Glass Science & Technology Research @ UPM

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2012/2013

Hindawi Publishing CorporationAdvances in Condensed Matter PhysicsVolume 2013, Article ID 783207, 6 pageshttp://dx.doi.org/10.1155/2013/783207

Research ArticleEffect of ZnO on the Thermal Properties of Tellurite Glass

H. A. A. Sidek, S. Rosmawati, B. Z. Azmi, and A. H. Shaari

Glass Ceramic and Composite Research Group, Department of Physics, Faculty of Science, Universiti Putra Malaysia,43400 Serdang, Selangor, Malaysia

Correspondence should be addressed to H. A. A. Sidek; [email protected]

Received 16 August 2012; Accepted 29 January 2013

Academic Editor: Nigel Wilding

Copyright © 2013 H. A. A. Sidek et al. This is an open access article distributed under the Creative Commons Attribution License,which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

Systematic series of binary zinc tellurite glasses in the form (ZnO)𝑥(TeO2)1−𝑥 (where 𝑥 = 0 to 0.4 with an interval of 0.05

mole fraction) have been successfully prepared via conventional melt cast-quenching technique. Their density was determinedby Archimedes method with acetone as buoyant liquid. The thermal expansion coefficient of each zinc tellurite glasses wasmeasured using L75D1250 dilatometer, while their glass transition temperature (𝑇

𝑔) was determined by the SETARAM Labsys

DTA/6 differential thermogravimetric analysis at a heating rate of 20Kmin−1. The acoustic Debye temperature and the softeningtemperature (𝑇

𝑠) were estimated based on the longitudinal (𝑉

𝐿) and shear ultrasonic (𝑉

𝑠) wave velocities propagated in each glass

sample. For ultrasonic velocity measurement of the glass sample, MATECMBS 8000 Ultrasonic Data Acquisition Systemwas used.All measurements were taken at 10MHz frequency and at room temperature. All the thermal properties of such binary telluriteglasses were measured as a function of ZnO composition. The composition dependence was discussed in terms of ZnO modifiersthat were expected to change the thermal properties of tellurite glasses. Experimental results show their density, and the thermalexpansion coefficient increases as more ZnO content is added to the tellurite glass network, while their glass transition, Debyetemperature, and the softening temperature decrease due to a change in the coordination number (CN) of the network formingatoms and the destruction of the network structure brought about by the formation of some nonbridging oxygen (NBO) atoms.

1. Introduction

Tellurite glasses are at present the subject of intensive inves-tigations because the glassy phase can be formed over a widerange of concentrations. The application of these types ofglasses in areas of optoelectronics such as laser technologyand fiber optics and other fields is immense due to theirgood physical properties, high density, chemical stability,high homogeneity, and relatively high electrical conductivity[1–4].

Even tellurite glasses and glass ceramics are promisingchoices due to their high refractive index (larger than 2),wideband infrared transmittance (extending up to 6microm-eter), and large third-order nonlinear optical susceptibility.In addition, tellurite glasses combine the attributes of a shortwavelengthUV edge (about 350 nm), good glass stability, rareearth ion solubility, a slow corrosion rate, and relatively lowphonon energy (600–850 cm−1) among oxide glass formers[5]. Furthermore, their low transformation temperatures andabsence of hygroscopic properties limits the application of

phosphate and borate glasses. Based on the information, theuse of tellurite glassesmay bemore advantageous than silicateglasses [6, 7].

Another basic system that has good glass-forming abilityand used by many researchers is the ZnO-TeO

2system.

Tellurium (IV) oxide in combination with ZnO forms stablesglasses [7]. Zinc tellurite glasses are reported to be a suitablehost for optically active rare earth ions because of the wideglass-formation range which is close to the extremum forbinary tellurite glasses [8]. ZnO-TeO

2system was used as a

basis formulticomponent optical glass synthesis and has beenreported as a useful medium for ultralow loss (1 dB 1000m−1)optical fibers for wavelengths in the 3.5–4𝜇m region [9]. Itseems clear from the coverage above that tellurite glasses arestrategically important solid materials.

The above indicate undoubtedly the existence of a prac-tical interest in the zinc-tellurium-containing systems as achoice of compositions for super heavy optical flint glasses.Previous studies showed that the glass formation occurs inthe zinc tellurite system in the region of the eutectic (21mol%

2 Advances in Condensed Matter Physics

Table 1: The density, transition temperature, thermal expansion, acoustic Debye temperature, and softening temperature of zinc telluriteglasses.

Glasssample ZnO-TeO2 (mol%) Density

(g/cm3) (± 0.01)Thermal expansion coefficient

(× 10−6 K−1) (± 0.01)Temperature (K)

Transition Debye SofteningZT0 0–100 4.80 12.40 658 263 857ZT1 10–90 5.09 12.14 654 259 852ZT2 15–85 5.10 12.36 653 259 833ZT3 20–80 5.14 12.51 646 257 802ZT4 25–75 5.19 12.66 647 257 783ZT5 30–70 5.21 12.90 638 252 736ZT6 35–65 5.28 12.54 637 253 722ZT7 40–60 5.29 12.78 633 251 694

ZnO) on the TeO2-rich side of the phase diagram [10]. These

types of glasses are characterized by a high refractive indexwhich increases with TeO

2content [11–13].

Apart from their applications, there is a lack of data onstructural investigations as well as the thermal properties ofthese ZnO-TeO

2glass systems in the literature.Therefore, the

aimof this research is to study the effect of zinc on the thermalproperties of tellurite glass system in order to understand thefundamental origin of such properties.

2. Experimental and Materials

Systematic series of binary zinc tellurite glasses in the form(ZnO)

𝑥(TeO2)1−𝑥

(where 𝑥 = 0 to 0.4 with internal of0.05 mole fraction) have been successfully prepared viamelt quenching technique. The density of the glasses wasdetermined by Archimedes method with acetone as buoyantliquid. The preparation of the tellurite-based glass systemsand related experimental method has been discussed else-where [14–16].

To check the amorphous state, the X-ray diffraction wascarried out for each glass sample by using a computer-controlled X’pert Pro Panalytical set. Both longitudinaland shear ultrasonic velocities were measured in differentcompositions of the glass system by using the MBS8000Ultrasonic Data Acquisition System at 10MHz frequencyand at room temperature. The thermal expansion coefficientwas measured using L75D1250 dilatometer with the rectan-gular parallelepiped 3 × 3 × 6mm3 of each glass samples.The thermal expansion was obtained over a range of 30∘to 210∘C, while the glass transition temperature (𝑇

𝑔) was

determined by the differential thermogravimetric analysis(Setaram instrumentation Labsys DTA/6) at heating rate of20Kmin−1. The accuracy in the measurement of 𝑇

𝑔is ±2∘C.

3. Results and Discussion

Table 1 presents the density, transition temperature, thermalexpansion, acoustic Debye temperature, and softening tem-perature of (ZnO)

𝑥(TeO2)1−𝑥

zinc tellurite glasses.

2𝜃

Inte

nsity

(a.u

.)

10 20 30 40 50

ZT7ZT6ZT5ZT4ZT3ZT2

ZT1

ZT0

Figure 1: The XRD patterns of zinc tellurite and pure tellurite glass.

3.1. XRD Analysis. The XRD patterns of the present glasssamples depicted in Figure 1 were found to show no discreteor continuous sharp peaks but broad halo at around 26∘–30∘, which reflected the characteristic of amorphous glassstructure. This indicates the absence of long-range atomicarrangement and the periodicity of the three-dimensionalnetwork in the glassy materials.

3.2. Density. The density of a glass is an important prop-erty capable of evaluating the compactness. The density isaffected by the structural softening/compactness, change incoordination number, and dimension of interstitial spacesof the glass. The increase in the density (as depicted inTable 1) can be related to two reasons: The first reason isthe replacement of TeO

2by ZnO which has high relative

molecular weight where the molecular weight of TeO2and

ZnO is 159.6 and 81.38, respectively. The second reason maybe due to the transformation of TeO

3to TeO

4where the

formation of TeO3+1

polyhedron has one nonbridging oxygenatom. This increase can be attributed to the zinc ions thatoccupy the interstitial position, and therefore the three-dimensional structure of tellurite glass is not destroyed.Thesebehaviors of the studied glasses are agreed with reported dataelsewhere [4, 7].

Advances in Condensed Matter Physics 3

3.3. Thermal Expansion. Thermal expansion is one of thevery important properties of materials for many technolog-ical and practical applications. The aim of the present workis to characterize glass thermal expansion coefficient, andtheir transition temperatures. There is a strong dependencebetween the glass transition temperature, 𝑇

𝑔, thermal expan-

sion coefficient and the kind of the modifier [17].The thermal expansion coefficient of a material is a

measure of the rate of change in volume and thereforedensity with temperature. Although the thermal expansioncoefficient is actually defined in terms of the volume of thesubstance, this value is somewhat difficult to measure. As aresult, the expansion coefficient for glasses is usually onlydetermined in one direction; that is, the measured value isthe linear thermal expansion coefficient, 𝛼

𝐿:

𝛼𝐿=

Δ𝐿

𝐿𝑥Δ𝑇

, (1)

where 𝐿 is the original length of the sample,Δ𝐿 is the increasein length, andΔ𝑇 is the increase in temperature. Since glassesare usually isotropic materials with relatively small thermalexpansion coefficients, 𝛼V = 3𝛼𝐿 can be used to approximate𝛼V with very little error in calculation.

All reported thermal expansion coefficients for glassesare actually average linear thermal expansion coefficientsover some specified temperature ranges. The particular tem-perature ranges from 0 to 300∘C, 20 to 300∘C, or 25 to300∘C. The data for experimental studies may be reportedfor almost any temperature range. Since most linear thermalexpansion coefficients lie between 1 and 50 × 10−6 K−1,metallurgists, ceramists, and other material scientists usuallyreport values with units of ppmK−1. Traditionally, however,glass technologists used 10−7 K−1 as the basis for reportingthermal expansion coefficients.

An understanding of how the thermal expansion coeffi-cient varies as a function of the glass composition is needed.The linear 𝛼th of any solid material depends strongly on theanharmonic nature of interatomic forces.

Figure 2 shows the plot of thermal expansion coefficientversus chemical composition of binary zinc tellurite glasses atvarious temperatures and summarized in Table 1.

The thermal expansion coefficient indicates the relationbetween the volume of a glass and its temperature. Thisproperty is a strong function of glass composition. In therange between room temperature and 𝑇

𝑔, the expansion

coefficient of the glass was often assumed to be independentof temperature and was defined as 𝛼 = (Δ𝐿/𝐿

0)Δ𝑇. From

Figure 2 it can be seen that the thermal expansion coefficientincreases as the ZnO content added from 0.10 mole fractionto 0.40 mole fraction.

The substitution of ZnO might be due to the change ofthe coordination number of TeO

2from 4 to 3. This change is

associatedwith the creation of non-bridging oxygens (NBOs)that caused the decrease in rigidity. Further substitutionof ZnO, that is, 0.35 mole fraction, decreases the thermalexpansion coefficient which can be supported by the earlierwork [18]. Further increasing the modifier content stabilizesTeO4units with NBO. The decrease of thermal expansion

14

12

10

8

6

4

20 10 20 30 40

14

13

12

11

10

9

8

7

6

𝑦 = 0.0001𝑥2 + 0.0078𝑥 + 12.278

𝑅2 = 0.5521

𝑅2 = 0.9541

ZnO content (mol%)D

ensit

y (g

cm−3)

𝑦 = −0.0003𝑥2 + 0.0221𝑥 + 4.8296

Ther

mal

expa

nsio

n (×10−6

K−1)

Figure 2:The density and linear expansion coefficient of TeO2-ZnO

glasses.

coefficient increased the tightness of the structure. Above 0.35mole fraction, the conversion of TeO

4to TeO

3occurs again

which causes the decrease in rigidity.

3.4. Thermal Stability. Thermal stability is defined as theresistance to permanent change in properties caused solelyby heat. Glass stability is defined in terms of resistance tocrystallization of a glass during heating. Glass stability ismost important during processes involving re-forming of anexisting glass. The glass forming ability automatically leadsto glass stability [19]. Thermal stability is frequently charac-terized by the difference in temperature between the onsetof the glass transformation range (𝑇

𝑔) and the occurrence

of crystallization (𝑇𝑥) for a sample heated at a specified

linear rate. These measurements are routinely carried outusing a differential scanning calorimeter (DSC) or differentialthermal analyzer (DTA).

The exact definitions of 𝑇𝑔and 𝑇

𝑥are subjected to

the preference of the experimenter, as is the choice of theappropriate heating rate used in the study. Typical thermalspectra may contain one or more exothermic peaks due tocrystallization of different phases, but the lowest temperaturepeak is considered in discussing glass stability. Once asignificant number of crystals are formed, subsequent eventsat higher temperatures are not considered important in glassstability.

It is known that the glass transition temperature (𝑇𝑔)

is affected by the alteration of the glass structure, and thestructure of the thermally stable glasses is close-packedstructure. The glass transition temperature 𝑇

𝑔helps to reveal

the close or loosely packed structure of the glass [20], wherethe higher single-bond energy in glass network, the morestable the glass-forming system.

3.5. Glass Transition Temperature. Glass transition temper-ature, 𝑇

𝑔, plays a vital role in understanding the physical

properties of glass [21]. DTA curves for the studied glass


900

800

700

600

500

400

300

200

1000 10 20 30 40

ZnO content (mol%)

900

800

700

600

500

400

300

200

100

𝑦 = −0.0758𝑥2 − 1.4117𝑥 + 863.81

𝑅2 = 0.9787

𝑦 = −0.0068𝑥2 − 0.3747𝑥 + 658.59

𝑅2 = 0.9614

𝑦 = −0.0003𝑥2 − 0.2841𝑥 + 262.79

𝑅2 = 0.9348

Deb

ye te

mpe

ratu

re (K

)

Tran

sitio

n/so

fteni

ng te

mpe

ratu

re (K

)

Figure 3:The acousticDebye, transition and softening temperaturesof ZnO-TeO

2glasses.

samples with different ZnO contents have been obtained todetermine the glass transition temperature (𝑇

𝑔) values. Gen-

erally, glasses with close-packed structure will have thermalstability, while those with loose-packed structure will haveunstability [17]. In the present investigation, all of the glasseshave an endothermic change between 385.21 and 360.52∘C,which attributes to the glass transition temperature, 𝑇

𝑔.

The Zn2+ is incorporated into the glass structure as anetwork modifier, resulting in loose packing of the glassstructure. As a result, a continuous decrease in 𝑇

𝑔with the

increase in network modifier content has been observed(Figure 3). In the present glasses, the decrease in 𝑇

𝑔values

with increase in Zn2+ content contributes to a decreasein thermal stability of the glasses leading to loose-packedstructure as discussed elsewhere [21].

The glass transition reflects a change in the coordinationnumber of the network forming atoms and destruction ofthe network structure brought about by the formation ofsome non-bridging atoms [22]. The decrease in the glasstransition temperature values implies that the number ofbridging oxygen groups decreases. This is mainly due to theaddition of ZnOwhich weakens the bond between each atomsample (increases the number of NBOs atom). The bond iseasier to break and hence the 𝑇

𝑔of the sample decreased.

Furthermore, it also implies a decrease in rigidity of the glassnetwork.

3.6. Acoustic Debye Temperature. Acoustic Debye tempera-ture (𝜃

𝐷) is a characteristic property of a solid lattice related

to its acoustic phonon spectrum [23] where it representsthe temperature at which nearly all modes of vibration in asolid are excited [24]. Debye temperature is a characteristictemperature of glass; any modifier added to the host networkaffects this temperature [25]. Also the increase in the rigidity

of the glass is associated with an increase in the lattice vibra-tions. The observed acoustic Debye temperature, obtainedfrom the ultrasonic velocity data, [13, 24, 25] is

Θ𝐷= (

ℎ

𝑘

)𝑀𝑆(

9𝑍𝑁𝐴

4𝜋𝑉𝑚

)

1/3

, (2)

where𝑀𝑆, the mean velocity, is given by

𝑀𝑆= [(

1

𝑉𝐿

3) + (

2

𝑉𝑠

3)]

−1/3

(3)

ℎ the Planck’s constant, 𝑘 the Boltzmann’s constant, 𝑁𝐴the

Avogadro’s number, and 𝑍 the number of atoms given by

𝑍 = ∑𝑥𝑖𝐿𝑖, (4)

where 𝑥 and 𝐿 are the mole fraction and number of atoms inthe 𝑖th oxide.

Regarding the compositional dependence of the Debyetemperature, it can be seen that Debye temperature decreasesfrom 263K to 251 K as ZnO content increases (Figure 3).It decreases when the ultrasonic velocity decreases. Theobserved decrease in 𝜃

𝐷indicates a monotonic decrease in

the total vibrational energy of the system. This is becauseany of the conceivable vibrational units resulting from thesubstitution will be of lower energy. Also, the observeddecrease in Debye temperature is mainly attributed to changein the number of atoms per unit volume and also the existenceof non-bridging oxygen. It also indicates the loosing packingstructure of the glasses with creation of NBOs as discussedabove. In general the acoustic Debye temperature of thepresent glasses is particularly sensitive with the addition ofZnO content.

3.7. Softening Temperature. Softening temperature (𝑇𝑠) is

another important parameter defined as the temperaturepoint at which viscous flow changes to plastic flow. In actualpractice, it plays a crucial role in determining the temperaturestability of the glass. Softening temperature (𝑇

𝑠) is related to

the ultrasonic velocity of shear waves (𝑉𝑠) by the equation

𝑇𝑠=

𝑉𝑠

2

𝑀

𝐶2𝑍

, (5)

where𝑀 is the effective molecular weight, 𝑍 is the numberof atoms in the chemical formula, and 𝐶 is the constantof proportionality and has the value 507.4 (ms−1 K1/2) foralumina-silicate glasses and is assumed to be the same for theglasses under investigation.

The higher the value of softening temperature of a glass,the greater the stability of its elastic properties [24, 25].Values of softening temperature for ZnO-TeO

2glasses were

calculated and presented in Table 1. Figure 3 shows thaton addition of ZnO to TeO

2, the softening temperature

decreases from 857K to 694K with increasing ZnO content.This shows that the stability of the glasses decreases as thenetwork modifier content increases. This change is what maybe expected from the decrease of elastic moduli.

Advances in Condensed Matter Physics 5

Table 2: Nonlinear regression analysis of the variables (�� = 𝛿��2 + 𝛽�� + 𝛼) for various properties of ZnO–TeO2 glass.

Variables (��) 𝛿 𝛽 𝛼 𝑅2 % change

Density −0.0003 0.033 4.830 0.954 10Thermal expansion coefficient 0.0001 0.008 12.278 0.552 3.06Transition temperature −0.0068 −0.375 385.59 0.961 −6.41Acoustic Debye temperature −0.0006 −0.275 262.76 0.910 −4.56Softening temperature −0.1037 −0.609 861.08 0.978 −19.02

3.8. Regression Analysis. All the current experimental datawere analyzed using Microsoft Excel, by fitting regressioncurves, and the results of the regression coefficients arepresented in Table 2. The regression coefficients obtainedfrom each curves are shown in Figures 2 and 3. In Table 2,�� stands for the variables shown in the first column and �� isthe ZnO concentration. As can be seen in previous figures,for most of the variables a nonlinear polynomial (�� = 𝛿��2 +𝛽�� + 𝛼) gives the best fit.

Except for the softening temperature, the overall resultsfrom Table 2 show that the addition of ZnO with less than40mol% into the tellurite glass system causes small effect (lessthan 10%) on their thermal properties.

4. Conclusion

A number of experimental techniques have been employedto determine a number of important thermal properties ofzinc tellurite glass system. The thermal properties of telluriteglasses such as the linear thermal expansion coefficient, theacoustic Debye temperature, glass transformation tempera-ture,𝑇

𝑔, and softening temperature were studied with respect

to ZnO content. The addition of ZnO content increases thedensities of ZnO-TeO

2glasses due to a change in crosslink

between TeO2chains and coordination number of Te2+

ions. The addition of ZnO on TeO2network also causes

the decreasing values of the acoustic Debye, transition, andsoftening temperatures of ZnO-TeO

2glasses probably due

to the change in Te2+ coordination number. The increase ofZnO in the tellurite glass system results in lower networkrigidity, which in turns results in decrease of most of theirthermal properties. Experimental data shows that the densityand thermal properties are greatly a strong function ofglasses composition. The changes in microstructure glassynetwork can have an effect on the physical as well as thermalcharacteristics of zinc tellurite glass.

Acknowledgment

The authors like to thanks the Universiti Putra Malaysia(UPM) who funded this research project under the ResearchUniversity Grant Scheme (RUGS 2-2012) Project no. 05-02-12-1838RU.

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The Effect of Remelting on the Physical Properties of Borotellurite Glass Doped with Manganese

Syed Putra Hashim Syed Hashim, Haji Abdul Aziz Sidek *, Mohamed Kamari Halimah,

Khamirul Amin Matori, Wan Mohamad Daud Wan Yusof and Mohd Hafiz Mohd Zaid

Glass and Ultrasonics Studies Centre (GUSC), Department of Physics, Faculty of Science, Universiti Putra

Malaysia, 43400 UPM Serdang, Selangor, Malaysia; E-Mails: [email protected] (S.P.H.S.H.);

[email protected] (M.K.H.); [email protected] (K.A.M.);

[email protected] (W.M.D.W.Y.); [email protected] (M.H.M.Z.)

* Author to whom correspondence should be addressed; E-Mail: [email protected];

Tel.: +603-8946-6682; Fax: +603-8943-2508.

Received: 7 October 2012; in revised form: 26 December 2012 / Accepted: 29 December 2012 /

Published: 7 January 2013

Abstract: A systematic set of borotellurite glasses doped with manganese (1–x)

[(B2O3)0.3(TeO2)0.7]-xMnO, with x = 0.1, 0.2, 0.3 and 0.4 mol%, were successfully

synthesized by using a conventional melt and quench-casting technique. In this study, the

remelting effect of the glass samples on their microstructure was investigated through

density measurement and FT-IR spectra and evaluated by XRD techniques. Initial

experimental results from XRD evaluation show that there are two distinct phases of glassy

and crystallite microstructure due to the existence of peaks in the sample. The different

physical behaviors of the studied glasses were closely related to the concentration of

manganese in each phase. FTIR spectra revealed that the addition of manganese oxide

contributes the transformation of TeO4 trigonal bipyramids with bridging oxygen (BO) to

TeO3 trigonal pyramids with non-bridging oxygen (NBO).

Keywords: borotellurite glass; density; bridging oxygen; FTIR spectra

1. Introduction

Tellurite glass is an extremely promising material for laser and nonlinear applications in optics due

to some of its essential characteristic features, such as high density, high refractive index, low phonon

OPEN ACCESS

Int. J. Mol. Sci. 2013, 14 1023

maxima, low melting temperature and excellent transparency in the far infrared region [1,2].

Furthermore, tellurite glass has a low melting point and is nonhygroscopic, which is an advantage

when compared to borate and phosphate glasses. These types of glasses are extremely stable against

devitrification, nontoxic and resistant to moisture for long periods of time [2]. It is widely recognized

that the refractive index, n, and density, ρ, of many common glasses can be varied by changing the

base glass composition [3]. In binary tellurite glasses, the basic structural unit of TeO4 is trigonal

bipyramid (tbp) with a lone pair of electrons, and the structural units permit Te–O–Te bonding for

glass formation [4].

The addition of tellurite to any other glass former or network modifier, such as B2O3, is of scientific

and practical interest and may lead to the formation of interesting structural units that affect the

physical properties of the glass network [5]. As reported earlier, the boron coordination number in the

borate glass changed from three to four as more alkaline content was added into the system where the

network linkage was increased. In contrast, the Te coordination number changed from four to three by

the cleavage of the tellurite glassy matrix [6,7]. In fact, the presence of TeO2 in the matrix of alkali

borate glasses decreases its hygroscopic nature; however, it improves the quality and enhances the IR

transmission [8,9]. The role of alkali, alkaline earth, and transition metal oxides (TMO) in the

borotellurite network is to modify the host structure through the conversion of the structural units

of the borate system from [BO3] to [BO4] and the tellurite network from trigonal bipyramid [TeO4]

to trigonal pyramid [TeO3] [10–13]. The elastic moduli of borotellurite glasses (TeO2–B2O3) have

been reported and discussed based on the bond compression model [10,12].

In this work, borotellurite glasses doped with manganese oxide (MnO) in the form of (1–x)

[(B2O3)0.3(TeO2)0.7]-xMnO, with x = 0.001, 0.002, 0.003 and 0.004, were prepared by using a

conventional melt and quench-casting technique. The main objective of this work is to determine the

optimum concentration of manganese needed to prepare the glass by examining amorphous

characteristics using X-ray diffraction. The effect of remelting each glass sample is also studied.


The XRD patterns for the various compositions of (1–x)[(B2O3)0.3(TeO2)0.7]-xMnO glasses are

shown in Figure 1, and the existence of a peak for MnO concentrations of 0.1 mol% and 0.2 mol%,

which is related to the existence of a crystalline phase in the samples, is shown in Figure 1a. The

addition of the MnO has disturbed the borotellurite glass system. In general, crystal growth can occur

at any temperature if a seed crystal is available. It may establish and enhance crystal growth inside a

system where a detectable growth rate can occur at any temperature below the Tm. These crystalline

peaks correspond to cubic manganese telluride borate, Mn3B7O12.65Te0.85, with the reference number

00-026-1255 [11–13]. At 0.3 mol% and 0.4 mol% MnO, the glass systems are in the amorphous state

due to the optimum concentration of MnO, where it acts as a stabilizer for the glass system. Figure 1a

also shows that no sharp peaks exist at 0.3 mol% and 0.4 mol% of MnO [10–15].

Manganese ions seem to exist in the Mn2+ and Mn3+ states in the glass network. However, at lower

concentration of MnO, a majority of the manganese ions are in the Mn2+ state. The linkage of the Mn2+

and Te4+ ions is expected to be extremely weak because the difference in ionic radii of the Mn2+ (0.8 Å)

and Te4+ (0.84 Å) is high when compared to that of Mn3+ (0.58 Å) and Te3+ (0.52 Å) ions [15].

Int. J. Mol. Sci. 2013, 14 1024

To study the remelting effect on the glass structure, all of the glass samples were then remelted. In

general, the features of the XRD patterns confirm that all of the remelted glasses are in the amorphous

state (Figure 1b), as indicated by the broad hump that occurs at approximately 2θ = 20°–30° for all of

the remelted glass samples.

All of the prepared glasses are free of bubbles, purple in color and of good quality. The density (ρ)

and molar volume (Vm) of the glass samples are shown in Figure 2 and Table 1. The density of the pure

borotellurite increased steadily with the addition of TeO2 into the glass structure, as depicted in Figure 2a.

It can be observed that the density decreases gradually with the compositions for both glasses before

and after remelt with an addition of MnO (see Figure 2b,c).

Figure 1. The XRD pattern of (a) the original sample and (b) the remelted MnO–B2O3–TeO2

glass samples.

0.1%

0.2%

0.3%

0.4%

0.1%

0.2%

0.3%

0.4%

(a) (b)

The density results as depicted in Table 1 show that as the manganese cation concentration

increases, the glass structure becomes more open, allowing for the likely creation of more nonbridging

oxygen (NBO) [15,16]. Additionally, Figure 2 shows that the molar volume increases with the

introduction of manganese content. In the present samples, the glass densities vary from 4.57 to

5.56 gcm−3 and 3.23 to 4.38 gcm−3 after remelt, revealing a rather linear relationship with the

manganese content. However, there are slight differences in the density between the before and after

remelt (Figure 2c) samples due to the existence of a crystalline phase inside the glass system. The

occurrence of crystal growth causes a decrease of NBO. The remelting effect of this glass network is

the reconstruction of the structure of the glass system and an increase of NBO, causing a decrease in

the density.

The composition dependence of the molar volume gives information about the coordination state of

the manganese cations. The density and molar volume for these glasses are compatible with the ionic

size, atomic weight, and amount of different elements in the glasses.

Int. J. Mol. Sci. 2013, 14 1025

Table 1. The glass composition (mol%), density and molar volume of (100–x)

[(B2O3)30(TeO2)70]-xMnO.

MnO B2O3 TeO2 ρ (g/cm3) ρremelt (g/cm3) Vm (cm3) Vm,remelt (cm3)

0.1 29.97 69.93 5.555 4.3773 22.734 27.597 0.2 29.94 69.86 5.370 3.885 23.507 29.668 0.3 29.91 69.79 4.764 3.5798 26.483 30.649 0.4 29.88 69.72 4.569 3.2253 27.604 32.300

Figure 2. The density and molar volume of (a) borotellurite glasses; (b) the original

(100–x) [(B2O3)30(TeO2)70]-xMnO glass samples and (c) the remelted samples.

Int. J. Mol. Sci. 2013, 14 1026

Figure 2. Cont.

The experimental FTIR spectra for the borotellurite glasses doped with manganese (100–x)

[(B2O3)30(TeO2)70]-xMnO, with x = 0.1, 0.2, 0.3 and 0.4 mol%, are presented in Figure 3a,b. The FTIR

spectral bands of the glasses and their assignments are summarized in Table 2. The data were analyzed

following the method proposed by Condrate [17], comparing the experimental data of the glasses with

those of their corresponding crystalline compounds.

The present study shows that the quantitative evolution of these glass structures are greatly

influenced by the MnO concentration. The addition of MnO to the glass matrix leads to a drastic

reduction in intensity between the ~520 and ~650 cm−1 absorption bands due to the Te–O bond

between the trigonal bypiramidal unit [TeO4] and bridging oxygen and also contributes to the specific

vibration of the Mn–O bond [18–20]. If we take into account the Mn–O bond vibrations’ contribution

to the ~520 cm−1 absorption band, it seems that the controlled addition of manganese ions constricts, to

a large degree, the bending motion of different boron-oxygen bonds and gradually increases the

number of Mn–O linkages.

The vibration of the B-O arrangement in the infrared region of 400–1400 cm−1 is more

profound [20,21]. The medium absorption observed at ~1200 cm−1 is attributed to the B–O asymmetric

stretching of the tetrahedral BO4 [18] and orthoborate group [21,22]. The intensity of this band

decreases for the original samples from x = 0.1% to x = 0.4%. As for the remelted sample, the

intensities of these bands remain the same as the concentration of manganese ions is increased.

The band intensity at ~1400 cm−1 is due to the asymmetric stretching of the B–O bond from the [BO3]

trigonal unit in varied borate rings [18,22,23]. The band intensity decreases with the increase in

concentration of manganese ions for the original samples.

The band intensities between ~1600 cm−1 and ~3200 cm−1 are assigned to the bending of O–H and

the asymmetric stretching of O–H, respectively [24,25]. The occurrence of the O–H bond inside the

glass for the x = 0.1% to x = 0.3% original samples corresponds to the existence of a crystal structure

of manganese telluride borate, which is highly soluble and simply reacts with H2O. As the band

Int. J. Mol. Sci. 2013, 14 1027

intensity at ~1200 cm−1 and ~1400 cm−1 decreases, the band intensity at ~1600 cm−1 and ~3200 cm−1

also decreases and eventually disappears for the x = 0.4% original sample. There is no bending of O–H

or stretching of O–H for the remelted sample because no crystal structure exists in the glass matrix, as

affirmed by Figure 1b.

Table 2. Frequencies and their assignments for the FT-IR spectra of (100–x)

[(B2O3)30(TeO2)70]-xMnO.

Peak positions (cm−1) Assignments

~520 Corresponds to the Mn–O bond

~650 The Te–O bond of the trigonal bypiramidal unit [TeO4] with NBO and the

contribution of the specific vibration of the Mn–O bond

~1200 The asymmetric stretching vibration of the B–O bond for the tetrahedral and

orthoborate group

~1400 The asymmetric stretching of the B–O bond from the [BO3] trigonal unit in

diverse borate rings ~1600 The bending of O–H ~3200 The asymmetric stretching of O–H

Figure 3. Selected FT-IR spectra of (a) the original sample; (b) the remelted sample.

% T

3. Experimental Section

A 13 g batch of the (1–x) [(B2O3)0.3(TeO2)0.7]-xMnO glass system, with x = 0.1, 0.2, 0.3 and 0.4 mol%,

was prepared by mixing all of the components together. The mixture was mechanically ground and

homogenized using an agate mortar for 15 min. The mixture was then preheated inside an alumina

crucible in an electrical furnace for half an hour at a temperature of 400 °C. The preheated mixture was

then transferred to the second furnace for one hour at a temperature of 950 °C. To improve

homogeneity, the crucible was constantly shaken inside the furnace. The melt was then poured into a

stainless steel cylindrically shaped split mold, which was preheated at 350 °C before being transferred

Int. J. Mol. Sci. 2013, 14 1028

to an annealing furnace for two hours at 350 °C. After two hours, the furnace was allowed to cool to

room temperature.

The cylindrically shaped samples obtained were then cut using low speed diamond blade to make

parallel fine surfaces of 6 mm thickness. The unused part of the glass was taken and ground in to a fine

powder. The fine powders were then remelted, and the entire procedure above was then repeated to

examine the remelting effect. The entire procedure for the remelted sample preparation was the same,

including the preheating, melting and annealing temperature, so that the conditions of the samples

could be maintained as the conditions of the original samples.

The amorphous nature of the glasses was ascertained from XRD analysis using an X-ray

Diffractometer (PAnalytical (Philips) X’Pert Pro PW 3040/60).

The density of each glass was measured using the Archimedes method with distilled water as the

immersion liquid. A bulk glass was weighed in air (Wair), immersed in distilled water and then

reweighed (Wdw), where the density of the distilled water was 1.00 g cm−3. The relative density is given

as ρs = ρdw (Wair/Wdw) [26].

4. Conclusions

Borotellurite glass doped with manganese oxide was prepared by the melt-quenching technique.

The XRD, density and molar volume of the glasses were discussed. The overall features of the XRD

curves showed that the occurrence of peaks for the 0.1 mol% and 0.2 mol% samples are due to the

existence of crystal seeds inside this glass structure. However, for the 0.3 mol% sample, the XRD

pattern confirms the amorphous nature of the glass. The remelting effect, on the other hand,

reconstructs the glass structure and avoids the nucleation of crystal growth. This result was confirmed

by the XRD patterns showing the amorphous nature of all of the remelted glasses. Density was

observed to decrease with the increase of MnO in the glass and the remelting effect. However, a

slightly different density was shown for both glass systems due to the existence of crystallites inside

the remelted glass network. This effect reconstructed the structure of the glass system and increased

the amount of nonbridging oxygen inside the system, causing the density to decrease after it was

remelted, which was proven by the FT-IR spectral analysis.

Acknowledgments

The researchers gratefully acknowledge the financial support from Universiti Putra Malaysia for

this study through the Research University Grant Scheme (RUGS) no. 91748.

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© 2013 by the authors; licensee MDPI, Basel, Switzerland. This article is an open access article

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Study of the elastic properties of (PbO)x(P2O5)1−x lead phosphate glass using anultrasonic technique

Khamirul Amin Matori a,b,⁎, Mohd Hafiz Mohd Zaid a, Sidek Hj. Abdul Aziz a,Halimah Mohamed Kamari a, Zaidan Abdul Wahab a

a Department of Physics, Faculty of Science, Universiti Putra Malaysia, 43400 UPM Serdang, Selangor, Malaysiab Materials Synthesis and Characterization Laboratory, Institute of Advanced Technology, Universiti Putra Malaysia, 43400 UPM Serdang, Selangor, Malaysia

a b s t r a c ta r t i c l e i n f o

Article history:Received 31 July 2012Received in revised form 18 September 2012Available online 24 November 2012

Keywords:Glasses;Ultrasonic measurement;Elastic properties

Fabrication of a series of binary (PbO)x(P2O5)1−x lead phosphate glasses with various mole fractions (x=0.1to 0.6) was carried out using a conventional melt-quenching method. The glass density was measured byusing Archimedes principle. The ultrasonic wave velocities (Vl and Vt) of the glasses were determined atroom temperature by using a nondestructive test: the digital signal processing technique of the UltrasonicData Acquisition System (Matec 8020, Matec Instruments, USA). The experimental data for the wave veloci-ties and densities were then used to determine the elastic properties in each series of lead phosphate glasssystems: the longitudinal, shear, bulk and Young's moduli; Poisson's ratio; and the Debye temperature.Based on the results obtained, the longitudinal, shear, bulk and Young's moduli of the glasses increasedwith the addition of PbO content. The Poisson's ratio obtained remains almost constant, while the Debye tem-perature shows a continuous decrease with the addition of PbO content.

© 2012 Elsevier B.V. All rights reserved.

1. Introduction

Recently, phosphate glasses have become technologically impor-tant materials, primarily because of their relatively large thermal ex-pansion coefficients, low optical dispersions and low glass transitiontemperatures [1–5]. Phosphate glasses have been extensively investi-gated due to their transparency in a wide spectral range fromUV to IR,which makes them suitable for the fabrication of optical fibers, detec-tion, sensing and laser technologies (laser host glasses) [6]. Phosphateglass is also important for the study of hazardous waste immobiliza-tion [7,8]. Phosphate glasses have unique characteristics and bondlattices, which include a low melting point, a high thermal stability,a high gain density, a low refractive index and a low dispersive power[9–12].

Pure phosphate glass has a very viscous hygroscopic nature [13–15],so many studies were conducted to improve its chemical resistance.Various studies related to the preparation and use of phosphate glasshave been widely carried out in areas such as bioceramics and glasswith metal connectors for semiconductor materials. Phosphate glass isimportant in glass technology because the pure phosphate glass

viscosity is low at the melting point [16], and it is suitable for use as ahost of the networkmodifier ions in the glass matrix. Phosphate glassescontaining rare-earth ions have important applications in optical fibers,sensors, and radiation shield glasses. These glasses also have been usedin optoelectronics technology for the fabrication of solid-state lasers[7,17].

Research in the field of glass and crystal using ultrasonic methodshas been carried out for many years. Anderson et al. [18,19] studiedsilica glass at different pressures, temperatures and frequencies.After that, the study of silica glass using ultrasonic methods was con-tinued by Cantrell et al. [20], and the study of other types of glass,such as borate, phosphate and others, followed [21–24].

In a study of the elastic constants of materials using ultrasonicmethods, the main point to note is the ultrasonic wave propagationvelocity and the density [25]. From the data for the ultrasonic wavepropagation velocity mode, the series elastic constants can be deter-mined. Due to the elastic constants of this second-order differenceof the total energy of the strain, the constant values can be used toexplore the bonding forces between the atoms in the material [26].Typically, when the material undergoes a phase change, the value ofthe elastic constants will also change.

In the study of glass samples, the ultrasonic method is greatlyinfluenced by the density of the sample. The phosphate density islower than the density of the glassmodifier. The addition of phosphateglass modifier in the glass will increase the density of the glass pro-duced. This is proved by the addition of V2O5 by Farley et al. [27],Fe2O3 by Brassington et al. [28], Sm2O3 by Mierzejewski et al. [29],

Journal of Non-Crystalline Solids 361 (2013) 78–81

⁎ Corresponding author at: Department of Physics, Faculty of Science, Universiti PutraMalaysia, 43400 UPM Serdang, Selangor, Malaysia. Tel.: +60 3 89466653; fax: +60 389454454.

E-mail addresses: [email protected] (K.A. Matori),[email protected] (M.H.M. Zaid), [email protected] (S.H.A. Aziz),[email protected] (H.M. Kamari), [email protected](Z.A. Wahab).

0022-3093/$ – see front matter © 2012 Elsevier B.V. All rights reserved.http://dx.doi.org/10.1016/j.jnoncrysol.2012.10.022

Contents lists available at SciVerse ScienceDirect

Journal of Non-Crystalline Solids

j ourna l homepage: www.e lsev ie r .com/ locate / jnoncryso l

and ZnO by Higazy [30] to phosphate glass, which increases the densi-ty. In the present work, PbO has been added to the phosphate glassnetwork to improve its chemical durability as a glass modifier. Theelastic properties of (PbO)x(P2O5)1−x glasses have been discussed bylooking at the structural modifications that take place in the glassnetwork.

2. Experimental

A series of (PbO)x(P2O5)1−x lead phosphate binary glasses (wherex=0.1, 0.2, 0.3, 0.4, 0.5 and 0.6 mol%) were prepared by the conven-tional melt quenching method. The starting materials, phosphoruspentoxide (P2O5) with a purity of 97% and lead oxide (PbO) with apurity of 98% were weighed in appropriate quantities according tothe mol% of the samples. The powdered mixture was placed in a cru-cible and melted in an electrical furnace to obtain a homogenous meltat 1100 °C for 1 h. A special mold was made to obtain samples with acylindrical shape and dimensions of 10 mm×20 mm. The glass meltwas poured into the stainless steel mold. All of these glass sampleswere annealed at 400 °C (below Tg) for 1 h to remove the thermalstrain.

The glass samples were later cut and polished to obtain flat, paral-lel end faces that were suitable for ultrasonic measurements. The den-sity measurement was performed using the Archimedes method withacetone as the buoyant liquid. The room-temperature ultrasonic mea-surements were carried out at 10 MHz using x-cut and y-cut quartztransducers. A pulse superposition technique was employed using anUltrasonic Data Acquisition System (Matec 8020, Matec Instruments,USA). Burnt honey was used as a bonding material between the glasssamples and the transducers. Bymeasuring the thickness of the sample(d), longitudinal (Vl) and transverse (Vt) wave velocities were calculat-ed using the relation V=2d/t. The absolute accuracy in the measure-ment of the velocity is ±5 m s−1, and the relative error is ±0.1%.

In an amorphous solid, the elastic strain produced by a small stresscan be described by two independent elastic constants, C11 and C44.Elastic moduli were calculated using the following standard relations.

Longitudinal modulus C11 ¼ L ¼ ρV l2; ð1Þ

Shear modulus C44 ¼ G ¼ ρV t2; ð2Þ

Bulk modulus K ¼ L– 4=3ð ÞG; ð3Þ

Young’s modulus E ¼ 1þ σð Þ2 G; ð4Þ

Poisson’s ratio σ ¼ L–2Gð Þ=2 L–Gð Þ; ð5Þ

3. Results and discussion

The density of the binary lead phosphate glasses (PbO)x(P2O5)1−x

together with the molar volume, sound velocities (both longitudinaland transverse), the calculated elastic constants (C11 and C44), thebulk modulus (K), Young's modulus (E), Poisson's ratio (σ) and the

Debye temperature obtained from the experimental results are givenin Table 1.

Lead phosphate glasses are interesting systems to study becausethe glass phase can be formed over a wide concentration. Moreover,PbO can enter the glass network both as a network modifier and alsoas a network former [23]. It was suggested that the addition of PbOto phosphate networks results in the formation of P\O\Pb bonds,leading to a dramatic improvement in the chemical durability of thephosphate glasses [31].

The variation of density versus mol% of PbO as shown in Fig. 1,suggests that the addition of PbO to phosphate glass networks causesa nonlinear increase in density. As shown in Fig. 1 the densities of(PbO)x(P2O5)1−x glasses increase with an increase of the PbO content.This change in density by the addition of PbO is related to the changein the atomic mass and the atomic volume of the constituent ele-ments. The atomic masses of the Pb and P atoms are 207.20 and30.87, and their atomic radii are 1.75, and 1.28 Å, respectively. Thisexplains the increase in density with the increase in the PbO content.

There is a nonlinear increase in density up to 40 mol%. For higheramounts of PbO, the increase in the density is highly pronounced.The addition of PbO and the decrease in the P2O5 concentration inthe glass network caused the densities to increase, which indicatesthat the Pb2+ acts as a network modifier, altering the structure ofthe glass by reducing nonbridging oxygens (NBOs) in the network,so that the structure turns out to be more compact. The possible reac-tions in the glass network can be represented as follows:

P2O5≡2 POO3=2

h iand PbO≡ PbO2=2

h i

PbO2=2

h iþ 2 POO3=2

h i↔ PbO4=2

h i2− þ 2 PO4=2

h iþ

The additional oxygen sharing and charge-balance requirementsare met by the conversion of P_O in [POO3/2] units to form P\O in[PO4/2]+ units. It is therefore suggested that the P_O bonds are titrat-ed continuously to incorporate Pb into the network [32]. This processgives rise to the formation of P\O\Pb linkages.

Fig. 2 shows that the molar volume of the glasses increases with in-creases in the PbO content. Both the density and the molar volume of

Table 1The values of density, molar volume, sound velocities, elastic moduli, Poisson's ratio and Debye temperature of (PbO)x(P2O5)1−x glasses.

Glass sample(x)

ρ(kg/m3)

Vm

(cm3/mol)Vl

(m/s)Vt

(m/s)C11

(GPa)C44

(GPa)K(GPa)

E(GPa)

σ ƟD

(K)

0.1 4136 36.3 2706 1576 30.3 10.3 16.6 25.5 0.238 2560.2 4275 37.0 2857 1664 34.9 11.8 19.1 29.4 0.242 2530.3 4470 37.1 2912 1702 38.0 12.9 20.6 32.1 0.243 2510.4 4521 38.6 3011 1765 41.0 14.1 22.2 34.9 0.244 2490.5 4706 38.8 3141 1788 46.4 15.0 26.4 37.9 0.256 2440.6 4856 39.3 3237 1844 50.9 16.5 28.9 41.6 0.261 240

3800

4000

4200

4400

4600

4800

5000

5200

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

PbO mol%

Fig. 1. Variation of density versus PbO mol%.

79K.A. Matori et al. / Journal of Non-Crystalline Solids 361 (2013) 78–81

the glasses increase with an increase in PbO. The density andmolar vol-ume increase by replacing P2O5 by PbO. As shown in Fig. 1 and Fig. 2 thedensity of (PbO)x(P2O5)1−x glasses varies from 4136 to 4856 kg m−3

and that themolar volume varies from 36.3 to 39.3 cm3 mol−1. Gener-ally, the density and the molar volume show opposite behaviors, but inthis study, different results were obtained. In this glass, the substitutionof phosphorus by lead causes an expansion of the network. Similartrends for densities and molar volumes have already been reportedelsewhere for other glass systems [16,24,33,34].

It is clear that by increasing PbO, the molar volume increases,which is similar to the variation density that occurs with increasingPbO content. The Pb ions may enter the glass network interstitially;hence, some network P\O\P bonds are broken and replaced byionic bonds between Pb ions and singly bonded oxygen atoms. There-fore, if one assumed that the only effect of adding Pb cations was tobreak down the network P\O\P bonds, then an increase in themolar volume with PbO content would be expected for the entirevitreous range of the studied glass system. Experimentally, this effectincreases the molar volume, and as a consequence, the values of thedensity are increased. The addition of PbO increased the values ofthe density, which is most likely attributable to simultaneous fillingup of the vacancies in the network by the interstitial Pb ions withan atomic mass of 207.20. This increase in density indicates a struc-tural change in the glass network, which is accompanied by an in-crease in the molar volume [35].

The addition of PbO in glass interstices causes more ions to fill upthe network, thus compacting the glass structure and increasing therigidity of the network. As a consequence, both velocities Vl and Vt

increase with the addition of PbO, as shown in Fig. 3. An increase ofthe ultrasonic velocities with an increase in the PbO concentrationhas been observed, which indicates that PbO plays a dominant rolein the velocities. In this (PbO)x(P2O5)1−x glass system, PbO plays

the role of a network modifier. It will modify the glass structure,thus causing the glass to become harder. Although the glass is harder,this does not mean that the glass is dense.

The independent elastic constants for isotropic solids and glassesare the longitudinal modulus (C11) and the shear modulus (C44). Thecalculation of other elastic constants and Poisson's ratio depends onthe values of the density and on both of the velocities. The sound ve-locities also determine Young's modulus, which is defined as a ratioof the linear stress over the linear strain and is related to the bondstrength. Additionally, the bulk modulus is defined as the change involume when a force is acting upon it in all directions.

Fig. 4 shows the variation of the elastic moduli; C11, C44, K and Eversus mol% of PbO. It can be observed that for every type of glass,there is a similar pattern in the elastic moduli with increases of thePbO content in the composition. The values of the elastic moduli in-crease linearly with increases in the PbO content. In general, the addi-tion of PbO to a phosphate glass network increases the rigidity, thevelocity and hence the elasticity of the glass. The increase in the rigid-ity of the glass contributes to the increase in the velocity and elasticmoduli.

Poisson's ratio obtained from elastic moduli remains almost con-stant and linearly increases from 40 mol% PbO and above, while theDebye temperature obtained from ultrasonic velocities shows a con-tinuous decrease with the addition of PbO, as shown in Figs. 5 and 6.Although the modulus is sensitive to structural changes in glassesbecause it depends on the change in the cross-link density, the obser-vationmade about Poisson's ratio supports the finding that there is noappreciable change in cross-link densities with the addition of PbO.The change in the cross-link density of the glass network is wellunderstood from the variation in Poisson's ratio. In general, a highcross-link density has a Poisson's ratio on the order of 0.1 to 0.2 [29],while a low cross-link density has a Poisson's ratio between 0.3 and0.5. In the present system, Poisson's ratio (Fig. 5) is almost constant

35

36

37

38

39

40

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

Mo

lar

Vo

lum

e (c

m3 /m

ol)

PbO mol%

Fig. 2. Variation of molar volume versus PbO mol%.

1400

1900

2400

2900

3400

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

Vel

oci

ty (

m/s

)

PbO mol%

Longitudinal velocity

Transverse velocity

Fig. 3. Variation of sound velocities versus PbO mol%.

0

10

20

30

40

50

60

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

Ela

stic

mo

du

lus

(GP

a)

PbO mol%

C11

C44

E

K

Fig. 4. Variation of elastic moduli versus PbO mol%.

0.23

0.24

0.25

0.26

0.27

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

Po

isso

n's

rat

io

Pbo mol%

Fig. 5. Variation of Poisson's ratio versus PbO mol%.

80 K.A. Matori et al. / Journal of Non-Crystalline Solids 361 (2013) 78–81

(changes from 0.238 to 0.261) when the PbO content is increased. Theobserved Debye temperature, obtained from the ultrasonic velocitydata, is particularly sensitive to the PbO content (Fig. 6). The observeddecrease in the Debye temperature with the addition of PbO supportsthe finding that the addition of PbO content indicates the compactpacking structure of the glass structure and a reduction in the creationof NBOs, as discussed above.

4. Conclusions

The elastic properties of binary (PbO)x(P2O5)1−x lead phosphateglass systems have been studied to ascertain the role of Pb2+ ions inthese glasses. Based on the results obtained, the density and molarvolume increase with the addition of PbO in the (PbO)x(P2O5)1−x

glass system. The velocities (Vl and Vt) and elastic moduli (C11, C44,K, and E) show gradually increasing trends as PbO is being addedinto the lead phosphate glass network. Poisson's ratio remains almostconstant in the early stage and then increases, while the Debye tem-perature obtained from ultrasonic velocities shows a continuous de-crease with the addition of PbO content. The addition of PbO in theglass network caused the densities to increase according to the atomicmass and atomic radii of the Pb atoms. In addition, the addition of PbOand the decrease of the P2O5 concentration in the glass networkcaused the densities to increase, which indicates that the Pb2+ actsas a network modifier, altering the structure of the glass by reducingthe non-bridging oxygens (NBOs) in the network and causing thestructure to be more compact.

Acknowledgments

The researchers gratefully acknowledge the financial support forthis study from the Malaysian Ministry of Higher Education (MOHE)through the Fundamental Research Grant Scheme (5523753).

References

[1] N.H. Ray, C.J. Lewis, J.N.C. Laycock, W.D. Robinson, Glass Technol. 14 (1973) 50.[2] N.H. Ray, J.N.C. Laycock, W.D. Robinson, Glass Technol. 14 (1973) 55.[3] Y.B. Peng, D.E. Day, Glass Technol. 32 (1991) 166.[4] Y.B. Peng, D.E. Day, Glass Technol. 32 (1991) 200.[5] Y. He, D.E. Day, Glass Technol. 33 (1992) 214.[6] R. Praveena, V. Venkatramu, P. Babu, C.K. Jayasankar, Phys. B 403 (2008) 3527.[7] P. Stoch, M. Ciecinska, J. Therm. Anal. Calorim. 108 (2012) 705.[8] M.I. Ojovan, W.E. Lee, Metall. Mater. Trans. A 42 (2011) 837.[9] M.I. Ojovan, W.E. Lee, J. Non-Cryst. Solids 356 (2010) 2534.

[10] G. Le Saout, Y. Vaills, Y. Luspin, Solid State Commun. 123 (2002) 49.[11] B. Eraiah, S.G. Bhat, Phys. Chem. Solids 68 (2007) 581.[12] P. Jozwiak, J.E. Garbarczyk, Solid State Ionics 176 (2005) 2163.[13] K.V. Shah, V. Sudarsan, M. Gaswami, A. Sarkas, S. Manikandan, R. Kumar, B.I.

Sharma, V.K. Shrikhande, G.P. Kothiyal, Bull. Mater. Sci. 26 (2003) 715.[14] L.D. Burling, PhD thesis, University of Nottingham (2005).[15] K. Suzuya, D.L. Price, C.K. Loong, B.C. Sakas, L.A. Boatner, in: Proceeding of Mate-

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(2009) 1489.[17] H.M. Farok, H.B. Senin, G.A. Saunders, W. Poon, H. Vass, J. Mater. Sci. 29 (1994)

2847.[18] O.L. Anderson, H.E. Bommel, J. Am. Ceram. Soc. 38 (1955) 125.[19] O.L. Anderson, J. Phys. Chem. Solids 27 (1966) 547.[20] J.H. Cantrell, M.A. Breazeale, Phys. Rev. B 17 (1978) 4864.[21] A. Tawansi, I.A. Gohar, D. Holland, N.A. El-Shishtawi, J. Phys. D: Appl. Phys. 21

(1988) 607.[22] H.A.A. Sidek, S.P. Chow, Z.A. Talib, S.A. Halim, Turk. J. Phys. 28 (2004) 67.[23] M. Hamezan, H.A.A. Sidek, A.W. Zaidan, K. Kaida, A.T. Zainal, J. Appl. Sci. 6 (2006)

943.[24] Y.B. Saddeek, J. Alloys Compd. 467 (2009) 14.[25] K.A. Matori, M.H.M. Zaid, H.A.A. Sidek, M.K. Halimah, Z.A. Wahab, M.G.M. Sabri,

Int. J. Phys. Sci. 5 (2010) 2212.[26] M.H.M. Zaid, K.A. Matori, L.C. Wah, H.A.A. Sidek, M.K. Halimah, Z.A. Wahab, B.Z.

Azmi, Int. J. Phys. Sci. 6 (2011) 1404.[27] J.M. Farley, G.A. Saunders, Phys. Status Solidi 28 (1975) 199.[28] M.P. Brassington, A.J. Miller, J. Pelzl, G.A. Saunders, J. Non-Cryst. Solids 44 (1981)

157.[29] A. Mierzejewski, G.A. Saunders, H.A.A. Sidek, B. Bridge, J. Non-Cryst. Solids 104

(1988) 323.[30] A.A. Higazy, B. Bridge, A. Hussein, M.A. Ewaida, J. Acoust. Soc. Am. 86 (1989) 1453.[31] P.Y. Shih, S.W. Yung, T.S. Chin, J. Non-Cryst. Solids 224 (1998) 143.[32] B. Bridge, A.A. Higazy, Phys. Chem. Glasses 27 (1986) 1.[33] S.E. Van Kirk, S.W. Martin, J. Am. Ceram. Soc. 75 (1992) 1028.[34] Y.B. Saddeek, Mater. Chem. Phys. 83 (2004) 222.[35] V. Rajendran, N. Palanivelu, D.K. Modak, B.K. Chaudhuri, Phys. Status Solidi A 180

(2000) 467.

235

240

245

250

255

260

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

Deb

ye t

emp

erat

ure

(K

)

PbO mol%

Fig. 6. Variation of Debye temperature versus PbO mol%.

81K.A. Matori et al. / Journal of Non-Crystalline Solids 361 (2013) 78–81





Article

Structural and Optical Properties of Lead-Boro-Tellurrite Glasses Induced by Gamma-Ray

Iskandar Shahrim Mustafa 1, Halimah Mohamed Kamari 2,*, Wan Mohd Daud Wan Yusoff 2,

Sidek Abdul Aziz 2 and Azhar Abdul Rahman 1

1 School of Physics, Universiti Sains Malaysia, 11800 Minden, Pulau Pinang, Malaysia;

E-Mails: [email protected] (I.S.M.); [email protected] (A.A.R.) 2 Physics Department, Faculty of Science, Universiti Putra Malaysia, 43400 UPM, Serdang,

Selangor, Malaysia; E-Mails: [email protected] (W.M.D.W.Y.);

[email protected] (S.A.A.)


Tel.: +603-89466657; Fax: +603-89454454.

Received: 25 October 2012; in revised form: 9 January 2013 / Accepted: 12 January 2013 /

Published: 4 February 2013

Abstract: Spectrophotometric studies of lead borotellurite glasses were carried out before

and after gamma irradiation exposure. The increasing peak on the TeO4 bi-pyramidal

arrangement and TeO3+1 (or distorted TeO4) is due to augmentation of irradiation dose

which is attributed to an increase in degree of disorder of the amorphous phase. The

structures of lead tellurate contain Pb3TeO6 consisting of TeO3 trigonal pyramid connected

by PbO4 tetragonal forming a three-dimensional network. The decrease of glass rigidity is

due to irradiation process which is supported by the XRD diffractograms results. The

decreasing values of absorption edge indicate that red shift effect occur after irradiation

processes. A shift in the optical absorption edge attributed to an increase of the conjugation

length. The values of optical band gap, Eopt were calculated and found to be dependent on

the glass composition and radiation exposure. Generally, an increase and decrease in

Urbach’s energy can be considered as being due to an increase in defects within

glass network.

Keywords: tellurite glass; optical band gap; Urbach’s energy; irradiation

OPEN ACCESS

Int. J. Mol. Sci. 2013, 14 3202

1. Introduction

Glass in an amorphous (non-crystalline) solid material. Glasses are typically brittle and optically

transparent. The most familiar type of glass, used for centuries in windows and drinking vessels, is

soda-lime glass, composed of about 75% silica (SiO2) plus sodium oxide (Na2O) from soda ash, CaO,

and several minor additives. Some glasses that do not include silica as a major constituent may have

physico-chemical properties useful for their application in fibre optics and other specialized technical

applications. These include fluoride glasses, tellurite glasses, aluminosilicates, phosphate glasses,

borate glasses and chalcogenide glasses. Tellurite glasses contain tellurium oxide (TeO2) as the main

component. Tellurium dioxide is known as a conditional glass former, which it is, needs a modifier in

order to easily form the glassy state. The formation of glass on two glass formers interest both

scientific and practical locale. The structural network will be perturbed and may lead to the formation

of new structural units [1,2]. Glass forming substances are fall into two categories of inorganic

compounds containing bonds which are partially ionic and partially covalent, and, inorganic or organic

compounds which form chain structures with covalent bonds within the chains and van der Waals’

bonds between the chains. Glasses containing heavy metal oxide (HMO) have recently attracted the

attention of several researchers for the excellent infrared transmission compared with conventional

glasses. The γ-irradiation on glasses is found to affect the optical and physical properties [3,4]. Hence,

radiation damage caused by electrons, alpha particles and gamma rays has been thoroughly

investigated [5]. The structural and physical properties of PbO glasses are well described by Worrel

and Henshell [6]. In previous work, Atul et al. [7] have studied borate glasses containing heavy-metal

oxides and shown that it has potential applications in radiation shielding. The objective of the present

work is to study the effect of radiation on the structural and optical properties of lead borotellurite

glass system. To achieve this, a systematic study on optical properties has been performed to

understand the variation of irradiation dose as a function of PbO composition in borotellurite glasses.

In addition, X-ray diffraction patterns and Raman spectra measurements were also performed in order

to support the available data.


2.1. Raman Spectra

Marker labeling of Raman peak is shown in Table 1. Raman spectrum in Figures 1 and 2 were

corrected for baseline and normalized which allows for an effective comparison across a

heterogeneous set of samples. Eventually, the baseline correction utilized the multiple point level

method (Savitzgy-Golay) in which the baseline is leveled at a value that is the average of the baseline

points. Normalization of Raman spectral was performed based on the common normalization method

referring to min/max technique. The min/max (normalization) method is expressed by:

(1)

where I is the intensity after baseline correction was performed, Imin is the minimum intensity and Imax

is the maximum intensity on single spectral measured. Raman spectrum in Figure 1 shows significant

Int. J. Mol. Sci. 2013, 14 3203

peak at <100 cm−1 which indicate the strong presence of Pb and Te in the chemical bonding through

vibrational mode due to addition of PbO and glass network. PbO stands out as unique because of its

dual role [8], one as modifier, if Pb–O bond is ionic and the other as glass former with PbO4 structural

units, if Pb–O bond is covalent. Occasionally, PbO concentration deteriorates glass forming ability [9]

of (TeO2)y[(PbO)x(B2O3)1−x]1−y system. The addition of heavy metal oxide modifiers to pure TeO2

leads to the progressive formation of distorted TeO3+1 polyhedron followed by the creation of regular

trigonal TeO3 pyramids that contain non-bridging oxygen. In all compositions, the appearance of the

low-frequency Boson peak (<200 cm−1) affirms the presence of the glass structure. The increase broad

shoulders at 410 cm−1 indicate that new features to vibrations of one of the partially crystalline phase

of Pb3TeO6. The existence of Pb3TeO6 is confirmed by X-ray analysis. Clearly, the shoulders at

410 cm−1 were getting broader as the content of PbO increased, possibly due to the PbO unique ability.

At low portions of PbO (up to 0.2% mol), it enters the glass network by breaking up the Te–O–Te and

B–O–B bonds and introduces coordinate defects known as dangling bonds along with non-bridging

oxygen ions (Te–O−…Pb2+…−O–Te) which in turn neutralizes the negative charge of non-bridging

oxygens (NBO) by forming TeO3 and BO4 units. Normally, the oxygen of PbO breaks the local

symmetry while Pb2+ ions occupy interstitial positions. As PbO increases (from 0.3% to 0.5% mol),

a considerable portion may be acts as a double bridges between adjacent TeO4 such as

=Te–O–Pb–O–Te= which can formed besides the formation of PbO4 and TeO3 units. Therefore, for

PbO ≥ 0.3% mol, Pb2+ acts as glass forming agent and is incorporated in the glass network in the form

of PbO4 units. Decreasing on (ss) Te–O–Te and B–O–B bending shoulder at approximately 490 cm−1

and 450 cm−1 ascribe that the splitting of Te–O–Te and B–O–B bonds and hence, the bridging

oxygen’s (BOs) are converted into NBOs. The pure B2O3 was known to consist of the boroxol rings by

linking among trigonal-plane BO3 units, but the network structure was altered through the addition of

PbO. Some parts of the boroxol ring of BO3 units were changed into BO4 tetrahedral units [10].

TeO4 trigonal bipyramids is known to be the main structural unit of the network of all tellurite

glasses [11], as well as of the lattices of crystalline TeO2 polymorphs. All stable tellurite glasses are

multi-component and, what is important, cations on non-tellurite components have a coordination

number other than four. The tellurite structural units of two types are always present in

multi-component tellurite glass network, namely, fourfold coordinated Te atoms (TeO4 trigonal

bipyramids, where all O atoms from bridging bonds with the environment) and threefold coordinated

Te atoms (O=TeO2 trigonal pyramids, where one of O atoms is non-bridging, one forms O=Te double

bond and one O atoms form bridging bonds with the environment). The spectral features from

710 cm−1 to 730 cm−1 and 790 cm−1 in Figure 1 correspond to the TeO4 bi-pyramidal arrangement and

the TeO3+1 (or distorted TeO4) and TeO3 trigonal pyramids structures respectively. It can be clearly

observed that the evolution of TeO4 to TeO3+1 and TeO3 units which one of the Te sp3 hybrid orbital is

occupied by a lone pair of electron. This transformation causes increases in the number of

non-bridging oxygen (NBO) atoms. Legitimately, the increasing peak on the TeO3 trigonal pyramids

shows that modification of lattice and interstitial occurs in the system due to addition of PbO

and B2O3.

Int. J. Mol. Sci. 2013, 14 3204

Table 1. Marker labeling Raman peak according to chemical bonding and stretching [12].

Bonding Raman Shift (cm−1)

S—Strong (S) <100cm−1; Pb, Te W—Weak (S) 450 cm−1; B–O–B (b)

B—Bending 410 cm−1; G—Group (W) 490 cm−1; Te–O–Te (ss)

as—Asymmetry stretching (S) 710 cm−1 ; TeO4 bi-pyramidal arrangement ss—Symmetry stretching (S) 730 cm−1; TeO3+1 (distorted TeO4)

(S) 790 cm−1–860 cm−1; Te–O bending vibrations in TeO3

trigonal pyramids and TeO6

Figure 1. Raman spectra (at ambient temperature) of (TeO2)y[(PbO)x(B2O3)1−x]1−y glasses

with different compositions before irradiation.

Figure 2. Raman spectra (at ambient temperature) of (TeO2)y[(PbO)x(B2O3)1−x]1−y glasses

with different irradiation doses at x = 0.5% mol, y = 0.7% mol.

0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

0 200 400 600 800 1000

Nor

mal

ized

Ram

an I

nten

sity

(a.

u.)

Raman Shift (cm-1)

0 Pb

0.1 Pb

0.2 Pb

0.3 Pb

0.4 Pb

0.5 Pb

Boson frequency peak region

0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

0 200 400 600 800 1000

Nor

mal

ized

Ram

an I

nten

sity

(a.

u.)

Raman Shift (cm-1)

0 kGy

5 kGy

10 kGy

20 kGy

25 kGy

Boson frequency peak region

Int. J. Mol. Sci. 2013, 14 3205

The presence peaks at <100 cm−1 in Figure 2 do not indicate any strong changes of intensity due to

variation of gamma irradiation exposure. Eventually, the decrease in broad shoulders at 410 cm−1 also

indicates the existence of new features in the vibrations of the partially crystalline phase of Pb3TeO6.

The existence of Pb3TeO6 is confirmed by X-ray analysis. Clearly, the shoulders at 410 cm−1 were

getting lower as the irradiation dose increased, possibly due to the network compaction. The spectral

features from 710 cm−1 to 730 cm−1 and 790 cm−1 correspond to the TeO4 bi-pyramidal arrangement

and the TeO3+1 (or distorted TeO4) and TeO3 trigonal pyramids structures respectively. It can be

clearly observed that the evolution of TeO3+1 and TeO3 units occurs as the TeO2 concentration

decreases. Legitimately, the decreasing peak on the TeO4 bi-pyramidal arrangement and TeO3+1

(or distorted TeO4) is due to augmentation of irradiation dose which is attributed to an increase in

degree of disorder of the amorphous phase. According to El-Alaily and Mohamed [3], irradiation with

gamma rays are assumed to create displacements, electronic defects and/or breaks in the network bonds,

which allow the structure to relax and fill the relatively large interstices that exist in the interconnected

network of boron and oxygen atoms causing expansion followed by compaction of the volume.

Shelby [13] also suggested that the boron-oxygen bond is more likely to be affected by irradiation.

2.2. XRD

The XRD diffractograms result in Figure 3 shows the partially crystalline precipitation before and

after irradiation exposure for x = 0.5% mol; y = 0.7% mol. All the glass that was prepared proved to fit

the amorphous state. In addition to the enhanced amount of PbO, the glass had been acclimatized to the

partially crystalline phase from the full amorphous phase. It also shows the presence of a hunch for 2θ

around 20°–35°. All these XRD reflections were assigned to the two polymorphic phases; hexagonal

of tellurium (Te, PDF2 No. 00-036-1402) and monoclinic of lead tellurate (Pb3TeO6, PDF2

No. 00-033-0770). The intensity of the XRD reflections indicate that more monoclinic crystals are

reduced in quenched samples as the modifier content increases. The structures of lead tellurate contain

Pb3TeO6 probably consisting of TeO3 trigonal pyramid connected by PbO4 tetragonal, forming a

three-dimensional network. The decrease of glass rigidity is due to irradiation processes which are

supported by the XRD diffractogram results in Figure 3. The XRD result shows the presence of

crystalline precipitation before irradiation process. However, the partially crystalline peaks vanished

due to 5 kGy of gamma irradiation exposure. Eventually, modification of lattice and structural

arrangement along with network compaction occurs in the system due to irradiation exposure of more

than 5 kGy which comprises the presence of crystalline peaks. As the irradiation dose increases (above

20 kGy), the crystalline peaks once again begin to diminish. Obviously, the partially crystalline glass

still proved to fit the amorphous state as the irradiation dose increased.

Int. J. Mol. Sci. 2013, 14 3206

Figure 3. X-ray diffractogram patterns (at ambient temperature) of

(TeO2)y[(PbO)x(B2O3)1−x]1-y glasses (x = 0.5% mol; y = 0.7% mol) with varied gamma

irradiation dose.

2.3. Optical Absorption Spectra

The optical absorption spectra were taken in the ranges of 340 to 550 nm. The optical absorption is

one of the most productive tools to understand the band gap of optical materials. The optical properties

of a solid are governed by the interaction between the solid and the electric field of the electromagnetic

wave. The optical absorption measurements coefficient α() near the fundamental is calculated from

absorbance A, using the following Equation [14]:

2.303 ⁄ (2)

where d is the thickness of the samples. The rapid change in α(ω) against ω is called “the fundamental

absorption edge” and the corresponding energy is defined as “the optical energy gap (Eopt). In the

compound, a typical absorption edge can be broadly ascribed to any of the three processes: (i) residual

below-gap absorption; (ii) Urbach tails; and (iii) interband absorption. In the second process, the

absorption edge depends exponentially on the photon energy according to the Urbach relation. In

crystalline materials the fundamental edge is directly related to the conduction band and valence band,

i.e., direct and indirect band gaps, while in the case of amorphous materials a different type of optical

absorption edge is observed. Figure 4 illustrates the variation of absorption coefficient, α with incident

photon energy at different doses.

Urbach edge analysis is a useful way to parametrically characterize glass optical absorption edge

and potentially distinguished intrinsic contributions to absorbance. Urbach’s absorption edge is formed

in the region of photon energies below the forbidden gap. The interaction between lattice vibrations

and localized states in tail of band gap from the glass samples has a significant effect on the optical

properties. The plot of ln (α) against photon energy, ħω is linear for the absorption region near the

fundamental absorption edge. Thus, it is evaluated that the absorption coefficient near the fundamental

absorption edge is exponentially dependent on the photon energy and obeys the Ubach’s rule. Figure 5

illustrates the dependence of Urbach’s absorption edge with different irradiation of prepared glasses.

10 20 30 40 50 60 70 80

Relative

Counts (a.u.)

2θ

0 kGy

5 kGy

Tellurium, syn, Te

Tellurium, syn, Te

(114)

(114)

(440)

(440)(422)

Lead Tellurate, Pb3TeO6


(422)

25 kGyTellurium, syn, Te

(110)

(114)

(114)

Tellurium, syn, Te

Tellurium, syn, Te

(440)

(440)



20 kGy

10 kGy

Int. J. Mol. Sci. 2013, 14 3207

The absorption edge decreases with the increase of dose. Significantly, the decreasing values of

absorption edge indicate that after irradiation processes, there is red shift effect on the

(TeO2)y[(PbO)x(B2O3)1−x]1−y glasses. The red shift effect is a process when the absorbance band shifts

to longer wavelength and widens due to irradiation. A shift in the absorption edge can be attributed to

an increase of the conjugation length. The number of Te atoms and Pb atoms per conjugation length is

found to increase with increasing dose which create structure defect within the prepared

(TeO2)y[(PbO)x(B2O3)1−x]1−y glass.

Figure 4. Optical absorption edge, α of (TeO2)y[(PbO)x(B2O3)1−x]1−y glasses with

y = x = 0.5% mol; 0.7% mol at various irradiation exposure. The lines represent

the expolation.

Figure 5. The dependence of Urbach’s absorption edge on different irradiation dose for

(TeO2)y[(PbO)x(B2O3)1−x]1−y glasses with y = x = 0.5% mol; 0.7% mol.

The optical band gap energy is determined by using the following Equation [15]

(3)

0

1

2

3

4

5

6

7

8

1.5 2.0 2.5 3.0 3.5

(cm

‐1)

ħω (eV)

0 kGy

5 kGy

10 kGy

20 kGy

25 kGy

1.50

1.70

1.90

2.10

2.30

2.50

2.70

2.90

3.10

0 5 10 15 20 25 30

Absorpotion edge, ħω (eV)

Dose (kGy)

0 % mol Pb

0.1 % mol Pb

0.2 % mol Pb

0.3 % mol Pb

0.4 % mol Pb

0.5 % mol Pb

Int. J. Mol. Sci. 2013, 14 3208

where α is the absorption coefficient, ħω is the incident photon energy, A is a constant and Eopt is the

optical band gap. Values of n are 2 and 1/2 for direct and indirect transitions, respectively. Figure 6

shows the information of indirect band gap (αħω)1/2 against photon energy ħω of

(TeO2)y[(PbO)x(B2O3)1−x]1−y glasses with x = 0.5 mol %; y = 0.7 mol % at various irradiation exposure,

plotted in the absorption region. Indirect energy gap is determined from the linear regions of the plots

as shown in the figures and corresponding values presented in Table 2. The Eopt has been calculated

approximately from the linear region of the arc extrapolating to meet the ħω axis at (αħω)1/2=0.

Figure 6. Plot of (αħω)1/2 against photon energy for indirect band gap of

(TeO2)y[(PbO)x(B2O3)1−x]1−y glasses with y = x = 0.5% mol; 0.7% mol at various irradiation

exposure. The lines represent the pattern.

The variation of indirect optical band gap with mole fraction of PbO content before and after

irradiation is shown in Figure 7. The connected lines do not resemble any significant explanations

rather than to show the decreasing and increasing pattern of the graph. The optical band gap values of

the indirect process before irradiation decreases through the augment of PbO from 0% mol to

0.15% mol. This is due to the increase of the network disorder and consequently the extension of the

localized states within the gap. More likely, PbO enters the glass network by breaking up the Te–O–Te

and B–O–B bonds and introduces coordinate defects known as dangling bonds along with

non-bridging oxygen. Normally, the oxygen of PbO breaks the local symmetry while Pb2+ ions occupy

interstitial positions. PbO content >0.2% mol shows a slight increase before tends to decline slowly

for 0 kGy and increase for irradiated samples with 5 kGy up to 25 kGy (towards 0.5% mol).

Consequently, as PbO increases (from 0.2% to 0.5% mol), modification of lattice and interstitial

occurs in the system with the nearest neighboring atoms and arrangements, such as PbO4 and/or BO4.

A considerable portion may act as double bridges between adjacent TeO4 which can form in addition

to the formation of PbO4 and TeO3 units. Therefore, for PbO ≥ 0.2% mol, Pb2+ acts as glass forming

agent and is incorporated in the glass network in the form of PbO4 units. The optical band gap, Eopt

values for indirect transition decrease with increasing of irradiation dosage as the content of

PbO ≤ 0.2% mol due to increase in degree of disorder of the amorphous phase. Obviously, the Eopt

values for indirect transition increase with increasing of irradiation dosage as the content of

0

1

2

3

4

5

6

1.8 1.9 2.0 2.1 2.2 2.3 2.4 2.5 2.6 2.7 2.8 2.9 3.0 3.1 3.2 3.3 3.4 3.5

(αħω)1/2

(cm

‐1eV)1/2

ħω (eV)

0 kGy

5 kGy

10 kGy

20 kGy

25 kGy

Int. J. Mol. Sci. 2013, 14 3209

PbO > 0.2% mol. It is believed that the increasing of irradiation create displacements, electronic

defects and/or breaks in the network bonds, which allow the structure to relax and fill the relative large

interstices that exist in the interconnected network of boron and oxygen atoms causing expansion

followed by compaction of the volume.

Figure 7. Variation optical band gap, Eopt of (TeO2)y[(PbO)x(B2O3)1−x]1−y glasses for

indirect transition with x = 0%–0.5% mol; y = 0.7% mol at various irradiation exposure.

The lines represent the extrapolation.

In many crystalline and non-crystalline semiconductors, the α(ω) depends exponentially on the ħω.

This exponential dependence, known as the Urbach rule, can be written in the form [16]:

∆ (4)

where B is a constant and ΔE is the width of the band tails of the localized states which also known as

Urbach energy. The value of Urbach energy (ΔE) is calculated by taking the reciprocals of the slopes

of the linear portion of the ln α() against ħω curves in the lower photon energy regions. Figure 8

shows the reciprocal of the slopes of the linear portion from the ln α() against ħω curves in the lower

photon energy regions for (TeO2)y[(PbO)x(B2O3)1−x]1−y glasses with x = 0.5% mol; y = 0.7% mol at

various irradiation exposure. The reciprocal values will be used to calculate the value of Urbach

energy (ΔE) using the Equation 3. The values of (ΔE) of (TeO2)y[(PbO)x(B2O3)1−x]1−y glasses’ various

irradiation exposure with x = 0.0%–0.5% mol; y = 0.7% mol are visualized in Figure 9 and tabulated in

Table 2. Generally, an increase and decrease in Urbach energy can be considered as being due to

defects within the glass network.

0.50

1.00

1.50

2.00

2.50

3.00

0 0.1 0.2 0.3 0.4 0.5 0.6

E opt(eV)

PbO, x (mol %)

0 kGy

5 kGy

10 kGy

20 kGy

25 kGy

Int. J. Mol. Sci. 2013, 14 3210

Figure 8. Optical absorption coefficient of (TeO2)y[(PbO)x(B2O3)1−x]1−y glasses with

y = x = 0.5% mol; 0.7% mol at various irradiation exposure. The line represents the pattern.

Figure 9. Variation Urbach energy, ΔE of (TeO2)y[(PbO)x(B2O3)1−x]1−y glasses for indirect

transition with x = 0%–0.5% mol; y = 0.7% mol at various irradiation exposure. The lines

represent the pattern.

Urbach’s energy is a characteristic energy which determines how rapidly the absorption coefficient

decreases for below band gap energy. Urbach measured the absorption tail for different temperatures

and showed that the Urbach’s energy is approximately kT, the thermal energy. The temperature

dependence of the urbach tail led to the conclusion that the below-bandgap transitions are assisted

transition. The Urbach tail can be caused by mechanisms other than phonon-assisted absorption.

Considering Figure 9, before irradiation took place, the Ubach energy having a non-prominent increase

pattern when the content PbO in the network increase which illustrating the saturation of defects at

high content of PbO. Other than that, it is more likely due to the increase of the disorder and

consequently the more extension of the localized states. For every irradiation dose of between 5 kGy

and 25 kGy, Ubach energy increases as the content of PbO ≤ 0.2% mol suggesting the increasing of

disorder and consequently the further extension of the localized states. Nevertheless, Ubach energy for

every irradiation dose from 5 kGy up to 25 kGy decreases as the content of PbO > 0.2% mol

advocating the possibility of long range order locally arising from the minimum in the number

of defects.

‐1.5

‐1.0

‐0.5

0.0

0.5

1.0

1.5

2.0

2.5

2.6 2.7 2.8 2.9 3.0 3.1 3.2 3.3 3.4

ln α(cm

‐1)

ħω (eV)

0 kGy

5 kGy

10 kGy

20 kGy

25 kGy

0.0

0.5

1.0

1.5

2.0

0 0.1 0.2 0.3 0.4 0.5 0.6

ΔE (eV)

PbO, x (mol %)

0 kGy

5 kGy

10 kGy

20 kGy

25 kGy

Int. J. Mol. Sci. 2013, 14 3211

Table 2. Optical characteristic for (TeO2)y[(PbO)x(B2O3)1−x]1−y glasses, y = 0.7% mol.

PbO, x (mol %)

Dose (kGy)

Direct transition Eopt (eV)

Indirect transition Eopt (eV)

Urbach energy ΔE (eV)

0.00

0 3.22 2.78 0.27 5 3.13 2.48 0.36 10 2.82 2.10 0.65 20 2.76 1.84 0.68 25 2.74 1.80 0.73

0.10

0 3.22 2.78 0.27

5 2.97 2.28 0.43

10 2.82 1.96 0.81

20 2.72 1.70 1.06

25 2.62 1.60 1.09

0.20

0 3.18 2.72 0.28 5 2.98 2.38 0.38 10 2.54 1.16 1.58 20 2.62 1.36 1.17 25 2.66 1.60 0.94

0.30

0 3.16 2.73 0.25 5 2.94 2.30 0.41 10 2.66 1.64 1.23 20 2.68 1.68 1.00 25 2.76 1.80 0.79

0.40

0 3.16 2.66 0.30

5 3.04 2.40 0.35

10 2.74 1.59 1.17

20 2.78 1.62 1.00

25 2.80 1.80 0.87

0.50

0 3.10 2.58 0.28 5 3.00 2.38 0.34 10 2.80 1.82 0.87 20 2.82 1.92 0.70 25 2.92 1.90 0.70

With the increasing of irradiation dosage, the content of PbO ≤ 0.2% mol, the Ubach energy

increases significantly, this suggests the increase in degree of disorder of the amorphous phase. As the

content of PbO increases to be greater than 0.2% mol, the Ubach energy increases with 5 kGy and

10 kGy of irradiation dose and begins to decrease with irradiation dosage of 20 kGy and 25 kGy. It is

believed that the increasing of irradiation creates displacements, electronic defects and/or breaks in the

network bonds, which allow the structure to relax and fill the relative large interstices that exist in the

interconnected network, causing expansion followed by compaction of the volume.

Int. J. Mol. Sci. 2013, 14 3212


The ternary (TeO2)y[(PbO)x(B2O3)1−x]1−y glass system (x = 0.0%–0.50% mol and y = 0.7% mol)

were prepared using a conventional melt-quenching method [17]. All the glass samples arranged were

homogenous, transparent and bubble free. The glasses were prepared by mixing together specific

weights of Tellurium dioxide—TeO2 (Alfa Aesar 99.99%), Lead oxide, Litharge—PbO (99%) and

Boron oxide—B2O3 (Alfa Aesar, 97.5%). Appropriate amounts of TeO2, PbO and B2O3 were weighed

by using an electronic balance having an accuracy of 0.0001 g. The chemicals were then thoroughly

mixed in an Agate pestle mortar for half an hour and poured into an Alumina crucible. The crucible

was transferred to a furnace and heated at 950 °C for 2 h to aid the melting process. When the melting

process was complete, the molten liquid was cast into a stainless steel cylindrical shape mold which

had been preheated at 340 °C for 30 min. The produced glass samples were annealed at the

temperature range 340 °C for 2 h, and then the furnace was turned off for cooling process reaching the

atmospheric temperature. The glass samples were cut using Buhler ISOMET diamond cutter at a

thickness of approximately 2 mm for the required measurements.

The irradiation process had been conducted using 6°Co gamma rays (J.L. Sherperd & Associates,

model 109-68# 3044) with the dose rate of 5.52 kGyh−1 on March 2004. The 6°Co radioisotope

produced two gamma rays of energy 1.17 MeV and 1.33 MeV, often indicated in the machine as the

average energy of 1.25 MeV, is the most widely used gamma source not only for research, but also for

the food preservation, sterilization of medical equipment and pharmaceutical raw materials. The dose

rate at the day of irradiation was calculated using decay equation , where λ is the decay

constant, Dt is the dose rate at time t and D0 is the dose rate calibrated using Fricke dosimeter (Ferrous

Sulphate). The half-life of 6°Co is 5.3 years. The irradiation has been conducted at the building 41,

SINAGAMA at Malaysia Nuclear Agency. Raman spectra were measured using a Raman spectrometer

(RSI 2001 B, Raman system, INC) equipped with a 532 nm solid-state diode green laser. Grams/32,

version 6 software was used to analyze the spectra. All spectra were corrected for baseline; smoothed

and Fourier Transformed (FT). The baseline correction utilized the multiple point level method in

which the baseline is leveled at a value that is the average of the baseline points. A constant correction

factor of 80% of the degree of smoothing parameter was used throughout the data collection. The

Fourier smoothing was accomplished by the peak data, applying a triangular filter function at the

specified cut-off point of 40% and then reverse Fourier transforming the data. Optical absorption

measurements in the wavelength range of 190 to 800 nm with 0.1 nm internal spacing at slow scan

speed were performed at room temperature using SHIMADZU spectrophotometer model UV-1650PC

(absorption within ± 0.1 a.u. and wavelength within ± 2 cm−1). Phase identification of the samples will

be determined by X-ray Powder Diffractometer, Philip X’Pert Pro Holland, using Cu-Kα

monochromatized radiation (λ = 1.5418 mm). The operating generator tension is 40 kV while

generator current is 30 mA. X-Ray diffraction patterns of powders were recorded at room temperature

with a diffraction angle from 2θ = 5°–90° and at a rate of 0.01°/min. The XRD measurements were

held using bulk dimension since the results gained did not shows any significant difference between

bulk dimension and powder dimension.

Int. J. Mol. Sci. 2013, 14 3213

4. Conclusions

The structural and optical properties of lead borotellurite glass system were studied as model

glasses for radiation protection material. They depend not only on the content of Tellurium dioxide as

the network former, but rather on the total fraction of network modifying oxides (PbO, B2O3) in the

glass structure and effect of radiation towards the glass network. The increasing peak on the TeO4

bi-pyramidal arrangement and TeO3+1 (or distorted TeO4) is due to augmentation of irradiation dose

which is attributed to an increase in degree of disorder of the amorphous phase. The existence of

Pb3TeO6 in Raman spectra is confirmed by X-ray analysis. XRD shows the presence of a hunch for 2θ

around 20°–35°. The structures of lead tellurate contain Pb3TeO6 probably consisting of TeO3 trigonal

pyramid connected by PbO4 tetragonal forming a three-dimensional network. The decrease of glass

rigidity is due to irradiation processes which supported by the XRD diffractograms results.

Significantly, the decreasing values of absorption edge indicate that red shift effect occur after the

irradiation processes. Red shift effect is a process when the absorbance band shifts to longer

wavelength and widens due to irradiation. A shift in the absorption edge can be attributed to an

increase of the conjugation length. The values of optical band gap, Eopt were calculated and found to be

dependent on the glass composition and radiation exposure. The increasing of irradiation creates

displacements, electronic defects and/or breaks in the network bonds, which allow the structure to

relax and fill the relative large interstices that exist in the interconnected network causing expansion

followed by compaction of the volume.

Acknowledgments

The authors appreciate the financial support for the work from the Ministry of Higher Education of

Malaysia through RUGS (9199837) and Universiti Sains Malaysia through incentive grant and Short

Term grant (304/PFIZIK/6312046). A special gratitude addressed to Khairul Zaman Haji Dahlan,

Former Director of Radiation Processing Technology Division in Malaysia Nuclear Agency for his

ample support and permission in providing author with the usage of irradiation facilities.

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(http://creativecommons.org/licenses/by/3.0/).

Ultrasonic and optical properties and emission of Er3+/Yb3+ doped leadbismuth-germanate glass affected by Bi+/Bi2+ ions

Hamid-Reza Bahari Poor a,n, Hj.A.Aziz Sidek a, Reza Zamiri b

a Department of Physics, Faculty of Science, Universiti Putra Malaysia, 43400 UPM Serdang, Selangor, Malaysiab Department of Materials Engineering and Ceramic, CICECO, University of Aveiro, Campus Santiago, 3810-193 Aveiro, Portugal

a r t i c l e i n f o

Article history:Received 19 March 2013Received in revised form9 May 2013Accepted 30 May 2013Available online 11 June 2013

Keywords:Er3+/Yb3+ dopingBi-dopingUltrasonic measurementPhotoluminescenceJudd–Ofelt theoryMultiphonon relaxation.

a b s t r a c t

Rare earth doped heavy metal oxide glasses instead of silicates are interesting research area, especially,for their potential application in optoelectronics. In addition, contribution of different bismuth ionizationstates in photoluminescence spectra is still an open question. In this research work, [GeO2]60–[PbO]40−x–[½Bi2O3]x, glass hosts where x¼0, 10, 20, 30, and 40 mol% and 0.5 wt% of Er2O3 and 1.5 wt% of Yb2O3 asdoping agents were studied. The activated heavy metal oxide glass samples were synthesized byconventional melt quenching method. The optical properties were studied by refractive index, UV–visibleabsorption and photoluminescence (PL) measurements, and explained in terms of the Judd–Ofelt theory.The glass was also studied by ultrasonic measurements and showed that the velocity of sound is lower inPb-rich samples. Our results also, showed that emission intensities are higher in host glasses with lowersound velocities, which is attributed mainly to multiphonon relaxation. In addition, variation of PLintensities with increase of bismuth composition was related not only to the variation of Debyetemperature and refractive index; but also, to the increase of Ω6 in Pb-rich samples due to the ligandfield and existence of Bi2+/Bi+ ions in Bi-rich glasses.

& 2013 Elsevier B.V. All rights reserved.

1. Introduction

Rare earth doped glass and ceramics are widely used in opticsand optoelectronics. Among rare earths (RE), erbium is one of themost interesting elements with a wide range of applications inoptoelectronics and telecommunication, especially Er and Er/Ybdoped glasses with broad near-infrared emissions [1] which havepotential applications in tunable lasers, optical amplifiers andoptical fibers, waveguides and devices. Bi-doped glasses are alsointeresting for technology because of its unique emission between1200 nm and 1400 nm which could not be generated by rareearths. Heavy metal oxide glasses, containing bismuth and/or leadin the glass structure are one of the best alternatives for RE dopedhosts. The optical fibers from heavy metal oxide glasses have goodenvironmental stability and medium optical loss in comparisonwith fluorides, making them suitable for photonics applications atshort distances [2].

Among the heavy metal oxide glasses, bismuth-germanate glassesare of growing interest, because of their high density, high linear andnonlinear refractive index [3], high thermal expansion, low glasstransition temperature, and excellent infrared transmission [4] which

make them suitable for applications in non-linear optics, opticalswitching and second harmonic generation (SHG) [5]. In addition,lead-germanate glasses show one of the lowest cut-off opticalphonon energy among oxides (≤700 cm−1). It has been found thatthe upconversion efficiency is strongly influenced by the maximumphonon energy of the host. This indicates low nonradiative relaxationrates to make upconversion easily observable [4]. It is also, inter-esting to note that different ionic states of bismuth ions (Bi+, Bi2+ orBi+5) affect on photoluminescence in glasses containing bismuthbut the origin of this luminescence is still a challenge [6]. Germanteglasses are of particular interest for their capacity to show bismuthemission because they can present two different co-ordinations andcan host large bismuth ions. Therefore, they are better alternativethan silicate glasses for bismuth luminescence which show only oneco-ordination.

To design a glass for specific applications, basic understandingof host material and dopants is necessary. Recent IR and Ramanstudies on (100−x)GeO2–xBi2O3 system showed that bismuth actsas a network-former in low composition of Bi2O3, but has modifierbehavior at higher amounts [7]. According to EXAFS and vibra-tional study on xPbO–(1−x)GeO2 [8], at low lead content, Pb ionsact as a modifier in the germanate network, however, in PbOcompositions higher than 40%, lead ions increasingly play anetwork former role.

IR and ultrasonic studies are powerful techniques to studythe vibrational modes and elastic properties of glasses and are

Contents lists available at SciVerse ScienceDirect

journal homepage: www.elsevier.com/locate/jlumin

Journal of Luminescence

0022-2313/$ - see front matter & 2013 Elsevier B.V. All rights reserved.http://dx.doi.org/10.1016/j.jlumin.2013.05.053

n Corresponding author. Tel.: +60133087610.E-mail addresses: [email protected],

[email protected] (H.-R. Bahari Poor).

Journal of Luminescence 143 (2013) 526–533

motivating researchers to understand the amorphous host [7,9,10].Optical properties of rare earths are modified by the glass host;hence, a study on the glass hosts and dopants is of great interestfor science and technology [11,12]. Several research works, forexample, have been conducted on tellurite glasses and showed anenhancement of elastic properties by increasing of TeO2 as well asmodification of upconversion efficiency of Er doped tellurite glassby increasing of rare earth concentration [13–16]. Multiphononrelaxation and nonradiative loss of upconversion of Er3+/Yb3+

co-doped glass is smaller in glasses with low phonon energy. Mostof the papers in this area are concentrated on optical properties andemission enhancement abilities of co-doped materials specificallyEr3+/Yb3+ co-doped glasses [1,17]. But, studies on the photo-luminescence (PL) of Er3+/Yb3+ co-doped glass in one hand andthe relation of the phonon spectrum to PL spectrum in the otherhand plays an important role, especially for photonic applications.Multiphonon relaxation and nonradiative loss of upconversion ofrare earth doped glass is related to phonon spectrum of glass hostand is drastically smaller in glasses with low phonon energy [18]. Inthis paper refractive index, ultrasonic velocities, FTIR, UV–visibleabsorption and photoluminescence measurements of Er3+/Yb3+

doped GeO2–PbO–Bi2O3 glass were used to probe a range of physicalproperties of the glass and the Debye temperature, refractive index,Judd–Ofelt parameters and bismuth ionization states of glassymedium were linked to rare earth dozed glass photoluminescencespectrum. Data of density, ultrasonic velocities and FTIR spectro-scopies were published previously in [19,20].

2. Experimental work

The conventional melt-quenching method was used to synthe-size the glass samples. Precursors: GeO2, (Alfa Aesar, 99.999%),PbO, (Alfa Aesar, 99.9%), Bi2O3, (Alfa Aesar, 99.975%), Er2O3,(Aldrich, 99.9%), Yb2O3, (Alfa Aesar, 99.998%) were used to synthe-size Er and Er/Yb active lead–bismuth–germanate vitreous samples.Er/Yb active glasses were prepared in the form of [GeO2]60–[PbO]40−x–[½Bi2O3]x, where x¼0, 10, 20, 30, and 40 mol% and all the glasssamples were doped with 0.5 wt% of Er2O3 and 1.5 wt% of Yb2O3.Also, two Yb-free binary glass specimen with molar formula:[GeO2]40–[PbO]60 and [GeO2]40–[½Bi2O3]60, doped with 1 wt% ofEr2O3 were prepared. After mixing and grounding GeO2, PbO, Bi2O3,Er2O3 and Yb2O3 raw materials with above concentrations anddrying the mixture at about 300 1C, the mixture was heated inalumina crucible and air atmosphere at 1100 1C for 1 h. During theprocess, the melt was stirred inside the furnace to achieve opticallyhomogenous samples. Then, the melt was quenched into a pre-heated metal mold to obtain a transparent glassy cylinder with12 mm diameter, and annealed at 420 1C for 1 h. The glassiness andamorphous nature of samples were confirmed by XRD. Then thesamples were cut and optically polished for measurements.

The density of the glasses was determined by Archimedes'method as described in Ref. [21]. Ultrasonic wave velocities at5 MHz frequency were measured by a MATEC MBS-8000 at roomtemperature from 1 cm cylinder polished glasses with possibleerror of about 5 m/s. The absorption spectrum was probed with anOcean-Optics UV–visible absorption setup from 400 nm to 800 nmwith 0.25 nm resolution. Refractive indexes were measured by aCK40 Olympus microscope using the apparent depth method.Photoluminescence (PL) from powder samples was measured bya Perkin-Elmer-LS55 fluorescence spectrometer. The specimenswere excited at low intensity absorption peak, 490 nm, and highintensity absorption peak, 523 nm, and the emission spectrumcollected from 200 nm to 900 nm with 0.5 nm resolution. Excita-tion spectra in the range of 450 nm–600 nm also, were probedin this paper by adjusting emission of the Perkin-Elmer-LS55

fluorescence spectrometer at 376 nm with 0.5 nm resolution.Density, ultrasonic, and FTIR results related to these samples werepublished in Ref. [19,20] previously.

3. Results and discussion

3.1. Density, ultrasonic velocities and elastic moduli

Table 1 presents the measured values of density, molar volumeand ultrasonic wave velocities. By increasing bismuth content, thedensity increases and molar volume shows a slight increase in fullrange and a greater decrease between GPB631 and GPB622, whichindicates that the material network slightly expands by substitu-tion of Pb with Bi, but almost contracts from x¼0.1 to x¼0.2(Fig. 1a and b). This may be due to changes of coordination numberof GeO4 and GeO6 units [19,22].

Longitudinal and shear moduli (L and S respectively) of eachglass sample were derived from measured longitudinal and shearvelocities and density by L¼ρVL

2 and S¼ρVS2 relations [9]. Bulk

and Young's moduli (B and Y respectively) and Poisson's ratio (s)were determined from longitudinal and shear moduli usingstandard relations [9].

Fractal bond connectivity (d) [10] and Debye temperature (ΘD)were calculated using the following equations:

d¼ 4SB

¼ 61−2s1þ s

ð1Þ

ΘD ¼ hKB

3N4π

� �1=3

Vm ð2Þ

Vm ¼ 1=V3L þ 2=V3

S

3

!−1=3

ð3Þ

N¼ ZNA

VMð4Þ

where h, KB, NA, are Plank, Boltzman, and Avogadro's constants,and Vm, VM, Z, and N are mean ultrasonic velocity, molar volume,number of atoms in chemical formula, and volume density ofvibration units, respectively.

Based on Table 1, the longitudinal velocity increases from 3422to 3511 m/s and the shear velocity increases from 1912 to 2094 m/s.Variations of VL and VS depends on bond strength, packing density,coordination number and cross-linking of units [22]. Increase inlongitudinal ultrasonic velocity with x is attributed to the increasein the stretching force constant [23] and implies that the rigidity ofthe glass network has increased. This may be due to the mechanismsuggested previously for the modifier role of Pb and Bi [7]. Increaseof modifiers can increase the coulomb contribution of the modifiersto the lattice energy [7].

Elastic moduli are also given in Table 2. Longitudinal, shear andYoung's moduli increase with increase in bismuth content, yet thebulk modulus remains constant (Table 2 and Fig. 2).

Table 1Measured values of density, molar volume, longitudinal and shear ultrasonicvelocities and mean ultrasonic velocity of (GeO2)0.6(PbO)0.4−x(1/2Bi2O3)x glass.

x Density Molar volume VL VS Vm

g/cm3 (70.021) cm3 (70.08) m/s (72)

0 5.901 26.08 3422 1912 21290.1 5.912 26.20 3445 1963 21820.2 5.986 26.04 3475 2027 22480.3 6.012 26.09 3482 2037 22590.4 6.046 26.11 3511 2094 2318

H.-R. Bahari Poor et al. / Journal of Luminescence 143 (2013) 526–533 527

Increase in Young's modulus values with x is attributed to theincrease in the stretching force constant (F) [23]. This may be dueto the mechanism suggested previously for the modifier role of Biand Pb [7,8,24]. Increase of modifiers (Pb and Bi) can increase thecoulomb contribution of the modifiers to the lattice energy[18,25,26]. Higazy et. al. [17] suggested that variation of shearmodulus is proportional to the bond bending force constant (Fb).Hence, the increase in shear modulus with respect to x canindicate an increase in Fb.

The ring deformation model [23,26] was used to relate theaverage bond bending force and atomic ring size of 3D structure

of A–O bonds (A¼cation, O¼anion) to bulk modulus with thefollowing equation:

B¼ 0:0106� Fbl3:84

ð5Þ

where average bond bending force (Fb) is in Nm−1 and ring size (l)in nm and bulk modulus (B) in GPa. Ring size is supposed as thesmallest closed circle of A–O bonds or the diameter of the atomicrings. In a first approximation, the average bond bending force canbe replaced by average bond stretching force (F) of 3D polycom-ponent glass:

F ¼ ∑iðxnf f Þi∑iðxnf Þi

ð6Þ

According to this theory, Poisson's ratio is related to averagecross-link density:

scal ¼0:28

nc0:25 ð7Þ

nc ¼1η∑iðxncNcÞi ð8Þ

η¼∑iðxNcÞi ð9Þ

where Nc is the number of cations in unit glass formula, and nc isthe number of cross-link per unit cation and is equal to thenumber of bonds minus two.

Ring size values are given in Table 2. All the ring size values areat 0.45270.001 nm and by consideration of 0.2% relative error,ring size values are constant in all samples. It is known that atomicring size decreases by increasing the cross-link density [26] andconstant values of ring size indicate that the cross-linking densityof samples is unchanged.

3.2. Poisson's ratio, fractal bond connectivity and Debye temperature

Compositional variation of experimental values of Poisson'sratio and calculated values from ring deformation model areshown in Fig. 3. Poisson's ratio decreases from 0.27 at x¼0 to

Fig. 1. (a) Density and (b) molar volume of (GeO2)0.6(PbO)0.4−x(1/2Bi2O3)x glassversus bismuth composition.

Table 2Measured and calculated values of longitudinal, shear, bulk and Young's moduli,ring size, Poisson's ratio and fractal bond connectivity and Debye temperature of(GeO2)0.6(PbO)0.4−x(1/2Bi2O3)x glass.

x L S B Y l sexp scal d ΘD

(GPa) (nm) (70.002) (70.001) K(70.6)

0 69.10 21.57 40.33 54.93 0.451 0.273 0.268 2.14 248.40.1 70.16 22.78 39.79 57.39 0.453 0.260 0.262 2.29 255.90.2 72.28 24.59 39.49 61.10 0.453 0.242 0.257 2.49 265.80.3 72.89 24.95 39.63 61.86 0.452 0.240 0.253 2.52 268.50.4 74.53 26.51 39.18 64.90 0.453 0.224 0.249 2.71 277.2

Fig. 2. Elastic moduli of (GeO2)0.6(PbO)0.4−x(1/2Bi2O3)x glass.

H.-R. Bahari Poor et al. / Journal of Luminescence 143 (2013) 526–533528

0.22 at x¼0.4 but shows deviation from decreasing rate fromx¼0.2 to x¼0.3. According to Bridge et al. [27], cross-link densityof two, one and zero are related to Poisson's ratio of 0.15, 0.3 and0.4 respectively. Specifically, the cross-link density increases fromnearly one to two. Previous studies [28] indicate that Poisson's ratiois proportional to cross-link density and also to the ratio of bondbending to bond stretching force constant Fb/F (i.e. bond ionicity).Calculated Poisson's ratio which is related to estimated cross-linkdensity decreases with increasing x, but shows a deviation fromexperimental values. The constant values of ring size are alsoattributed to the invariance of cross-link density. As such, thedecrease of Poisson's ratio may be caused by an increase in ionicitydue to the presence of more polar bonding in the modifiers. Thegreater Poisson's ratio of GPB613 and GPB604 relative to thedecreasing rate may be related to existence of Bi2+/Bi+ ions inBi-rich samples (as discussed in our previous paper based on FTIRresults, [20]) which produce lesser cross-link density in the networkbecause of lower valence of Bi2+/Bi+ ions in comparison to Bi3+.

A useful parameter that provides information about the dimen-sionality of the glass structure is fractal bond connectivity (d)proposed in Eq. (1) [10]. One dimensional chain structure isrepresented by d¼1, 2D layer structure by d¼2, and 3D tetra-hedral co-ordination polyhedral structures by d¼3 [29]. Table 2and Fig. 4 show that fractal bond connectivity increases from 2.14to 2.71 with respect of bismuth content. Hence, the dimensionalityin x¼0 is predominantly 2D, and by increasing of Bi compositiontends to be 3D structure.

The temperature at which nearly all modes of vibrations in asolid are excited is represented by Debye temperature [30]. TheDebye temperature extracted from measured ultrasonic velocitiesincreases from 248 K to 277 K by increasing of bismuth content(Table 2) and its increase implies an increase in the compactness inthe glass structure [30].

3.3. UV–visible absorption

Fig. 5 shows the UV–visible absorption spectrum of the sam-ples. The net amount of absorption band increases in higherbismuth contents, due to the generation of Bi2+ and/or Bi+ colorcenters [20,31]. In addition to the glass absorption edge, a numberof sharp absorption peaks are distinguishable, especially in Pb-richsamples. The most intensive peak occurred in 523 nm and twoother strong peaks in 490 and 649 nm where detected in allsamples except sample GPB604 in which the glass absorption edgecovered most of the peaks. A weak absorption peak can beobserved in 455 nm, especially in high lead content. These peaks,455, 490, 523, 544 and 649 nm are related to Er3+ excitations from

ground state (4I15/2) to 4F3/2+4F5/2, 4F7/2, 2H11/2, 4S3/2 and 4F9/2states, respectively [5].

Three absorption peaks of Pb-rich samples (x¼0, 0.1, and 0.2)are suitable for least square fitting for Judd–Ofelt parameters.According to Judd–Ofelt theory [32,33], the oscillator strength of4f–4f electric dipole transition from (SL)J energy level to (S′L′)J′energy level is given by:

f calðaJ; bJ′Þ ¼8π2mc

3ð2J þ 1Þhe2n2λχEDSEDðaJ; bJ′Þ ð10Þ

Where h is the Plank's constant, m and e are mass and charge ofelectron, c is speed of light, λ is the mean wavelength of theabsorption band, and n is measured refractive index of themedium. χED is the local field correction factor for electric dipoletransition, equal to χED¼n(n2+2)2/9. J and J′ are initial and finaltotal angular momentum quantum number, respectively. SED(aJ,bJ’)is electric dipole line strength, and is obtained by:

SED ¼ e2 ∑t ¼ 2;4;6

Ωt j⟨aJjjUðtÞjjbJ′⟩j2 ð11Þ

Where Ωt are the J–O intensity parameters related to coordinationstatus of the matrix or crystal field and radial integral of 4f and 5delectrons. oaJ||U(t)||bJ’4 are the reduced matrix elements of unittensor operator calculated in the intermediate-coupling approx-imation and are invariant of environment. Matrix elements forEr3+ ion can be obtained from Judd study [32]. Regarding that the

Fig. 3. Dependence of experimental and calculated Poisson's ratio to the composi-tion of (GeO2)0.6(PbO)0.4−x(1/2Bi2O3)x glass.

Fig. 4. Dependence fractal bond connectivity to the composition of(GeO2)0.6(PbO)0.4−x(1/2Bi2O3)x glass.

Fig. 5. UV–visible spectrum of (GeO2)0.6(PbO)0.4−x(1/2Bi2O3)x glass.


matrix elements are characteristics of rare earths, Eq. (11) showsthat more than three absorption peaks are necessary to obtainthree intensity parameters Ωt (t¼2, 4, and 6).

The experimental oscillator strength can be obtained fromintegrated absorption coefficient spectra in the area of absorptionband, and is given by:

f exp ¼mc2

πe2N

ZαðλÞλ2

dλ ð12Þ

α(λ) is absorption coefficient (cm−1) and N is concentration of Er3+

ion (ion/cm−1). J–O parameters can be evaluated from the fitbetween fcal and fexp and are given in Table 3.

By increasing of x, Ω2 increases from 3.67�10−20 to9.38�10−20 cm2 and Ω6 from 3.36�10−20 to 11.97�10−20 cm2,however, Ω4 remains reasonably unchanged. Ω2 is related to theinteraction of RE to the ligands and correlates with covalencybetween rare earth and oxygen ions and local symmetry of theglass host [34,35]. Thus, increase of Ω2 indicates more symmetricbehavior of crystal field and stronger Er–O interaction. In addition,an increase observed in Ω6 is proportional to rigidity of mediumand is controlled by composition and structure of glass host[34,35]. Dependence of J–O parameters to mean ultrasonic velocity(similarly, to the bismuth composition) is shown in Fig. 6. Lineardependence of Ω6 to mean ultrasonic velocity demonstrated inthis figure, confirms that an increment in rigidity of glass hoststrongly affected on Ω6 parameter.

Also, the radiative life time of each state can be calculated fromJ–O parameters [17,36]:

τrad ¼ ∑bJ0AðaJ; bJ′Þ

" #−1¼ ∑

bJ′

64π4e2

3ð2J þ 1Þhλ χEDSEDðaJ; bJ′Þ" #−1

ð13Þ

where A(aJ, bJ′) is the probability of transition between aJ and bJ’.Eq. (13) shows linear dependence of radiative life time to J–O

parameters. Increasing value of Ω2 and Ω6 and almost unchangedvalue of Ω4, in this study, results in an increase in the radiative lifetime with increased x.

3.4. Photoluminescence

Fig. 7 shows the room temperature photoluminescence spec-trum of Er3+/Yb3+ co-doped GeO2–PbO–Bi2O3 glass samplesexcited with 490 nm light. Two intense emission peaks at 547and 647 nm, three weak emission peaks at 279, 392 and 582 nm,and three shoulders at 355, 447 and 525 nm were observed. Themechanism of emission is demonstrated graphically in Fig. 8.Excitation from ground state of Er3+ ion to 4F7/2 level and thennonradiative relaxations to 2H11/2, 4S3/2 and 4F9/2 followed byradiative relaxation to ground state cause emissions at 525, 547and 657 nm wavelengths, respectively. The emission at 657 nmcan be enhanced by cross relaxation (CR1) process that depopu-lates the 4F7/2 and 4I11/2 from different Er3+ ions to their 4F9/2 [37]:

4F7=2 þ 4I11=2-4F9=2 þ 4E9=2 ð14Þ

The existence of emission peaks at 271 and 392 nm indicatethere are transitions from4F9/2 to 4G9/2 and from 4I9/2 to 2K13/2+

2P1/2+

4G5/2 due to excited state absorption (ESA) process. The energyof emission from the long-lived state 4S3/2 of an Er3+ ion can alsotransfer to other Er3+ ion to populate the 4G9/2 level [37]. Theseexcited levels will decay to populate 4G9/2, 4G11/2, 4F5/2+4F3/2 and2G9/2+

2K15/2+2G7/2 states nonradiatively,and then their relaxations

to ground state will yield the emission peaks at 271, 392, and447 nm and a shoulder at 355 nm, respectively (Fig. 7).

Photoluminescence spectrum of samples excited by 523 nm lightis shown in Fig. 9. A very strong peak at 547 nm exists close to theexcitation wavelength, and is generally excited due to the spectralwidth of the excitation source (and therefore is not mentioned inFig. 9, because of accuracy of the spectrum), two intense emissions at376 and 685 nm, two weak emissions at 419 and 623 nm, two veryweak emissions or shoulders at 476 and 653 nm are distinguishablein the spectrum. The 523 nm excitation beam excites ground stateelectrons of Er3+–2H11/2 level, and after decay to long-lived level 4S3/2,the electrons relax to ground state with a very strong emission at547 nm. The energy of this transition can transfer to other Er3+ ions(ET) and with assistance of excited state absorption from theexcitation beam (ESA) and can populate 2P3/2 and 4G11/2 from 4I9/2to 4I13/2, respectively. The population in 4G11/2 can also increase bythe following cross relaxation (CR2) process:

2P3=2 þ 4I11=2-4G11=2 þ 4F9=2 ð15Þ

Table 3Measured and calculated values of longitudinal ultrasonic velocity, energy bandgap, refractive index and Judd–Ofelt intensity parameters of Er3+/Yb3+ co-doped(GeO2)0.6(PbO)0.4−x(1/2Bi2O3)x glass.

x Sample code ΘD (K) Refractive index (70.01) Ω2 Ω4 Ω6

10−20 cm2

0 GPB640 248.4 2.25 3.29 1.34 3.020.1 GPB631 255.9 2.20 6.33 1.9 6.890.2 GPB622 265.8 2.07 9.78 0.96 12.480.3 GPB613 268.5 2.10 – – –

0.4 GPB604 277.2 2.08 – – –

Fig. 6. Dependence of J–O intensity parameters to mean ultrasonic velocity withdifferent bismuth content (x¼0, 0.1, and 0.2).

Fig. 7. Photoluminescence emission spectrum of (GeO2)0.6(PbO)0.4−x(1/2Bi2O3)xglass, excited by 490 nm light.


Relaxation of the populated 4G11/2 to the ground state willproduce strong upconverted emission (UC) at 350–390 nm. 4G11/2

can also decay nonradiatively to 2H9/2, then relax to the groundstate giving the 419 nm emission peak. Another strong emissionoccurred at 685 nm related to the transition 2H9/2-

4I11/2. Exited2H11/2 and 4S3/2 levels can also decay nonradiatively to populate4F9/2 state. Then 4F9/2 relaxes to the ground state giving just ashoulder at 653 nm. The existence of a weak shoulder in 653 nm incontrast to the strong peak at 657 nm when excited by 490 nmlight indicates the process which depopulates 4F9/2 level. Popu-lated 4F7/2 relaxes to 4I11/2 and its energy may excite 4F9/2 torepopulate 2H9/2 state of another Er3+ ion (CR3) or transfer to Yb3+

ion (ET) and retransfer to the Er3+ ion to excite the 4F9/2 to 2H9/2

state. These mechanisms populate 4I11/2 and depopulate 4F9/2 leveland make CR2 process possible. Such a three-ion (Er3+/Yb3+/Er3+)transition was proposed previously by Gouveia-Netoa et al. in a

glass system triply doped by Nd3+/ Yb3+/Tm3+ [38]. The mechan-ism of photoluminescence of rare earth doped samples at 523 nmexcitation is demonstrated graphically in Fig. 10.

Excitation spectra collected from GeO2–PbO and GeO2–Bi2O3

binary glasses, singly and/or doubly doped with Er3+ or Er3+/Yb3+

(abbreviated with −Er and −Er–Yb) at a 370 nm fixed emissionwavelength are demonstrated in Fig. 11. This figure shows anexcitation band in a range from 500 to 550 nm which is responsiblefor the emission at 376 nm of Fig. 9. In the same doping, leadgermanate binary glasses (GPB460-Er and GPB640-Er–Yb) showedhigher intensities in comparison with corresponding glasses. Inaddition, it is observed in Fig. 5 that the excitation intensity at550 nm is clearly modified by Yb3+ co-doping. According to theprevious discussion, relaxation of the populated 4G11/2 to theground state is responsible of 370 nm emission when glass samplesexcited at 523 nm. This state is populated throw CR2 process whichis dependent to CR3 process and consequently to Yb3+ ion dopants.

3.5. Dependence of photoluminescence intensities to glass host

The main emission peaks of the tow excitation (λex¼490 and523 nm) versus bismuth composition parameter, x, are plotted in

Fig. 8. The Er3+ energy levels scheme with the excitation and relaxation transitionsfor excitation energy of 490 nm.

Fig. 9. Photoluminescence emission spectrum of (GeO2)0.6(PbO)0.4−x(1/2Bi2O3)xglass, excited by 523 nm light.

Fig. 10. The Er3+/Yb3+ energy levels scheme with the excitation and relaxationtransitions for excitation energy of 523 nm.

Fig. 11. Photoluminescence excitation spectrum of Er doped and Er/Yb co-dopedgerminate-lead oxide and germinate bisbuth oxide binary glasses with emission at370 nm.


Fig. 12. By increasing the bismuth content of the glass host, all ofthe emission intensities decrease, however an increase is observedbetween GPB622 to GPB613 (the arrow in Fig. 12 drawn to guidethe eye). This decrease is strongly related to nonradiative decaydue to the multiphonon relaxation depending on the maximumphonon energy [39]. The relationship between total life time of anexcited state (τ) and radiative, multiphonon and energy transferdecay rate (WR, WMP and WET respectively) is given by:

τ−1 ¼WR þWMP þWET ð16ÞIn low doping conditions, with a low space distribution of rare

earths, theWET is negligibly small, so the nonradiative decay rate ispredominated by WMP. The WR is related to the enhanced localfield effect and is proportional to linear combination of J–Ointensity parameters (especially Ω6). Therefore, life time is influ-enced mainly byWMP, which is largely affected by the host [39,40].In the case of weak-coupling of rare earth ions, the multiphonondecay rate is given by [36,39,40]:

WMP ¼ C½1þ nðTÞ�pexpð−αΔEÞ ð17ÞWhere n(T)¼[exp(ħω/KT)−1]−1 is the average vibrational quantumnumber, p is effective number of phonons due to the nonradiativedecay (phonon order), ΔE is energy gap to the next level, C is a hostdependence constant, and α¼−ln (ε)/ħωmax is function of ħωmax

and electron phonon coupling constant. WMP and τ−1 increaseexponentially with an increase in ħωmax. This is relevant to thehigher mean ultrasonic velocity or Debye temperature (ΘD) in highbismuth content samples. ΘD is the temperature at which thehighest frequency mode and, therefore, nearly all vibrationalmodes are excited [41]. And so, we can consider KBΘD as a phononcutoff energy or maximum phonon energy. Therefore, in smallerDebye temperature nonradiative decay is unlikely to occur andluminescence intensity will be greater. So, excellent agreement ofphotoluminescence intensity with Debye temperature (reversedependence in Fig. 12 and Fig. 13) implies that the photolumines-cence intensity is governed drastically by nonradiative relaxation.

On the other hand, the increase in the PL intensities betweenGPB622 to GPB613 may be also related to existence of more Bi2+/Bi+

ions in Bi-rich samples. Bi2+/Bi+ ions show a wide range of photo-luminescence emissions by various excitation wavelengths and this,namely white light emission, probably increases the PL intensities ofBi-rich samples [42]. In addition, upconversion efficiency is affectedby Ω6 which is increased by increasing of Er3+ ionicity [40]. Also, thePL efficiency is enhanced for higher refractive index by increasing oflocal field for Er3+ ion [40]. By increasing of x, refractive index

decreases from 2.25 to 2.08 which is related to a decrease in polarityof the medium (Table 3). Therefore variation of PL intensities withincrease of bismuth composition is related to the variation ofmultiphonon decay rate and refractive index in one hand and tothe increase in Ω6 in Pb-rich samples and the existence of Bi2+/Bi+

ions in Bi-rich samples on the other hand.

4. Conclusion

Longitudinal and shear ultrasonic velocities showed the max-imum values of 3422 and 1912 m/s respectively at the highest leadoxide containing sample. This is related to the increase in modifica-tion behavior of bismuth and lead and increase in the stretchingforce constant. Increase of ultrasonic velocities and elastic moduliimply that the rigidity of sample increases and is relevant to anincrease in Debye temperature with increasing bismuth content.

UV–visible absorption experiment showed that the band gapdecreases and as a result NBO concentration increases with increas-ing of bismuth content. Three main absorption peaks were observedat 490, 523 and 649 nm and were attributed to excitation of Er3+ ionfrom ground state to 4F7/2, 2H11/2 and 4F9/2, respectively. Judd–Ofeltintensity parameters were obtained from the absorption bands.Increase in Ω2 and Ω6 by increasing of bismuth content indicated amore covalent Er–O band in a more rigid glass medium. Thephotoluminescence properties of (GeO2)0.6(PbO)0.4−x(1/2Bi2O3)x glasssamples were measured under 490 and 523 nm excitation, and up/down conversion luminescence was observed. The strong emissionpeak at 657 nm arising from 490 nm excitation was correlated totransition from 4F9/2 level to the ground state. Excitation at 523 nmresults in two main emission peaks at 376 nm and 685 nm. Emissionat 376 nm is due to the transition from 4G11/2 level which ispopulated by ESA process to the ground state. The strong emissionpeak at 685 nm is related to the transition of 2H9/2 level to 4I11/2excited state which is populated not only by the ESA process fromexcitation light, but with a cross relaxation of two Er3+ ion (CR3) anda three-ion (Er3+/Yb3+/Er3+) transition process. Emission intensitiesshow larger values in the samples with higher lead content. This canbe attributed to the decrease of nonradiative decay due to smallersound velocity in Pb-rich samples and increase in local electric fielddue to larger refractive index. More study on electron-phononinteraction and comparison of J–O parameters with other matrices[11,12] are recommended for future study.

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Fig. 13. Dependence of Debye temperature to bismuth composition. The dashedlines and arrow is drawn to guide the eye.


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Entropy 2013, 15, 1528-1539; doi:10.3390/e15051528

entropy ISSN 1099-4300

www.mdpi.com/journal/entropy

Article

Characterization and Synthesis of Silver Nanostructures in Rare Earth Activated GeO2-PbO Glass Matrix Using Matrix Adjustment Thermal Reduction Method

Hamid-Reza Bahari 1,2,*, Reza Zamiri 3,*, Hj. A. A. Sidek 1, Azmi Zakaria 1 and

Faisal Rafiq M. Adikan 2

1 Department of Physics, Faculty of Science, University Putra Malaysia, 43400, UPM Serdang,

Selangor, Malaysia; E-Mails: [email protected] (H.A.A.S.);

[email protected] (A.Z.) 2 Department of Electrical Engineering, Faculty of Engineering, University of Malaya,

Kuala Lumpur 50603, Malaysia; E-Mail: [email protected] (F.R.M.A) 3 Department of Materials Engineering and Ceramics, CICECO, University of Aveiro, Campus Santiago,

Aveiro 3810-193, Portugal

* Authors to whom correspondence should be addressed; E-Mails: [email protected] (H.-R.B.);

[email protected] (R.Z.); Tel.: +60-13-308-7610 (H.-R.B.); Tel.: +351-916-759-671 (R.Z.).

Received: 5 February 2013; in revised form: 24 April 2013 / Accepted: 25 April 2013 /

Published: 25 April 2013

Abstract: This paper reports matrix adjustment thermal reduction method to synthesize

silver nanostructures in Er3+/Yb3+ activated GeO2-PbO glass matrix. The GeO2-PbO glass,

the medium of nanoparticle formation, doped with Er2O3, Yb2O3 and AgNO3 was prepared

by a melt quenching method. Annealing of the glass for different times was utilized, not

only due to thermally reduce Ag+ ions to Ag nanostructures, but also to influence the

glassy network. This is because, the glass structural transformation temperature is near to

435 °C and heating at more than this temperature can cause some structural changes in the

glass matrix. According to TEM images, samples that tolerate 450 °C annealing

temperature for one hour show the formation of basil-like silver nanostructures with a

mean length of 54 nm and mean diameter of 13 nm embedded in the glass matrix, whereas

with annealing at 450 °C for 5 to 20 h, silver nanoparticles of about 3–4 nm mean diameter

size are formed. Annealing for 30 h causes silver nanoparticles to aggregate to form larger

particles due to an Oswald ripening process. Observation of the characteristic Ag-NP SPR

band at 400–500 nm in the UV-visible absorption spectra confirms the existence of silver

OPEN ACCESS

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nanoparticles. The SPR band widens to longer wavelengths in one hour annealed samples,

which relates to the existence of nanostructures with different size or fractal shapes. In

addition, an increment in the peak of the SPR band by increasing the duration of annealing

indicates the formation of more nanoparticles. Furthermore, the existence of a peak at 470

cm–1 in the FTIR spectra of annealed samples and its absence in the samples not exposed to

an annealing process suggests that the glass matrix is polymerized by Pb-O chains during

the 450 °C annealing process. This is the main source of different nanostructures because

of the dissimilar stabilizing media. The tighter media cap the particles to form small and

dense nanoparticles but a loose environment leads to the creation of basil-like particles in

the glass matrix.

Keywords: thermal reduction; silver nanoparticles; glass; germanate; PbO

1. Introduction

The unusual physical properties of metal nanostructures embedded in solid matrices show potential

application in optical devices. Fabrication and characterization of metal nanostructures embedded in

glass matrices have been performed by many researchers [1,2]. Photoluminescence can be enhanced by

metal nanostructures dispersed in an active medium through local field enhancement and/or energy

transfer. Surface Plasmon Resonance (SPR) arising from metallic nanoparticles results in giant and

highly localized electric fields around nanoparticles. This intense localized electric field can enhance

the yield of optical transition of nearby rare-earth ions, a phenomenon known as local field

enhancement (LFE) [3]. Metal nanoparticles (NPs) can also enhance the luminescence spectrum by

energy transfer to emitting ions. In addition to isotropic spherical NPs, optically coupled NPs and also

anisotropic NPs with sharp edges by confinement of the local surface electric field at their sharp edges

are fascinating for nanometal enhanced fluorescence (NMEF) studies [4,5]. The synthesis of

directional nanostructures such as nanorods, nanowires and nanobelts of various materials have

attracted material scientists and led to many challenging studies [6]. However, fabrication of

non-spherical metallic nanostructures in glass matrices has had remarkably little success and remains

a challenge.

Various fabrication techniques such as chemical processes, ion implantation, irradiation or thermal

reduction have been employed to produce nanometals inside solid matrices [1,7,8]. In situ fabrication

of metal nanostructures in a glass medium is affected by temperature, and annealing or irradiation

time. In this research, thermal treatment of glass specimens has been performed to mutually adjust the

glass matrix. Hence the thermally reduced silver ions dissolve in the glassy medium, leading to

fabrication of silver nanostructures embedded in Er3+/Yb3+ activated GeO2-PbO glass matrix. In

addition, annealing of the glass at a temperature higher than 435 °C, the glass transition temperature [9],

was used to adjust the glass matrix. This technique represents a novel method to produce promising

(isotropic and anisotropic) nanostructures embedded in a glass matrix in a more controlled fashion. The

UV-Visible absorption, transmission electron microscopy (TEM), X-ray diffraction (XRD) analysis,

energy-dispersive X-ray (EDX) spectroscopy, Fourier transform infra-red (FTIR) spectroscopy, and

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X-ray photoelectron spectroscopy techniques have been employed to study the formation of

nanostructures in the glassy network.

2. Results

2.1. UV-Visible Absorption

Figure 1 shows the UV-visible absorption of NP-Ag/glass composite samples where a number of

sharp absorption peaks are observed in the spectra of all samples. The most intense peaks occur at 490,

523, and 649 nm, which are related to Er3+ excitations from the ground state 4I15/2 to 4F7/2, 2H11/2, and

4F9/2 states, respectively [10]. The characteristic SPR peak of silver nanoparticles at about 410 nm is

overlapped by a glass band edge at about 400 nm and forms some shoulder-like peaks seen in the

absorption spectra. The absorption at about 400 nm for all samples are higher than in the reference

sample, but for the samples prepared by annealing at 450 °C for one hour, it is widened to longer

wavelengths. The intensity of the SPR peak increases gradually with increasing of the duration of

annealing from 5 h to 30 h.

Figure 1. UV-visible absorption spectra for Ag-NP doped Er3+/Yb3+ activated GeO2-PbO

glasses with different annealing timing. The blue spectrum, specified with an arrow, is

related to a specimen annealed at 450 °C for one hour and shows a higher SPR absorption

at longer wavelengths. The SPR peak at about 400–500 nm also increases with increasing

annealing time.

2.2. TEM

Figures 2 and 3 combine TEM images of nanoparticle-doped glass samples and apparently shows

the formation of nanostructures embedded in the glass matrix. In Figure 2 basil-like nanostructures

with a mean length and width of 54 nm and 13 nm, respectively, are shown, with neither specific shape

nor dense appearance. The TEM picture reveals that the nanobasils consist of many rod-like

nanostructures which are joined each other to form a wider basil-like nanostructure. Figure 3a–c shows

the nanoparticle formation in the glass matrix by annealing at 450 °C for 5, 10, and 20 h, respectively.

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Figure 2. TEM images for Ag-NP doped RE activated GeO2-PbO glasses annealed at 450 °C

for one hour. basil-like silver nanostructures is shown in the figure. More details from this

sample are shown in the intercept.

Figure 3. TEM images for Ag-NP doped RE activated GeO2-PbO glasses with (a) 5 hour,

(b) 10 h, (c) 20 h and (d) 30 h annealing time. The Ag nanoparticles embedded in glass

matrix are shown in the figure. Increasing of annealing duration causes aggregation of small

nanoparticles (mean diameter of 3.6 nm) turn to larger particles (5.6 nm) as seen on Figure 3d.

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Figure 3. Cont.

These figures clearly represent pieces of glass-embedded 3–4 nm size spherical silver nanoparticles

in their matrix. Particle size distribution and Lorentzian fit for Figure 3c are depicted in Figure 4a. This

figure shows that particles are uniformly distributed around a mean diameter of 3.6 nm. Based on

Figure 4b, the mean size of nanoparticles does not change considerably with an increase of annealing

time up to 30 h. Moreover, when the annealing time increases to 30 h, the silver nanoparticles

aggregates turn into bigger, but slightly deformed ones, as shown in Figure 3d and Figure 4b.

Figure 4. (a) The size distribution of NP particles and Lorentzian fit for TEM image of

Figure 3 b,c. The relation between NP’s mean size and annealing duration. The error bar

represents standard deviation of size measurement.

2.3. XRD

Figure 5 shows X-ray diffraction patterns collected from annealed samples. In addition to the

amorphous background, sharp peaks at 38.3°, 44.6° and 64.8° were detected in the XRD patterns

which show crystallizations in the glass matrix. All peaks were indices of a cubic silver single phase

which is in good accordance with XRD reference number (JCPDS 1-87-720). Comparison between

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two patterns related to minimum and maximum annealing time in Figure 4 reveal that nanostructures

after one hour annealing growth more in the [111] direction due to the higher intensity of the peak at

38.3° in comparison with the 450 °C-30 h sample. In addition, no evidence was obtained from XRD

measurements to show the existence of Ag2O (JCPDS 76-1393) or AgO (JCPDS 84-1108)

crystalline phases.

Figure 5. XRD spectra of Ag-NP doped Er3+/Yb3+ activated GeO2-PbO glass for samples

annealed for 1 and 30 h. The sharp peaks were indices according to the crystalline

directions of silver.

2.4. EDX

Figure 6 shows energy-dispersive X-ray (EDX) spectra collected from Ag-NP doped Er3+/Yb3+

activated GeO2-PbO glass.

Figure 6. EDX spectra of Ag-NP doped Er3+/Yb3+ activated GeO2-PbO glass.

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Figure 6. Cont.

These spectra show the elements participate in the glass material and are summarized in Table 1 in

terms of weight percent. According to EDX data, contribution of all materials which make up the glass

matrix (Ge, Pb, and O) and the dopants (Ag, Er, and Yb) is confirmed.

Table 1. EDX data of Ag-NP doped Er3+/Yb3+ activated GeO2-PbO glass. All values are in

weight percent.

Sample C O Ge Ag Er Yb Pb Total

Ref 7.94 19.49 26.53 0.34 0.36 1.17 44.17 100.00

450-1h 6.83 20.79 25.36 0.35 0.24 1.47 44.96 100.00

450-5h 7.09 14.56 30.68 0.32 0.45 1.71 45.19 100.00

450-10h 6.96 20.26 25.09 0.27 0.39 1.51 45.51 100.00 450-20h 6.6 21.02 29.72 0.34 0.32 1.46 40.54 100.00

450-30h 7.2 21.67 23.94 0.34 0.46 1.27 45.12 100.00

2.5. FTIR

The FTIR absorption spectra of silver nanoparticles actively doped in heavy metal oxide germanate

glasses is shown in Figure 7. A wide and intense band at 650–850 cm–1 is deconvulated into two peaks

at 710 and 770 cm–1. These peaks are related to the stretching vibration mode of Ge-O-Ge bonds in

GeO6 octahedral units [11] and antisymmetric stretching vibrations of Ge-O-Ge bonds in GeO4

tetrahedral units [12], respectively. This shows the existence of both of four- and six-coordination. The

absorption at 500–600 cm–1 is attributed to the antisymmetrical bending vibration of a Pb-O-Pb chain [13]

and also related to bending, and the symmetric stretching vibration of Ge-O-Ge in the GeO4 tetrahedral

unit [12]. The weak peak observed at 470 cm–1 is related to the symmetrical bending vibration of Pb-O

in a PbO4 tetragonal pyramid [14]. In the germanate-lead oxide system lead ions have a modifier role

Entropy 2013, 15

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at low PbO content, but participate in the glass network as PbO chains at higher compositions [15].

The peak in 470 cm–1 is related to a PbO covalent bond which is formed when Pb acts the former role

in the glass matrix. This peak was not observed in a reference sample, which means that the Pb-O chains

did not originally exist in the glass and form in the glass matrix because of the applied heat treatment.

Figure 7. FTIR spectra of Ag-NP doped Er3+/Yb3+ active GeO2-PbO glass. The arrow

in the picture shows the weak peak related to a PbO covalent band that arises after

the annealing process.

2.6. XPS

X-ray photoelectron spectroscopy was used as a powerful technique to study the chemical

surrounding of Pb atoms and their role in the glass network. XPS collective spectra of Pb 4f5/2 and Pb

4f7/2 spin-orbit doublet of selective lead oxide germanate glasses actively doped with silver

nanoparticles are shown in Figure 8. Binding energy of Pb 4f5/2 and Pb 4f7/2 in all selected samples are

about 143.0 and 138.1 eV, respectively. These binding energies are in good agreement with data

previously reported for PbO [16]. Two Pb 4f peaks of the sample annealed for 10 h are symmetric and

are fitted with two single Gaussian-Lorentzian curves. The other samples show a clear asymmetry in

their Pb 4f5/2 and Pb 4f7/2 spectra and are deconvoluted into two peaks for each spin-orbit doublet (the

sample annealed for 1 hour shows more asymmetric behavior). The appearance of these new peaks at

quite higher binding energies (144.8 and 140.1 for Pb 4f5/2 and Pb 4f7/2 respectively) is related to the

Entropy 2013, 15

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existence of a new chemical state due to a change in the Pb-O bonding nature from covalent to ionic

form [17]. A change in the role of cation from former to modifier is typically associated with the

formation of an ionic bond with a transfer of electrons from the cations to neighboring oxygens. Owing

to the inverse relation of binding energy to local electron density at the site of atoms, the appearance of

new peaks at higher energies in samples annealed for 1 and 30 h, are due to the formation of Pb-O

ionic bonds and the modifier role of lead in the glass network. In addition, the sample annealed for 10

h shows higher intensities of both Pb 4f peaks relative to other samples. This reveals a more covalent

nature of the Pb-O bond in this sample in comparison with the others, because of higher average Pb

electron density. The XPS results clearly show that ionic Pb-O bonding that appeared in the glass

network during 1 hour of annealing, is transformed into covalent Pb-O chains after heat treatment for

10 h. Some Pb-O chains then, are broken into ionic bonds after 30 h of heat treatment, but to a lesser

extent than the sample exposed to 1 hour of annealing duration.

Figure 8. Experimental XPS spectra of Pb 4f transition and the deconvoluted curves of

Ag-NP doped Er3+/Yb3+ active GeO2-PbO glass annealed for 1, 10, and 30 h from bottom

to top.

Entropy 2013, 15

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3. Discussion

One of the most well-known modifiers in glass science, silver, dissolves in the glass matrix in ionic

form (Ag+) when the melt of glass forming material cools down rapidly to form a rigid glassy

specimen [13]. These ions can be reduced automatically by heat treatment because the reduction

potential (E0) of the Ag+/Ag0 redox system is positive (E0 = 0.8 mV) [1]. During the annealing process,

the thermally reduced silver atoms collect with each other to form nanoparticle seeds. Continued

exposure of the samples at 450 °C leads not only to the seeding process, but also to the growth of

nanoparticles based on the existing seeds from additional reduced silver atoms [6]. When the glass is

annealed for just one hour and cooled down slowly in the furnace, the silver atoms do not have enough

time to create a large number of seeds, but the growth process still persists even during cooling down

to create larger particles. Moreover, our FTIR results showed that Pb acts as a modifier of the glass

network before the annealing process took place. According to the XPS data, 1 hour of heat treatment

causes the formation of Pb-O ionic bonds and opens the glass network, so silver atoms can move more

freely in a glass with an open structure to form more random nanostructures. This is in reasonable

agreement with the UV-visible absorption spectra and TEM images. The weak and uniformly wide

SPR absorption of the sample annealed for one hour illustrates perfectly the smaller number of

particles of various sizes and shapes. The large, uncompressed, basil-like nanostructures observed in

the TEM images are well matched with previous results and hence support this evocative mechanism.

Based on the XPS and FTIR data, the Pb-O chains form to recover the glass network during the

annealing process. An increase of the duration of annealing causes silver atoms to form more

nanoparticle seeds in an environment that is tightened by the formation of Pb-O chains. This leads to

the formation of several small nanoparticles which are capped by the glass network to control the NPs’

size. Holding the samples at 450 °C for a longer annealing time causes an increase in the number of

nanoparticles with almost the same size, but the highest annealing duration results in aggregation of

small NPs to form slightly larger nanoparticles. This mechanism can describe the increase of SPR

absorption by increasing of annealing time which indicates an increase in the number of nanoparticles.

Slightly larger particles observed in the sample annealed for 30 h may lead to a slight red-shift of the

SPR peak which may make the SPR band easy to observe, but it cannot be certainly proven by its

absorption spectrum. Enhancement of the seeding process in a tight amorphous glassy environment by

increasing of annealing time, results in uniformly distributed cultivation of small NPs embedded in the

glass matrix. Holding the samples in 450 °C for 30 h not only enhances NP cultivation, but also gives

the small nanoparticles the opportunity to aggregate into larger particles and breakdown the glassy

network, which is confirmed by the photoelectron data.


In order to study the silver NP-glass material, one set prepared by a melt and quenching method

with different annealing processes. A sample with a composition of (GeO2)0.5(PbO)0.5 triply doped

with 0.5, 1.5, 1 wt% of Er2O3, Yb2O3 and AgNO3, respectively, was prepared by the melt and

quenching method. The sample was divided into several pieces that were separately annealed in

450 °C for 1, 5, 10, 20 and 30 h which is higher than glass transition temperature of this glass

Entropy 2013, 15

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(435 °C), and cooled down inside the furnace to room temperature. Then, they were finely grounded for

TEM, FTIR, EDX, XPS and UV-Vis absorption (reflection setup) characterization.

The glass samples’ crystallinity or amorphous nature was investigated by an X’Pert Pro PW 3040

MPD X-ray powder diffraction instrument (PANalytical B.V., Almelo, The Netherlands). The

absorption spectrum was probed with an UV-3600, UV-VIS-NIR spectrometer (Shimadzu Scientific

instrument Inc., Columbia, MD, USA) from 300 nm to 1,200 nm using a reflection setup. FTIR results

were extracted from a Spectrum-100, FT-IR spectrometer (Perkin-Elmer Inc., Waltham, MA, USA)

using a UATR accessory as the sampling method. The spectra were taken from 400 cm−1 to 1,000 cm−1

and deconvoluted into Gaussian component bands. Material compositions were probed by energy-

dispersive x-ray spectroscopy. The EDX data were collected using a JSM-6400 electron scanning

microscope (JEOL Ltd., Tokyo, Japan) working at an accelerating voltage of 20 kV. Morphological

evaluation and measurement of size and size distribution was performed with an H-7100 Transmission

Electron Microscope (Hitachi, Chula Vista, CA, USA) operating at an accelerating voltage of 120 kV.

For TEM experiments the powder sample was dissolved in isopropyl alcohol and a drop of the solution

was deposited onto carbon coated copper grids and left to air dry for one day at ambient temperature.

The photoelectron and binding energies study was performed using an AXIS-Ultra XPS instrument

(Kratos Analytical Ltd., Manchester, UK). The glass samples were measured in an ultra-high vacuum

chamber (4.5 × 10−10 Torr) and using a monochromatic Al Kα X-ray source. The binding energy data

were calibrated by referencing the main peak of C 1s peak to 284.6 eV.

5. Conclusions

Matrix adjustment thermal reduction fabrication method was introduced as a cheap, easy and swift

technique to fabricate various silver nanostructures in GeO2-PbO glass matrices. Annealing the glass

sample for one hour provided a loose environment to thermally reduce the Ag+ ions to form basil-like

nanostructures with a mean length of 54 nm and mean width of 13 nm according to TEM images. Pb-O

chains didn’t appear in the FTIR results for the glass before annealing, but arose in annealed samples.

Also, Pb-O ionic bonds appear in the XPS spectrum of the Pb 4f transition for sample annealed for a

short time and disappear in the samples exposed to longer annealing. This provides an open

environment for fractal nanostructure formation after a short time of annealing. Longer annealing

times resulted in formation of 3–4 nm size nanoparticles in a tight glass network that acted as a

capping agent. The SPR band of NP silver showed higher values in the extinction spectra against the

increase in annealing time, which reveals the formation of more Ag nanoparticles. Holding the glass

specimen at 450 °C for 30 h caused aggregation of smaller NPs to form larger but deformed

nanoparticles and some loosening in the glass network which was confirmed by the appearance of Pb-

O ionic bonds in the XPS data.

Acknowledgments

The financial support from Universiti Putra Malaysia (UPM) and the Institute of Bioscience (IBS)

Zahidah Muhamed for TEM and EDX services, and Center for Research Instrumentation and

Management (CRIM) for XPS services are gratefully acknowledged.

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© 2013 by the authors; licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution license (http://creativecommons.org/licenses/by/3.0/).





Article

Preparation and Elastic Moduli of Germanate Glass Containing

Lead and Bismuth

Hj A. A. Sidek *, Hamid R. Bahari, Mohamed K. Halimah and Wan M. M. Yunus

Department of Physics, Faculty of Science, Universiti Putra Malaysia, 43400 UPM Serdang,

Selangor, Malaysia; E-Mails: [email protected] (H.R.B.);

[email protected] (M.K.H.); [email protected] (W.M.M.Y.)


Tel.: +603-89466682; Fax: +603-89432508.

Received: 31 January 2012; in revised form: 5 March 2012 / Accepted: 30 March 2012 /

Published: 11 April 2012

Abstract: This paper reports the rapid melt quenching technique preparation for

the new family of bismuth-lead germanate glass (BPG) systems in the form of

(GeO2)60–(PbO)40−x–(½Bi2O3)x where x = 0 to 40 mol%. Their densities with respect of Bi2O3

concentration were determined using Archimedes’ method with acetone as a floatation

medium. The current experimental data are compared with those of bismuth lead borate

(B2O3)20–(PbO)80−x–(Bi2O3)x. The elastic properties of BPG were studied using the ultrasonic

pulse-echo technique where both longitudinal and transverse sound wave velocities have

been measured in each glass samples at a frequency of 15 MHz and at room temperature.

Experimental data shows that all the physical parameters of BPG including density and

molar volume, both longitudinal and transverse velocities increase linearly with increasing

of Bi2O3 content in the germanate glass network. Their elastic moduli such as longitudinal,

shear and Young’s also increase linearly with addition of Bi2O3 but the bulk modulus did

not. The Poisson’s ratio and fractal dimensionality are also found to vary linearly with the

Bi2O3 concentration.

Keywords: glasses; bismuth; lead; germanate; elastic moduli

OPEN ACCESS

Int. J. Mol. Sci. 2012, 13 4633

1. Introduction

Bismuth germanate-based glasses doped with rare-earth oxides gain much attention from

researchers due to their potential applications in non-linear optics devices such as optical

communication fibers, solid-state lasers, light converters, waveguide, sensors and scintillation detector

in positron camera [1–3]. Some optical properties of bismuth in borate, phosphate and silicate glasses

were also extensively studied [4].

The lead germanate glasses are also of growing interest, because their promising application

including new lasing materials, upconverting phosphors and optical waveguides. All are due to their

high density, refractive index, thermal expansion, mechanical strength, high chemical durability,

temperature stability as well as with excellent transmission in the infrared (IR) region up to 4.5 micron [5].

Recent studies by Rabukhin and Belousova [6] found that the coordination numbers of cation present in

the composition study range does not significantly affect the structural grouping character of the

bismuth–containing gallate glasses. However, they noticed that the elastic constants of bismuth-containing

borate and germinate are strongly dependent on the change of coordination number of B3+

and Ge4+

.

The photoelastic constants of germanate glasses containing lead and bismuth oxides were studied

by Rabukhin [7,8] where he suggested that the photoelastic and elastooptic data could be used for the

light and acoustic lines of acoustooptical devices. He proposed that the high photoelastic constants of

lead bismuth germanate glass can be achieved through high refractive index and low modulus of elasticity.

For lead borate glass, at low PbO content the formation of four-coordinated boron proceeds at the

rate of two tetrahedral for each added oxygen. The lead enters the glass as modifiers Pb2+

ions.

However, at 15–20 mol% PbO, lead enters the network in the form of PbO4 pyramids (with Pb at the

apex of the pyramid). These PbO4 units bridge prefentially to BO3 rather than BO4 units [9].

Hamezan et al. [10] studied the elastic constants and thermal properties of lead bismuth borate

glasses. They found the density and molar volume of such glasses increase with glass modifier content

which attributed to the replacement of Bi2O3 and PbO; both have larger density and molar volume than

B2O3 in the glass networks.

So far, the sound velocity in lead and bismuth containing germanate glass and their elastic constants

have not been widely studied. In this research the glassy network of GeO2 will be added with PbO and

Bi2O3 glass modifiers. Their elastic properties will be investigated by the ultrasonic pulse-echo

technique [11–14]. The present work reports the preparation and ultrasonic characterization of

bismuth-lead germanate (BPG) glass systems.


Glass composition of the ternary bismuth lead germanate glasses (BPG) of the ternary form

(GeO2)60–(PbO)40−x–(1/2Bi2O3) together with density, molar weight and molar volume are given in

Table 1. The experimental data of bismuth lead borate (BPB) (B2O3)20–(PbO)80−x–(Bi2O3)x [14] is

presented for comparison purposes.

It can be seen from Figure 1 that, in contrast to BPB, the density and molar volume of BPG glasses

vary monotonically and slowly with the increase of Bi2O3 content. This change in density by the

addition of Bi2O3 is much related to the change in the atomic mass and atomic of volume constituent

Int. J. Mol. Sci. 2012, 13 4634

elements as given in Table 1. The atomic mass of the Bi, Pb, Ge and B atoms is 208.98, 207.20, 72.63

and 10.811 and their atomic radii are 1.56, 1.75, 1.22 and 0.90 Å respectively. This explains the

increase in density with increasing Bi2O3 content.

Table 1. Glass composition of the ternary lead-bismuth germanate glasses of the form

(GeO2)60–(PbO)40−x–(1/2Bi2O3)x together with density, molar weight and molar volume.

The experimental data of (B2O3)20–(PbO)80−x–(Bi2O3)x [14] is included for comparison.

Samples Glass Composition (mol%) Density (g cm−3

) Molar Weight (g/mol) Molar Volume (cm3)

GeO2–PbO–1/2Bi2O3

A1 60–40–0 5.90 152.06 25.77

A2 60–30–10 5.91 153.04 25.89

A3 60–20–20 5.99 154.01 25.73

A4 60–10–30 6.01 154.99 25.78

A5 60–0–40 6.05 155.97 25.80

B2O3–PbO–Bi2O3

B1 20–70–10 6.69 216.76 32.40

B2 20–60–20 6.90 241.04 34.93

B3 20–50–30 7.41 265.31 35.80

B4 20–40–40 7.26 289.59 39.89

B5 20–30–50 7.39 313.86 42.47

B6 20–20–60 7.36 338.14 45.94

B7 20–10–70 7.24 362.42 50.06

Figure 1. Density of the ternary lead-bismuth germinate glasses of the

form (GeO2)60–(PbO)40−x–(1/2Bi2O3)x (A) and compared with those of

(B2O3)20–(PbO)80-x–(Bi2O3)x (B) [14]. The line is drawn to guide the eye.

Figure 1 shows that densities of all the glass examined are increased with gradual substitution of

PbO with Bi2O3. The ternary borate containing lead and bismuth (B) show higher density changes as

compared with those of BPG glasses.

Int. J. Mol. Sci. 2012, 13 4635

Both density and molar volume of the glasses increase with an increasing of Bi2O3. The density and

molar volume increases by replacing PbO by Bi2O3 in Bi2O3–PbO–GeO2 glass system. It can be seen

from Figure 1 and Figure 2 that the density of BPG glasses varies from 5.90 to 6.05 g cm−3

and the

molar volume varies from 25.77 to 25.80 cm3 mol

−1. Generally, the density and the molar volume

show opposite behavior but, in this study, this was not the case. In this glass, substitution of lead by

bismuth causes the expansion of network. Similar trends for densities and molar volumes have already

been reported elsewhere for other glass systems [15–18].

Figure 2. Molar volume of the ternary lead-bismuth germinate glasses of the

form (GeO2)60–(PbO)40−x–(1/2Bi2O3)x (A) and compared with those of

(B2O3)20–(PbO)80−x–(Bi2O3)x (B). The line is drawn to guide the eye.

The molar volume of the ternary BPG glasses of the form (GeO2)60–(PbO)40−x–(1/2Bi2O3)x (A) is

shown in Figure 2 and compared with those of (B2O3)20–(PbO)80−x–(Bi2O3)x (B). It is clear that by

increasing Bi2O3, the molar volume increases which is similar with the variation density with

increasing Bi2O3 content. For the ternary BPG glasses (A), the change is quite small as compared with

the ternary BPB glasses and this might due to the compactness of the glass structure.

The Bi ions may enter the glass network interstitially, hence, some network bonds Ge–O–Ge or Pb–O–Ge

are broken and replaced by ionic bonds between Bi ion and singly bonded oxygen atoms. So if one

assumed that only effect of adding Bi cations was to break down the network bonds Ge–O–Ge and Pb–O–Ge

then an increase in the molar volume with Bi2O3 content would be expected for the entire vitreous

range of the studied glass system. Experimentally, this effect increases the molar volume in the glass

compositional range from 0 to 40 mol% Bi2O3 content (see Table 1) and as a consequence the values

of the density are increased.

The addition of Bi2O3 increased the values of density, which is probably attributable to

simultaneous filling up of the vacancies amidst the network by the interstitial Bi ions with atomic mass

208.98. This increase in density indicates a structural change in the glass network which is

accompanied by an increase in the molar volume [18,19].

The elastic moduli are proportional to the square of velocity and a plot of sound velocities vs.

composition is indicative of relative structure. The compositional dependence of longitudinal (Vl) and

Int. J. Mol. Sci. 2012, 13 4636

transverse (Vs) sound velocities are depicted in Table 2 and Figure 3. Figure 4 shows the variation of

elastic moduli as a function of Bi2O3 content. All the elastic moduli show the same trend as the

ultrasonic wave velocities.

Table 2. The room temperature ultrasonic wave velocities and elastic moduli of the ternary

lead-bismuth germanate glasses together with fractal dimensionality and Poisson’s ratio.

Glass sample x

Velocity (m/s) Elastic Moduli (GPa)

d = 4G/K

Poisson’s

Vl Vs L E K G Ratio

A1 0 3422 1912 69.10 54.93 40.34 21.57 2.14 0.27

A2 10 3445 1963 70.16 57.39 39.79 22.78 2.29 0.26

A3 20 3475 2027 72.28 61.10 39.49 24.59 2.49 0.24

A4 30 3482 2037 72.89 61.86 39.63 24.95 2.52 0.24

A5 40 3511 2094 74.53 64.90 39.18 26.51 2.71 0.22

Figure 3. Longitudinal and shear ultrasonic velocities at room temperature in ternary

(GeO2)60–(PbO)40−x–(1/2Bi2O3)x glass systems. The line is drawn to guide the eye.

Figure 4. Variation of elastic moduli of ternary (GeO2)60–(PbO)40−x–(1/2Bi2O3)x glass

systems with composition. The line is drawn to guide the eye.

Int. J. Mol. Sci. 2012, 13 4637

As can be seen from Figures 3,4, both Vl and Vs, and elastic moduli increase with the addition of

Bi2O3 content over the entire composition studied. The results in Table 2 indicate that the elastic

moduli increase linearly with Bi2O3 content. Longitudinal modulus changes from 69.10 GPa to

74.53 GPa, shear modulus from 21.57 GPa to 26.51 GPa, Young’s modulus from 54.93 GPa to

64.90 GPa and bulk modulus decrease from 40.34 GPa to 39.18 GPa as depicted in Figure 4. The

increase in elastic moduli is due to an increase in the cross-link density [20] and therefore an increase

in the rigidity of glass samples.

This variation of ultrasonic wave velocities and elastic moduli can be explained on the basis of the

structural consideration of germanate glassy network. In the present glass system, their longitudinal

ultrasonic velocity increases from 3422 to 3511 ms−1

and shear velocity increases from 1912 to 2094 ms−1

by increasing of bismuth content. According to Higazy and Bridge [21] the longitudinal strain changes

directly with bond stretching force constant. On the other hand, as reported by Reisfeld et al. [4] the

shear strain changes with bond bending force constant. The increase in velocities is attributed to the

increase in rigidity of the glass network.

Glass is considered as elastic substance and, thus, can be characterized through a modulus of

elasticity [11,12]. This modulus increases as the lengthening at a certain applied stress diminishes.

That will be the case if the glass structure is rigid and therefore contains the fewest possible non

bridging oxygen. When an oxide is introduced to germanate, the strength of the structure depends on

the field strength of the cation. With increasing Bi2O3 content in the germanate glass, the structure

becomes more rigid and so the density also increase and hence the modulus of elasticity increases [17–19].

It may also be noted from Figure 4 that the rate of change of elastic moduli is more pronounced in

longitudinal modulus (L) and least in case of transverse or shear modlulus (G). This indicates

resistance to deformation and it is most probably due to presence of large number of covalent bonds.

Bulk modulus (K) is the elastic property of material that can be derived most easily from the glass

structure [20]. GeO2 posses an open structure characterized by many open spaces. The addition of

Bi2O3 will occupy such spaces and this should leads to an increase in bulk modulus. With low Bi2O3

contents, the bulk modulus is small since many open spaces are present which will be filled most

quickly by the large Bi3+

ions (ion radius is 1.56 Å so that the bulk modulus increases. With higher

Bi2O3 contents, the ability for deformation of the cations becomes influential. This is clear from Figure 4,

which shows the relation between bulk modulus and Bi2O3 content.

The compositional dependence of Poisson’s ratio and fractal dimensionality parameter as a function

of Bi2O3 content is given in Figure 5. For BPG glasses their Poisson’s ratio, the ratio of transverse and

linear strains for a linear stress, decrease from 0.27 to 0.22. Poisson’s ratio has also been discussed in

terms of the dimensionality of glass network and it is observed that the Poisson’s ratio for a three

dimensional network is less than that of a two dimensional structure, which in turn is less than that of a

one dimensional structure. The decrease in Poisson’s ratio implies to the increase of crosslink density

of the glass as proposed by Higazy and Bridge [21].

Int. J. Mol. Sci. 2012, 13 4638

Figure 5. Variation of fractal dimensionality and Poisson’s ratio of ternary

(GeO2)–(PbO)–(1/2Bi2O3) glass systems with composition. The line is drawn to guide the eye.

The fractal dimensionality parameter (d = 4G/B) listed in Table 2 suggest that in all glasses the

dimensionality of glass structure lies between 2.14 and 2.71. As it can be seen from Figure 5, the

Poisson’s ratio is found to nonlinearly decrease with an increase of Bi2O3 content. Further, the values

of Poisson’s ratio are believed to be that of a covalently bonded structure. With an increase of Bi2O3

concentration, an increase in fractal dimensionality is observable. This suggests that, by increasing

bismuth cations, dimensionality of BPG glass samples increase.

3. Experimental

The new family of bismuth lead germanate (BPG) glasses in form of (½Bi2O3)x–(PbO)40−x–(GeO2)60

where x = 0 to 40 wt% have been successfully prepared by a rapid melt quenching technique [22]. The

reagent grade raw materials (99.9% purity of GeO2, PbO and Bi2O3) were weighed out the ratio of the

specified composition (in batches of approximately 15 g for each composition) as presented in Table 1.

The raw materials were mixed carefully in an agate mortar and then the mixtures in alumina

crucible were preheated in the first furnace at a temperature of 300 °C for 30 min. The mixtures were

then transferred in a second furnace and melted in 1100 °C for 1 h under atmospheric conditions. The

mixtures were stirred regularly to obtain a better mixing of the composition and to improve the

homogeneity of the glass samples.

The melt was then poured quickly on a preheated split metal mold and transferred to the first

furnace for annealing at 420 °C before cooling down to room temperature for 24 h. The glass samples

with 12 mm diameter and 30 mm thickness were bubble free and homogeneous with good optical

quality. A similar procedure was employed to prepare other glass samples. The color of glass sample

varies from light to dark violet as more Bi2O3 were added into germanate glassy networks.

The density (ρ) of the glasses were determined by Archimedes method with acetone as buoyant

liquid using the relation

Int. J. Mol. Sci. 2012, 13 4639

a

a b

W=

W Wb

where Wa is the glass sample weight in air, Wb the glass sample weight in buoyant and ρb the density of

the buoyant. All the glass sample’s weights were measured with a digital balance (±0.0001 g accuracy).

Their molar volume was calculated from the molecular weight (M) and density (ρ). The accuracy in the

measurement of the density is ±0.01 g cm−3

and the relative error is ±0.05%.

The chemically estimated elemental composition values present in the glass samples with the

Atomic Absorption Studies (Perkin–Elmer, Model 1372, USA) were found to be slightly smaller than

the corresponding elemental nominal composition values (before glass formation), which we

considered to be due to evaporation losses and uncertainties in the chemical analysis.

All the glass samples were checked by X-ray diffraction for their amorphous nature using

X’Pert Pro Panalytical PW 3040 MPD X-ray powder diffractometer by employing Cr-Ka radiation.

The absence of any crystalline peaks in the XRD patterns of the present glass samples indicates the

amorphous nature.

For the measurements of ultrasonic velocity in bismuth lead germanate (BPG) glass samples, the

samples were shaped into a circular disc of 12 mm diameter and 10–12 mm thickness. The opposite

faces of the disc shaped glass samples were highly polished using very fine lapping papers to achieve a

good surface finish with plane parallelism having accuracy of ±5 micron.

Ultrasonic velocity measurements were carried out at a frequency of 10 MHz using x-cut and y-cut

quartz transducers. A pulse superposition technique was employed using Ultrasonic Data Acquisition

System (MATEC 8020, Matec Instruments, USA) [12,22]. Burnt honey was used as a bonding

material between the glass samples and transducers. By measuring the thickness of the sample (d),

longitudinal (Vl) and transverse (Vt) wave velocities were calculated using the relation,

V = 2d/t [12,13]. The absolute accuracy in the measurement of the velocity is ±5 ms−1

and the relative

error is ±0.1%.

Glasses are isotropic and have only two independent elastic constant of L and G which are obtained

from longitudinal and shear sound wave velocity and density of the present glass samples. The various

elastic properties of the glasses were calculated using the following relations [22]:

Longitudinal modulus: L = ρV12 (1)

Shear modulus: G = ρVs2 (2)

Bulk modulus:

22

3

4sl VVK

(3)

Young’s modulus: 22

22243

sl

sls

VV

VVVE

(4)

Poisson’s ratio: 22

22

2

2

sl

sl

VV

VV

(5)

Fractal dimensionality: d = 4G/K (6)

Int. J. Mol. Sci. 2012, 13 4640

4. Conclusions

New ternary germanate based glass samples added with bismuth and lead oxides were successfully

prepared and their glassy natures were confirmed by the XRD method. Based on the result obtained,

it demonstrated that the density and molar volume increase with glass modifier content, which more

attributed to the replacement of Bi2O3 and PbO; both had larger density and molar volume than

germanate glass networks. Such increase in their density and volume with composition are due to the

compactness of the structure. The observed higher values in velocity at high modifiers content for this

glass types, confirmed a substantial change in glass structure. Their longitudinal, shear and Young’s

modulus for bismuth lead germanate glasses are also found to increase with the addition of Bi2O3.

Meanwhile, there was a similar pattern in elastic moduli with the increasing of Bi2O3 content, where

the values of both ultrasonic wave velocities increased subsequently.

Acknowledgement

The financial support from Universiti Putra Malaysia (UPM), under the Research University Grant

Scheme (RUGS), vote no. 91748 is gratefully acknowledged.

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(http://creativecommons.org/licenses/by/3.0/).

JOURNAL OF OPTOELECTRONICS AND ADVANCED MATERIALS Vol. 15, No.3 - 4, March – April 2013, p. 239 - 243

Effect of thickness on structural, optical and magnetic properties of Co doped ZnO thin film by pulsed laser deposition

A. KAMALIANFARa,b, S. A. HALIMa,c*, KASRA.BEHZADa, MAHMOUD GOODARZ NASERIa, M. NAVASERYa, FASIH UD DINa, J. A. M. ZAHEDId , K.P.LIMa, S.K.CHENa, H.A.A.SIDEKa

aDepartment of Physics, Faculty of Science, Universiti Putra Malaysia, 43400 UPM, Serdang, Selangor , Malaysia. bDepartment of Physics, Faculty of Science, University of Jahrom, Jahrom, 74137-66171 , Iran. cInstitute for Mathematical Research, Universiti Putra Malaysia, 43400 UPM, Serdang, Selangor, Malaysia dDepartment of Physics, Karaj Branch, Islamic Azad University, Karaj, Iran. Zn0.97Co0.03O thin films were deposited on n-type silicon substrate (111) using pulsed laser deposition (PLD) technique with different thicknesses of 200 20 , 320 20 , and 480 20 . The Zn0.97Co0.03O powders prepared by sol-combustion were formed as pellets. A neodymium-doped yttrium aluminum garnet (Nd:YAG) class 4 laser (λ = 266 nm, 5-8 pulse duration, 2 Hz duration rate) were used for the PLD of the nanostructures. The structure of the prepared thin films was investigated using different techniques such as X-ray diffraction (XRD) and field emission scanning electron microscope (FESEM). The optical properties of prepared thin films were investigated using UV-vis spectrometry. In addition, the magnetic properties of the samples were studied by vibrating sample magnetometer (VSM). The results showed that the properties of the thin films altered by thickness changes. (Received August 24, 2012; accepted April 11, 2013) Keywords: Co-doped ZnO, Thin Film, PLD, Dilute magnetic semiconductor

1. Introduction Recently, the potential for some applications in

spintronics has attracted wide research on dilute magnetic semiconductor (DMS). The main idea of DMS is to dope transition metal ions (e.g. Mn, Co or Fe) into a semiconductor host to prepare magnetic semiconductor at room temperature [1],[2],[3],[4]. It is supposed to have good control on the spin as well as electron charge for next generation spintronics devices. Some properties of ZnO such as wide band gap energy (~3. 37ev), transparency to visible light and chemical compatibility with other oxidizing materials provide a potential for the usage in magneto optical devices [5]. A variety of methods have been used to synthesize doped and undoped ZnO semiconductors including sol-gel [6], Spin-Coating Technique [7], pulsed laser deposition (PLD) [8], [9], [10] and vapor phase transport (VPD [11]. Among these methods, pulsed laser deposition method is a versatile technique and allows deposition different materials, e.g. high-temperature superconductors and oxides with high deposition rates.

Here, we study the effect of thickness on the structural, optical and magnetic properties of Zn0.97Co0.03O thin films. PLD technique was chosen and three Co doped ZnO thin films were prepared.

2. Experimental Zn0.97Co0.03O powder was prepared by sol-combustion

method [12]. The powder was pressed in a form of pellet and sintered at 600 ˚C for 1h to be used as a target. The silicon substrates were cleaned by HCl, NH3 and acetone before being placed onto the sample holder. A Nd:YAG laser system (λ=266 nm, 5-8 ns of duration pulse, 2 Hz of repetition rate and 0.5 J/cm2 fluence of energy) was used to deposit three Zn0.97Co0.03O thin films with different thicknesses. The background pressure of the vacuumed chamber was 10-5mbar. The deposition was carried out in an oxygen atmosphere about 0.15 mbar instead of nitrogen due to the incorporation of nitrogen affecting the properties of the DMS material [13],[14]

The temperature of the substrates was fixed at 600 ˚C while the holder of the target was rotated slowly to obtain better thin films crystalline and homogeneity. A post-deposition annealing at 550 ºc was applied to decrease oxygen vacancies and the samples were cooled to room temperature at 8 ◦C/min rates. Using a high surface profilometer (Ambios, XP-200), the thickness of the films was estimated about 200 20 . 320 20 , 48020 . The structural analysis of the samples was performed using X-ray diffraction (XRD, Pw 3040 MPD) with energy dispersive X-ray spectrometer (EDX). The morphology of the samples was observed by field emission scanning electron microscopy (FESEM, JSM

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tures and mofilms

ows the XRDposited on silie 600°C. The mbe seen in al

s and Table 1me causes more thickness of

XRD patterns sses of (a) 200

esponding pepatterns beca

e. A typical Ewn in Figure 2e detected in a

While the EDXbsence of Co), indicates tha

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nfar, S. A. Hali

optical propevis, Perkin Elmmples in the emperature. ne using vake Shore

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orphology of

patterns of thcon substratemain peaks ofl of the patte1, it can be fre intensity ofthe film.

of the preparednm (b) 320 nm

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the Co doped

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d films with (c) 480 nm.

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Goodarz Naseri

Fig.2. EDX of t

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the thin film wi

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Fasih Ud Din,

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Effect of thickness on structural, optical and magnetical properties of Co doped ZnO thin film by pulsed laser deposition 241

Table 1. Comparison of position, height, FWHM, crystallite size and dislocation density of (002)

diffraction peaks.

No. Position angle (2θ)

Height (CTs)

FWHM (θ)

Crystallite size (nm)

Dislocation density (10-4

nm-2 )

(a) 20020

35.51

20.93

0.5723 15.22

43.14

(b) 32020

34.50

21.98

0.4149 20.94

22.80

(c) 48020

34.42

22.04

0.3046 28.52

12.30

For magneto-optic applications, surface properties are significant factors that influence some optical and magnetic properties. Fig. 3 shows the morphology of the surface of thin films after annealing at 550˚C for 1 h under O2 gas flow. All the images were obtained by field emission scanning electron microscopy (FESEM). The size distribution diagram for the particles is shown below the FESEM images. As observed in the images, for the thin film with 200 20nm thickness, the mean size of particles is 14.75±0.03 nm, whereas for 320 20nm and 480 20nm thicknesses, the mean size of particles are 19.01±0.03nm, 28.13±0.03 nm respectively. It can be stated that the particles have probably more freedom directions for growth when the thickness of film increases, so the size of particles will be larger.

Fig.3. FESEM images their histogram particle size diagrams of the films with thicknesses of (a) 200 20 (b) 320 20

(c) 480 20

3.2 Optical characteristics The variation of thickness of Co doped ZnO layers

and the growth temperatures have significant effect on the optical properties and crystalline structure. To investigate the film thickness effect on the energy band gap values, the (αhν) 2 versus (hν) of absorption spectra in the range of 200 and 800 NM was plotted, where α is the absorption coefficient and hν is the photon energy. Figure 4 presents variation of the energy band gap for different film thicknesses. It indicates that by increasing the thickness of the films, the band gap energy values decrease from 3.27 to 3.32 eV. This variation of the band gap values can be related to decreasing the crystalline size of thin films because the crystalline size rises with increases of the film thickness (quantum size effect). This change can also be attributed to the improvement in the crystals, changes of

structural defects, atomic distances and grain size in the films. The inset shows the absorption spectra of the samples.

Fig.4. Variation of (αhν) 2 versus hν for Zn0.97Co0.03O thin films with different thicknesses.

242 A. Kamalianfar, S. A. Halim, Kasra.Behzad, Mahmoud Goodarz Naseri, M. Navasery, Fasih Ud Din,…

3.3 Magnetic characteristics Fig 5 shows the field-dependent magnetization (M–H)

of the Zn0.97Co0.03O samples with different thicknesses at room temperature. As a point of comparison, the magnetization of two other ZnO films with 6 and 9 percent of cobalt was shown as well. The diamagnetic ratio due to Si substrate was subtracted out of the obtained data. As can be seen, the Zn0.97Co0.03O films exhibit weak ferromagnetic behavior. The curves of (d) and (e) shows the field-dependent magnetization of the Zn0.94Co0.06O and Zn0.91Co0.09O films with a similar thickness (~44020 and 550 20 ) but with more concentration of Co. Hence, the films show a stronger ferromagnetic behavior. The question here is : what causes a room temperature ferromagnetism in a DMS material? Based on ZnO semiconductors with defects, the Ruderman–Kittel–Kasuya–Yosida (RKKY) model was employed by Story et al. [17] to explain the exchange coupling energy. Thus, the effective magnetic Hamiltonian is given by this equation:

∑ · ∑ · ∑ ′ ·

(3)

In this equation, JF is the parameter of ferromagnetic coupling between the magnetic ions, JAF is the antiferromagnetic coupling parameter between them and

is the ferromagnetic coupling parameter for the impurities and localized carriers. The spatial distance between randomly located substitutional magnetic pair in the lattice (rij) and the impurity spins (Si, Sj) are other parameters considered.

Therefore, the room ferromagnetic behavior originates from three parts. One of them is direct interaction between magnetic ions (antiferromagnetic coupling) and two others are indirect couplings between the bound magnetic polaron mode and delocalized carriers or weak localized carriers. These two couplings provide ferromagnetic alignment inside the lattice. According to [18], for the samples with an absolute moment less than 5 µemu, the moment comes from extrinsic sources including substrate effect. Liu et al. [19] reported that the oxygen vacancies contribute to an indirect interaction with the bound magnetic polarons model and provide the ferromagnetic behavior of Co doped Zno. They also mentioned that the ferromagnetism is not due to metalic cobalt clusters.

Although, the Co ions of Co doped ZnO can cause the small amount of absolute moment, but the possibility of the mismatch coupling between Co spin orientation and donor electron states for the room temperature ferromagnetic cannot be ruled out. In addition, the Co doped ZnO has this ability to present strong room temperature ferromagnetism [20].

Fig.5. Room temperature field-dependent magnetization (M–H) for Zn0.97Co0.03O thin films with different thicknesses (a) 300 20 (b) 320 20 (c) 480 20 (c) Zn0.94Co0.06O with 440 20 thickness (d) Zn0.91Co0.09Owith 455 20 thickness

4. Conclusion The Zn0.97Co0.03O thin films were prepared on Si

substrate using a PLD method with different thickness. The XRD patterns of the films showed ZnO hexagonal wurtzite structure and no cobalt oxide peaks were observed. It also showed that by increasing the thickness of the films, the (200) peak position is shifted a little to lower degrees, and the grain size parameter increased showing better crystallization of films with the increase of the thickness. From FESEM pictures and histogram particle size diagrams, it was estimated that the size of the particles on the surface rose in the range of 14 to 30 nm with increasing the thickness of the films. The optical measurement of the films indicated that the band gap of samples was reduced (~0.01 ev) by increasing the thickness. The films with 3% cobalt doped showed a superparamagnetic behavior. By increasing the concentration of cobalt in ZnO lattice, the thin films showed a weak ferromagnetic behavior with saturation magnetization about o.oo6 emu/g.

References

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Effect of thickness on structural, optical and magnetical properties of Co doped ZnO thin film by pulsed laser deposition 243

[5] K. Sato and H. Katayama-Yoshida, Jpn. J. Appl. Phys. 39, L555 (2000). [6] T. Shi, S. Zhu, Z. Sun, S. Wei, and W. Liu, Appl. Phys. Lett. 90, 102108 (2007). [7] F-Z. Ghomrani, S.Iftimie, N. Gabouze, A. Serier, M.Socol, A. Stanculescu, F. Sanchez, S.Antohe, M.Girtan , J Optoelectron. Adv. Mater. 5, 257 (2011). [8] S. Pilban Jahromi, N.M. Huang, A. Kamalianfar, H.N. Lim, M.R. Muhamad, R. Yousefi, Journal of Nanomaterials 2012, 173825 (2012). [9] K. Samanta, P. Bhattacharya, R. S. Katiyar, W. Iwamoto, P. G. Pagliuso,and C. Rettori, Phys. Rev. B 73, 245213 (2006). [10] J. H. Kim, H. Kim, D. Kim, Y. E. Ihm, W. K. Choo, Physica B. 327, 304 (2002). [11] A. Kamalianfar, S.A. Halim, S.P. Jahromi, M. Navasery, F.U. Din, K.P. Lim, S.K. Chen, J.A.M. Zahedi, Chinese. Phys. Lett. 29(12), 128102 (2012). [12] A. Khorsand Zak, M. Ebrahimizadeh Abrishami, W. H. Abd Majid, Ramin Yousefi, S. M. Hosseini,Cer. Inter. 37, 393 (2011).

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Glass Science & Technology

2012/2013

Glass Science & Technology Research @ UPM

Education

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