic and 
optical devices and elimination can be done by optimization of the proce parameters [ 19). 
These SEM images give a rough indication of the surface morphology but better imaging wa , 
obtained with AFM analysis. 
88 
Figure 6.3 : SEM images of the thin fi lms ablated in a) vacuum, b) 20 mtorr, c) 60 
mtorr and c) 140 mtorr 0 2 ambient at 300 °C with a fluence of 0.767 ± 0.1 Jcm-2 (5 kV 
beam energy, magnification of x 20 000 and a scale of l µm (FOV: 2 x 1µ  m). As 
insets: 30 Height AFM images done in contact mode for the thin fi lms ablated in a) 
Vacuum, b) 20 mtorr and c) 140 mtorr oxygen ambient. 
6.3.3 Atomic Force Microscopy (AFM) 
Figures 6.3 also shows AFM images of the samples deposited in (a) vacuum, (b) 20 mtorr and 
( c) 140 mtorr as insets. It is clear that almost hexagonally-shaped nanoparticles were 
deposited during the deposition process. The particle sizes of the different fi lms varied from 
60 to 70 nm depending on the oxygen partial pressure. An average particle size of 70 run was 
calculated for the fi lm grown at 20 mtorr, 300 °C, and the particle size of 58 nm was 
calculated fo r the film grown at 140 mtorr, 300 °C. The particles were also less agglomerated 
at the higher oxygen partial pressures. The mean free path of the particles in a low ambient 
pressure is longer compared to the mean free path at higher ambient pressures. 
More collisions between the ultra fine particles (vaporised particles close to the target) at a 
higher ambient pressure lead to nucleation and growth of smaller nanoparticles when arriving 
at the substrate. In vacuum there are virtually no colli sions between the particles before 
reaching the substrate. Longer residence time of the particles in the plume, as is the case at 
higher ambient pressures, lead to more evenly di stributed particles. Light emission from the 
spherical shaped phosphor particles as excited by the electron beam is more intense due to the 
fact that much less photons encounter total internal refl ection [20]. The increase in the 
deposition pressure is reported to have caused an increase in the connectivity (agglomeration) 
between particles due to sintering of small particles [21]. This would eventually lead to grain 
growth at high enough pressure. In this case, we also fo und that more agglomeration occurred 
at higher oxygen partial pressure. 
6.3.4. Photoluminescence spectra 
Figure 6.4 indicates the excitation spectra of Y20 2S:Eu
3 3
+ thin fi lms as well as Y 20 2S; Eu + 
powder (inset). Excitation spectra were recorded keeping the emission wavelength at 6 19 nm. 
These spectra consist of several excitation bands of the f-f transitions, which are ascribed to 
89 
di fferent transitions fro m ground state 50 0 to the van ous excitation states of 
50 1 (J= 
1,2,3,4,. .... ) electronic configuration of the Eu3+ ions. 
Figure 5 shows the emission spectra of Y20 2S: Eu
3
+ thin films and Y20 2S:Eu3+ powder 
(inset) red-emitting phosphor. Films deposited in vacuum and those deposited at di fferent 
oxygen partial pressures as well as the powder Y20 2S:Eu3+ were used for the PL spectrum 
measurements. All the samples were excited at a wavelength of 230 nm. For the powder 
sample, the PL emission spectra consist of several peaks corresponding to different energy 
transitions of the Y 320 2S :Eu + phosphor. Totally, there are three groups of di stinctive emission 
peaks between 590 and 629 nm, which are related to the 5D 70- F1 (J= I, 2, 3) transitions of 
Eu3+, respecti vely. In the case of the thin films the PL reveals three peaks only which are 
assigned to 5D 7F 5 7 5 7 30- 1, Do- F2 and Do- F3 transitions of Eu + ions. The strongest emission peak 
fo r all these samples both thin fil ms and powder Y20 2S:Eu
3
+ is from the 50 70- F2 transition of 
Eu3+. The fact that the dominant emission is from the parity forbidden electric d ipole 
transition rather than fro m the magneti c dipole transition (5 D 70- F1) indicates that Eu
3
+ is 
located at the site w ith no inversion symmetry in the Y 20 2S latti ce (C2 site) (22]. Compared 
with the 5Do-7F2 trans ition, the intensity of the 
50 70- F 1 transition corresponding to the orange 
colour is much lower, which makes the deposited Y20 2S:Eu
3
+ a purer red phosphor. This is 
also confi rmed by the inset o f Fig.5, which shows that the PL emission intensity of the 
powder is more than 6 times higher than the intensity of Y20 2S:Eu
3
+ thin films. 
The reason behind observing the intense red emission from Y20 2S:Eu
3
+ in the case of 
hexagonal and also cubic phase can be understood by considering the structure of Y 20 2S. The 
coordinate number o f Y 20 2S is twelve and forms hexagonal structure with two di fferent s ites 
(C2 and C3i) fo r rare-earth (RE) ions substitution. The C2 is a low symmetry site without an 
inversion centre whereas C3i is a high symmetry site having an inversion centre. When Eu
3
+ is 
located at a low symmetry (C2) , the red emission is dominant whereas the orange emission is 
dominant w hen Eu3+ is located at high symmetry (C3i). In the present case, red emission is 
dominant suggesting that the location of Eu3+ is more favorable at C2 site. As the C2 site does 
not have an inversion centre, electric dipole trans ition from Eu3+ ions attached to this site are 
more favorable than the magnetic di pole transitions. The similarity of the ionic radii of Eu3+ 
and y 3+ ions for both hexagonal as well as cubic phase allows the easy substitution of y 3+ 
ions w ith Eu3+ ions at C2 sites giving rise to intense red emission in all the samples. The 
emission peaked at 590 nm due to 5D 70- F1 transition of Eu
3
+ is quenched at higher 0 2 partial 
pressure. The most intense peak at 6 19 nm due to 5D 70- F2 transition is totally quenched at 
90 
higher 0 2 partial greater than 20 mtorr. Simi larly, the peaks at 630 nm due to 
5Do-7F3 
transition are quenched at higher 0 2 pressure. In the case of vacuum atmosphere as well as 60 
mtorr 0 2 partial pressure, the peaks are not visible in all the three transition , while 20 mtorr 
0 2 partial pressure provides the highe t intensity at all the transitions. Thi. can be explained 
by visiting the oxidation states or kinetics of the Eu3+ ions. The most intense peak at 619 nm 
is ass igned to the 5D 70- F2 transition for 20 mtorr 0 2 partial pressure confi rms the presence of 
Eu3+ , making 20 mtorr to have highe t intensity in all three transi tions. 
1000 
---- 0 
- 20mtorr 
BOO - • • - 60 mtorr - 100 mtorr ::J . . . 
- 140 mtorr ci:s J';OO >- 600 
!:::: --:- :ml 
:I 
-"c:'  ~ 2500 4> :I c: .~ 2000 Ill 400 c 
-4> .!! 1500 .~ .E ci:s 
a4:>  
soo 
200 
- -- 250 3'JO 360 400 450 , ..... WaYelength (nm) 
0 -:..-.;'..":.,,_- ··-·~· ...·..·...·... . :..·-~-:-::-::l':-:;m.-...;::;...;;;.,,,....,._ _ _ _ _ ..... 
200 250 300 350 400 
Figure 6.4. Excitation spectra for films deposited in vacuum and at di fferent oxygen 
partial pressure. The inset show excitation spectrum of Y20 2S:Eu
3
+ powder phosphor. 
91 
1000 
5CXX) V acuum 
- PO\'\der 
20 mtorr 
800 4000 60 mtorr 
100 m torr 
'3' 3000 140 m to rr 
~ 
:::J 600 
rd i 2CXX> ->- ~ V> 4 00 1000 
-c: Q) c: 
2 00 500 600 
Wavelength (rwn) 
0 
500 520 5 4 0 560 580 600 6 2 0 6 4 0 6 60 
Wavelength (nm) 
Figure 6.5. Emission spectra fo r films deposited in vacuum and at different oxygen 
partial pressure. The inset show emission spectrum of Y 320 2S:Eu + powder 
phosphor. 
Figure 6.6 shows a plot of maximum peak intensity as a function of oxygen partial pressure. 
The fi gure indicates that the maximum peak intensity initially increases monotonously 
between 0 mtorr and 20 mtorr, then decreased at higher oxygen partial pressure. The 
optimum luminescence intensity was achieved at 20 mtorr due to the fact that posible 
oxidation of Eu~+ to Eu2+ did not take place at this particular 0 2 partial pressure. This 
optimum value of 20 mtorr 0 2 partial pressure can be uti lized in future research but under 
different deposition conditions. 
92 
Oxv~en partial pressure (mtorr) 
Figure 6.6. The plot o f maximum peak intensity versus oxygen 
partial pressure. 
3.5 Afterglow decay curves of the red phosphors 
The a fte rg low properties of thin film abla ted a t vacuum and various oxygen pres ures are 
compared, a hown in Fig. 7. It can be seen that the decay curve of the fil m ab lated at 20 
mtorr ha highe t afterg low and brightness, whil e the film ablated in vacuum and 60 mtorr 
has the lowe t afte rg low and brightncs . This indicates that vacuum and lower oxygen 
pressure favo r long afte rg low and higher intensities and vice versa. The decay times of the 
phosphor can be estimated by using the fo llowing double exponentia l equation; 
where I is the phosphorescence intensity, A 1, and A2, a re constants, t is ti me, T1 and T2. are 
decay Limes fo r exponential componenls, rcspccLi vcly. T he fi lling resul ts of parameLer · L1 and 
t2 are listed in Table 6.3 below and the expected experimental error is± 0.00 17 m/s. 
93 
- 0 
- - - 20 mtorr \ ......... 60 rntorr :::l 
-m \ · · · 100 rrtorr > \ - · - 140 rrtorr ;!:::: 
-"c:' 
\ 
Cl) \ 
c: 
\ 
Cl) 
:> \ ;: 
m \ 
Cl) 
a: \ 
\ ' 
' ' 
' ' ' ' 
3 4 5 
Tirre(ms) 
Figure 6.7: Decay curves for the thin films deposited in vacuum atmosphere and 
at di ffe rent oxygen ambient. 
Table 6.2: Decay constants for the fitted decay curves of the thin 
films ablated in vacuum and various oxygen partial ambient. 
0 2 pressure (mtorr) Vacuum 20 60 100 140 
Components (ms) Decay constants(<, s) 
Fast (T1) 0.346 0.400 0.352 0.388 0.372 
Medium (T2 ) 0.542 0.600 0.558 0.597 0.582 
6.3.6 Optical properties 
6.3.6.1 Absorbance spectra 
The UV-vis reflectance spectra of the samples are given in Fig. 8. The spectra of all the 
samples show good optical quality in the visible range due to reflectance in the 200- 500 nm 
range. The sharp absorption edge is characteri stic of a homogeneous structure [23]. The 
figure shows that the absorption edge shifted to highe r wavelength for lower 0 2 partial 
pressure and then reduced to lower wave length for higher 0 2 partial pressure. Furthermore, 
94 
absorption bands corresponding to the forbidden Eu3+ 4f-4f transitions were detected for 
higher 0 2 partial pressure. The band at around 344 run is attributed to the exciton absorption, 
which is red-shift compared with powder Y20
3
2S:Eu + [24). The absorption peaks at around 
290 and 340 run are assigned to 5Do-7F1and 
5Do-7F2 transitions of Eu
3
+ ions, respectively 
[25). 
6.3.6.2 Determination of band gap from reflectance spectra 
The Kubeika- Munk equation was used to calculate the band gap of the as- deposited thin 
film using a diffuse reflectance spectrum. On the other hand, the band gap (Eg) and 
absorption coefficient a of the indirect band gap semiconductor is re lated through the well-
known Tauc relation. The average Eg value for the thin films was found to be 4.03 eV, which 
is almost in agreement with the literature values by other researchers [26-27). The observed 
optical band gap for Y20 2S: Eu
3
+ thin films has increased to 4.24 eV as the 0 2 partial 
pressure increased to I 00 mtorr as shown in Fig. 9 (a). The change in optical band gap values 
may also be due to the change of crystal structure of the Y 20 2S thin films. This is al so 
confirmed by the fact that the PL emission intensity of the powder is more than 6 times 
greater than the intensity of Y20 2S:Eu
3
+ thin films 
The dependence of the band gap energy of the Y 20 2S on the 0 2 partial pressure is shown in 
Fig. 6.9 (b ). It can be seen from this graph that the band gap of the Y 320 2S; Eu + thin films 
increased with the amount of 0 2 partial pressure except for film deposited in 60 mtorr 
ambient. The decrease in the band gap energy and the shift of the absorption edges to higher 
wavelengths might be due to the presence of defect states and disorder due to the 0 2 partial 
pressure (28-30). The 0 2 partial pressure might have introduced new states close to the 
conduction band of the Y 320 2S:Eu +. A new defect band is therefore formed below the 
conductions which lead to reduction in the effective band gap [31-32). It is clear that at 20 
mtorr 0 2 partial pressure the estimated band gap increased to 4.40 eV. 
95 
70 
--Vacuum 
- - 20 mtorr 
60 I\ - · - 60 mtorr 
\ 
1. .. - • • • 100 mtor 
ti \\ 
~ 50 • 'A 0 
Cl> ...~ / \ 
0 
-c: (IJ 40 0 
Cl> 
~ 
Cl> 
a: 30 
20 . 
I 
10 ----· -·-·-·-· 
200 300 400 500 600 700 800 
Wavelength (nm) 
Figure 6.8. UV-vis diffuse reflectance spectra of nanocrystall ine 
Y 20 2S:Eu
3
+ thin film deposited in vacuum and different oxygen 
pressure. 
4.5 
0.20 (a) -- Vacuum (b) - • - Band gap energy 
- - 20 rrtorr 
0.18 .. . 4.4• 60 rrtorr ..I•   •  \ 
0.16 - • • - 100 mtorr "\ 4.3• I ..I   \ 
0.14 I . 
C'll 4.2• .I  . ~ > \ 0.12 . / 
~ I 
!..... -
cc 0.10 >Q) 
4.1• ..  \ 
 / 
i..:L
I 
.......  ~ 4.0• . 0.08 .. 
\ 
w .I 
/ 
0.00 3.9•   \ 
I ./  
0.04 3.8• . \ ' I 0.02 
3.7• ii 
0.00 
3.6 . . . . . . 
2 3 4 5 0 20 40 60 80 100 
Eg (eV) Oxygen partial pressure (rrtorr) 
Figure 6.9. (a) Graph of F[(R)*hv]2 as a function of band gap energy, (b) 
Dependance of band gap energy on partial oxygen pressure 
96 
6.4 Conclusion 
Y 20 2S: Eu
3
+ thin fi lms were successfully grown onto the Si (I 00) substrates at different 
oxygen ambient. XRD show mixed phases of cubic and hexagonal crystal structures. The 
fi lms grown at low and high oxygen partial pressure are predominantly cubic, while that 
grown at moderate pressure ( I 00 mtorr) is predominant)y  hexagonal. The thin films were 
composed of nanoparticles. The size of the particles depended on the oxygen partial pressure. 
The decrease in the oxygen partial pressure resulted into big particles which are actually 
piling up of smaller particles and therefore a rougher surface. The PL measurements showed 
red emission of Y20 2S:Eu
3
+ thin films as well as the powder Y20
3
2S: Eu + with the most 
intense peak appearing at 6 19 nm, which is assigned to the 50 70- F2 transition of Eu
3
+. This 
intense peak is quenched at higher 0 2 partial pressure greater than 20 mtorr. The emission 
peak at 590 nm due to 50 70- F 1 transition of Eu
3+ is also quenched at higher 0 2 partial 
pressure. Uv-vis measurement revealed an average band gap of 4.03 eY. 
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35 16. 
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[12] H.T. Cui, G.Y. Hong, H.P. You, J. Colloid Interface Sci. 252 (2002) 184. 
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J. Korean Phys. Soc. 46, 1193 (2005). 
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(2001) 921. 
[19] L. Chen, Particles generated by pulsed laser ablation, in D. B. Chrisey, G. K Hulber 
(Eds), Pulsed Laser Deposition of Thin Films, John Wiley & Sons, Inc, New York, (1994) 
184. 
[20] S. Yi, J.S. Bae, B.C. Choi, K.S. Shim, H.K. Yang, B.K. Moon, J.H. Jeong, 
J.H. Kim, Opt. Mater. 28 (2006) 703. 
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Suveges, J. Appl. Phys. 79 (1996) 2834. 
[22] H. Zhang, J. Liu, J. Wang, C. Wang, L. Zhu, Z. Shao, X. Meng, X. Hu, Opt. 
Lasers Eng. 38 (2002) 527. 
[23] V. Buissette, A. Huignard, T. Gacoin, J.-P. Boilot, P. Aschehoug, 8. Viana, 
Surf. Sci. 532-535 (2003) 444. 
[24] R.W.G. Wyckoff, Crystal Structures, Vol. 3, Interscience, New York, 
1963, p. 17. 
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99 
Chapter 7 
Energy transfer and material properties of Y 0 3 32 3: Eu +:Ho +nanophosphors 
synthesized by sol- combustion method 
7.1 Introduction 
Europium-activated Y 20 3 (Y 20 3:Eu
3+) has attracted much attention as a red -emitting 
phosphor for commercial use in fluorescent lighting and screen due to its high luminescence 
emission around 626nm. Because of the fast development of nano technology, the optical 
properties of nanocrystalline Y20 3:Eu
3
+ have also been extensively investigated for its 
potential app lication in high resolution imaging for flat plasma displays and in fu ndamental 
research [ I, 2]. Therefore, many researchers have already perfonned investigations related to 
this nanophosphor, such as the study of the influence of particle size on the intense 
luminescence of nanocrystalline Y 20 3:Eu
3
+ [3- 5]. In recent years, luminescent nanocrystals 
(NCs) doped with rare earth ions were paid more attentions because o f their interesting 
luminescent properties. They can be as components in displays [6], light emitting diodes [7], 
biological assays [8], and optoelectronic devices [9]. Cubic Y20 3:EuH is one of the most 
important commercial red phosphors, which can be used in fluorescent lights, cathode ray 
tubes (CRTS), plasma display panel (PDP), and field emission display (FED) [I O]. Yttrium 
oxide (Y20 3) has been investigated widely as a host material for rare-earth ion doping in 
optical appl ications [ 11] on account of its excellent chemical stabi lity, broad transparency 
range (0.2 to 8µm) with a band gap of 5.6eV, high refractive index, and low phonon energy 
[12-14]. Furthennore, the sim ilarities in the chemical properties and ionic radius of RE ions 
and Y20 3 make it an attractive choice as a host material [ 15, 16]. Many methods fo r synthesis 
of Y 20 Eu
3
3: + :Ho
3
+ phosphor has been reported, including sol-gel method [ 17], spray 
pyrolysis method [ 18], hydrothennal method [ 19] , and precipitation method [20]. Wet 
chemical methods have been widely developed to prepare the luminescent materials [2 l ], 
since these processes have advantages of good homogeneity through mixing the starting 
materials at the molecular level in solution, a lower calcination temperature and a shorter 
heating time. However, sol-gel process often requi res expensive (and environmentally 
unfriendly) organic precursors and solvents. Then a simple technique, so l- combustion 
synthesis, is beginning to attract a great deal of interest. Sol- combustion is one of the 
simplest and most versatile approaches avai lable to obtain single-phase powders at low 
temperatures with shorter reaction times and little residual impurities as compared with 
conventional solid-state reactions [22]. 
In the present study, a series of red -emitting phosphors Y 0 :Eu3 Ho32 3 - : "'" was prepared by 
sol- combustion method. The luminescence, excitation, and optical absorption, structural and 
morphological properties of Y20 3:Eu
3 3
+: Ho + has been studied. To investigate the effect of 
100 
Ho3+ on the morphology and luminescence properties of the phosphor, the percentage 
concentration of Ho3+ ions were varied while Eu3+ ions were kept constant. To the best of our 
knowledge, there is no publication on the synthesis of Y20 3: Eu3 3+: Ho + by sol-combustion 
method and got similar results. 
7.2 Experimental 
7.2.1 N anocrystal synthesis 
All the chemicals used for the preparation of the powders were of analytical grade. In the 
present study Y (N03).6H02 is used as an oxidizer, thiourea employed as a fuel and Eu
3+ and 
Ho3+ as activators. Stoichiometric compositions of metal nitrates and fuel are mixed. 
Precursors were dissolved in a very minimum amount of di stilled water and ethanol to obtain 
homogeneous solution. The solution was then introduced into a muffle furnace preheated to 
500°c. After about 7 minutes the solution boiled and was ignited to a self-propagating flame. 
The fluffy masses obtained were crushed into a fine powder and was ready for 
characterization. 
7.2.2 Characterization 
To determine the average crystallite diameter and the phase of the samples, X-ray powder 
diffraction (XRD) spectra were measured with a 08 Bruker Advanced AXS GmbH X-ray 
diffractometer using Cu Ka radiation at a wavelength of 0.154056 nm. The measurement 
conditions were as follows: the wavelength was 0.1541 9 nm {pure Cu Ka radiation), tube 
voltage/tube current was 40 kV/50 mA, and machine scanning range was 28 = 1o 0- 90°. The 
size and morphology of the as-prepared particles were carried out by using a scanning 
electron microscope (SEM), Tescan VEGA 3 SEM. The photoluminescence (PL) spectrum as 
well as decay curves for all the samples were investigated by Cary Eclipse fluorescent 
spectrophotometer equipped with a 150 W xenon lamp at an excitation source w ith the slit of 
1.0 nm and scan speed of240 nm min-1• 
101 
7 .3 Results and discussions 
7.3.1 X-ray diffraction study 
Fig. 7. l illustrates the X-ray diffraction patterns of Y20 :Eu33 +: Ho
3
+ powder samples, 
synthesized by the sol-combustion method. Six main peaks are observed, which can be 
assigned to (2 11 ), (222), (400), (440), and (622) of cubic Y20 3. All diffraction peaks are 
attributed to the pure body-centered cubic (bee) structure for Y 20 3 phase and match well with 
JCPDS (25- 1200) card with space group Ia3. All samples present (222) as the preferential 
orientation. At O. l % Ho3+, the product is already crystall ized and presents some three peaks. 
At higher concentrations more peaks are detected. This difference in crystallization is due 
probably to the difference in the Ho3+ ions concentrations. Increasing Ho3+% ions leads to the 
increase of the main peaks intensity with a sharpening, indicating the improvement of the 
crystallinity quality of the powders. This phenomenon can be explained by the fact that the 
y 3+ ions are accommodated in two different symmetries, i.e., the C2 without inversion center 
and the S6 with inversion center [23). The Eu
3
+ ion might be replacing the y 3+ ions in 
theY20 3 host lattice because no diffraction pattern characteristi c of Eu20 3 phase is detected . 
The radii of Y3\ Eu3.. . and Ho3+ ions are 0.89 A0 , 0.947 A0 and 0.9 A0 respectively,[24). 
According to the ion radius approximation principle in the crystal field theory, Eu3 3+ and Ho + 
ions enter into crystal lattices by replacing y 3+ ions other than by occupyi ng the slot between 
crystal lattices. Compared w ith Y3 ... ionic radius (0.89 A), Eu3 3+ and H0 + have larger ionic 
radii of 0.947 A and 0.9 A respectively. Therefore, the lat tice constant would increase w ith 
increa ing atomic content of Ho3- ions in the powder. 
The crystallite size of these materials has been estimated from the same XRD patterns. As a 
first approx imation, it is suggested that the appearance of broad XRD peaks is due to the 
nano-sized particles formation. The average crystallite size ' D' of Y 320 3:Eu +: Ho
3+ powders 
can be estimated using Scherrer's fo rmula [25), 
0.9-1 
D= 
{JcosfJ 
where A.= J .5406 A is the wavelength of the X-ray radi ation used, fJ is the fu ll width at half 
maximum (FWHM) of the (222) diffraction peak in the XRD patterns in radians, B is the 
Bragg diffraction angle. 
102 
- 0.1 % Ho3 
- 0.2% Ho3 
-. 
-. ca 
l-o.5% Ho3~ 
,..,........, .,.....  N.,...'  1- # 24-8756[ 
~ 
10 20 30 40 50 60 70 80 90 
2 theta (degrees) 
Fig.7.1. X -ray di ffraction pattern of Y 20
3
3: Eu +: Ho
3
+ phosphor 
The average crystallite si/e and crystallographic unit ce ll parameters arc calculated and l isted 
in T able 1. It is found that the diffract ion peaks shift to longer angles w ith i ncreas ing Ho% , 
indicati ng a decrease o f the l;lltice constant and a dc,·elopmcnt of internal tensile strain. In 
1
addi tion. the crystall ite size o f "Ynthes i7ed Y ~o, : Eu +: Ho-'+ powder increa-;ed with 
increa..., ing Ho~+ ion<.,. I t is known that FWI IM (/1 in the Scherrer formula) can be i nterpreted 
1
in terms of lall ice st ra in and crystal l ine size. The crystal lallice stra in generated hy the Ho + 
ions is determined from the W i ll ia1n....on- Hall relationsh ip [26] : 
cos8 1 
(J - A- = 0 + 77sin8 
w here (1 i s the full w idth at half maximum. i. i <> the X -ray wave length. (} is the diffraction 
angle, D is the effecti ve crystallite " i1.c and 11 i ~ the effecti ve strain. The strai n i., ca lculated 
from the slope of the plot of /1 ((co" 0)/i.) agai n.,t sin O/i. and the effective crysta l l i te size is 
calculated from the i ntercept to /1 ((cos 0)//. ) ax is. T hey shm\ a difference in angu lar 
dependence or the line w idth for Ho I+ ions. Indeed, the best lit is obta ined when the Ho I+ ion 
percen t is incrca:-.cll rrom 0 . 1 to 0.51/r and w hen the slope becomes negative. In the same 
tahk. we al '>O present the effecti ve cry .... rallite size and the effecti ve stra in 11 again:-.! the Ho>+ 
ions percent. T he effective cry'>tal l i te si1.e values arc very c lose to tho'ic ex trac ted from the 
Scherrer rormula and increase \i\ hen the Ho i- ions percent increa.,es. The tensile strai n 
decreases from 0 .1%  to the negat ive value corresponding to O.S <fr . It has been reported that a 
negati ve value indicates the presence o r compress ive lalli ce strain [27] . Th is -;trai n behaviour 
103 
as a function of Ho3- ions is in a good agreement with that of the lattice parameters. T he 
reduction of the lattice parameter and the increase of crystallite size with increasing of the 
Ho3+ ions signify the closeness of lattice plane which can lead to an increase of the density 
and a decrease in the dislocation density. The average nanocrystallite sizes fo r Y 20 3: Eu3+ 
with different concentration of H03+ is compared as shown in the table I below; 
Table 7. 1. The crystallite sizes as a function of % concentration of H03+ ions ions. 
% ofHo3+ 0. 1 0.2 0.3 0.4 0.5 
Crystal sizes (nm) 6.5 7.5 8.9 9.2 9.8 
Table I lists the crystallite sizes of synthesized particles estimated from the well-known 
Scherrer's equation. The diffraction data of the strongest peak (222] plane was used to 
calculate the mean crystallite sizes. The crystall ite sizes showed a slightly increasing 
tendency, which can be attributed to the effect of the increased dopant concentration (28]. 
Figure 2 shows the ro le of Ho3+ mole concentration on crystallite sizes and lattice parameter. 
A monotonous decrease of the lattice constant and increasing crystallite size is established by 
means of X-ray analysis for a series of Y20 3: Eu
3 Ho3+: + powders. lt is shown that the effect 
cannot be explained by impurity or intrinsic defects in the powders. The expressions derived 
are based on the size dependence of the internal stress and the intracrystalline pressure 
stipulated by the interaction of the elements of the crystal charge lattice. Calculations made 
for the ion charge lattice of quartz evidence the monotonous increase of the lattice parameters 
with decreasing particle size and agree with the results of the X-ray experiment. 
104 
10.0 ------k- Crystallite sizes * 
• 10.9 ·~ Lattice constay 9.5 -E 9.0 • ---------* 10.8 -c: * -c: t/) 8.5 ns 
Q) 
N -"c:'  
-t/) 
10.7 
8.0 0 ·~ 0 Q) Cl> - 0 ca 7.5 -t/) 10.6 ns ... ...J >a 7.0 
(.) ·~ 
6.5 • 10.5 
6.0 
0.1 0.2 0.3 0.4 0.5 
0/o mole concentration 
Fig. 7.2. Crystallite sizes and Lattice constant as functions of 
Ho3+ concentration 
7.3.2 Scanning electron microscopy 
SEM was carried out us ing Tescan VEGA 3 SEM. Fig. 3 shows the SEM image of the 
Y 0 :Eu3 32 3 +: Ho + co-doped phosphor. The image shows the characteristics surface 
morphology of combustion product. Particles are highly agglomerated and form a connected 
network type with some vacant space among the m, which is expected due to the evo lution of 
different gases during combusti on of the gel. lt also reveals inhomogeneous di stribution of 
the particles (a variation and shape of the particles). The shape o f some o f the particle are 
spherical and the size lie between 5 - I0  nm (fi g.2a) which is well in agreement with the 
crystalline s ize calculated us ing XRD. However ome particles are of sub-micron ize due to 
agglomeration of many crystallites and lac k of eparate boundary (fi g.2b-2e). The shapes of 
the particles are irregular and seem to be polygonal. 
105 
-------~------------------------------
Fig.7.3. SEM micrographs ofY20 3: Eu
3 3
+: Ho + samples with (a) 0. 1 (b)0.2 (c) 0.3 (d) 0.4 (e) 
0.5 % of Ho3+ ions. 4.77µm field of view. 
106 
7.3 Photoluminescence 
7.3.l Excitation 
Fig. 7.4 shows room-temperature optical ab orption spectra of Y2 0 3:Eu
3+: Ho3+ phosphor. 
Ho3+ absorption bands of 3H 5G5 5G6 56, , , F3, 
5F4 (5S and 
5
2) F5 were pre ent in Y2 0 :i : 
Eu3+:Ho3+. The Ho3+ doped Y 20 3:Eu
3+ showed green (544 nm) and red (626 nm) emiss ion 
peaks, which are assigned to 5F 5 5 5 34 (5S2) - 18 and F5- I8 transitions of Ho + ions re pectively. 
F a!17', -0.1 % 
2 5 0 - 0 .2 % 
- 0 .3 % 
- 0.4 % 
200 
- - 0. 5% ::J 
-«I 150 >. 
.~ 
VI 
-c: Q) 100 c: 
50 
0 
200 250 300 350 4 00 
Wavelength (nm ) 
Fig.7.4 Excitation spectra of Ho3+ co-doped Y 0 3 32 3 : Eu + phosphor when Ho + ion concentration was 
varied from 0.1 to 0.5% 
7 .3.2 Emission 
Fig. 7.5 shows the emission spectra of Y 20
3
3: Eu +:Ho
3+ powder when the percentage 
concentration of Ho3+ ions was varied. Three main emission spectra appear at 540, 588 and 
626 nm for the samples with lower mole concentrations (0.1 and 0.2%) of Ho3+ ass igned to 
50 o-7F 5 7 5 71, · 0 0- F 2 an d 0 - F 0 3 trans.it i.o ns o f E u
3 + . The se em1. ss1. on spectra d ue to Eu ·:i + .i ons are 
also confirmed by the inset of Fig. 3. Other minor peaks also appear around 440 and 500 nm 
107 
speculated to be from Ho3+. At higher mole concentration of Ho3+, the three main peaks are 
to tally quenched. Y 20
3
3:Eu:i+:Ho + phosphor shows a red-emitting afterglow phenomenon, and 
the Eu3+ ions are the luminescent center during the decay process. The bright red e mission 
near 626 nm has been noticeable due to the 50 70- F2 transition of Eu
3+. The Intensity of the 
luminescence has decreased with an increase of concentration of Ho3+. In sufficient quantities 
of Eu3+ to Ho3+, the bright red emission near 626 nm has been predominant due to 5D 70- F2 
transition of Eu3+. From the figure, the intensities of spectra are quenched as the % 
concentration of Ho3+ ions are increased. Since Eu3+ ions are kept constant, it is speculated 
that Ho3+ ions are the ones which caused the quenching o f the intensities of the spectra. 
300 
' o ,-' F, - Y o : Eu 32 3 •only 0.1 % 
8 00 0.2% 
250 " ~ 600 0.3% 
.~..  
c 4 00 0.4% 
"c  .... : ... 
. 
- 200 .' . 0.5% . 200 0  :i. - m 150 450 500 550 600 650 7 00 7 50 800 ->- Wavelength (nm) ·-
-"c:Q' )  100 c: 
50 
400 450 500 550 600 650 700 
Wavelength (nm) 
Fig.7.5. Photoluminescence emission spectra of Ho3 co-doped Y 0 Eu3+ 2 3: + phosphor 
when Ho3+ ion concentration was varied from 0.1 to 0.5% 
108 
0.9--~~~~~~~~~~~~~~~~~~-.-1~----. 
• 2 
4 
es 
0 .6 
>-
-w (.) 
0.3 
0 0.2 0 .4 0.6 0.8 
CIE X 
Fig.7.6. Chromatic ity colour coordinate of the Y20 3:Eu
3+:Ho3+ powder under 325 nm UV exc itation. 
7 .3.3 Decay characteristics 
The afterg low properties of Y 0 Eu32 3: +: Ho
3+ powder with different mole concentration of 
Ho3+ ions are shown in Fig. 7 . It can be seen that decay curve o f the powders with lowest 
concentrations has highest intensities and afterglow, while those with higher mole 
concentration of Ho3+ ions has the lowe t inten ities and . This ind icates that lower mo le 
concentra tion of Ho3+ favo rs Jong afterglow and higher intensities and vice versa. The decay 
times of the phosphor can be estimated by us ing the fo llowing double expone ntia l equation; 
109 
where I is the phosphorescence intensity, A 1, and A1, are constants, t is time, T1 and T2. are 
decay times for exponential components, respectively. The fitting results of parameters t1 and 
t2 are listed in Table 2 below. 
o.1°1c 
8000 0.2°/c 
0.3°/c 
0.4°/c 
.-..  6000 0.5°/c 
::s.  
.c_a.  
>a 
·..-, 4000 "c .Q.,')   
-c 
2000 
0 1CXX> 2000 4000 
lime(ms) 
Fig.7.7. Decay curves of Ho3+ co-doped Y20
3 3
3 : Eu + phosphor when Ho + ion 
concentration was varied from 0.1 to 0.5% 
Table 7.2: Decay constants for the fitted decay curves of the Y 20 Eu
3+:Ho33: + powder 
% concentration 0.1 0.2 0.3 0.4 0.5 
Components Decay constants(T, s) 
Fast (T1) 0.368 0.352 0.345 0.339 0.326 
Medium (T2) 0.680 0.668 0.641 0.625 0.599 
110 
The graph of maximum peak intensity of the as-prepared powders as a function of% Ho3+ 
concentration is shown in Fig.8. The emission peak intensity quenched gradually with 
increase in the % concentration of Ho3+ ions. It can be seen from the curve that the powders 
showed differences in initial intensity and medium persistence when the powders were 
efficiently activated by UV lamp. The results indicate that the initial luminescence intensity 
and the decay time of phosphors are enhanced with a decrease in the % concentration. 
3000 ...- -------....- --M-a-x-.- P-e-a-k- I-n-te-n-s-it-y- V-s- -%- -C-o-n-c-en-t-r-a-ti-on- ---~ 
-:::J.  2500 
-«S 
~ 2000 
t/) 
c: 
..Q..),  
c: 
1500 
~ 
«S 
Q) 
a. 
E 1000 
::::::s 
·E>-
 
< 
«S 500 
:i 
0.01 0.02 0.03 0.04 0.05 
0/o Concentration 
Fig.7.8. Concentration of Ho3+ ions vs. maximum peak intensity graph of 
Y20  3:Eu 3+ :Ho 3+ p h osp h or 
7 .3.4 Optical properties 
7.3.4.1 Absorbance spectra 
The UV- vis absorbance spectra of the samples are given in Fig. 9. The spectra of all the 
powder samples show good optical quality in the visible range due to the complete 
111 
absorbance in the 200- 700 run range. The sharp absorption edge is characteristic of a 
homogeneous structure [29). The figure shows that the absorption edge shifted to higher 
wavelength for higher mole concentration and then reduced to lower wavelength for lower 
mole concentration. Furthermore, absorption bands corresponding to the forbidden Eu3+ f- f 
transitions were detected for higher mole concentration. The band at around 220 and330 run 
is attributed to the exciton absorption, which is red-shift compared with powder 
Y 3 320 3:Eu +:Ho + [30]. The absorption peaks at around 290 and 340 nm are assigned to 
5Do-
7F 1and 
5Do-7F2 transitions of Eu
3
.,. ions, respectively [31]. 
7.4.2 Determination of Eg. from reflectance spectra 
The Kubeika- Munk equation was used to calculate the band gap of the as- prepared powder 
as fo llows; 
(1-R%) 2 
F ( Roo) = =k/s--------------------- ( I) 
2R% 
Where Roo= Rsampi./ Rrcfcrences, K is absorption coefficient and S is scattering coefficient. 
On the other hand the bang gap Eg, and absorption coefficient a of indirect band gap 
semiconductor are related through the well- known Tauc relation; 
2 
a hv = C1 ( hv - E9 ) ---- ---------------- ----------- (2) 
Where hv is the photo energy and C 1 is proportionality constant, n=2 for indirect transition. 
When the material scatters in a perfectl y manner, the absorption coeffi cient K becomes equal 
to 2a (K=2a). Considering the scattering coeffi cient S as constant with respect to wave length, 
and using equations ( I) and (2), the fo llowing expression can be written: 
[F( Roo) * hv]2 = C2 (hv - £9 )----------------- -- (3) 
By plotting [F(R)*hv) 2 against hv and fit the linear region with a line and extend it to the 
energy axis, then one can easily obtain Eg by extrapolating the linear regions to [F (R·) hv] 
2=0. The average Eg. value for the powders was fo und to be 4.43 eV, which is in a good 
agreement with the literature values by other researchers [26-27). The observed optical bang 
gap fo r Y20 3: Eu3+:Ho3+ samples have decreased to 4.0 eV at lower mole concentration of 
112 
Ho3+ shown in Fig. 10. The change in optical band gap values may also be due to the change 
of crysta l structure of the Y2 0 3 powders. This change in the optical band gap materials can be 
explained on the basis of quantum size effect. 
100 0.1°/oHO 
0.2%HO3 + 
90 0.3°/oHO3 + 
00 0.4°/oHo3+ 
.- - 0.5°/oHO3 + 
.0~. _ . 70 
Q) 
0 60 
c: 
.cco  50 
I.. 
0 
."c'  40 ct 
30 
20 
10 
200 300 400 500 600 700 
Wavelength {nm) 
Fig.7.9. Uv-vis absorbance spectra of Y2 0
3 3
3: Eu +: Ho + red- emitting phosphor with % mole 
concentration of Ho3+ from 0.1 to 0.5%. 
The dcpc11uc111.,,c v1 LllC ua11u bal-' c11c1by v1 Lile 1 2v 3.Lu .11v vu Lile pc11.,,c11L 111v1c 
concentration of Ho3+ ions is shown in Fig. I I . It can be seen from the graph that the band 
gap of the Y 3 3 32 0 3; Eu +: Ho +powders increased with the percent mole concentration of Ho +. 
The decrease in band gap energy and the shift of the absorption edges to higher wavelengths 
might be due to the presence of defect states and disorder due to the percent mole 
concentration. The mole concentration of Ho3+ might have introduced new states close to the 
conduction band of the Y2 0 3:Eu
3
+: Ho3+. A new defect band is therefore formed below the 
113 
conduction which lead to reduction in the effective band gap [32]. It is clear that at 0.5% 
mole concentration of Ho3+ the estimated band gap increased upto 4. 7 e V. 
C,\.I. ..., 
.>c  0.10 
.-tc- .. 
.a._:. . 
L....L... . 
0.05 
2 3 4 5 6 
hv(eV) 
Fig.7.10. Graph o f F [(R)*hv] 2 as a function of band gap energy 
114 
4.8 ~Band gap Vs o/o Ho3+ mole concentration 
4.7 ~ 
4.6 
- 4.5 -> Cl> . 4.4 
C) 
w 4.3 
4.2 
4.1 
4.0 
0.1 0.2 0.3 0.4 0.5 
0/o mole concentration 
Fig. 7.11. Band gap energy as a function of Ho3+ mole concentration 
7 .4 Conclusion 
Y20 3:Eu
3+:Ho3+ red-emitting phosphor was successfull y synthesized and the influence of 
mo le percent of Ho3+ ions has been investigated. The powders composed of nanoparticles. 
T he ize of the particles depended on the Ho3+ mole concentration. An increase in the Ho3+ 
percent resulted in mailer crystallite size . SEM show particles are highly agglomerated and 
fo rm a connected network type with some vacant space among them, which is expected due 
to the evo lution of different gases during combu tion of the gel. At lower mo le of Ho3+ all the 
peaks for both Eu3+ and Ho3+ were observed . T he PL measurements showed red emis ion of 
Y20 3:Eu
3+: Ho3+ powder with the most intense peak appearing at 626 nm, which is ass igned to 
the 5 D 70- F2 transition of Eu
3+. At higher mole concentration of Ho3+, energy trans fer 
115 
occurred from Eu3 3+ to Ho +. The band gap energy increased with increase of Ho3.,. mole 
concentration. 
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117 
Chapter 8 
Temperature dependence of structural and luminescence properties of Eu3+ 
-doped Y 20 3 red- emitting phosphor thin films by Pulsed Laser Deposition 
8.1 Introduction 
In an earl y development of low-voltage cathodoluminescent (CL) phosphors for applications 
in fi eld emission display (FED) devices, the cathode-ray tube (CRT) phosphors have readily 
been tested as candidates. In the case of the red phosphor materials, unfortunately, 
Y20 2S:Eu
3
+ which is used as the red primary color in the CRT has known to be degraded 
under electron bombardments with high current densities and the sulphur containing volatile 
gases escaping from the surface under electron bombardment contaminate the cold cathodes, 
resulting in a fatal damage to FED devices [ I] . The oxide based thin-film phosphors are 
highly attractive in the use of the FED devices because of the advantages such as higher 
lateral resolution from smaller I:T:, fains, better thermal and mechanical stabi lity, and reduced 
outgassing over conventional powder phosphors [2]. Among the ox ide phosphors, Y 0 3: Eu32 + 
is currently one o f the leading red phosphor materials for FEDs [3]. Y 320 3:Eu + fi lms have 
been grown using various deposition techniques [ 4-6]. However, because of its high melting 
po int of about 2400 °C, Y 320 3:Eu + thin fi lms require a post-annealing process at high 
temperatures above I 000 °C to crystallize the deposits [7-1 OJ. Therefore, the high 
temperature process is inevitable in order to obtain high efficient and bright Y 320 3:Eu + thin 
film phosphors and Y 20 3:Eu
3
+ fi lms had been grown onl y on the heat-resistant substrates 
such as Si wafers [ 11 , 12], Ni based alloys [9] and sapphire plates [1 3- 15]. However, 
annealing at high temperature is definitely a concern for the fabrication of the current FED 
devices which adapt low temperature glass substrates. 
Journal ofA pplied Physics A, A P YA-D-15-01 792.1, DOI: J0.1007/s00339-016-9946-5 
118 
By replacing the rastered electron beams in the CRT with an array of cathodes, FEDs promise 
to be significantly thinner and lighter, have higher brightness, better power efficiency, and 
viewing angle, and operate over a large temperature range as compared to liquid crystal 
displays (LCDs) [ 16-1 8]. While the cathode array enables FEDs, the light from this emissive 
display comes from the phosphor anode. Phosphor anodes are currently powders screened 
onto glass plates using a variety of techniques including electrophoretic, dusting, and slurry 
methods [19]. The slurry method is the most common method, in which phosphor powder is 
mixed with photosensitive chemicals and is patterned using photol ithography techniques. For 
operating voltages greater than 2 kV, the screen is back coated with a thin layer of aluminium 
which acts as an optical refl ector as well as a charge dissipater [20-22]. At operating voltage 
below 2 kV, the electron penetration depth is so small that the screens are left un-coated [23-
25]. Even though powder phosphors are very efficient, the particle size may limit resolution 
of the display. Thin films are an alternative to powder phosphors, and have both advantages 
and disadvantages as compared to powders [26-29]. 
The intention of the present research is to investigate whether modifying the surface of 
deposited thin film phosphors by annealing at different temperatures would resu lt in 
increased efficiency. We therefore report on the study of the annealing temperature on pulsed 
laser deposited thin films, the consequent crystalline and surface morphology structures, and 
photoluminescence (PL) characteristics ofY20 3:Eu
3
+ thin films. 
8.2 Experimental details 
8.2.1 Powder synthesis 
Y 20 3:Eu
3
+ nanocrystals were synthesized using the so l- combustion route. The method of 
synthesis essentially comprises of mixing the precursors in appropriate stoichemetric ratios, 
followed by firing in an air tube furnace at a temperature of 400 °c. The white foamy product 
was then grounded and left to dry in an enclosed oven for 24 hours. 
119 
8.2.2 Pulsed Laser Deposition (PLD) 
The Si ( l 00) wafers used as substrate were first chemically cleaned. The phosphor was 
pressed with binders to prepare a pellet that was used as an ablation target. The deposition 
chamber was evacuated to a base pressure of 8 x 1o -6 mtorr. The Lambda Physic 248 nm KrF 
excimer laser was used to ablate the phosphor pellet in a constant 20 mtorr 0 2 atmospheres. A 
Baratron Direct (Gas Independent) PressureN acuum capacitance Manometer ( 1.33 x 10-2 
mtorr) was used for the high pressure measurements. The laser energy density, number of 
pulses and laser frequency were set to 0.74 J/cm2, 12000 and 10 Hz respectively. The 
substrate temperature was fixed at 300 °C, and the target to substrate distance was 6 cm. The 
ablated area was l cm2. The films were then annealed at the temperatures of 600, 700, 800 
and 900 °c. 
8.2.3 Characterization 
The Shimadzu Superscan SSX-550 system was used to co llect the Scanning Electron 
Microscopy (SEM) micrographs. Atomic Force Microscopy (AFM) micrographs were 
obtained from the Shimadzu SPM - 9600 model. X-ray diffraction (XRD) data was collected 
by using a SIEMENS D5000 diffractometer using CuKa radiation of/... = l .5405 nm. PL 
excitation and emission spectra were recorded using a Cary Eclipse fluorescence 
spectrophotometer (Model: LS 55) with a bui lt-in xenon lamp and a grating to select a 
suitable wavelength for excitation. The excitation wavelength was 209 nm and the sli t width 
was 10 nm. The afterglow curves for the films were also obtained with the Cary Eclipse 
spectrophotometer. 
8.3 Results and discussions 
8.3.1 Structur a l and morphological analysis 
8.3.1.1 X-ray d iffraction ana lysis (XRD) 
XRD data were analyzed for the identification of phase and crystallite size. Figure 1 shows 
the X-ray diffraction patterns of the un-annealed and thin films annealed at the temperature of 
600 to 900 °c for 2 hours. The un-annealcd thin film was amorphous, while those annealed 
were crystalline. Two different phases were obtained at low and high annealing temperatures. 
120 
The thin film annealed at lower temperatures were indexed to cubic bixbyte phase [space 
group Ia-3(206)] with average lattice constant a= 10.60A which is in good agreement with the 
standard value for bulk cubic Y20 3 (JCPDS No. 72-0927). At low annealing temperature 
(600 to 700 °c), there are mainly two diffraction peaks at 2 theta values of 29.4 and 44.3. It 
can be clearly seen that by increasing the annealing temperature to 800 °c more peaks 
emerged at 2 theta values of 29.4, 33.0, 44.5, 47.8, 54.6, 56.3 and 62.0. This seven diffraction 
peaks indexed to hexagonal structure of Y 20 3 are in agreement with data from JCPDS card 
No. 24-1424 with average lattice constants a=3.779 and c=6.590. The thin film annealed at 
900 °c also revealed two peaks at 2 theta values of 29.4 and 44.5 similar to the peaks 
annealed at 600 and 700 °c, indicating same phase of cubic. This change in the structure 
when the annealing temperature is increased to 800 °c could be due to reduction in the lattice 
stress. The films annealed at 600, 700 and 900 °c have large lattice stress; whi le the stress of 
film at 800 °c is small. The role of crystal fi eld change in the fonnation of a new phase 
should not be neglected as well. The broadening of the peaks in both the phases suggests 
small particle size. No other impurities peaks have been found which means that dopant ion 
completely occupies the Y 20 3 host lattice. Also, the ionic radii of y
3
+ (0.90A) and Eu3+ (0.89 
A) are very close, and hence it is possible to substitute y 3+ with Eu3+ ions. 
121 
Card llo 24·1424 [222) ~I I~ 
·- Unannealed 
.-...- soo0c 
-
Card llo:n .0927 
1 2 3 40 60 
2t heta (degree) 
Figure 8.1. X-ray diffraction pattern of un-annealed and annealed Y2 0 3: Eu
3
+ 
thin film deposited on a ( I 00) Si substrate after firing at temperatures between 
600 and 900 °c in air fo r 2 hours. 
Figure 8.2 shows the role of annealing temperature on crystalli te sizes and latt ice parameters. 
A slight increase in the crystallite size and a decrease in lattice parameters are establ ished by 
means of X-ray diffraction analysis for a eries of thin films annealed at temperatures 
between 600 and 900 °c. 
122 
80 12 
70 - • • 10 
- 60 .. *.. . 8 -E c -* 50 .Ii.. (.1.)  (1) tn 
(1) E 6 
·N- ca Ii.. 
·.
t.n.  aca (1) . 
-..-. ..._ Crystallite sizes (nm) 
(1) 
4 
ca 0 
tn - Lattice pararreter ......  ca 
>a. ...I 
Ii.. 
(..) 2 
0 
0 ...- -....- --.--....~ --....- -..---...- --.--.... 
600 700 ED> 900 1COO 
Te11 peratLre (°C) 
Figure 8.2. The c rystallite sizes and lattice parameters as a function 
of temperature. 
Further, it is marked that as the sample is annealed at higher temperatures the FWHM (full 
width at half max imum) o f the diffraction peaks decreased and peaks became sharper. This 
suggest for an increase in the crystalline size of the annealed sample . To confi rm this, 
crystalline size of both samples was calculated using the Scherrer fo rmula (30-3 1] . 
D= o.9..1 ---------- - - ------------- (8. 1) 
{3cos8 
Where, D is crystallite size, A. is the wave length of incident x-ray [Cuk ( 1.54056)), ~ 
is the FWHM and e is the diffraction angle for (hkl) plane. To calculate crystallite s ize, three 
most intense peaks (29.24°, 32.9 1° and 44.51 °) were selected. The FWHM of these peaks 
were taken by their Gaussian peak fitting. The average crystallite size for the annealed fi lms 
123 
calculated through aforementioned procedure comes out to be 64 nm. The lattice parameters 
and crystallite sizes are shown in table 8. 1. 
Table 8.1. Showing lattice parameters and crystalline sizes ofY 320 3 : Eu + thin films. 
----------------------------------------------------------------------------------------------------
Annealing temp. (°C) Lattice parameters Crystallite size (nm) 
a c a 
Unannealed 
600 3.7 12 6.575 10.603 70 
700 3.790 6.630 10.625 67 
800 3.796 6.645 10.647 62 
900 3.8 15 6.649 10.654 58 
8.3.1.2 Scanning Electron Microscopy (SEM) 
The morphologies of the nanostructures of the Y20
3
3: Eu + thin fi lms annealed at different 
temperatures were studied by the SEM patterns and presented in Figure 8.3. It shows that the 
obtained fi lms are composed of particles which are nearly spherical in shape. SEM 
micrograph shows the aggregated nature of the secondary particles which were made up of 
the agglomeration of many primary particles. The films which were un-annealed and those 
annealed at lower temperatures appeared with the same morphology and crystalline sizes. At 
higher annealing temperature, the morphology of the film changed drastically to finer one 
probably due to change in the structure as also revealed by XRD. 
124 
Figure 8.3. SEM micrographs of (a) un-annealed and annealed samples 
(b) 600 (c) 800 and (d) 900 °c. 
8.3.1.3 Atomic Force Microscopy 
Figure 8.4 show the AFM images of the Y ~o.~ : Eu, ... thin film~ (a) unannealed, (b) annealed at 
600 °C and (c) annealed at 900 °c in an open-air furnace. The images were obtai ned in 
contacting mode taken over a scale of 5x5 µm 2. Better crystal gain can be obtained by 
annealing at different temperatures, and the ~u rface with different features can also be 
observed. The crystal grains of Y 20 , :ELr'+ film annealed at higher temperature i~ more even 
and the grain si1.e is lesser than that of un-annealcd sample. The uneven grains arc distributed 
on the ~urfacc of Y20 3:Eu'+ un-annealed film , large 72 nm and sma ll 40 nm. 
125 
(a) 
288.45 
nm 
a 
5.00 x 5.00 [um] Z 0.00 - 288.45 [nm] 
0 
c 
126 
(b) 
354.60 
[nm] 
0 
4 
354.60[nm] 
2 
2 
4 
0 
0.00 
5.00 x5.00 [um] Z 0.00 - 354.60 [nm] 
127 
(c) 
367.34 
nm 
0 
4 367.34 
nm 
2 
2 
4 
5.00 x 5.00 [um] Z 0.00 - 367.34 [nm] 
0.00 
Figure 8.4. 30 Height AFM images done in contact mode for the thin films a) un-
annealed, b) annealed at 600 and c) annealed at 900 °c 
8.3.2 Optical properties 
8.3.2.1 Absorption band 
The UV-vis ible reflectance spectra of the as prepared samples are illustrated in Figure 8.5. 
The spectra of all the samples show good optical quality in the visible range due to the 
complete reflectance in the 200-400 nm range. It clearly indicates that, firstly as the 
annealing temperature increased the optical absorption edge shi ft to a higher wavelength 
while the reflectance intensity decreased [32]. Furthermore, absorption bands corresponding 
to the forbidden 4f-4f transitions are usually weak and therefore not detected. From the 
128 
graph, Figure 8.5, (indicated by arrows) the material absorbs between 320 and 350 nm 
depending on the annealing temperatures. Except for the thin film annealed 900 °c, the 
absorption wavelength of the fi lms increased with increasing annealing temperature. 
70 
Unannealed 
600 °C 
60 
700 °C 
800 °C 
--
50 900 °C 
0~  
C1) 
0 
-c 40 ca 0 C1) 
:;::: 
C1) 30 
cc 
20 
10 ....- ----.----~. ........., .._ ___. .,.._ ___~ . ...- ---~ 
200 300 400 500 
Wavelength (nm) 
Figure 8.5. Diffuse reflectance measurements for unannealed and 
those annealed at different temperatures for Y 0 :Eu32 3 +thin films. 
8 .3.2 .1.1 Determination of Eg. from reflectance spectra 
Figure 8.6 shows graph of [F(R)*hv]2 versus hv for the un-annealed and thin fi lms annealed 
at various temperatures. In order to calculate the band gap of the fi lms, the following 
Kubelka- Munk equation was used; 
F (Rx:) = (1 -R"') 2/2R,, = K/5------------------------------------------- (8.2) 
Where R.o= Rsamp1elRrefercnces. K is absorption coefficient and S is scattering coefficient. 
On the other hand the bang gap Eg, and absorption coefficient a of direct band gap 
semiconductor are related through the well- known Tauc relation: 
129 
a hv = C (h v-Eg) 112 1 -------------------------------------------------------(8.3) 
Where hv is the photo energy and C 1 is proportionality constant. When the material scatters 
in a perfectly manner, the absorption coefficient K becomes equal to 2a (K=2a). Considering 
the scattering coeffic ient S as constant with respect to wave length, and using equations (8. l) 
and (2), the following expression can be written: 
[F (RC/))* hv ]2= C2(hv-Eg)-------------------------------------(8.4) 
By plotting [F(R)*hvf against hv and fit the linear region w ith a line and extend it to the 
energy axis, then one can easily obtain Eg by extrapolating the linear regions to [F(R,, 
)hvf=O. The arrows in the figure indicate these ex trapolations for the unannealed and the thin 
films annealed at different temperatures. Eg is the band gap at n=2 fo r direct transitions. 
From the figure, the band gap of Y20
3
3: Eu + was found to be in the range of 4 .6-4.8 eV. It 
can be seen clearl y that the band gap energy of the Y 20 3 decreases linearly with increasing 
temperature. 
0.25 
Unannealed 
600°C 
0.20 700 °C 
800°C 
900°C 
N, ......, 0.15 
.>c  
.~.- ... 
.a._:. . 0.10 
.L..L... . 
0.05 
2 3 4 5 6 
hv (eV) 
Figure 8.6. Graph of F[(R)*hv]2 as a function of band gap 
130 
8.3.4 Photoluminescence 
8.3.4 .1 Excitation spectra 
Figure 8.7 gives the PL excitation spectra of the Y20 :Eu
3
3 + thin fi lms for un-annealed and 
those films annealed at temperatures between 600 and 900 °c. These spectra consist of 
several excitation bands which are ascribed to different transition. In the excitation spectrum 
monitored by 50 70 - F2 transition of Eu
3
+ at 6 12 nm for Y 20
3
3: Eu + thin films, the broad band 
with a maximum at 209 nm originates from the excitation of the oxygen-to-europium co2-
-+Eu3"'") charge transfer band (CTB) and some very weak peaks in the longer-wavelength 
region of 300 nm are ascribed to the f-+ f transitions of the Eu3+ ions. The band located at 
234 run is due to host absorption band. 
8.3.4.2 Emission spectra 
Figure 8.8 shows the emission spectra of un-annealed and films annealed between 600 and 
900 °c. The emission spectra were recorded at excitation wavelength of 209 nm. The most 
intense peaks appearing at 6 12 nm are recorded fo r the un-annealed and films annealed at 
lower temperatures o f 600 and 700 °c. At higher annealing temperature of 800 °c, this peak 
is tremendously quenched possibly due to structural change. At annealing temperature of 900 
0c, all these peaks are quenched but the peak appearing at 6 12 run is red-shifted and 
enhanced which is a confirmation of formation of a new phase. Also, this peak is 
predominant at 588 and 7 15 nm wavelengths. This phenomenon can be explained by the fact 
that the 209 nm radiation excites the Eu3+ ions from ground state to the higher excited state 
5D 70- F2 and quickly relaxes to 
5D 70- F4 level by emitting non-radiative transition. The strong 
red emission band centred at 6 12 nm corresponds to the hypersensitive transition 5D 70- F2. 
Another feeble yellow emission band at 588 nm corresponds to 50 70- F 1 transition, which is 
less sensitive to the host. The reason behind observi ng the intense red emission from 
Y20 3:Eu3+ can be understood by considering the structure of Y20 3. The coordi nate number of 
Y20 3 is six and forms cubic bixbyi te structure with two different sites (C2 and S6) for RE ions 
substitution. The C2 is a low symmetry site without an inversion centre whereas C2 is a high 
symmetry site having an inversion centre. 
131 
400 
Unanneal 
600°C 
700°C 
300 209nm 
- aoo 0c • :::J 210nm 
..c_a
• 
.. . 
.>. 200 ·t-.  ~~ n 
c ~'ti-.C.l.: >  ~ ~o 
c: o..._q ~t"f> - · 0 100 N~ ~'!} ~<!·~)  
M ~'ti-
N Ot'!) ~'ti-
~ tc.:' 
o r ........... ._...:::::::=-;;::::::..- ...- .......... ------1 
200 250 300 350 
wavelength (nm) 
Fig. 8.7. The excitation spectrum o f Y20
3
3: Eu + thin films for un-annealed 
and those annealed at 600, 700, 800 and 900 °c. 
When Eu3+ is located at a low symmetry (C2), the red emission is dominant whereas the 
yellow emission is dominant when Eu3+ is located at high symmetry (S6) [34] . In the present 
case, red emission is dominant suggesting that the location of Eu3+ is more favourable at C2 
site. As the C2 site does not have an inversion centre, electric dipole transition from Eu
3
+ 
ions attached to this site are more favourab le than the magnetic dipole transitions. The 
similarity of the ionic radii of Eu3+ and y 3+ ions allows the easy substitution of y 3+ ions with 
Eu3+ ions at C2 sites giving rise to intense red emission in the un-annealed and films annealed 
at lower temperatures [35]. 
132 
Figure 8.9 represents the chromacity co-ordinates of the PL spectra for the samples (a) Un-
annealed and those annealed at (b) 600, (c) 700, (d) 800 and (e) 900 °c, which are determined 
using the CIE (International Commission on Ill umination) Coordinate Calculator software . 
According to the software the position of the color coordinates lie well in the red and orange 
region. The detailed analysis of the phosphor finds its suitability in making the red/orange 
producing phosphor for display applications and light emitting diodes. 
1000.------------------------------------------------
Emission ---Unannealed 
---600 °C 
800 ---100 °C 
---aoo 0c 
- ---900 °C ::::..s -  
600 
m 
.>...  ·- 400 
."c..' Cl>   
c 
450 500 550 600 650 700 750 800 
Wavelength (nm) 
Figure 8.8. The emission spectrum of Y 20 3 : Eu
3
+ thin films for un-annealed 
and those annealed at 600, 700, 800 and 900 °c. 
133 
0.9----------------------------------------------..... 
a< Sample 1 
ac Sample 2 
>< Sample 3 
fl. Sample 4 
0 .6 
>-
-w (..) x 
0.3 
0-+-----
0 0 .2 0.4 0.6 0.8 
CIE X 
Figure 8.9. The CIE co-ordinates for samples that were un-annealed and 
those annealed at various temperatures 
8.3.4.3 Decay curve 
Figure 8. 10 shows the decay characteristics of the thin films un-annealed and those annealed 
at temperatures between 600 and 900 °c. The films which are un-annealed and those 
annealed at lower temperatures has the highest initial intensities. This is also consistent with 
the PL spectra shown in Figure 8. The fi lms were characterized by the fast and medium 
decays characteristics, s ince they are indicati ve of the different rates of decay for the films. 
The decay curves were fitted according to the equation (3) and gave the decay constants 
listed in Table 2. 
134 
20 • Unanneal 
• • E300 oC 
• A - 700 °C . 15 • T 800 °C :J.  
-co • • ~oc 
~ • 
+-' 
en • 
c 10 • 
Q) • 
+c- ' ••••
  •• 
Q) •••  •
  
> •• • 
+-' 
co 5 ••.•..•. -  Q) l • 
a: 
8 
lirre (rrs) 
Figure 8.10. Showing decay characteristics of Y 20 3: Eu
3
+ phosphor thin films 
for un-annealed and films annealed at 600, 700, 800 and 900 °c. 
I= A1exp (-tlr1) + A2exp (-ti r2) - - ------------ (3) 
Where I is phosphorescence intensity, A 1, and A2 are constants, t is time, r 1 andr 2 are decay 
constants, deciding the decay rate fo r rapid and slow exponentially decay components, 
respectively. The fitting results of parameters r 1 and r 2 are listed in Table 2 below. 
135 
Table 8.2. Results for the fitted decay curves of the un-annealed and fi lms annealed at 
di fferent temperatures. 
Temperature Unannealed 600 700 800 900 
(oC) 
Components 
Fast ('t 1) 0.5896 0.5524 0.5275 0.51 65 0.4962 
Medium ('t2) 0.8564 0.8 124 0.7856 0.7522 0.7050 
Two components namely fast and slow are responsible for the luminescence properties of the 
as- synthesized phosphor. A trend can be observed (Table 2) that the decay constants of the 
phosphors decrease gradually with the increasing annealing temperature. 
8.4 Conclusion 
Y 20
3
3 : Eu + thin films have been successfull y deposited using PLO. The average crystallite 
size o f the films after annealing was 64 nm. The study showed that the annealing temperature 
positively affected the crystalline phase, the morphology and the PL effi ciency of the thin 
films. The un-annealed thin fi lms were amorphous, while the annealed films consist of cubic 
and hexagonal phases depending on the annealing temperatures. SEM show similar 
morphology for the un-annealed and films annealed at lower temperatures which changed at 
higher temperature. The luminescence results show that the PL intensities were quenched and 
red-shifted at higher annealing temperatures showing formation of a new phase. UV-vis 
measurement indicated a band gap in the range of 4.6 to 4.8 eV. 
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138 
Chapter 9 
The influence of different species of gases on the luminescent and 
structural properties of pulsed laser ablated Y20 2S:Eu
3
+ thin films 
9.1 Introduction 
Nanostructure materials are being actively explored nowadays because of their size induced 
novel characteristics and applications ranging from micro tips to optoelectronics. To exploit 
the optical properties of these materials and to meet the ever- increasing demands of energy, 
tremendous emphasis is being placed on alternate sources of energy conservation and low 
power driven display devices. Afterglow phosphors are materials that are readi ly excited by 
ambient lights and continuously radiate visible light for many hours at high brightness levels 
[ I]. Nanoscale Y20 2S:Eu3+ has remarkably different luminescent properties from those of 
bulk samples: such as emission line broadening, lifetime changes and its spectral shift [2]. 
There are several methods to synthesize nanocrystall ine Y20 2S:Eu
3
+ , such as sol-gel [3], 
combustion [ 4), micro emulsion [5], and spray pyrolysis method [6] , but these methods are 
limited in the complexity of the preparation methods. Solid state reaction at room temperature 
is a good method to synthesize nanoparticles [7-9]. Recently, the pulsed laser deposition 
(PLO) technique, which provides a unique process for stoichiometric evaporation of target 
materials and contro l of fi lm morpho logy [ 10, 11], has been used for the deposition of 
oxysul fide fi lms [ 12- 14). In most of the reported works [ 15-17), the Y 20 2S: Eu
3
- phosphors 
have been prepared and investigated in the form of powders. 
Journal of Applied Physics A. •DOI: J0.1007/s00339-0J6-0062-3 
139 
However, for various industrial applications such as device fabrication and surface coatings it 
is also important to investigate the performance of these phosphors in the form of thin films. 
Moreover, it is well documented that thin film phosphors have several advantages over 
powders, such as higher lateral resolution from smaller grains, better thermal stability, 
reduced out gassing, and better adhesion to solid substrates [ 18]. Amongst the techniques 
used to prepare luminescent thin films, PLO has several attractive features, including 
stoichiometric transfer of the target material, generation o f quality plume of energetic species, 
hyper thermal reaction between the ablated cations and molecular gas in the ablation plasma 
and compatibility with background pressures ranging from UHV to 100 Pa [19]. The plasma 
fabricated during pulsed laser ablation is very energetic, and its mobility can be easily 
controlled by changing processing parameters [20]. The presence of a background gas in the 
chamber has a strong influence on the quality of the plasma produced by the laser. The gas 
can modify the kinetic energy and the spatial distribution of the ejected species present in the 
plasma, and it may also induce compositional changes in the deposited films [21 ]. Plume 
colli sions may also provide an increase in the vibrational energy of molecular species [22]. 
Thus, the gas affects the spatial distribution, the deposition rate, the energy and distribution 
of ablated particles thereby controlling the cluster formation, cluster size, cluster energy and 
particle distribution [23]. Oauscher et al. [24] investigated the influence of argon atmosphere 
on the microstructural properties of CaxC0 4Sb 12 films prepared by pulsed laser deposition 
and found that the films deposited in argon atmosphere were more delaminated than those 
prepared in vacuum. The report from Supab et al. [25] on laser ablated ZnO nanorods showed 
that the growth rate of nanostructures in oxygen atmosphere was slower than in argon 
atmosphere. 
In the current work, we report on the influence of vacuum, argon and oxygen atmospheres on 
the morphology, structural and photoluminescence (PL) properties of pulsed laser deposited 
(PLO) Y2 30 2S: Eu + thin films. A detailed report on the influence of working atmosphere on 
the properties of the pulsed laser deposited Y2 0 2S: Eu
3+ thin films are presented. 
9.2 Experimental procedures 
9.2. l Powder preparation 
Eu3+-doped yttrium oxysufide nanocrystals were synthesized using the sol- combustion route. 
The method of synthesis essentially comprises of mixing the precursors in appropriate 
140 
stoichemetric ratios, followed by firing in an air tube furnace at a temperature of 400 °c. The 
stoichemetric ratio between the fuel (thiourea) and the flux (ethanol) is very important 
otherwise ylttrium oxide formation is inevitable. The white fo rmy product was then grounded 
and left to dry in an enclosed oven for 24 hours. 
9.2.2 Pulsed laser deposition (PLO) 
The Si ( 100) wafers used as substrate were fi rst chemical cleaned. The powder was pressed 
without binders to prepare a pellet that was used as an ablation target. The target was 
annealed at 700°C for 12 hours in air to remove water vapour and other volatil e compounds 
that might be trapped in the pellet before placing it on the target holder of the PLO system. 
The deposition chamber was evacuated to a base pressure of 8.2 x I 0-6 mtorr. The Lambda 
Physic 248 nm KrF excimer laser was used to ab late the phosphor pellet in vacuum, argon 
and oxygen atmospheres. A Baratron Direct (Gas Independent) PressureNacuum capacitance 
Manometer (1.33 x 10·2 mtorr) was used for the 50 mtorr pressure measurements. The laser 
energy density, number of pulses and laser frequency were set to 0.74 J/cm2, 12000 and 1O  
Hz respectively. The substrate temperature was fixed at 300 °C, and the target to substrate 
distance was 5 cm. The ablated area was I cm2. 
9.2.3 Characterization 
The Shimadzu Superscan SSX-550 system was used to collect the Scanning Electron 
Microscopy (SEM) micrographs. Atomic Force Microscopy {AFM) micrographs were 
obtained from the Shimadzu SPM - 9600 model. X-ray diffraction (XRD) data was co llected 
by using a SIEMENS 05000 diffractometer using CuKa radiation of A. = 1.5405 nm. PL 
excitation and emission spectra were recorded using a Cary Eclipse fluorescence 
spectrophotometer (Model: LS 55) with a built-in xenon lamp and a grating to select a 
suitable wavelength for excitation. The excitation wavelength was 230 nm and the slit width 
was 10 nm. The afterglow curves for the films were also obtained with the Cary Eclipse 
spectrophotometer. 
141 
9.3 Results and discussion 
9.3.1 X-ray diffraction analysis 
Figure 9.1 shows the XRD patterns of Y 320 2S:Eu + thin films grown at 300 °C in vacuum, 
argon and oxygen atmospheres. The patterns show mixed phases of cubic and hexagonal 
crystal structures. The fi lm grown in 0 2 atmosphere has a cubic phase, while that grown in 
vacuum and argon atmospheres are hexagonal. The average lattice parameters for the 
hexagonal phase a=3.785 nm and c=6.589 nm, are very close to the standard values provided 
in the powder diffraction file PDF #24- 1424. The average lattice parameters for the cubic 
phase a=6.026 nm according to card# 72-0927. The intensities of the XRD peaks were fo und 
to increase in the order argon, vacuum and oxygen. This may be attributed to the enhanced 
oxidation kinetics and improvement in crystalline nature of the fi lms. It is also clear that the 
thin fi lm consist of both nanoparticles and microparticles as seen from the narrow and broad 
XRD peaks. The crystallite size of the films with hexagonal and cubic phase ranging between 
80 and 480 nm were estimated by using equations (I) and (2) respectively; 
1 4(hz +kz +hk) 1z 
- = 3a z + -cz- ------------- (9. I) d z 
1 hz+k z+1z 
- = a z ------------------------ (9 .2) d z 
There is a marginal decrease (-0. 17%) in crystallographic unit-cell that tends to contract due 
to the increase in surface area of the layers for the Y 20 2S: Eu
3
+ nanostructures. This may lead 
to a decrease in the lattice constant. Eu20 3 diffraction peaks from XRD patterns were not 
detected indicating that the Eu3+ was incorporated into the Y 20 2S host lattice homogeneously 
[26]. 
142 
,..., ,..., 
r N 
Card # 72-0927 0 r 
.r.. ... .r.. ... 
...... 
I Vacuum! 
:J 
I 
..r_a, , 
.).',l  Argon I ·e-n 
c: 
.Q.,>  Oxygen I 
-c: ,..., ,..., N 0 ,..., ,.. ,..., ,..., N N N (") r 
Card # 24-1424 0 Or r 0 (") rr N N,.. r r 
.r.. ... r .._. .....
 . .r...... ..
 ... .N...... ..., .r.. ... .(."..)..  
10 15 20 25 30 35 40 45 50 55 60 65 
2 theta (degrees) 
Figure 9.1. The XRD spectra of the Y20 2S: Eu3+ thin films deposited In 
vacuum and different gas atmospheres. 
9.3.2 Scanning electron microscopy (SEM) 
Figure 9.2 shows the SEM pictures o f the fi lms prepared in different working atmospheres. 
The fi lm deposited in vacuum, shown in Figure 9.2 (a), had a smooth surface with numerous 
bigger spherical particles . The film deposited in Ar atmosphere (Figure 9 .3(b)) had a much 
rougher surface packed with larger nu mber of spherical particles than that of the film 
deposited in the 0 2 atmosphere. The film that was depos ited in the 0 2 atmosphere, in Figure 
9.3(c), had a rougher surface with small cracks on the surface. The c rysta ll ite sizes as shown 
by SEM micrographs consist of both nano and micro particles. The SEM micrographs show 
that the surfaces of the films prepared in the gas atmosphere are much rougher than that 
deposited in vacuum. The increase in surface roughness in the 0 2 and Ar atmospheres is due 
to the enhanced particulate fo rmation in the plume, which is a typical characteristic of high-
pressure laser ablation [27]. 
143 
Figure 9.2. SEM micrographs fo r thin films 
deposited in (a) vacuum, (b) argon and (c) 
oxygen atmosphere 
9.3.3 Atomic force microscopy (AFM) 
Figures 9.3 shows AFM images of the samples deposited in (a) vacuum, (b) argon and (c) 
oxygen atmosphere. It is clear that almost hexagonall y-shaped nanopartic les were deposited 
during the deposition process. The particles were also less agglomerated for films grown in 
gas atmospheres (Fig. 9.3 (b) versus Fig. 9.4 (c)). ln vacuum, the plume does not undergo 
scattering (collisions) due to gas particles hence the arrival of the particle is more compared 
to case of argon and oxygen and thus higher PL intensity. Argon be ing a heavier gas scatters 
more particles as compared to oxygen; however the speed of argon particles is less than those 
of oxygen. The arrival of the particles is higher in case o f oxygen, hence intensity. 
144 
(a) 
360.61 
[nm] 
0 
360.61[nm] 
0 
0.00 
5.00 x 5.00 [um] Z 0.00- 306.61 [nm] 
145 
{b) 
354.60 
[nm] 
0 
4 
354.60[nm] 
2 
4 
0 
0.00 
5.00 x5.00 [um] Z 0.00 - 354.60 [nm] 
146 
(c) 
288.45 
[nm] 
\ 
0 1 
288.45 
4 
[nm] 
2 
4 
0 
0.00 
5.00 x 5.00 [um] Z 0.00 - 288.45 
Figure 9.3. AFM image of thin films deposited in (a) vacuum, (b) argon and (c) 
oxygen atmospheres 
Collisions between the vaporised particles close to the target in the case of films deposited in 
gas atmosphere lead to nucleation and growth of smaller nanoparticles when arri ving at the 
substrate. In vacuum there are virtuall y no collisions between the particles before reaching 
the ubstrate. Longer res idence time of the particles in the plume, as i the ca e of fi lm 
deposited in gas atmosphere, lead to more evenl y dis tributed pa rticles (F ig . 9 .3 (c)). Light 
emis ion from the spherical shaped phosphor particles as excited by the e lectron beam i 
more inten e due to the fact that much less photons encounter total internal reflecti on [28-30]. 
The increase in the deposition pressure i reported to have caused an increase in the 
connectivity (agglomeration) between particles due to s intering of small particles (31-32]. 
147 
This would eventuall y lead to grain growth at high enough pressure . In this case, we also 
fo und that more agglomeration occurred during deposition in gas atmosphere than in vacuum. 
9.3.4 Photoluminescence results 
Figure 9.4(a) indicates the excitation spectra of Y 20 2S: Eu
3
+ thin films grown under vacuum, 
argon and oxygen atmosphere. The inset indicates the magnified spectrum for sample 
deposited in oxygen atmosphere. Excitation spectra were recorded keeping the emission 
wavelength at 612 nm. These spectra consist of two charge trans fer bands (CTB). The band 
located at 237 nm is due to 0 2 3---+ Eu + CTB, while that at 312 nm is due to Eu3 2+ ---+ s - CTB. 
Figure 9.4(b) show the deconvulated excitation spectra of Y 0 S:Eu32 2 + th in films deposited in 
vacuum, argon and oxygen. The figure indicates two transitions at 237 and 312 nm 
wave lengths. The transition located at 237 nm is due to charge transfer band between 0 2- and 
Eu3+ ions, while that located at 312 and 315 nm is associated with Eu3+ ---+ s2- charge trans fer 
band. 
100 0 2 3-- Eu • 
en 
80 8 f.-- Oxygen 
- 237nm ::::J 
-ca 60 i: 4 >- ·;;; 
.'!::: ."!!  
t/) .5 
-c Q) 40 c 
300 350 400 45 
20 Wavelength (nm) 
0 
200 250 300 350 400 450 
Wavelength (nm) 
Figure 9.4(a). Excitation spectra fo r Y 20 2S: Eu
3
+ thin films deposited in vacuum, 
argon and oxygen atmosphere. Inset: Excitation spectrum for sample depos ited 
in oxygen atmosphere 
148 
o2-:_Eu3+ Fit 1 vacuum 100 Fit 2 vacuum 
Vacuum 
Fit 1 argon 
- 80 Fit 2 argon . 237nm Argon ::::.:s  
_ca.  60 
.~ ·-.. 
."c..'   40 
N 
(1) j 312nm c 315nm 
20 ~/ 
200 250 300 350 400 450 
wavelength (nm) 
Figure 9.4(b). Deconvo luted excitation spectra for Y20
3
2S: Eu + thin films 
deposited in vacuum, argon and oxygen atmosphere 
Figures 9.4(c) to 4(e) show excitation wavelengths for each film deposited in different 
atmospheres. In each case, the exc itation wavelengths were found at different emissions 
wavelengths in order to find the best exc itation for each film. Figure 9.4(c) indicate the 
exc itation wavelength for the film deposited in vacuum. At all the emiss ion wavelengths of 
590, 612, 625 and 655 nm, no reasonable excitation wavelength have been achieved 
indicating this situation does not favour vacuum atmosphere. Figure 9.4(d) indicate excitation 
wavelengths for the film deposited in argon atmosphere. The best excitation wave length was 
achieved when the emission wavelength was at 612 nm, followed by 625 and 590 nm 
respectively. The exc itation wavelength with the least intensity is achieved at 655 nm 
emission wavelength. Similar results are obtained for film deposited in oxygen atmosphere as 
shown in Figure 9.4(e). 
149 
30..---------------------~~~~~~~ 
Y 2o2S:Eu
3+ deposited in vacuum --A. em 590 nm 
--A. em 612 nm 
--A. em 625 nm 
20 --A. em 655 nm 
-. ::J
-.
  
m 10 
.>..  
."c: C.l.')   0 
c: 
-10 
200 250 300 350 400 
Wavelength (nm) 
Figure 9.4(c). Excitation spectra of Y2 0 2S:Eu
3
+ thin fi lm deposited in vacuum 
recorded at 590, 6 12, 625 and 655 nm emission wavelengths 
150 
aoo ..... -o---~-3-------------~~~~~~!!!!!!'!!!I Y 2 2S:Eu + deposited in argon -- /.... em 590 nm 
700 -- /.... em 612 nm 
-- /....em 625 nm 
600 -- /....e m 655 nm 
. 
~- 400 >-·"-c:'  300 
Q) 
c 200 
100 
0 
200 250 300 350 400 
Wavelength (nm) 
Figure 9.4(d). Excitation spectra of Y 20 2S:Eu
3
+ thin film deposited in argon 
recorded at 590, 6 12, 625 and 655 nm emission wavelength 
Y 0 S:Eu + deposited in oxygen 
600 2 2 - A. em 590 nm 
-- A. em 612 nm 
-- A. em 625 nm 
500 -- A. em 655 nm 
- 400 ::::J 
-. ca ->- 300 fl) 
-c: CV 200 c: 
100 
0 
200 250 300 350 400 
Wavelength (nm) 
Figure 9.4(e). Excitation spectra of Y 0 S:Eu32 2 + thin film deposited in oxygen 
recorded at 590, 612, 625 and 655 nm emiss ion wavelengths 
151 
Figure 9.5 (a) and (b) show the room temperature PL emission spectra for the Y20 2S: Eu
3
+ 
films deposited in vacuum, argon and oxygen. All the samples were excited at 245 nm using 
a monochromatized xenon lamp. The intense red emission peak that is observed at 612 nm 
from the film deposited in vacuum is ascribed to the 5Do - 7F2 transition of Eu3+ [33). A 
shoulder at 595 nm for the vacuum sample is also visible, which is ascribed to Eu2+ 
transitions. Minor emission peaks associated with residual Eu3+ [34, 35) was also observed at 
402, 469, 496, 514 and 540 nm for the films deposited in the vacuum atmosphere. The 
relative ratios of the 595 nm to 613 nm peaks (Eu2+ to Eu3+ ratio) clearly support this 
observation. It should be mentioned that 0 2 atmosphere should favour the stabilization of 
Eu3+ cations (oxidation), while vacuum should play the opposite role and favour Eu2+ 
(reduction). The intensity of the film deposited in the vacuum atmosphere was the highest 
followed by the one prepared in the argon atmosphere, though greatly quenched. All the 
peaks for the sample deposited in 0 2 are totally quenched due to the smallest intensity as 
compared to that of vacuum and argon, hence not visible. But when PL emission spectra for 
sample deposited in 0 2 is drawn alone, the peaks are clearly visible as shown in the inset of 
Figure 9.5 (a). The most intense peak for this sample also appears exactly at 612 nm just as 
that of vacuum. In the case of film deposited in argon atmosphere, the peak at 612 nm is 
quenched while that at 315 nm is enhanced. In all the three cases, the peak at 612 nm is 
associated with 0 2 3 2 3- - Eu + charge transfer band (CTB) between 0 • and Eu + ions, while the 
peak at 315 nm is due to Eu3+ - s2• charge transfer band. These peaks are clearly shown by 
Figure 9.5 (b) inset. A shift of the Y20 2S: Eu
3
+ main peak toward shorter wavelengths was 
measured for sample deposited in argon atmosphere. This is due to the quantum size effect of 
the nano-sized particles, compared with those deposited in vacuum and oxygen as confirmed 
by XRD results. 
152 
Vacuum 
250 I Otygeri 
Argon 
0.15 Oxygen 
200 
- -;- 0.10 
-::::J .!!!. C'tl 150 ->- f o.~ 
-"c:Q')   100 c: 0.00 
50 
200 300 400 500 600 700 800 
Wavelength (nm) 
Figure 9.5 (a). Emission spectra for Y20 2S: Eu
3
+ thin films deposited in vacuum, 
argon and oxygen atmosphere. Inset. Emission spectrum for thin film deposited 
in oxygen. 
153 
250 Vacuum 
Argon 
Oxygen 
- 200 .. 3' ::::s  .!!!- 1 -ca 150 -~ f U'J 
-c: Cl> 100 c: 
50 :m 
200 300 400 500 600 700 800 
Wavelength (nm) 
Figure 9 .S(b). Emission spectra for Y 320 2S: Eu • thin films deposited in vacuum, argon and 
oxygen atmosphere. Inset . Emission spectrum for thin fi lm deposited in argon. 
Figures 9.5(c) to 9.5(d) show emis. ion wavelengths for each film deposited in different gas 
atmospheres. In each case, the emis. ion wavelength were found at different excita ti on 
wave lengths in order to optimize emission for each fi lm. Figures 9.5(c) indicate the emission 
wavelength fo r the fi lm deposited in vacuum. At all the excitation wavelengths of 237, 245, 
260 and 290 nm, high quality emiss ion wavelength with high intensit ies have been achieved 
with 245 and 260 nm excitation providing the highe t intensity. Figure 9.5(d) indicate 
e mission wavelengths for the fi lm depo ited in oxygen atmosphere. Similarly, the be t 
emission wavelength was achieved when the excitation wavelength was at 245 and 260 nm, 
followed by 290 nm respectively. T he excitation wavelength with the least intensity is 
achieved at 237 nm excitation wavelength. The film deposited in oxygen atmosphere 
indicates the lowe t inten ity as shown in Figure 9.5(d) at all excitations. 
154 
700..-----------------------.....m !!!!!!!!!!!!!!!!!!!!~--------· 
--A. ext 237 nm 
600 --A. ext 245 nm 
-- A. ext 260 nm 
500 --A. ext 290 nm 
-. 
-::J.  400 C'O ~ 
·..e-. 300 n 
.cQ.. )  200 
c 
400 500 600 700 800 
Wavelength (nm) 
Figure 9.S(c). Emission spectra o f Y 20
3
2S:Eu + thin fi lm deposited in vacuum 
atmosphere recorded at 237, 245, 260 and 290 nm exc itation wavelength. 
155 
3
Y  0 S : E u •  d e p o s i t e d  i n  o x y g e n  
- - A .  e x t  2 3 7  n m  
2 2
2 0  
- - A .  e x t  2 4 5  n m  
- - A .  e x t  2 6 0  n m  
- - A .  e x t  2 9 0  n m  
1 5  
.  
-
: : : : J  
.  
« J  
-
> -
1 0  
. ~  
c :  
" '  
Q )  
c :  
-
5  
0  '  
4 0 0  
5 0 0  6 0 0  7 0 0  8 0 0  
W a v e l e n g t h  n m )  
3
F i g u r e  9 . S ( d ) .  E m i s s i o n  s p e c t r a  o f  Y  0 S : E u +  t h i n  f i l m  d e p o s i t e d  i n  o x y g e n  
2 2
a t m o s p h e r e  r e c o r d e d  a t  2 3 7 ,  2 4 5 ,  2 6 0  a n d  2 9 0  n m  e x c i t a t i o n  w a v e l e n g t h s  
3
I t  i s  w e l l  k n o w n  t h a t  t h e  6 1 3  n m  p e a k  d u e  t o  t h e  E u +  i s  m u c h  m o r e  p r o n o u n c e d  t h a n  t h e  4 0 2 -
5 9 5  n m  p e a k  f o r  t h i n n e r  f i l m s  ( l e s s  l a s e r  p u l s e s )  [ 3 3 ] .  W i t h  a l l  t h e s e  i n  m i n d  i t  m i g h t  b e  
3
s p e c u l a t e d  a t  t h i s  s t a g e  t h a t  t h e  s h i f t  i n  t h e  e m i s s i o n  w a v e l e n g t h  o f  t h e  E u +  e m i s s i o n  i n  a r g o n  
3 2
a t m o s p h e r e  m i g h t  a l s o  b e  d u e  t o  r e d u c t i o n  o f  E u +  t o  E u + .  T h e  c r y s t a l  f i e l d  d u e  t o  d i f f e r e n t  
s t r u c t u r e s  o f  c u b i c  a n d  h e x a g o n a l  m u s t  a l s o  b e  k e p t  i n  m i n d .  T h e  n a n o  s i z e  o f  t h e  p a r t i c l e s  
f o r m e d  d u r i n g  d e p o s i t i o n  a l s o  c a n  p l a y  a  r o l e .  T h e  h i g h e r  P L  i n t e n s i t i e s  f r o m  t h e  f i l m s  
d e p o s i t e d  i n  t h e  v a c u u m  a t m o s p h e r e s  c a n  b e  a s c r i b e d  t o  t h e  f i l m s '  r e l a t i v e l y  r o u g h e r  s u r f a c e  
a s  o b s e r v e d  f r o m  t h e  A F M  p i c t u r e s  i n  F i g u r e  3 .  I t  i s  w e l l  k n o w n  t h a t  r o u g h  s u r f a c e s  i n c r e a s e s  
t h e  p r o b a b i l i t y  o f  l i g h t  e m i s s i o n  f r o m  t h e  s u r f a c e  b y  l i m i t i n g  t h e  c h a n c e s  o f  t o t a l  i n t e r n a l  
r e f l e c t i o n  a t  t h e  f i l m - s u b s t r a t e  i n t e r f a c e  [ 3 4 - 3 7 ] .  
9 . 3 . 5  D e c a y  c u r v e s  
F i g u r e  9 . 6  s h o w s  t h e  d e c a y  c h a r a c t e r i s t i c s  o f  t h e  t h i n  f i l m s  d e p o s i t e d  i n  v a c u u m ,  0 a n d  A r  
2  
a t m o s p h e r e s .  C o n s i s t e n t  w i t h  t h e  P L  d a t a  i n  F i g u r e  5 ,  t h e  f i l m  p r e p a r e d  i n  v a c u u m  
1 5 6  
atmosphere has the highest initial intensity fo llowed by the film prepared in the argon 
atmosphere. The fluorescence of Y20 2S: Eu3+ is believed to originate from the photo-
oxidation of Eu2+ cation under UV-irradiation [38] . According to this model, an electron from 
the 4 f7 ground state is excited to the 
4 f6_5d 21 level of Eu "'" followed by an electron capture from 
the valence band reducing Eu2+ to Eu+. 
The films were further characterized by the fast and slow decays characteristics [38], since 
they are indicative of the different rates of decay for the films. The decay curves were fitted 
according to the equation (3) and gave decay constants listed in Table 9. 1. 
I= A 1 exp (-t/t1) + A1exp (-ti t2) ---------------------------------- (9 .3) 
where I is the phosphorescence intensity, Ai, and A2, are constants, tis time, 1 1 and 12, are 
decay times for exponential components, respectively. The fitting results of parameters t1a nd 
t2 are listed in Table 9. 1 below. 
Table 9.1: Decay constants for the fitted decay curves of the thin 
fi lms ablated in vacuum, argon and oxygen atmospheres. 
Atmosphere Vacuum Argon Oxygen 
Components Decay constants(t , s) 
Fast (t 1) 0.48±0.002 0.47±0.007 0.47±0.001 
Medium (t 2) 0.66±0.006 0.62±0.006 0.59±0.005 
157 
--Vacuu 
Argon 
Oxygen 
....-.. 
• 
U) 
.cQ.. )  
c 
2 4 6 
lime(ms) 
Figure 9.6. The decay curves for PLOY2 0 S: Eu
3
2 + thin films deposited in 
vacuum and different gas atmospheres. 
It is clear trom Table Y.1 that the decay constants tor the vacuum and argon samples are 
longer than that of the oxygen sample. According to these results the gas atmospheres did not 
play any role in the trap level concentration and also in the possible trap types. 
9.3.6 Optical properties 
9.3.6.1 Reflectance spectra 
The UV-vis reflectance spectra of the samples are given in Figure 9.7. The spectra of all the 
samples show good optical quality in the visible range due to the complete reflectance in the 
200- 500 nm range. The sharp absorption edge is characteristic of a homogeneous structure 
[39]. The figure shows that the absorption edge shifted to higher wavelength in the order 
vacuum, argon and oxygen. Absorption bands corresponding to the forbidden Eu3+ 4f-4f 
transitions were also detected for film deposited in 0 2 . The band at around 343 nm is 
158 
attributed to the exci ton absorption, which is red-shift compared with powder Y 0 S:Eu32 2 + 
(40]. The absorption peaks at around 290 and 340 nm are ass igned to 5Do-7F1and 
50 70- F2 
transitions of Eu3+ ions, respectively [4 1] . 
9.3.6.2 Determination of band gap from reflectance spectra 
The Kubeika- Munk equation was used to calculate the band gap of the as- deposited thin 
film using a diffuse reflectance spectrum. The bang gap (Eg), and absorption coefficient a of 
direct band gap semiconductor are related through the Tauc relation. By plotting [F(R)*hv] 2 
against hv and fit the linear region with a line and extending it to the energy ax is, then one 
can easi ly obtain Eg by extrapolating the linear regions to [F (R)*hv] 2=0. 
65 -Vacuum 
60 - Argon 
55 - Oxygen 
- 50 -0~  Q) 45 
(,) 
-c cu 40 (,) Q) 35 
;;:: 
Q) 
a: 30 
25 
20 
15 
200 300 400 500 600 700 800 
Wavelength (nm) 
Figure 9.7. UV-vis diffuse reflectance spectra of nanocrystalline 
Y20 2S:Eu
3
+ thin film deposited in vacuum, argon and oxygen atmo pheres 
159 
Figure 9.8. Graph of F[(R)*hv] 112 as a fu nction of band gap 
energy 
The dependence of the band gap energy of the Y20 2S on the vacuum and di fferent gas 
spec ies is shown in Figure 9.8. The observed optical band gap for Y20 2S: Eu
3
+ thin films 
increased in the order vacuum (3.4 eY), argon (4.3 eY) and oxygen (4 .4 eY) as shown in 
Figure 8. The average Eg. value for the thin films was found to be 4.07 eY, which is in a 
good agreement with the literature values by other researchers [42]. The change in optical 
band gap values may also be due to the change of crystal structure of the Y20 2S th in films 
(Hexagonal verse Cubic). This i · also confi rmed by the fact that the PL emission intensity of 
the fi lms deposited in vacuum is more than S times greater than the intensity of Y 20 2S:Eu
3
+ 
thin fi lms deposited in gas atmosphe re. 
The increase in band gap ene rgy and the shi ft of the absorption edges to higher wavelengths 
in the case of film deposited in gas atmosphere might be due to the presence of defect states 
and disorder due to the different depos ition atmospheres [43-44] . The gas atmospheres might 
have introduced new states close to the conduction band of the Y 20 2S:Eu
3
+. A new defect 
160 
band is therefore formed below the conductions which lead to reduction in the effective band 
gap [44). 
9.4 Conclusion 
Red- emitting Y 20 2S: Eu
3
+ thin film phosphors were successfu ll y ablated on Si (I 00) 
substrates by the pulsed laser deposition technique. The X-ray diffraction patterns showed 
mixed phases of cubic and hexagonal crystal structures. The SEM micrographs show that the 
surfaces of the films prepared in the gas atmosphere are much rougher than that deposited in 
vacuum. Intense red emission, with a maximum peak at 612 run associated with the 5D 70- F2 
transitions of Eu3+ was detected from the fi lms deposited in vacuum and argon atmosphere. 
The working atmosphere has a severe influence on the PL properties of the films. The films 
deposited in the vacuum atmospheres gave better PL and afterglow properties than the film 
prepared in gas atmospheres. The crystallinity and PL properties of the sample ablated in 
vacuum gave the best bright red emission. UV-vis measurement gave an average band gap of 
4.69 eV. 
Reference 
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164 
Chapter 10 
Summary and suggestions for future work 
10.1 Thesis summary 
A simple and efficient chemical route to prepare a promising red- emitting phosphors by sol-
combustion synthesis is presented. Ethanol has the effect of decreasing the water needed, 
hence simplifying the experimental procedure by dissolving rare earth ni trate and sulfur-
contained organic fuel into an even solution. This prompts the formation o f rare earth oxide 
by ign iting first during heating that leads to a combustion decomposition reaction. Y 320 3 : Eu + 
microcrystalline structures were obtained using thiourea as organic fuel. The increase in the 
ratios of fuel to the host decreased the grain size and the maximum PL intensity of the 
phosphors. Also, trivalent-europium doped La20 2S microcrystals powders were prepared by 
sol- combustion method and their material properties were studied at room temperature. The 
phosphors emit bright red phosphorescence originating from the 50 70- F2 transition of Eu
3+. 
The XRD pattern of the as-synthesized powder revealed the presence of the hexagonal 
La20 2S phase with an average crystalline sizes of 178 run with a space group P3m I [ 164] and 
cell constants of (a) 0.4128 nm and (c) 0.6985 run. The lower ratio of fue l to oxidizer favors 
the formation of the La20 2S rather than La20 3 phosphor. De-hydration of the reactants was 
crucial fo r the successful synthesis of the oxysulfide phase. SEM images of the as-
synthesized powders showed that the morphology consisted of a foamy agglomeration and a 
continuous three-dimensional network. At higher La/S molar ratio the morphology appears as 
regular crystalline lattices. In order to study the effect of co-dopant, Y 0 3 32 3:Eu +:Ho + red-
emitting phosphor was synthesized and the influence of mole percent of Ho3+ ions has been 
investigated. The powders composed of nanoparticles. The size of the particles depended on 
the Ho3+ mole concentration. An increase in the Ho3+ percent resulted in smaller crystallite 
sizes. SEM show particles are highly agglomerated and form a connected network type with 
some vacant space among them, which is expected due to the evolution of different gases 
during combustion of the gel. At lower mole of Ho3+ all the peaks for both Eu3+ and Ho3+ 
were observed. The PL measurements showed red emission of Y20 3:Eu
3+:Ho3+ powder with 
the most intense peak appearing at 626 run, which is assigned to the 5D 70- F2 transition of 
165 
Eu3+. At higher mole concentration of Ho3+, energy transfer occurred from Eu3+ to Ho3+. The 
band gap energy increased with increase of Ho3+ mole concentration. 
To compare the material properties of the phosphors, the as- prepared powder samples were 
deposited on Si ( lOO) substrates by PLO under various conditions. In the initial sample, 
Y 20 2S:Eu
3+ thin films were successfully grown onto the Si ( I 00) substrates at different 
oxygen ambient. XRO show mixed phases of cubic and hexagonal crystal structures. The 
fi lms grown at low and high oxygen partial pressure are predominantly cubic, while that 
grown at moderate pressure ( I 00 mtorr) is predominantly hexagonal. The thin films were 
composed of nanoparticles. The size of the particles depended on the oxygen partial pressure. 
The decrease in the oxygen partial pressure resulted into big particles which are actually 
piling up of smaller particles and therefore a rougher surface. The PL measurements showed 
red emission of Y20 2S:Eu
3
+ thin films as well as the powder Y20 2S: Eu3+ with the most 
intense peak appeari ng at 6 19 nm, which is assigned to the 50 70- F2 transition of Eu
3
+. This 
intense peak is quenched at higher 0 2 partial pressure greater than 20 mtorr. The emission 
peak at 590 nm due to 5D 70- F 1 transition of Eu
3+ is also quench~d at higher 0 2 partial 
pressure. Uv-vis measurement revealed an average band gap of 4.03 eV. To investigate effect 
of annealing temperature, the next sample, Y 20 3: Eu
3
+ thin films have been successfully 
deposited using PLO. The average crystallite size of the films after annealing was 64 nm. The 
study showed that the annealing temperature positively affected the crystalline phase, the 
morphology and the PL efficiency of the thin films. The un-annealed thin fi lms were 
amorphous, whi le the annealed films consist of cubic and hexagonal phases depending on the 
annealing temperatures. SEM show similar morphology for the un-annealed and films 
annealed at lower temperatures which changed at higher temperature. The luminescence 
results show that the PL intensities were quenched and red-shifted at higher annealing 
temperatures showing formation of a new phase. UV-vis measurement indicated a band gap 
in the range of 4.6 to 4 .8 eV. Finally, a Red- emitting Y 20 2S: Eu
3
+ thin film phosphors were 
successfully ablated on Si ( I 00) substrates by the pulsed laser deposition technique and effect 
of different species of gases were studied. The X-ray diffraction patterns showed mixed 
phases of cubic and hexagonal crystal structures. The SEM micrographs show that the 
surfaces of the fi lms prepared in the gas atmosphere are much rougher than that deposited in 
vacuum. Intense red emission, with a maximum peak at 6 12 nm associated with the 50 70- Fi 
transitions of Eu3+ was detected from the films deposited in vacuum and argon atmosphere. 
The working atmosphere has a severe influence on the PL properties of the films. The fi lms 
166 
deposited in the vacuum atmospheres gave better PL and afterglow properties than the film 
prepared in gas atmospheres. The crystallinity and PL properties of the sample ablated in 
vacuum gave the best bright red emission. UV-vis measurement gave an average band gap of 
4.69 eV. 
10.2 Suggestions for future work 
The work presented in this thesis is very valuable to the literature of powder synthesis as well 
as laser ablated Y20 2S:Eu
3
- thin films. Other methods of synthesising powder Y20 2S: Eu3+, 
Y20
3
3 : Eu + and La20 2S: Eu
3
+ should be investigated. However, there are some investigations 
which need more attention to get more light on some of the findings . With regards to the 
deposition conditions, it is important to investigate more on the changes of oxygen, argon and 
nitrogen with the PL properties of Y 20
3
2S: Eu + thin film phosphors. A wider range of the gas 
pressures (0.1-2 Torr) is needed for better optimizations. Other deposition parameters which 
are of interest are the target-substrate, the laser-target distance and the substrate type. It is 
important to do some more investigations on the in-situ post deposition analysis on the films 
and compare with the ex-situ annealing to determine the better annealing process for better 
PL and afterglow properties of Y20 2S: 
3
+ films. Annealing in different atmospheres such as 
argon, nitrogen can also be interesting. In the current work, all the depositions were done by 
using a 248 nm KrF laser source. It will be interesting to use another source of laser with 
higher energy such as 193 nm ArF in the course of seeing how the properties of Y 20
3
2S: Eu + 
thin films wi ll change. 
Publications 
( I) Abdub G. Ali, Francis B. Dejene, Hendrik C. Swart, Reinhardt J. Botha, Kittesa Roro, 
Liza Coetsee and Mart M. Biggs, Optical properties of ZnO nanoparticles synthesized by 
varying the sodium hydroxide to Z inc acetate molar ratios using a sol- gel process, Cent. Eur. 
J. Phys. 9(2).2011.1321-1326. DOI: J0.2478/Sl1534-0JJ-0050-3 
(2) A.G. Ali, B.F. Dejene, H.C. Swart, Effect of Mn3+ doping on structural and optical 
properties of Sol-gel derived ZnO nanoparticles, Cent. Eur. J. Phys. JO (2).2012.478-
484.DOI: 10.2478/SJ 1534-011-0106-4 
(3) A.G. Ali, B.F. Dejene, H.C. Swart, Synthesis and Characterization of Y20 3S:Eu
3
+ 
phosphors Using Sol-combustion Method, Physica B: Physics of Con densed Matter, 
Volume 439, pp. 181-184. DOI: J0.1016/j.physb.2013.11.051 
167 
(4) A.G. Ali, 8 .F. Dejene, H.C. Swart, The influence of oxygen partial pressure on material 
properties of Eu3+- doped Y 20 2S thin film deposited by Pulsed Laser Deposition, Physica B: 
Physics of Condensed Matter (2016), pp. DOI information: 10.1016/j.physb.2015.10.005 
(5) A.G. Ali, B.F. Dejene, H.C. Swart, Temperature dependence of structural and 
luminescence properties of Eu3+-doped Y20 3 red-emitting phosphor thin fi lms by pulsed laser 
deposition, Journal ofA pplied Physics A, APYA-D-15-01792.1, DOI: I0.1007/s00339-016-
9946-5 
(6) A.G. Al i, B.F. Dejene, H.C. Swart, The influence of different species of gases on the 
luminescent and structural properties of pulsed laser ablated Y20 2S:Eu
3+ thin fi lms, Journal 
ofA pplied Physics A. •DOI: 10.1007/s00339-016-0062-3 
(7) A.G. Ali, B.F. Dejene, H.C, Swart, Energy transfer and material properties of Y20 3: 
Eu3+:Ho3+ nanophosphors synthesized by sol- combustion method, This paper has been 
accepted by Journal of Chemistry and Engineering 
(8) A.G. Ali , B.F. Dejene, H.C, Swart, Structural, thermal and optical properties of ZnO, 
Si02 and ZnO: Si02:Ce
3+ nanoparticles by sol-gel process. This paper was submitted to the 
Journal of Rare- earths 
(9) A.G. Ali , B.F. Dejene, H.C, Swart, Structural and Luminescent Properties of 
Y20 3:Eu
3+:Ho3+ Nanocrystals Prepared by Sol- Combustion Synthesis" at the South African 
Institute of Physics (SAIP) conference July 2015. This paper is under review 
( I 0) A.G. Ali, B.F. Dejene, H.C, Swart, Synthesis and characterization of europium 
activated yttrium oxysulphide by sol- combustion method. This paper was submitted to the 
Journal ofR are- earths 
( 11 ) A.G. Ali, B.F. Dejene, H.C, Swart, Synthesis and characterization of europium acti vated 
lanthanum oxysulphide by sol- combustion method, This paper was submitted to the Journal 
of Luminescence 
( 12) A.G. Ali, B.F. Dejene, H.C, Swart, Synthesis and characterization o f europium activated 
lanthanum oxide by so l- combustion method, This paper was submitted to the Journal of 
Luminescence 
168 
Papers presented at conferences 
(I) A paper entitled "Optical properties of ZnO nanoparticles synthesized by varying the 
sodium hydroxide to Zinc acetate molar ratios using a sol- gel process" at the South African 
Institute of Physics{SAIP) conference in Durban in July 2009. This paper was published in 
Cent. Eur. J. Phys. 9(2).20 11.1321- 1326. DOI: 10.2478/Sl 1534-011-0050-3. 
(2) A paper entitled "Effect of Mn3+ doping on structural and optical properties of Sol-gel 
derived ZnO nanoparticles"at the South African lnstitute of Physics (SAlP) conference in 
Johannesburg in July 20 IO.This paper was published in Cent. Eur. J . Phys. I 0(2).2012.478-
484.DOI: 10.2478/Sl 1534-011-0 106-4 
(3) A paper entitled "Synthesis and Characterization of Y20 3S:Eu
3
+ phosphors Using Sol-
combustion Method" at the South African Institute of Physics conference in Kariega in July 
20 11. This paper is in press at the Journal of Condensed Matter Physics- Physica B. 
( 4) A paper entitled "Structural, thermal and optical properties of ZnO, Si02 and ZnO: 
Si02:Ce
3
+ nanoparticles by sol-gel process" at the African Material Society {AMRS) 
conference in Victoria Falls Zimbabwe December 20 11. This paper was submitted to the 
Journal of Rare- earths. 
(5) A paper entitled "Synthesis and characterization of europium activated yttrium 
oxysulphide by sol- combustion method" at the African Material Society conference(AMRS) 
in Victoria Falls Zimbabwe December 20 11. This paper was submitted to the Journal of 
Rare- earths. 
(6) A paper entitled "Synthesis and characterization of europium acti vated lanthanum 
oxysulphide by sol- combustion method" at the South African Institute of Physics conference 
in University of Pretoria in July 2012. This paper was submitted to the Journal of 
Luminescence. 
(7) A paper entitled "Synthesis and characterization of europium activated lanthanum oxide 
by sol- combustion method" at the African Laser Center{ALC) conference in University of 
Namibia in Windhoek in November 20 12. This paper was submitted to the Journal of 
Luminescence. 
169 
(8) A paper entitled "Synthes is and Characterization of Y 20 3:Eu
3+ phosphors Using Sol-
combustion Method" at the South African Institute of Physics conference in Kariega in May 
2013. This paper i. in Press in the Journal of Condensed Matter Physics- Physica B. 
(9) A paper entitled "Structural and Lu mine cent Properties of Y 0 :Eu3+: Ho32 3 + Nanocrystals 
Prepared by Sol- Combustion Synthesis" at the South African Institute of Physics (SAIP) 
conference in Richards Bay in July 20 13. This paper is under review. 
( I 0) The influence of different species of gases on the luminescent and structural propertie 
of pulsed laser ablated Y 32 0 2S:Eu + thin fi lms at the conference of laser ablation (COLA) at 
Cairns, Australia. This paper was submitted to the Journal of Applied Physics A. 
( 11 )Temperature dependence of structural and luminescence properties of Eu3+ -doped Y2 0 3 
red- emitting phosphor thin film by Pulsed Laser Deposition at the conference of laser 
ablation (COLA) at Cairns, Australia. This paper was submitted to the Journal of Applied 
Physics A. 
Appendix: Emission spectra for fi lms deposited in vacuum and at different 
oxygen partial pressure. The inset show emis ion spectrum o f 
Y20 2S:Eu
3
+ powder phosphor. 
1000 
500'.) - v acuum 
- FW.del" 
- 20 mtorr 
400) 
800 - 60 mtorr 
- 100 mtorr 
"3' JOO) 
~ - 140 mtorr 
:::J 600 
n; I 200) ->-U) 
c: 400 
100) 
-Q) c: 
2 00 500 600 
VVavel~h (rm) 
0 
500 520 540 560 580 600 620 640 660 
Wavelength (nm) 
170