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Sensors and Actuators B 146 (2010) 395–402

Contents lists available at ScienceDirect

Sensors and Actuators B: Chemical

journa l homepage: www.e lsev ier .com/ locate /snb

he influence of sintering temperature on the photoluminescence propertiesf oxyapatite Eu3+:Ca2Gd8Si6O26 nanophosphors

. Seeta Rama Rajua, Hong Chae Junga, Jin Young Parka, Byung Kee Moona,∗, R. Balakrishnaiaha,ung Hyun Jeonga, Jung Hwan Kimb

Department of Physics, Pukyong National University, Busan 608-737, Republic of KoreaDepartment of Physics, Dong Eui University, Busan 614-714, Republic of Korea

r t i c l e i n f o

rticle history:eceived 16 January 2010ccepted 21 February 2010vailable online 26 February 2010

eywords:a2Gd8Si6O26:Eu3+

intering temperature effecthotoluminescence properties

a b s t r a c t

Eu3+ activated oxyapatite Ca2Gd8Si6O26 (CGS) nanophosphors were synthesized using a solvothermalreaction method. The structural and luminescent properties of these nanophosphors were investigatedas a function of sintering temperature and Eu3+ ion concentration. The SEM images of the prepared phos-phors reveal spherically shaped particles in the nanometer range and the XRD patterns confirm theirhexagonal structure. The photoluminescence excitation (PLE) spectra of Eu3+:CGS showed the chargetransfer band (CTB) and intense f–f transitions of Eu3+ and Gd3+. The intensity of the f–f transitions ofGd3+ increases with increasing the sintering temperature and decreases with increasing the Eu3+ con-centration. It was observed that the CTB of Eu3+ shifted to a shorter wavelength region with an increase

ecay profileseddish-orange emission

the crystallite size due to the variation of coordination environments. The photoluminescence (PL) spec-tra of Eu3+:CGS exhibit two emission lines corresponding to the 5D0 → 7F0 transition which results fromoccupation of Eu3+ ions in two different low symmetry local sites in CGS host lattice. The optimized sin-tering temperature and concentration of Eu3+ were observed for Eu3+:CGS nanophosphors based on thedominant red (5D0 → 7F2) emission intensity under NUV (395 nm) excitation. The decay curves of 5D0

level show that the lifetime decreases with increasing the crystallite size. These luminescent powdersntial

are expected to find pote

. Introduction

In recent years, rare-earth (RE) ions doped phosphors havettracted a great deal of attention for their wide range of potentialpplications, including optoelectronic devices, temperature sen-ors, solid-state lighting (SSL), optical communication systems anduminescent probes or labels in biological systems and so on [1–5].mong them, SSL is one of the most focused areas because the solid-tate white light emitting diodes (WLEDs) have more advantageshan the traditional incandescent or fluorescent lamps [1]. Manyesearch reports suggested that the WLEDs would reduce globallectricity consumption by about 50% by taking advantage of theonversion from direct electricity to light rather than the processes

n which light is the by-product of another conversion, as withraditional lamps [6–8].

Different methods are available to mimic white light fromlue/near-UV (NUV) LEDs. The combination of a blue chip and

∗ Corresponding author at: Department of Physics, Building No. 5, Pukyongational University, Daeyon 3 Dong, Busan 608-737, Republic of Korea.el.: +82 629 5569.

E-mail address: bkmoon@pknu.ac.kr (B.K. Moon).

925-4005/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.snb.2010.02.056

applications such as WLEDs and optical display systems.© 2010 Elsevier B.V. All rights reserved.

yellow phosphor has already been developed and is commer-cially available, but tricolor white LEDs consisting of a soft UVchip emitting at a 400 nm wavelength with red, green, and blue(RGB) phosphors continues to be a challenge [9–11]. In the tricolorwhite LED, inorganic oxides would be the best candidates for RGBphosphors in terms of both chemical stability and luminescenceefficiency [12,13]. This development of WLED generated the neces-sity for the development of tricolor phosphors, which can be excitedefficiently, typically in the range of 380–480 nm. The phosphor-converted color LEDs show some advantages, for example, a widelytunable color range from blue to red region and a stable color qual-ity under different operating conditions, especially under differentforward-bias currents and during the industrial mass production[1]. Thus, the NUV pc-LEDs are expected to have great potentialapplications in the field of solid-state lighting. Nowadays, full colordisplays made of RGB LEDs or WLEDs, which are energy efficient,reliable and having long working lifetime (∼100,000 h) have beengiven much attention for future applications [12].

Development of novel phosphors is, therefore, one of the mostimportant aspects in designing luminescent devices. It is knownthat the phosphors with oxyapatite structure (space group P63/m)have been effective host lattices for luminescent materials. Amongthe many synthetic oxyapatites, the ternary rare-earth-metal sil-

3 s and Actuators B 146 (2010) 395–402

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cate with the oxyapatite structure of the form Ca2RE8Si6O26RE = Y/Gd) is the efficient host lattice for the luminescence of var-ous RE3+ (Sm3+, Eu3+, Tb3+ and Dy3+) ions and mercury like ions.his oxyapatite host lattice consists of two low symmetry cationicites, that is, the 9-fold coordinated 4f sites with C3 point symme-ry and 7-fold coordinated 6 h sites with Cs point symmetry [14,15].oth sites are very suitable for the luminescence of RE3+ ions owingo their low symmetry features.

Among the RE3+ ions, red luminescence of trivalent europiumEu3+) ions due to intra-configurational f–f transitions are ofechnological importance because it has been widely applied tohosphors for color display systems and SSL industry. It is wellnown that the characteristic emission of Eu3+ strongly dependn their local environment in the host lattice. That is, the orange5D0 → 7F1) emission corresponding to the magnetic-dipole tran-ition with the selection rule, �J = 0, ±1, which is insensitive to therystal field symmetry around the Eu3+ ion and the red (5D0 → 7F2)mission belongs to the hypersensitive (forced electric-dipole)ransition with the selection rule, �J = 0, ±2, which is stronglynfluenced by the outside surrounding environment. When Eu3+ isocated at a low symmetry local site (without inversion symmetry),he red emission is often dominant in the emission spectrum andhen Eu3+ is at a high symmetry local site (with inversion symme-

ry center), the orange emission is stronger than the red emissionnd is dominant in the emission spectrum [16].

Generally, phosphors are prepared by traditional solid-stateeaction (SSR) method. Due to insufficient mixing and low reac-ivity of raw materials, several impurity phases easily co-exist inhe product. In addition, phosphor particles prepared through thisSR method have large in size (generally in micrometer range)17]. However, in recent years, several wet chemical techniquesuch as co-precipitation method, sol–gel, combustion, hydrother-al, solvothermal and spray-pyrolysis were used to prepare the

hosphor precursor [17–23]. These wet chemical processing meth-ds offer an intimate mixing of starting materials and hence havehorter diffusion distances between reactants, requiring relativelyow temperatures for the formation of final products with excellenthemical homogeneity [24]. However, to achieve high crystallinitynd transform the powder into a dense ceramic, sintering at ele-ated temperatures becomes mandatory.

To the best of our knowledge, upto now, a very few reportsith fewer luminescent properties of Eu3+:Ca2Gd8Si6O26 (CGS)

hin films and phosphors by sol–gel and reverse co-precipitationethods have been found [14,15,25–27]. However, no reports have

een found on the detailed luminescent properties of Eu3+:CGSanophosphors using solvothermal process, and it is necessary

or the phosphors with a narrow size distribution, spherical-likehape and fine particle in order to offer high resolution and highuminescence efficiency. In the present work, we report the struc-ural and detailed luminescent properties of spherically shapedu3+:Ca2Gd8Si6O26 using solvothermal synthesis at various sinter-ng temperatures and at various Eu3+ concentrations. It was foundhat the Eu3+:CGS nanophosphors exhibit reddish-orange emissionith seven intense emission peaks centered at 579, 586, 591, 598,

15, 622, and 627 nm.

. Experimental

Different concentrations of Eu3+ activated CGS nanophos-hors were prepared using solvothermal process with the

omposition Ca2(Gd1−xEux)8Si6O26 (x = 0.5–10 mol%). The sto-chiometric amounts of high purity grade calcium nitrateetrahydrate (Ca(NO3)2·4H2O), gadolonium nitrate hex-hydrate (Gd(NO3)3·6H2O), Europium nitrate pentahydrateEu(NO3)3·5H2O), and tetraethyl orthosilicate (Si(OC2H5)4) were

Fig. 1. TG/DTA curves of Eu3+:CGS powder precursor.

dissolved in 40 ml of 2-propanol. All reagents were taken with-out any further purification and stirred vigorously by using themagnetic stirrer until the formation of homogeneous solution andwas transferred into a stainless steel autoclave with a Teflon liner(80 ml capacity and 50% filling). It was then heated to 230 ◦C at arate of 2 ◦C/min and maintained for 5 h with magnetic stirring (at180 rpm) to make stable networks between the reactants. Aftergradually cooling down to room temperature, the precipitate wasseparated by a centrifugal separator with 3000 rpm for 3 min andthen dried at 50 ◦C for a day in the ambient atmosphere. The driedpowder was sintered at different temperatures for 5 h and wasbrought to room temperature.

Thermogravimetric/differential thermal analysis (TG/DTA) ofthe dried powder precursors was carried out with a Material Analy-sis and Characterization TG-DTA 2000. This experiment was carriedout at a heating rate of 5 ◦C/min and the samples were heated fromroom temperature to 900 ◦C. X-ray diffraction (XRD) patterns of sin-tered powders were recorded on X’PERT PRO X-ray diffractometerwith CuK� = 1.5406 Å. The morphology of the sintered particles wasexamined by field emission scanning electron microscope (FESEM),model JEOL JSM-6700 FESEM. Osmium coating was sprayed on thesample surfaces by using Hitachi fine coat ion sputter E-1010 unitto avoid possible charging of specimens before FESEM observationwas made on each time. The reflectance spectra of Eu3+:CGS phos-phors at different sintering temperatures were measured by usingJASCO V-670 UV–visible spectrophotometer. The Fourier transforminfrared (FTIR) spectra of these sintered samples were measured byusing the PerkinElmer (Spectrum GX) spectrophotometer. The roomtemperature photoluminescent spectra were recorded on a Pho-ton Technology International (PTI) fluorimeter with a Xe-arc lampof power 60W and the lifetimes were measured with a phospho-rimeter attachment to the main system with a Xe-flash lamp (25 Wpower).

3. Results and discussion

Fig. 1 illustrates the TG/DTA curves of the powder precursorof Eu3+:CGS obtained by the solvothermal method (range fromroom temperature to 900 ◦C). The TG curve shows two distinct

◦

weight loss steps up to 700 C; no further weight loss was regis-tered up to 900 ◦C. The weight loss is related to the decomposition ofthe organic matrix. Simultaneously, three exothermic peaks wereobserved with maxima at 386, 499 and 830 ◦C on the DTA curveduring the process mass loss. The first two exothermic peaks indi-

G. Seeta Rama Raju et al. / Sensors and

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The particles acquire sphere-like morphology formed from agglom-

ig. 2. XRD patterns of Eu3+:CGS phosphors at different sintering temperatures.

ate that the thermal events can be associated with the burnout ofrganic species involved in the precursor powders of the residualitrogen and the third exothermic peak is due to crystallization ofGS powder from the amorphous component. These results were

ell in agreement with the XRD analysis.

Fig. 2 shows the XRD patterns of Eu3+:CGS nanophosphor sam-les sintered at different temperatures from 750 to 1300 ◦C. Thehosphor was found to be amorphous until the precipitate powders

Fig. 3. SEM images of Eu3+:CGS phosphors at various sintering temperatures (

Actuators B 146 (2010) 395–402 397

were sintering at 750 ◦C and crystallized as pure CGS at 850 ◦C with-out any intermediate phase, indicating that the powders preparedby solvothermal process is pure in both chemistry and crystallinephase. It is well known that the pure phase is favorable for lumi-nescent properties of phosphors. The diffraction peak intensity andthe crystallite size increases significantly with increasing the sin-tering temperature. At all sintering temperatures (850–1300 ◦C),the diffraction peaks are attributed to the hexagonal crystal formof oxyapatite based on the space group P63/m and, which are inwell agreement with the standard JCPDS card [PDF (28-0212)]. Ingeneral, the crystallite size can be estimated by using the Scher-rer equation, Dhkl = k�/ˇ cos �, where D is the average grain size,k(0.9) is a shape factor, � is X-ray wavelength (1.5406 Å), ˇ is thefull width at half maximum (FWHM) and � is the diffraction angleof an observed peak, respectively. The strongest diffraction peaksare used to calculate the crystallite sizes of Eu3+:CGS nanophos-phors sintered at 850, 900, 1000, 1100, 1200 and 1300 ◦C, whichyield average values of about 38.6, 42, 63, 71.4, 80.9 and 94 nm,respectively. The crystallinity of the CGS phosphors with differentconcentrations of Eu3+ was the same, because the energy supplywas constant.

The morphological features of the Eu3+:CGS nanophosphors sin-tered at different temperatures were investigated by taking theSEM images. The typical morphological images are shown in Fig. 3.

eration between the CGS particles during the period of samplesintering. It is well known that the spherical-shaped particles are ofgreater importance because of their higher packing density, lowerscattering of light and brighter luminescence performance. From

a) 850 ◦C, (b) 900 ◦C, (c) 1000 ◦C, (d) 1100 ◦C, (e) 1200 ◦C and (f) 1300 ◦C.

398 G. Seeta Rama Raju et al. / Sensors and Actuators B 146 (2010) 395–402

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ig. 4. Reflectance spectra of Eu3+:CGS phosphors sintered at different tempera-ures.

ig. 2, it is clear that, an average particle size is in nanometer rangend tends to increase with increasing the sintering temperature.ote that, when the sintering temperature increases to 1200 ◦C, theggregation between the CGS particles was almost same, and theintering temperature increases more than 1200 ◦C, the agglomera-ion between the CGS particles was also increased with the smallerGS particles. The particle sizes are in good agreement with thealculated crystallite sizes from XRD patterns.

Fig. 4 shows the diffused reflectance spectra as a function ofintering temperature. It is evident that by increasing the sinteringemperature up to 1200 ◦C the reflectance in the range 375–700 nmecreased, and the sintering temperature above 1200 ◦C, theeflectance in the range 375–700 nm increases. This is in goodgreement with the SEM studies because larger particles give lesscattering and agglomeration between the particles causes thencrease of scattering.

The FTIR spectra with different sintering temperatures ofGS:Eu nanophosphors have been shown in Fig. 5. First of all, thereppear no absorption bands that are characteristic of OH groups

ith stretching (3572 cm−1) and vibration (630 cm−1) modes in

patite structure, indicative of the successful formation of the oxya-atite phase in our samples. The IR absorption due to SiO4 units cane assigned based on four kinds of modes consisting of symmetric

ig. 5. FTIR spectra of Eu3+:CGS phosphors at various sintering temperatures.

Fig. 6. (a) Excitation spectra of Eu3+:CGS phosphors at different sintering tem-peratures, inset shows the excitation spectra of amorphous sample at 750 ◦C. (b)Excitation spectra of Eu3+:CGS phosphors at various Eu3+ concentrations.

stretching (�1), symmetric bending (�2), anti-symmetric stretch-ing (�3), and anti-symmetric bending (�4). As seen from Fig. 5,the stretching vibrations �1 and �3 give fairly intense and broadabsorption bands at wavenumber between 1200 and 800 cm−1.This result is consistent with that reported in the literature aboutIR absorption of the SiO4 unit in the apatite structure [28]. How-ever, it is impossible to report IR spectra below 560 cm−1 with ourinstrument measurement limit. So, we cannot report the bendingvibrations (�2 and �4) of SiO4 unit because those bands are pre-sented at lower frequencies between 600 and 400 cm−1. The IRspectrum exhibits weak bands between 1800 and 1200 cm−1 con-stituted by the absorption of C–H bending (1372 cm−1) and C Ostretching (1737 cm−1), which are almost disappeared by increas-ing the sintering temperature.

Fig. 6(a) and (b) shows the photoluminescence excitation (PLE)spectra for the Eu3+:CGS samples sintered at different temperatures

3+

and different concentrations of Eu ions, respectively, by monitor-ing the emission wavelength at 615 nm. The both PLE spectra revealthe broad excitation peaks overlapped with the f–f transitions ofGd3+ ion in the shorter wavelength region, and it also consists ofsharp excitation bands in the longer wavelength region due to the

s and Actuators B 146 (2010) 395–402 399

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Fig. 7. (a) Emission spectra of Eu3+:CGS phosphors at different sintering tempera-

G. Seeta Rama Raju et al. / Sensor

–f transitions of Eu3+ ion. The broad excitation band also calledharge transfer band (CTB) is due to the charge transfer betweenhe completely filled 2p orbital of O2− ion and the partially filledf orbital of the Eu3+ ion and the position of this band dependstrongly on the host lattice. The other sharp excitation peaks dueo intra-configurational f–f transitions of Gd3+ and Eu3+ ions aressigned to the electronic transitions of (8S7/2 → 6D9/2) at 252 nm,8S7/2 → 6I11/2) at 275 nm, (8S7/2 → 6I9/2) at 277 nm for Gd3+ ionsnd (7F0 → 5D4) at 362 nm, (7F0 → 5G3) at 376 nm, (7F0 → 5G4) at82 nm, (7F0 → 5L6) at 395 nm, (7F0 → 5D3) at 412 nm, (7F0 → 5D2)t 464 nm for Eu3+ ions. These Eu3+ sharp excitation peaks indi-ate that violet and blue laser diodes/LEDs are efficient pumpingources in obtaining Eu3+ emissions. In both the excitation spectra,he dominated excitation is at 395 nm.

It is worthy to note from Fig. 6(a) that the intensities of CTB at75 nm and (7F0 → 5L6) transition at 395 nm of Eu3+ ions are almosthe same in the amorphous state [shown in the inset of Fig. 6(a)],intered at 750 ◦C. While increasing the sintering temperature from50 to 1200 ◦C the broadness (or strength) of the CTB decreases and

7F0 → 5L6) transition in the longer wavelength region increasesnd also the CTB maxima shift towards the higher energy side,hich indicate that the quantum efficiency of Eu3+ increases with

ncreasing the sintering temperature [26]. As the sintering temper-ture increases above 1200 ◦C the band maxima shift towards theower energy side. The CTB is related to the stability of the elec-ron of the surrounding O2− ion, i.e., the CT transition is sensitiveo a ligand environment (the bonding energy between the centralon and the ligand ions). The fact that it occurs in the solid-stateionic oxides) is explained by the strong potential field at the anionO2−) sites due to the surrounding ions. If this potential increase,he energy required transferring an electron from the O2− ion to theation (Eu3+) increases too, and the charge transfer band moves toigher energy side. This means that the mixing of Eu3+ and O2−

rbitals decreases, as a result the intensity of the CTB decreases29–31]. In nanoscale samples, especially very tiny ones, the hostattice is imperfect and incompact and Eu–O distance is long, indi-ates that the Eu–O bonds become weaker and less covalent andigh ionicity, weakening the binding energy of an electron to O2−.herefore, the electron needs less energy to transfer from O2− tou3+, resulting in the CTB shift to lower energy. On the other hand,s particle size increases the surface to the volume ratio decreasesnd degree of disordered of the nanostructured system decreasesnd Eu–O distance is short. As a result, it requires more energyo remove an electron from an O2− ion; therefore, the CTB shiftedowards higher energy. It can also be noted that, at higher tempera-ures the CTB shifts towards the lower energy side, this effect mighte associated with an appearance of Eu3+–Eu3+ pairs in the systemecause of the decrease in the average distance between Eu3+ ions.hese results are also consistent with the Hoefdraad’s study on theTB of Eu3+ in oxides, where he concluded that the CTB band posi-ion of Eu-doped oxide depends on the bond length of Eu–O and theoordination environment around Eu3+ [32]. Note that the variationf the CTB described in different literatures is quite incompatible.or example, in Y2O3:Eu3+ nanocrystals, some authors observed alue shift where as the others observed a red shift [30,33]. As it isell known, the resonant excitation cross-sections of f–f transitions

or trivalent rare-earth ions are generally small and play an impor-ant role in the excitation of Eu3+ ions. Furthermore, from Fig. 6(b),t can also be clear that by increasing the Eu3+ concentration, the–f transitions of Gd3+ ions are almost disappeared and the strengthf the CTB increases due to increase in the covalency between the

u3+ and O2−.

Upon irradiation of nanophosphors with commercial UV lamps,ed light emission was so strong as to be visible to the naked eye.ecause the intensity of emissions depends on many factors such

tures, inset shows the relationship between sintering temperature, crystallite sizeand 7F2 emission intensity. (b) Emission spectra of Eu3+:CGS phosphors at variousEu3+ concentrations, inset shows the relationship between the concentration andemission intensity.

as the porosity, scattering of light, UV transmittance, reflectanceand surface roughness of the phosphors. The photoluminescencespectra at different sintering temperatures and at different concen-trations of Eu3+ activated CGS nanophosphors by exciting at 395 nmare shown in Fig. 7(a) and (b). Obviously, these phosphors can beexcited with the intense f–f transition at 395 nm (7F0 → 5L6) andCTB (not shown here). In any cases, it is the good sign that thesephosphors can strongly absorb NUV light and transfer the excita-tion energy to the red radiation, indicating that the phosphors aresuitable to be excited by a NUV InGaN chip.

All these phosphors exhibit a stronger emission band with apeak at 615 nm, which can be ascribed to the forced electric-dipole5D0 → 7F2 transition of the Eu3+ ions located at the sites with-out inversion symmetry. The moderate emission bands at 591,

5 7

594, 598 nm are due to the magnetic-dipole D0 → F1 transitionand the other two emission bands at 579, 586 nm are due to5D0 → 7F0 transitions of Eu3+ ions. The position of the hypersensi-tive transition is well matched with the Isaacs reported result thatthe Ca2Gd8Si6O26:Eu3+ showed the strongest emission at around

400 G. Seeta Rama Raju et al. / Sensors and Actuators B 146 (2010) 395–402

Table 1Asymmetric ratio, lifetime and color coordinates of Eu3+:CGS nanophosphors fordifferent sintering temperatures at 1 mol% Eu3+ concentration.

Sinteringtemperature (◦C)

I(5D0 → 7F2)/I(5D0 → 7F1) Lifetime(ms)

Color coordinates

750 2.09 2.96 (0.551, 0.402)850 3.73 3.04 (0.589, 0.395)900 3.89 2.96 (0.592, 0.389)

1000 4.12 2.86 (0.601, 0.385)

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Table 2Asymmetric ratio, lifetime and color coordinates of Eu3+:CGS nanophosphors fordifferent concentrations of Eu3+sintered at 1200 ◦C.

Concentration(mol%)

I(5D0 → 7F2)/I(5D0 → 7F1) Lifetime(ms)

Color coordinates

0.5 3.77 2.00 (0.591, 0.385)1 4.38 2.04 (0.611, 0.369)2 4.78 2.87 (0.619, 0.364)

peratures have been listed in Table 1. The results demonstratedthat the decay profiles of all the samples were well fitted to sin-gle exponential. From Fig. 8 and Table 1, it is clear that the lifetimedecreases with increasing sintering temperature at above 850 ◦C. Insome literature, the luminescence lifetime for Eu3+ decreased with

1100 4.23 2.52 (0.605, 0.375)1200 4.38 2.04 (0.611, 0.369)1300 3.98 1.53 (0.607, 0.382)

14.7 nm [27]. Considerable emissions were not observed from theigher 5D1,2 levels in the shorter wavelength range of 400–550 nm.his is ascribed to the fact that the smaller energy gaps betweenD2 and 5D1 or 5D1 and 5D0 of Eu3+ can be bridged by the vibrationnergies of the silicate groups present in the oxyapatite. The higherD2 and 5D1 excited levels are then relaxed once to the lowest 5D0xcited level.

Generally, the 7FJ energy levels of Eu3+ split into some compo-ents under crystal field effects caused by surrounding ions [33,34].nder lower CS symmetry of the 6 h site, 7F1 and 7F2 split into threend five components, respectively, due to completely lifted “2J + 1”egeneracy, if overlapping of higher 5DJ emissions can be excluded,hough they cannot be with complete certainty. The emissions at79 and 586 nm due to 5D0 → 7F0 are particularly noted. Since Starkplitting of either 5D0 or 7F0 is not possible, the presence of twoeaks must be due to separate emissions from the two differentites occupied by the Eu3+ ion [34]. The 5D0–7F0 transition is for-idden both as a magnetic-dipole and an electric-dipole, and it isften very weak or altogether absent, but both sites in this crys-al have low symmetry, which relaxes the selection rules. Theseesults suggest a strong contribution to overall luminescence fromu3+ ions on both the sites.

From Fig. 7(a), a spectral structure for the phosphor sinteredt 750 ◦C is typical PL of Eu3+ present in the amorphous host. It islear that an unresolved broader and stronger emission band withpeak at 615 nm is ascribed to the electric-dipole 5D0 → 7F2 transi-

ion of the Eu3+ ions located inhomogeneously at the sites withoutnversion symmetry. A weaker emission band ranging from 583o 604 nm is due to the magnetic-dipole 5D0 → 7F1 transition andminor single emission peak for the 5D0 → 7F0 transition is also

bserved at 579 nm due to its amorphous nature. This observationoincides with the XRD results as described above. However, phos-hors sintered at above 850 ◦C exhibit an enhanced, sharpened, andell resolved Stark splits for 5D0 → 7F1 and 5D0 → 7F2 emissions

nd two 5D0 → 7F0 transitions are observed.The sintering temperature at above 850 ◦C the intensity of

he 5D0 → 7F2 transition increases greatly and the intensity ofD0 → 7F0 transitions increases as compared with the intensity ofD0 → 7F1 transitions. The relationship between sintering temper-ture, crystallite size and emission intensity has also been shown inhe inset of Fig. 7(a). The intensity ratio of 5D0 → 7F2 to 5D0 → 7F1,lso called the asymmetric ratio (R), is close to being related to theocal environment of Eu3+. Generally, larger the intensity ratio ofD0 → 7F2 to 5D0 → 7F1, lower the local symmetry. The asymmetricatios of Eu3+:CGS nanophosphors with various sintering temper-tures were also calculated, and the results are shown in Table 1.he results show that the asymmetric ratio increases slightly withncreasing sintering temperature up to 1200 ◦C due to increase ofrystallite size, which confirms the decrease in local symmetry and

ence an increase in red emission (5D0 → 7F2 at 615 nm). The sin-ering temperature increases above 1200 ◦C, the asymmetry ratioecreases due to increased agglomeration between CGS particlesith smaller CGS particles and increased reflectance, as a result

4 4.82 3.17 (0.624, 0.361)6 4.89 3.42 (0.629, 0.353)8 5.35 3.02 (0.632, 0.351)

10 5.65 2.86 (0.635, 0.349)

increase in local symmetry. From the above results, the optimizedsintering temperature is 1200 ◦C. So, we measured the PLE and PLspectra with different concentrations at 1200 ◦C.

From Fig. 7(b), clearly, the shapes of emission spectra with dif-ferent Eu3+ concentrations upon sintering at 1200 ◦C are similar.When the Eu3+ concentration increases from 0.5 to 6 mol% the emis-sion intensity also increases. When the Eu3+ concentration furtherincreased above 6 mol% the emission intensity decreased due toconcentration quenching. The concentration quenching might beelucidated by the following two factors, (i) the excitation migra-tion due to resonance between the activators is enhanced when thedoping concentration is increased, and thus the excitation energyreaches quenching centers, and (ii) the activators are paired orcoagulated and are changed to quenching center. Furthermore, theasymmetric ratio increases with increasing the concentration ofEu3+ presented in Table 2. The above results suggest that the phys-ical and chemical properties of Eu3+:CGS nanophosphors are stablewhen the phosphors sintered at 1200 ◦C and the optimum concen-tration of Eu3+ is 6 mol%. The correlation between concentration ofEu3+ and emission intensity has been shown in the inset figure ofFig. 7(b).

The decay profiles of the luminescence of 5D0 level of Eu3+:CGSnanophosphors at different sintering temperatures have beenrecorded under excitation at 395 nm and emission at 615 nm areshown in Fig. 8. The measured lifetimes for different sintering tem-

Fig. 8. Decay curves of Eu3+:CGS phosphors at various sintering temperatures, insetshows the relationship between the sintering temperature, lifetime and crystallitesize.

G. Seeta Rama Raju et al. / Sensors and Actuators B 146 (2010) 395–402 401

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phosphor–liquid crystal display, J. Appl. Phys. 88 (2000) 4660–4665.

ig. 9. Relationship between the 7F2 emission intensity, concentration and lifetimef Eu3+:CGS phosphors at various concentrations.

article size [35,36], on the contrary, in the other literature, theifetime increased, which depends on various factors such as theost and the preparation method, etc. [37,38]. The increase in 5D0

ifetime with decreasing particle size is attributed to a radiativeecay rate [38], whereas the decrease of 5D0 lifetime was gen-rally attributed to the increased non-radiative rate due to thenergy transfer from excited states of rare-earth ions to surfaceefect states [35,36]. The relationship between sintering temper-ture, crystallite size and lifetime has been shown in the inset ofig. 8. In addition the lifetimes of Eu3+:CGS nanophosphors havelso been studied as a function of Eu3+ concentration and, whichere presented in Table 2. It is clear that when the Eu3+ concentra-

ion increases to 6 mol% the lifetime also increases from 2 to 3.42 msnd when the concentration further increases above 6 mol% the life-ime decreased from 3.42 ms. Based on our results, we conclude theossible facts of the observed results as follows: when the Eu3+ con-entration is below 6 mol% the luminescent centers can be thoughtf as isolated, energy transfer from the interior to surface center isneffective owing to its longer distance between Eu3+ ions, a rela-ively longer lifetime could be observed. Also due to larger energyap (around 12,500 cm−1) between 5D0 and its immediate lowerevel (7F6), non-radiative mechanism through phonon relaxation isenerally almost absent in Eu3+-doped materials. However, a higheru3+ concentration would lead to form a resonance energy transferet, which acts as an additional channel to the non-radiative cen-ers on the surface, there by a shorter lifetime can be observed. Thisrocess is well correlated with that of the emission process. Theoncentration dependent emission intensity and lifetime investi-ations are presented in Fig. 9.

The Commission International De I-Eclairage (CIE) chromaticityoordinates for Eu3+:CGS nanophosphors at different sintering tem-eratures and at different concentrations were calculated and listed

n Tables 1 and 2, respectively. The CIE chromaticity coordinates ofptimized concentration and optimized sintering temperature haslso been represented in Fig. 10. From Tables 1 and 2 it is clear thatu3+:CGS exhibits excellent CIE coordinates. According to Wang etl. [41] two factors are considered as the reason for the phenomena:1) the asymmetric ratio increases with increasing crystallinity, and2) the Eu doping effect becomes stronger in the case of particlesith higher crystallinity, resulting in an improved activation degree

f Eu3+ (hence the improved red color emission). The CIE chromatic-ty coordinates of 6 mol% Eu3+:CGS upon sintering at 1200 ◦C showseddish-orange color because the intensities of 5D0 → 7FJ transi-ions are comparable to each other. It is useful for the production

Fig. 10. CIE diagram of Eu3+:CGS phosphor for the optimized concentration of6 mol% upon sinering at 1200 ◦C.

of artificial white light to be similar to those of natural white lightowing to its better spectral overlap [39,40]. The above observationshint at the promising application of Eu3+:CGS nanophosphor to pro-duce white light for NUV-LEDs as well as optical display systems.

4. Conclusion

The effect of sintering temperature and concentration on thestructural and PL characteristics of hexagonal oxyapatite Eu3+:CGSnanophosphors prepared by solvothermal reaction method weresystematically investigated. The morphology of the obtainedEu3+:CGS nanophophors were investigated using SEM analysis.The excitation spectra of these phosphors demonstrated a signif-icant blue shift of the CTB appearing in the shorter UV regionand the intensity of f–f transitions of Eu3+ increases in the NUVregion with increasing sintering temperature. Detailed analyticaldata revealed that the PL intensity increases under the excita-tion of NUV (395 nm) light with increasing sintering temperatureand concentration, respectively. Strong concentration quenching ofEu3+ emission was observed when the Eu3+ doping level exceeds6 mol%. Moreover, the decay profiles became single exponentialand the lifetime decreases with increasing sintering temperatureand crystallite size due to the increased non-radiative rate. From theabove detailed luminescent properties, we are able to conclude thatthe Eu3+-doped Ca2Gd8Si6O26 under the excitation of NUV 395 nmradiation is promising red phosphors for applications in WLEDs andoptical display systems.

Acknowledgement

This study was financially supported by Pukyong National Uni-versity in the 2009 Post-Doc. Program.

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Biographies

G. Seeta Rama Raju is a postdoctoral researcher in Prof. Byung K. Moon researchlaboratory at Pukyong National University. He specializes in structural and lumines-cence properties of rare-earth ions activated micro and nanophosphors. He earnedhis Ph.D. in 2008 from Sri Venkateswara University at Tirupati, India, under theguidance of Prof. S. Buddhudu.

Hong Chae Jung is a master student of Prof. Byung K. Moon at Pukyong NationalUniversity. He received his B.Sc. from the same university in 2008.

Jin Young Park is a Ph.D. student at Pukyong National University under the supervi-sion of Prof. Byung K. Moon. He received his B.Sc. in Physics in 2007 and his master’sdegree in 2009, both degrees from the Pukyong National University.

Byung K. Moon received his Ph.D. degree in physics in 1989 from Pusan NationalUniversity. He joined Pukyong National University, where currently he is a profes-sor of physics. His research interests are in the area of nano- and nanostructuredinorganic materials and their optical properties including their synthesis, structuraland microscopic analysis.

R. Balakrishnaiah is a postdoctoral researcher in Prof. Byung K. Moon research lab-oratory at Pukyong National University. He specializes in luminescence propertiesof lanthanide-doped glasses and phosphors. He received his Ph.D. (Physics) in 2008from Sri Venkateswara University, Tirupati, India under the guidance of Prof. C.K.Jayasankar.

Jung Hyun Jeong is a professor at the department of Physics, Pukyong NationalUniversity and the Chairman of BK21 research group. He published many researchpapers in highly reputed journals like Elsevier, ACS, AIP and IOP. Today, He is guid-ing research group of around 10 people at Pukyong National University aiming at

developing novel luminescent materials.

Jung Hwan Kim is a professor at the department of Physics, Dong Eui University. Hereceived his Ph.D. in Physics from Pusan National University. Today, he is supervisingthe research focused on the fluorescent materials within the group of Prof. Jung HyunJeong.