Structure and properties of low-phonon antimony glasses and nano glass-ceramics in K2O–B2O3–Sb2O3 system

Journal of Non-Crystalline Solids 356 (2010) 987–999

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Structure and properties of low-phonon antimony glasses and nano glass-ceramics in K2O–B2O3–Sb2O3 system Tirtha Som, Basudeb Karmakar * Glass Science and Technology Section, Glass Division, Central Glass and Ceramic Research Institute (Council of Scientific and Industrial Research), 196, Raja S.C. Mullick Road, Kolkata 700 032, India

a r t i c l e

i n f o

Article history: Received 6 March 2009 Received in revised form 11 January 2010 Available online 6 March 2010 Keywords: Phonons Heavy metal oxides FT-IR measurements Water in glass

a b s t r a c t A series of new low-phonon and low softening point monolithic antimony glasses and nano glass-ceramics (crystallite diameter 7–16 nm) containing as high as 90 mol% Sb2O3 were prepared in the ternary K2O– B2O3–Sb2O3 (KBS) system. Infrared spectra and X-ray analysis establish that their structures closely resemble valentinite form of Sb2O3. The phonon energy has been found to be 592–602 cm1. The softening point (Ts), glass transition temperature (Tg), coefficient of thermal expansion (CTE), dielectric constant (e) and optical band gap (Eopt) have been found to vary in the ranges 331–392 °C, 234–264 °C, 201– 222  107 K1, 12.4–14.5 and 3.15–3.22 eV, respectively. These properties are found to be controlled by the covalent character and optical basicity of prepared samples. The average electronic polarizability of the oxide ions and optical basicity are found to increase while interaction parameter and metallization criterion decrease with increase in Sb2O3 content. Ó 2010 Elsevier B.V. All rights reserved.

1. Introduction The enormous interest in heavy metal (atomic weight >100) oxide (HMO) glasses and glass-ceramics is due to their several inherent advantages over the conventional SiO2, GeO2, B2O3 and P2O5 containing glass systems. They possess attractive properties such as high density, high linear and non-linear refractive indices, high thermal expansion, low softening temperature, large transmission window (UV to IR), ultrafast response time and low-phonon energy [1–5]. This makes them suitable for several scientific applications like components of non-linear optics (NLO), rare-earth (RE) doped solid state upconversion laser materials, optical amplifiers, etc. [1–5]. Although much research has been performed on multi-component HMO glasses based on the oxides Ga, Bi, Pb and Te [1,2], but studies on antimony oxide based glasses and glass-ceramics are relatively very less. Antimony (atomic weight = 122) oxides are known to exist in wide range of compositions and display interconvertable polymorphism. The two common forms are senarmontite (cubic) and valentinite (orthorhombic) consists of Sb4O6 molecules and chains of SbO3 trigonal pyramids respectively [6]. The polymorphic forms of Sb2O4 are the orthorhombic a-phase (cervantite) and high temperature monoclinic b-phase [6]. It is difficult to obtain vitreous

* Corresponding author. Tel.: +91 33 2473 3469; fax: +91 33 2473 0957. E-mail address: [email protected] (B. Karmakar). 0022-3093/$ - see front matter Ó 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2010.01.026

Sb2O3 due to its low field strength z/a2 (0.73) which makes it a poor glass former [7]. A survey of the studies on antimony glasses concedes that pure Sb2O3 glasses are chemically unstable [8]. The addition of a second component such as B2O3 [8], SiO2 [9], P2O5 [10,11], or M2O (where M = Li, Na, K) [4,12] provides stabilization of the Sb2O3 network but the yield is mostly small quantities of pulverized glasses. Nevertheless persistent efforts have been made by several researchers to incorporate antimony within glasses either in form of oxides, halides, phosphates or sulphides and study their physical properties [9–16]. Sb2O3–SbPO4 glasses are suggested for technological applications as optical recording media [17] and NLO appliances due to their photosensitivity, non-centrosymmetric structure and fast response times as a consequence of the high polarisable Sb3+ ions with strongly localized stereochemically active lone pair of electrons (5s2) [5]. However, volatility of the melts, intense crystallization and devitrification while cooling the melts, and above all the difficulty in preparation of monolithic glass with very high Sb2O3 content (50–90%), essential for practical applications, have limited the study of antimony systems particularly in the areas of optics and photonics. It is only in recent times that we have first exploited the remarkable upconversion luminescence of rare-earth ions (Sm3+, Er3+, Nd3+) in antimony based (70 mol%) monolithic oxide glass [18–20]. Theoretically Sb2O3 containing glasses are expected to have low melting and glass transition temperature, and high linear thermal expansion to match those of certain metals and alloys. These have further sparked our interest in Sb2O3 based systems to provide the

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electrical industry with low melting insulating glass composition for sealing applications. To the best of our knowledge, the structural investigations of antimony based glasses have mainly been performed on binary Sb2O3–B2O3 systems [3,8,21–24] and some studies on vitreous Sb2O3 [24] and Sb2O3–SbPO4 systems [10], but no detailed comparable investigations of monolithic alkali–boron–antimony based glasses exists. The investigations on the structure of the pulverized Sb2O3–B2O3 glasses prepared by melt-quick(forced)-quench routes by EXASFS, 11B NMR, Mossbauer, and Raman spectra, reveal that antimony ions enter the glass network in the form of SbO3 trigonal pyramids, with each of the oxygens acting as a bridging oxygen to either another such pyramid or to a B–O group. The addition of Sb2O3 results in elimination of boroxal rings and conversion of some trigonal planar [BO3] units into three dimensional tetrahedral [BO4] units [21–24]. In this study, we demonstrate the preparation of a new series of monolithic antimony glasses and nano glass-ceramics in the K2O– B2O3–Sb2O3 (KBS) system, investigate their structure by X-ray diffraction and infrared spectral analyses and examine their various physical properties. Here we also evaluate the average electronic polarizability of the oxide ions comprising the systems, the covalent character and optical basicity of the glasses and glass-ceramics, interaction parameter and metallization criterion from the values of theoretical refractive index and optical band gap and correlate all properties with the fundamental properties that is, covalent character and optical basicity.

2. Experimental 2.1. Glass and glass-ceramic preparation The raw materials were potassium metaborate, KBO2xH2O (15.7% H2O, Johnson Matthey) and antimony(III) oxide, Sb2O3 (GR, 99%, Loba Chemie). They were used directly without any further purification. No separate method was followed to dehydrate KBO2xH2O. In order to mix the glass batch compositions homogeneously, all the raw materials were first mixed carefully in ethanol medium followed by drying. The nominal composition (mol%) of the investigated KBS glasses and glass-ceramics are: xK2O– xB2O3–(100  2x)Sb2O3, where x = 5–35 (details are given in Table 1). About 20 g of both glasses and glass-ceramics was prepared by simply melting the well mixed batches of calculated composition in a high purity silica crucible at 900 °C for 10 min with intermittent stirring for 0.5 min in air in a raising hearth electric furnace. All the molten samples were cast into carbon plates in air and annealed at 260 °C for 10 h to remove thermal stress followed by slow cooling to room temperature. The glasses and glass-ceramics thus obtained were cut and polished into desired shapes and sizes for the property measurements as described below.

2.2. Physical measurement The bulk densities of the samples were measured by the Archimedes method using toluene as immersion liquid. Density of toluene at the experimental temperature was found to be 0.861 g cm3. A sample of vitreous silica (density = 2.2 g cm3) was measured five times under similar conditions and the density was found to be 2.198 ± 0.0176 g cm3. An average of five separate measurements was calculated and chosen as representative value to minimize random error. Measurements were accurate to ±0.8%. Water may have some corrosive effect on the antimony glasses and also make them soluble to some extent. To avoid the solubility problem we used a non-polar hydrophobic solvent (e.g. toluene). The chemical durabilities of the KBS need to be investigated before their densities can safely be measured by water. The theoretically predicted respective glass properties and refractive indices were determined using SciGlass (Glass Properties Information System, Version 6.7) software following the Priven2000 method. The infrared transmission spectra (FT-IRTS) of the raw materials, glasses and glass-ceramics in the range 400–4000 cm1 were recorded by following the KBr pellet method with a Fourier transform infrared (FT-IR) spectrometer (Perkin Elmer, FT-IR 1615) at a resolution of ±2 cm1 and after 16 scans to identify the bond vibrations (absorption peaks) particularly in the fingerprint region. The FT-IRTS of the polished transparent monolithic samples (thickness: 1 ± 0.1 mm) in the range 1000–4000 cm1 were also recorded under similar conditions for estimation of OH content. The infrared reflection spectra (FT-IRRS) in the range 400–1500 cm1 were recorded with the same instrument at an incident angle of 15° and with the help of a specular reflectance attachment accessory at the resolution of ±1 cm1 after 256 scans. The instrument was calibrated with the peaks of a standard polystyrene film supplied by the manufacturer of the instrument and normalized. The UV–Vis transmission spectra were obtained with a double-beam spectrophotometer (Perkin Elmer, Lambda 20). The instrument was calibrated with the peaks of a standard holmium ion (Ho3+)-doped glass filter supplied by the manufacturer of the instrument to minimize the systematic error. The uncertainty of the band position is ±0.1 nm. XRD data of powder samples were recorded using X’pert Pro MPD diffractometer (PANalytical) with 2h varying from 10° to 80° with step size of 0.05° (2h) and step time of 0.5 s, using Ni filtered CuKa (k = 1.5406 Å) at 25 °C and generator power of 45 kV and 35 mA. An automated computer system was used to identify the patterns. The softening point (Ts) of each of the bulk samples was measured five times by a glass softening point system (Harrop/Labino, Model SP-3A) and the average values are reported. They are accurate to ±2 °C. The instrument was previously calibrated with a NBS (National Bureau of Standards, USA) standard glass of known softening point. The coefficient of thermal expansion (CTE, a), glass

Table 1 Nominal composition and some properties of glasses and glass-ceramics. Sample no.

B1 B2 B3 B4 B5 B6 B7

Composition (mol%) Sb2O3

K2O

B2O3

90 80 70 60 50 40 30

5 10 15 20 25 30 35

5 10 15 20 25 30 35

Color

Form

Nature

Deep yellow Yellow Yellow Pale yellow Colorless Colorless White

Translucent large broken pieces Transparent monolith Transparent monolith Transparent monolith Transparent monolith Transparent monolith but surface easily attacked by moisture Opaque small broken pieces

Glass-ceramic Glass Glass Glass-ceramic glass-ceramic Glass-ceramic Crystallized

T. Som, B. Karmakar / Journal of Non-Crystalline Solids 356 (2010) 987–999

transition temperature (Tg) and deformation temperature (Td) were measured using a horizontal vitreous silica dilatometer (Netzch, DIL 402C) with heating rate 5 K/min after calibration with standard alumina supplied with the instrument by the manufacturer. The CTE values reported here are in the temperature range 50– 150 °C. The CTE, Tg and Td values were reproducible with ±1 °C for all samples. Samples for dielectric constant measurements were cut and polished to the rectangular shape having cross-sectional area of 15  15 mm2 and thickness 2 mm. Either side of the large-faces of the samples thus prepared was coated with a thin layer of conducting silver paint and dried at 140 °C in order to serve as electrodes for dielectric measurements. A four probe method was employed and the dielectric constant of the samples was measured at room temperature using a LCR meter (Hioki 3532-50 LCR HiTESTER) at 1 MHz frequency. Dielectric constant was measured six times for each sample and the average values are reported. They are accurate to ±0.7%. The instrument was previously calibrated with a standard high purity, low OH content SUPRASIL-W silica glass (Heraus, Germany) having dielectric constant 3.8. Field emission scanning electron microscopic (FESEM) images were taken with a Gemini Zeiss Supra™ 35VP model (Carl Zeiss Microimaging GmbH) using an accelerating voltage of 4.9 kV. Fracture surfaces of samples having dimension 2  2  2 mm3 were examined by FESEM after etching with 5 vol.% HNO3 aqueous solution for 1 min. Standard wet-chemical analysis methods (error ±3%), such as flame photometry (for estimation of K2O) and titrimetry (for estimation of B2O3, Sb2O3 and SiO2) was used to measure the true compositions of the glasses and glass-ceramics (elemental composition analysis) [19]. The compositions of these glasses were also re-checked by energy dispersive X-ray (EDX) analysis following the method described by Metwalli and Brow [25] by an EDX analyzer attached with the FESEM instrument. A minimum of five spots on each sample were analyzed and the average composition (±5% relative) was considered. It was found that the loss due to volatilization of Sb2O3 varies in the range 2– 4 mol%. The loss increases with increasing Sb2O3 content. At the same time, inclusion of SiO2 from melting crucible was found to vary in the range 2–2.5 mol%.

3. Results

position values are used in the theoretical calculations for simplicity, as reported later in the paper. The positions of glass and glassceramic formation and crystallization are depicted in Fig. 1. All samples are obtained as transparent monoliths except B1 and B7. The visible appearance of the sample changes as the amount of Sb2O3 is increased. B1 is hazy, almost opalescent while B7 is opaque and obtained as tiny pieces. The samples having higher Sb2O3 content are yellow in color. Fig. 2 is a representative EDX spectrum of glass-ceramic B5 manifesting the incorporation of SiO2. 3.2. Physical properties 3.2.1. X-ray diffraction The formation of glasses and glass-ceramics can be explicitly determined by X-ray diffraction analysis. Fig. 3 shows the XRD spectra of the samples. Samples B2 and B3 show a hump between 2h = 22–35° while B1, B4, B5 and B6 shows gradual development of sharp peaks in addition to the hump. The XRD spectra of the raw material Sb2O3 (curve h) was also compared with those of the obtained samples. It shows major peaks at 2h = 19.3399° (d = 4.58586 Å), 27.9531° (d = 3.18932 Å), 28.3187° (d = 3.14897 Å) and 46.6931° (d = 1.94538 Å) due to diffractions from the (1 1 0), (1 3 0), (1 2 1) and (2 4 0) from the valentinite form of Sb2O3 (cell constants a = 4.914 Å, b = 12.468 Å, c = 5.421 Å; JCPDS Card No. 11-0689) and at 2h = 27.7656° (d = 3.21308Å), 32.1036° (d = 2.78583 Å) and 54.5713° (d = 1.68031 Å) due to diffractions from the (2 2 2), (4 0 0) and (6 2 2) planes of the senarmontite form (a = 11.152 Å; JCPDS 50534). In addition to the above peaks, the XRD spectra of the glassceramics (B1, B4–B6) and crystallized sample (B7) show prominent peaks at 2h = 14.5025° (d = 6.10783 Å), 24.8662° (d = 3.58077 Å), 35.9868° (d = 2.49569 Å) due to Sb2O5 crystalline phases (JCPDS 34-878). The XRD spectra of crystallized sample B7 also show peak at 16.4528° (d = 5.28796 Å) and 34.2995° (d = 2.61450 Å) due to (1 1 2) and (2 2 4) diffractions from the K4Sb2O75.6H2O (JCPDS 35-376). The peak at 19.8484° (d = 4.46950 Å) is attributed to (0 1 1) diffractions from the cervantite form (a = 5.436 Å, b = 4.810 Å, c = 11.76 Å; JCPDS 11-694). The average crystallite diameter calculated using Scherrer’s formula [26]:

d ¼ 0:9k=FWHM cos 2h ðpeakÞ;

3.1. Physical appearance

989

ð1Þ

Nominal composition and some physical properties of the investigated samples are indicated in Table 1. These nominal com-

Sb2O3 0 100 25

75

50

50

75

100 0

K2O

25

0 25 Glass-ceramic,

50 Glass,

75 100 Crystallized

B2O3

Fig. 1. Glass formation in the K2O–B2O3–Sb2O3 system (composition in mol%).

Fig. 2. EDX spectrum of B5 showing the inclusion of silica from melting crucible (for composition see Table 1).

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T. Som, B. Karmakar / Journal of Non-Crystalline Solids 356 (2010) 987–999

16000

⊗

14000

Intensity (a.u.)

♦ ♦

600

Intensity (a.u.)

10000 8000

♦

⊗

∇Δ

♦

◊

(a)

200 28

32 36 40 2θ (degree)

♦

⊗

♦

♦ = Valentinite

(b)

400

0

12000

♦ ♦ ⊗ ♦

The morphology of the glass-ceramics B4, B5 and B6 were investigated by FESEM. Their FESEM images (Fig. 4(a)–(c)) depict gradual change of microstructures from granule clustered to cellwall like to house-of-disks respectively which consists mainly of nano crystallites of valentinite Sb2O3 and have been left after elimination of borate-rich phase by acid.

⊗ = Senarmontite

800

♦

♦

◊ = Cervantite ∇ = Sb2O5 Δ = K4Sb2O7.5.6H2O ⊗ ⊗ ♦ ⊗

♦♦

♦ Δ∇♦ ⊗ ♦ ♦ ∇♦∇ ◊ ♦ Δ ∇ ◊

⊗⊗

h

g

6000

f 4000

e d c b a

2000 0 10

20

30

40 50 2θ (degree)

60

70

80

Fig. 3. X-ray diffractograms of (a) B1, (b) B2, (c) B3, (d) B4, (e) B5, (f) B6, (g) B7 (for composition see Table 1) and (h) Sb2O3 (used raw material is shown for comparison).

where k is the wavelength of X-ray radiation (CuKa = 1.5406 Å), FWHM is the full width at half maximum at 2h. The average diameter of the crystallites present in B4, B5, B6 and B7 are 7, 8, 11 and 16 nm (error ± 0.2 nm), respectively.

3.2.2. Infrared transmission spectra of KBr pellets (FT-IRTS, 400– 4000 cm1) The Fourier transform IR transmission (FT-IRTS) spectra of the raw materials and the synthesized KBS glasses are given in Fig. 5. The FT-IRTS spectra of reagent KBO2xH2O (curve a) can be divided into three major regions first group of bands lying between 1200 and 1500 cm1 is attributed as the asymmetric stretching vibration of the B–O–B bond of trigonal [BO3] units in KBO2xH2O. The second group of bands which occurs 800–1200 cm1 is due to asymmetric stretching vibration of the B–O–B bond of tetragonal [BO4] units. The third group around 646 and 708 cm1 is due to bending vibration of B–O–B linkages in the borate network [27,28]. In addition there is a wide band around 3462 cm1. This broad and very strong absorption band around 3462 cm1 is due to the hydroxyl groups and indicates the presence of the amount of water (15.7 wt% H2O) in this reagent (KBO2xH2O). Fig. 5 (curve b) depicts the FT-IRTS spectrum of raw material Sb2O3 powder used in glass preparation. It shows a small peak at 954 cm1 corresponding to Sb–O stretching vibrations of senarmontite form of SbO3 [5,10], a prominent sharp peak at 739 cm1 due to the symmetric stretching vibration mode of senarmontite form having C3V symmetry [5,10,29,30] and a number of small bands at 692, 592, 546 and 492 cm1 owing to symmetric stretch-

Fig. 4. FESEM images of fracture surfaces of a glass-ceramics (a) B4, (b) B5 and (c) B6 showing (a) granule clustered, (b) cell-wall like and (c) microstructures respectively comprised of nano crystallites. Scale bar for is 1 lm for all images (for composition see Table 1).

T. Som, B. Karmakar / Journal of Non-Crystalline Solids 356 (2010) 987–999

991

In addition, the dissolved water in the glasses produces sharp bands in this IR region. In case of the silica glass, the main band at 3662 cm1 (2.73 lm) due to the presence of free hydroxyl (OH) group while the two weak bands at 2645 cm1 (3.78 lm)and 2256 cm1 (4.43 lm) is probably due to hydrogen bonded OH. In case of the KBS antimony glass, the sharp band at 3259 cm1 (3.07 lm) is associated with hydrogen bonded hydroxyl groups (–O–Hd+Od–) [31]. As the concentration of KBO2 increases, the concentration of K2O increases and thereby the amount of nonbonding oxygen also increases. Consequently the two bonds shifted to higher wavelengths at 2541 cm1 (3.94 lm) and 2227 cm1 (4.49 lm) is due to hydrogen bonded O–H groups. The IR transmission spectra can provide much information about the hydroxyl group (OH) content in glass. The OH content of the glasses can be calculated by the method described by Ebendroff-Heidepriem et al. [32]. It must be noted that the OH bands obtained in the IR transmission spectra of KBS antimony glass (this study) are similar to that of fluorophosphates and calcium metaphosphate glasses [33–35]. These metaphosphate glasses display stretching vibrations of weakly hydrogen bonded OH at about 2900 cm1 and its corresponding absorption coefficient has been used as a measure of OH concentration. The 3259 cm1 band in KBS glass is also due to hydrogen bonded and is very close to that of metaphosphate glasses. So, the absorption coefficient, aOH at 3259 cm1, in KBS glass can be used to determine the OH concentration in KBS glass using the following relation [32]:

aOH ¼ logðT o =TÞ=d; Fig. 5. Infrared transmission spectra (FT-IRTS) of raw materials, glasses and glassceramics taken by the KBr pellet method: (a) KBO2xH2O, (b) Sb2O3, (c) B2, (d) B3, (e) B4, (f) B5 and (g) B6 (for composition see Table 1). Spectra of raw materials (a) and (b) are shown for comparison.

ing, asymmetric stretching, symmetric bending and asymmetric bending vibration modes of valentinite form of SbO3 trigonal pyramids with Cs or C2 symmetry, respectively [5,10,29,30]. This again insists that both forms of Sb2O3 are present in the starting material. The FT-IRTS spectra of the KBS antimony samples, B2–B6 (Fig. 5, curves c–g), the sharp peak at 592 cm1 and the small peaks at 685 and 931 cm1, and several minor peaks between 1000 and 1500 cm1 whose intensity increases as we proceed from B4 to B6 (curves e–g). The KBS antimony samples also display small peaks at 1630, 3277, 3392 and 3554 cm1. The intensity and number of peaks in the region 1000–1500 cm1 increases while the band around 3462 cm1 broadens and becomes stronger as we proceed down the series from B2 to B6 (curves c–g). The IR absorption bands of the samples with nature and assignments have been listed in Table 2 for convenience.

3.2.3. Infrared transmission spectra of bulk samples (FT-IRTS, 1000– 4000 cm1) We have measured the IR transmission spectra of the KBS glasses from 8000 to 1000 cm1, but interesting features due to presence of O–H group appears between 4000 and 1000 cm1. Consequently in Fig. 6, the IR transmission spectra from 1000 to 4000 cm1 of glass B2 and silica glass (both polished and 1.0 ± 0.1 mm thick) is demonstrated for comparison. The silica glass becomes completely opaque to the IR radiation at 2057 cm1 due to Si–O bond vibrations which are responsible for the strong IR absorption while in the case of KBS antimony glass, the IR absorption edge is red-shifted to 1492 cm1 due to the contributions from Sb–O bond vibration.

ð2Þ

where T0 is the highest transmission at 1400 nm (7142.85 cm1), T is the transmission of the glass at 3259 cm1 and d is the thickness (cm) of the glass sample. It must be mentioned here that %T of B2 throughout the range is almost 73.3% (±0.5%). The OH content can be calculated using above relation and according to the following equation [32]:

OH contentðppmÞ ¼ 30  aOH ðcm1 Þ:

ð3Þ

The estimated OH content for some of the samples under study is enlisted in Table 3. Estimated accuracy is ±0.4. Similar method was also followed by Lezal et al. [2] to estimate the OH content of heavy metal oxide glasses. In the present case, it is seen that the OH content increases with increase in K2O and B2O3 content in the glasses and glass-ceramics. It thus suggests that these O–H groups have been originated from KBO2 raw material which contains about 15.7% H2O. 3.2.4. Infrared reflection spectra (FT-IRRS) The phonon energy (hx) of glasses is the resonance vibration energy of the lattice and can be estimated by the IRRS. Theoretically, Sb2O3 based glasses are expected to have lower phonon energies due to lower stretching vibration of Sb–O–Sb bond (605 cm1) [29,30]. To establish the low-phonon energy of KBS antimony glasses, the IRRS of the samples B2–B5 in the region 490– 1500 cm1 is measured and represented in Fig. 7. Two major reflection bands are observed to be centered at 602 and 1207 cm1 due to Sb–O–Sb stretching vibration [29,30] and B–O–B stretching vibration of [BO4] unit [27,28], respectively. It is observed that the intensity of the 602 cm1 band decreases with decrease of Sb2O3 content and the intensity of the 1207 cm1 band increases indicating an increase in borate content. As a consequence, the IRRS curves intersect and display an isosbestic point at 777 cm1. 3.2.5. UV–Vis absorption spectra The UV–Vis spectra (Fig. 8) show that the position of the cutoff wavelength or fundamental absorption edge shifts towards higher

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Table 2 FT-IRTS band position in raw materials, glass and glass-ceramics along with band assignment. Sample identity/band position (cm1)a KBO2xH2O

3462 (s,b)

1646 (s) 1423 (w) 1377 (w) 1231 (s) 1008 (w)

Sb2O3

Band assignment B2

b

b

b

B4

B5

B6

938 (s)

3554 (w) 3392 (s,b) 3277 (w) 1654 (m) 1431 (w) 1315 (m) 1231 (m) 1008 (m) 931 (m)

3554 (w) 3369 (s,b) 3277 (w) 1654 (m) 1431 (w) 1315 (s) 1238 (m) 1008 (s) 931 (s)

3354 (w) 3369 (s,b) 3277 (w) 1646 (s) 1423 (m) 1315 (s) 1238 (m) 1008 (s) 931 (s)

3554 (w) 3368 (s,b) 3277 (w) 1654 (s) 1323 (m) 1338 (s) 1238 (m) 1008 (s) 931 (m)

685 (sh) 600 (s)

685 (m) 592 (s)

685 (m) 592 (s)

685 (m) 600 (s)

685 (sh) 600 (s)

462 (m)

439 (s)

439 (s)

454 (s)

454 (w)

1639 (w)

1223 (s)

954 (m) 739 (s) 692 (m) 592 (m) 546 (m) 492 (s)

b

B3

3508 (w)

1615 (w)

b

Free hydroxyl group O–H (s-s) O–H (s-v) Hydrogen bonded OH (–O–Hd+Od–) (s-v) H–O–H (b-v) B–O–B in [BO3] (as-s) B–O–B in [BO3] (as-s) B-O in [BO3] (as-v) B–O–B in [BO4] (as-s) O–Sb–O (s-v) (senarmontite) O–Si–O (s-v) O–Sb–O (s-s) (senarmontite) O–Sb–O (s-s) (valentinite) O–Sb–O (as-s) (valentinite) O–Sb–O (s-b) (valentinite) O–Sb–O (as-b) (valentinite)

s = Strong, b = broad, w = weak, sh = shoulder, m = medium. s-s = Symmetric stretching vibration, as-s = asymmetric stretching vibration. s-b = Symmetric bending vibration, as-b = asymmetric bending vibration. s-v = Stretching vibration, b-v = bending vibration, as-v = asymmetric vibration. a Band positions are correct within the range ±1 cm1. b For composition see Table 1.

where a(m) = (1/d)ln(I0/It). Here ln(I0/It) is absorbance (A), where I0 and It are the intensities of incident and transmitted light, respectively and d is the thickness of the sample. B is the proportionality constant and hm is the photon energy of incident radiation. The absorbance (A) can be expressed as a function of measurable quantity percent transmittance (%T) [31]:

A ¼ logð100=%TÞ;

Fig. 6. Infrared transmission spectrum (FT-IRTS) of polished transparent (a) glass B2 (for composition see Table 1). Spectrum of (b) silica glass is shown for comparison (thicknesses of both the glasses: 1.0 mm).

wavelength with increase in KBO2 content. The UV–Vis spectra also show an isosbestic point at 393 nm. The UV–Visible spectra can also be used to evaluate the optical band gap Eopt of the materials. According to Davis and Mott theory the relationship between absorption coefficient a(m) and Eopt is given by [36]: n

aðmÞ ¼ Bðhm  Eopt Þ =hm; or;aðmÞhm ¼ Bðhm  Eopt Þn ;

ð4Þ ð5Þ

ð6Þ

where %T can directly be determined from the transmission spectra (Fig. 8). The nature of the transition is determined by superscript n which takes the value 2, 3, 1/2 and 1/3 corresponding to indirect allowed, indirect forbidden, direct allowed and direct forbidden optical transitions that are involved in the occurrence of the UV absorption edge when electromagnetic radiation interacts with the electron in the valence band. In our materials, the transition between HOMO (Sb 5s + O 2pp) and LUMO (Sb 5p) is allowed direct transitions that is, n = 1/2. Consequently the studied materials are characterized by direct optical band transitions. Fig. 9 represents the Tauc’s plot [(ahm)2 vs. hm] for different KBS (B2–B5) samples. The optical band gap energies can be determined from the Tauc’s plot by extrapolating the linear portion of the curve to intersect the photon energy axis at zero absorption. The intersection of the slope of the linear part of the plot (ahm)2 vs. hm gives the values of the Eopt (in eV). It is seen that the Eopt value decreases from 3.22 eV down to 3.15 eV with decreasing concentration of Sb2O3 (inset). 3.2.6. Density, softening point, glass transition temperature, coefficient of thermal expansion and dielectric constant Fig. 10(a) shows variation of experimental and theoretical density (q) with Sb2O3 content of the glasses. Fig. 10(b) and (c) illus-

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T. Som, B. Karmakar / Journal of Non-Crystalline Solids 356 (2010) 987–999 Table 3 Some of experimental and calculated properties of the glasses and glass-ceramics. Sample identitya/property value

Property

B3

B4

B5

B6

125 4.7650

133 4.5560

141 4.3467

150 4.0985

3.8658

Calculated properties Covalent character, % (±0.03) Theoretical optical basicity, Kth (±0.002) Optical basicity from refractive index, K(n) (±0.005) Optical basicity from optical band gap, K(Eopt) (±0.005) Average molecular weight, Mav (gm/mol) (±0.004) Molar volume, Vm (cm3/mol) (±0.01) Refractive index, n (±0.0002) Molar refraction, Rm (cm3/mol) (±0.005) Total molar polarizability, am (Å3) (±0.007) P Total molar polarizability of cations, am (Å3) 3 Polarizability of oxide ions, aO2 (Å ) Interaction parameter, F(n) Å3 (±0.004) Metallization criterion (M) (±0.001)

48.55 1.072 1.176 1.27 249.55 52.37 2.0006 26.20 10.40 1.942 3.3832 0.013 0.4998

46.50 1.035 1.092 1.26 228.60 50.17 1.9477 24.19 9.60 1.8006 2.8887 0.029 0.5178

44.49 0.994 1.059 1.25 207.64 47.77 1.8925 22.10 8.77 1.662 2.7338 0.037 0.5376

42.50 0.950 1.023 1.21 186.68 45.55 1.8349 20.09 7.97 1.552 2.5792 0.046 0.5590

40.55 0.901 0.970 165.72 42.87 1.7746 17.89 7.10 1.382 2.3825 0.059 0.5827

For composition see Table 1.

35

90

-1 (602 cm )

B2 B3 B4 B5

25 20

80 70

Transmission (%)

Reflectivity (%)

30

Isosbestic point (777 cm-1) -1 (1207 cm )

15 10

90 80 Isosbestic point

60 50 40 30

B2 B3 B4 B5

20 5 10 600

800

1000

1200

400

Fig. 7. Infrared reflection spectra (FT-IRRS) of glass and glass-ceramics: B2, B3, B4 and B5 (for composition see Table 1).

trates the trend of the softening point (Ts), dilatometric softening point (Td) and glass transition temperature (Tg) with Sb2O3 content of the glasses while Fig. 10(d) demonstrates that effect of increase in Sb2O3 content of the glasses on the coefficient of thermal expansion (CTE) and dielectric constant (e). It is seen that the experimental density is found to vary in the range 3.87–4.77 g cm3 (±0.04), experimental softening point (Ts) and glass transition temperature (Tg) vary in the ranges 331–392 °C (±2 °C) and 234–264 °C (±1 °C), respectively while CTE and (r) of the glasses and glass-ceramics is found to vary in the range 201–222 (±1)  107 K1 and 12.4–14.5 (±0.1), respectively. It is found that density (q) follows the additivity rule and increases linearly with increase in Sb2O3 content. Similar trend is also found to be shown by the theoretically predicted density (Fig. 10(a)). The theoretically calculated physical property P can be calculated by the semi-empirical additive mathematical formula as given below [37]:

P¼

i¼1

g i;P mi ni

N X i¼1

mi ni ;

70 (393 nm) 60 50 40 30 20 10 0 360 370 380 390 400 410 420 Wavelength (nm)

0

1400

-1

Wavenumber (cm )

N X

Transmission (%)

a

B2 Measured properties OH content (ppm), (±0.4) Density (g cm3), (±0.04)

ð7Þ

600 800 Wavelength (nm)

1000

Fig. 8. UV–Visible transmission spectra of some glasses and glass-ceramics (thickness: 1.0 mm, for composition see Table 1). Inset shows isosbestic point at 393 nm.

where i is the index of the oxide; N is the number of types of oxides forming the glass in question; mi is a molar fraction of the ith oxide; ni is the number of atoms in the formula of the ith oxide and gi,P is a partial coefficient for the ith oxide. The softening point (Ts), dilatometric softening point (Td) and glass transition temperature (Tg) gradually decreases as the Sb2O3 content increases, but the effect is more pronounced on Ts and Td than that on Tg. Similar decreasing trends are observed in the theoretically predicted softening point (Mg) and glass transition temperature (Tg) as well. These facts are depicted in Fig. 10(b) and(c), respectively. It is seen that the experimental softening point (Ts) and glass transition temperature (Tg) vary in the ranges 331– 392 °C and 234–264 °C, respectively and they obey very well the empirical relation:

T s ¼ 3T g =2:

ð8Þ

Similar results have been reported for other Sb2O3 containing glasses [22,38].

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T. Som, B. Karmakar / Journal of Non-Crystalline Solids 356 (2010) 987–999

3.3. Theoretically evaluated physical properties Electronegativity of the constituting ions determines the covalent character while the electronic polarizability is intermittently linked to refraction and optical basicity. Correlating material properties with electronic charge distribution, that is, optical basicity and covalent character provides a better understanding of the underlying materials chemistry or physics and encourages the opportunity to make useful predictions. Conventional oxide glasses are found to possess low refractive index, low polarizability, low optical basicity, large metallization criterion and small third-order non-linear susceptibility. On the other hand, heavy metal oxide glasses have high refractive index, high polarizability, high optical basicity, small metallization criterion and large third-order nonlinear susceptibility and find promising application in NLO devises, telecommunications and upconversion solid state lasers. Consequently, the calculation of the above parameters in this new series of glasses (KBS) is crucial.

Fig. 9. (ahm)2 vs. hm plot of samples B2, B3, B4 and B5 for evaluation of optical band gap, Eopt (for composition see Table 1). Inset shows variation of Eopt with Sb2O3 content (lines are drawn to guide the eye).

3.3.1. Covalent character The extent of covalent bonding character of the resultant samples (glasses and glass-ceramics) can be calculated approximately using the formula [39]:

Covalent character ð%Þ ¼ exp½0:25ðDvÞ2   100; The CTE values are found to increases with increase of Sb2O3 concentration. These values and trend of the results agree well with those obtained in different Sb2O3 systems [3,38]. The dielectric constant values are also found to increase with increase of Sb2O3 content. It is interesting to note that their trend is very similar to that of CTE.

ð9Þ

where Dv is the electronegativity of the composite that is, the electronegativity difference (vA  vC) of the anions and the cations. The average electronegativity of the anions (vA) or cations (vC) can be evaluated by the following simple additive relation [18–20,40]:

vA or vC ¼

X

N i vi =

X

Ni ;

ð10Þ

5.0

(a) 450

Experimental Theoretical

4.6

Temperature (°C)

4.4 4.2 4.0

400 350 300 250

3.8 40

50

60

70

40

80

Sb2O3 content (mol %)

60

70

80

15.0

(c)

224

(d) Theoretical Experimental

400 350 300 250 200

50

Sb2 O 3 content (mol %)

Dielectric constant (ε)

Glass transition temperature, Tg (°C)

450

Theoretical (Mg ) Experimental (T s) Experimental ( T d )

(b)

14.5

Experimental (ε)

220

14.0

Experimental (CTE)

216

13.5

212

13.0

208

12.5

204 200

12.0 40

50

60

70

80

Sb2 O3 content (mol %)

CTE (α × 10−7)

Density (g.cm -3)

4.8

40

50

60

70

80

Sb2 O3 content (mol %)

Fig. 10. Variation of (a) density, (b) softening points (Mg, Ts and Td), (c) glass transition temperature (Tg), and (d) dielectric constant (e) and coefficient of thermal expansion (CTE) as function of Sb2O3 content in KBS antimony glasses (lines are drawn to guide the eye). For graph (b) and (c) the errors are comparable to the symbol size.

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T. Som, B. Karmakar / Journal of Non-Crystalline Solids 356 (2010) 987–999

3.3.2. Optical basicity The optical basicity, as proposed by Duffy and Ingram [41], is used as parameter to determine the acid–base properties of the glass in terms of the electron density carried by oxygen. It represents the average electron donation capacity of the oxide(II) species in the medium after the polarization of their electron charge clouds by constituent cations. The theoretical (ideal) optical basicity (Kth) is calculated according to the expression [42,43]:

Kth ¼ XðSb2 O3 ÞKðSb2 O3 Þ þ XðB2 O3 ÞKðB2 O3 Þ þ XðK2 OÞKðK2 OÞ; ð11Þ where X(Sb2O3), X(B2O3) and X(K2O) are the equivalent fractions based on the proportion of oxygen, each oxide contributes to the overall glass stoichiometry and K(Sb2O3) = 1.18, K(B2O3) = 0.42 and K(K2O) = 1.4 are the basicities assigned to individual oxides [44]. The calculated theoretical optical basicity is shown in Table 3. Fig. 11(a)–(d) displays that the properties Ts and Tg decrease while CTE and Eopt increase with increase in covalent character and theoretical optical basicity (Kth) of the glasses and glassceramics. The e values also increases with increase of covalent character and theoretical optical basicity (graph not shown here).

This indicates that there must exist a direct relation between these properties with covalent character and optical basicity of the materials which seem to regulate these properties. An inherent relation is known to exist between electronic polarizability of the oxide ions (aO2 ) and optical basicity of the oxide materials [41]. Increased oxide ion polarizability means stronger electron donor ability of the oxide ions, that is, increased optical basicity. Thus, alternatively optical basicity can be estimated based on refractive index K(n) and on optical band gap K(Eopt) values as follows [42,43,45]:

  KðnÞ ¼ 1:67 1  1=aO2 ðnÞ ;   KðEopt Þ ¼ 1:67 1  1=aO2 ðEopt Þ ;

3.3.3. Electronic polarizability of ions The average molar refraction (Rm) was calculated following the Lorentz–Lorenz equation [31]:

Rm ¼ ½ðn2  1Þ=ðn2 þ 2Þ  ðV m Þ;

ð14Þ

where Vm is the molar volume (Vm = Mav/q), Mav is the average molecular weight, q is the density, n is the refractive index evaluated by the Priven-2000 method using Sci Glass software. The molar polarizability (am) is calculated as [31,42,43]:

am ¼ 3Rm =4pNA ;

ð15Þ

3.23

224

(d)

3.22

215 210

-7

3.20 3.18

205

220

3.16

3.22 CTE, α E opt

CTE, α (× 10 )

CTE, α E opt

E opt (eV)

CTE, α ( × 10 -7 )

220

ð13Þ

where aO2 ðnÞ and aO2 (Eopt) are electronic polarizability of the oxide ions based on refractive index and optical band gap, respectively. The calculated K(n) and K(Eopt) for the present glass system is presented in Table 3, is seen to correlate well with the trend and values of Kth.

225

(c)

ð12Þ

3.21

216

3.20 3.19

212

3.18

208

3.17

204

3.16

E opt (eV)

where Ni and vi are the number of individual constituent atom per mole and its electronegativity respectively. Here, we have considered the electronegativity in Pauling’s scale as vO = 3.5, vB = 2.0, vSb = 1.9 and vK = 0.8. The calculated covalent character is shown in Table 3. It is found to vary in the range 48.55–40.55% (±0.05%) and decreased with decrease in Sb2O3 content.

3.15

200

200

3.14 40 41 42 43 44 45 46 47 48 49 50

Covalent character (%)

3.14 0.92

0.96

1.00

1.04

1.08

Theoretical optical basicity (Λth)

Fig. 11. Variation of experimental softening point (Ts) and glass transition temperature (Tg) as functions of (a) covalent character and (b) theoretical optical basicity (Kth) respectively, and variation of experimental coefficient of thermal expansion (CTE, a) and optical band gap (Eopt) as a functions of (c) covalent character and (d) theoretical optical basicity (Kth) of KBS glasses respectively (lines are drawn to guide the eye).

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T. Som, B. Karmakar / Journal of Non-Crystalline Solids 356 (2010) 987–999

where NA is the Avagadro’s number. For a ternary oxide glass system xAmOn  yBpOq  zCrOs, total molar polarizability of the cations P act can be calculated as:

X

act ¼ xmaA þ ypaB þ zraC ;

ð16Þ

where z = (1  x  y) and aA, aB and aB are the polarizabilities of the cations in Å3 (aKþ ¼ 0:002, aB3þ ¼ 0:821, and aSb3þ ¼ 1:111 Å3) [43]. Here cation polarizabilities are considered equal to their free polarizabilities because the cationic charge will hold their electrons despite their large size and that their deformation is considerably less than those of the O2 ions. Consequently, aO2 ðnÞ and aO2 (Eopt) are calculated as [42,43]:

aO2 ðnÞ ¼ ðam 

X

h

act Þ=ðxn þ yq þ zsÞ;

aO2 ðEopt Þ ¼ V m =2:521  ðEopt =20Þ

1=2



The calculated values of n, Rm, Vm, am, are provided in Table 3.

X

i

ð17Þ

act =ðxn þ yq þ zsÞ:ð18Þ

P

act, aO2 (n) and aO2 (Eopt)

3.3.4. Interaction parameter The interaction parameter represents the interaction along an average cation–oxide ion pair and provides an alternative way to estimate the ionic-covalent chemical bond formation in the glass. The interaction parameter F(n) (as shown in Table 3) is calculated as:

  AðnÞ ¼ XðSb2 O3 Þ ðaf  aO2 ðnÞÞ=2ðaSbþ þ af ÞðaSb3þ þ aO2 ðnÞÞ þ XðB2 O3 Þ½ðaf  aO2 ðnÞÞ=2ðaB3þ þ af Þða3þ B þ aO2 ðnÞÞXðK2 OÞ½ðaf  aO2 ðnÞÞ=2ðaKþ þ af ÞðaKþ þ aO2 ðnÞÞ; ð19Þ 3 where a f is the theoretical free oxide ion polarizability (3.921 Å ) estimated by Pauling [46]. Table 3 shows that the calculated values of F(n) decreases with increasing Sb2O3 content.

3.3.5. Metallization criterion Optical non-linearity is caused by electronic polarization of a material upon exposure to intense radiation. So, electronic polarizability is one of the most important properties that determine the non-linear response of the material. The metallization criterion (M) can be calculated as [42,43]:

M ¼ 1  Rm =V m :

ð20Þ

The calculated M is listed in Table 3. It is seen that M increases with decrease in Sb2O3 content. 4. Discussion 4.1. Physical appearance The yellow color of Sb2O3 glasses are essentially due to host absorption caused by the transition between HOMO (Sb 5s + O 2pp) and LUMO (Sb 5p) [5]. A small quantity SiO2 is picked up in the KBS samples during their melting from the silica melting crucible. This fact is evidenced in the qualitative energy dispersive X-ray (EDX) spectrum as shown in Fig. 2. 4.2. Physical properties 4.2.1. X-ray diffraction The XRD spectrum (Fig. 4, curve h) indicates that both valentinite and senarmontite forms of Sb2O3 are present in the starting raw material. For samples B2 and B3 (curves b and c), presence of the hump and absence of any peaks in samples B2 and B3 indicates that they are amorphous in character and are glasses while

the development of peaks in addition to the hump in samples B1, B4, B5 and B6 indicates their glass-ceramic nature. The decreasing broadness and the increasing intensity of the peaks as one proceeds from B4 to B6 indicate intense crystallization with decrease in Sb2O3 content. Although the XRD spectrum of B1 (curve a) noisy but the inset of Fig. 4 shows small peaks at 2h = 28.62° (d = 3.127 Å) and 38.92° (d = 2.311 Å) due to valentinite and Sb2O5 nano crystals (JCPDS card file No. 11-689 and 34-878, respectively) and are quite discernable to assign B1 glass-ceramic nature. Since no special method has been adopted for synthesis of nano glass-ceramic hence such intense crystallization leading to spontaneous glass-ceramic formation are probably caused by the large difference in field strengths of Sb3+ (F = 0.73), B3+ (F = 1.34) and K+ (F = 0.13) ions [7], when the natural and conditional glass former (B2O3 and Sb2O3, respectively) present in almost equal amounts. Orthorhombic Sb2O3 crystallizes in the space group Pccn (D10 2h ) and is built up of infinite polymeric Sb–O–Sb chains running along the c axis with Sb–O distances of 20 Å. [47]. The XRD results clearly support valentinite as the dominant phase with small amount of senarmontite and cervantite in glass-ceramics (B4– B7). In addition, there is evidence of presence of little crystalline potassium antimony hydroxide. The presence of sharp peaks and absence of any hump in sample B7 indicates that it is entirely crystalline in nature. 4.2.2. Infrared transmission spectra of KBr pellets (FT-IRTS, 400– 4000 cm1) The properties of a material is determined by its structure which is turn is regulated by the nature of chemical bonds and electronic properties. The Fourier transform IR transmission (FTIRTS) spectra, particularly in the fingerprint region (400– 1500 cm1), provide much information about structure (short range order) consequently the powdered samples were studied following the KBr pellet method and compared with those of raw materials. Within the KBS antimony samples, B2–B6 (Fig. 6, curves c–g), the sharp peak at 592 cm1 and the small peak at 685 cm1 are recognized as Sb–O–Sb asymmetric and symmetric stretching vibrations, respectively emphasizing the closeness of the glass structure to valentinite Sb2O3 [5,29,30]. The small peak at 931 cm1 probably arises as a consequence of Sb–O stretching vibrations of the senarmontite variety present in minor amount in the antimony glasses [5,10,30]. This 931 cm1 peak position also coincides with Si–O stretching vibrations of SiO4 tetrahedra which were contaminated from silica crucibles during the melting process. This is supported by the EDX spectrum of glass B5 (see Fig. 2). The peak at 1231 cm1 arises due to B–O stretching vibrations of (BO3)3 unit in metaborate chains [27,28]. The region between 1000–1200 and 1200–1500 cm1 is generally attributed to asymmetric stretching vibrations of the B–O–B bond of tetragonal [BO4] and trigonal [BO3] unit respectively. Presence of small peaks between 1000 and 1500 cm1 in the FT-IRTS spectra of B3 (curve d) suggest the presence of incipient crystallization in it. The intensity and number of peaks in this region 1000–1500 cm1 increases as we proceed down the series from B4 to B6 (curves e–g) indicating the presence of intense crystallization as the quantity of KBO2 increases. This is supported by the XRD spectra of the samples. The peak at 1630 cm1 is endorsed as bending vibrations of O– H group of H2O molecules. The broad band around 3392 cm1 corresponds to the stretching vibration of O–H group. While the small peak at 3277 cm1 is due to hydrogen bonded OH (–O–Hd+Od–) and the small peak at 3554 cm1 corresponds to free O–H group to the glass network [31]. These suggest the presence of water in the form of hydroxyl groups in the glass and glass-ceramics. Thus the study shows antimony based glass structure resembles valentinite form of Sb2O3 having SbO3 pyramids with the coordina-

T. Som, B. Karmakar / Journal of Non-Crystalline Solids 356 (2010) 987–999

tion number of Sb as 3. Oxygen atoms are present in three corners (Sb–O bond distance 2.0 Å) and the lone pairs of electrons of Sb3+ at the fourth corner [7]. The coordination polyhedra are joined by sharing corners to form double infinite chains with the lone pair pointing out from the chains. Such chains are held together by weak secondary Sb–O bonds with lengths greater than 2.6 Å. The third oxygen in each SbO3 unit takes part in Sb–O–B type bond formation [7,48]. The presence of common meta-centers and B–O–B vibrational bands in the IR spectra of the glasses supports the statement. Thus the structure essentially consists of long chains of entangled SbO3 units. The introduction of the glass modifier K2O probably breaks some linkages (B–O–B bonds, Sb–O–B bonds) and thereby the local symmetry of the glass network. This leads to the formation of non-bridging oxygen ions (NBOs) and introduction of dangling bonds (coordination defects). Earlier structural investigations of Sb2O3–B2O3 systems have shown that a small fraction of the Sb3+ ions are converted to Sb5+ ions, reaching a value of 15% of antimony in the higher oxidation state (Sb5+) for a glass containing 70 mol% Sb2O3. These Sb5+ ions enter into the glass as singly positive [SbO4] 4-coordinated units. The presence of positive [SbO4] would provide charge balance for the negative [BO4] units [5,21–24]. It is worth noting that nevertheless Sb3+ (as Sb2O3) is used as the raw material but a small part has been transformed into Sb5+ state during melting at high temperature (900 °C) by accepting oxygen from air due to its multivalent characteristics which is shown as follows:

Sb2 O3 þ O2 ¼ Sb2 O5 :

ð21Þ

The existence of Sb2O5 crystals has been confirmed by XRD analysis. 4.2.3. Infrared transmission spectra of bulk samples (FT-IRTS, 1000– 4000 cm1) Transparency in the longer wavelength (IR) region is restricted by the IR absorption edge because the propagating light becomes resonant with the fundamental and overtone vibrational frequencies of the molecules constituting the glass structure. The fundamental vibration frequency, m (in cm1) of a linear harmonic oscillator can be expressed by the Szigetti relation as follows [49]:

m ¼ ðk=4p2 c2 lÞ0:5 ;

ð22Þ

where (k) is force constant, (c) is speed of light and (l) is the reduced mass. The l can be defined as: 1 l ¼ m1 1 þ m2 ;

ð23Þ

where m1 and m2 are the masses of the two bond forming atoms or ions (e.g. m1 = Sb3+ and m2 = O2). Thus the IR absorption edge shifts to longer wavelengths for heavier atomic masses or due to reduced force constant (k) of the bond by the weak Columbic interaction. 4.2.4. Infrared reflection spectra (FT-IRRS) The phonon energy, thus, corresponds to the highest intensity stretching vibration bond of the network forming components of the glass, for example Si–O–Si bond (1060–1150 cm1) in silica and silicate glasses, Sb–O–Sb bond (605 cm1) in antimony oxide glasses, etc. [50]. Phonons can provide non-radiative decay pathways to suppress radiative upconversion luminescence of rareearth (RE) ions. Thus, glasses with lower phonon energy are desirable to reduce the multiphonon relaxation and non-radiative loss and to obtain high upconversion efficiency. It is reported that low hx of the host can be obtained from the phonon side band (PSB) spectra because the vibration of PSB originates from the local vibration around the RE ions [51]. Although IRRS represents the total vibrations of the whole matrix but it provides an alternative way to measure the hx of the matrix [50]. The average hx values of fluoride glasses obtained by PSB spectra have been found to be

997

very close to those obtained from the low frequency stretching vibrations of IRRS spectra [52]. In addition there are adequate literature reports that documents reflection spectra in the infrared (IR) region of 400–1200 cm1 were measured to establish the phonon energies of the glassy host where the wavenumber at the main peak was cited as the phonon energy of the glass matrix [52,53]. Similarly, we have denoted the main and highest intensity Sb–O– Sb stretching band at 602 cm1 as the phonon energy of the KBS antimony glass. The observation of isosbestic point at 777 cm1 is a good evidence of presence of antimony and borate as the two principal species. It manifests an existence of equilibrium between the antimony and boron containing species. 4.2.5. UV–Vis absorption spectra This red-shift of the fundamental absorption edge can be understood by the increasing influence of non-bridging oxygen ions (NBOs) with increasing K2O content [6]. The presence of isosbestic point at 393 nm indicates a decrease in transmittance of the glasses containing higher concentration of Sb2O3. This is because the highly polarisable Sb3+ cation with a lone pair of electrons causes high polarization the O2 anion which in turn increases the total electronic polarizability. Consequently the molar refraction and refractive indices increases causing a greater loss of transmittance due to scattering. The decrease in value of Eopt from 3.22 eV down to 3.15 eV with decreasing concentration of Sb2O3 can be interpreted in terms of structural changes in the glass system. In the present ternary system at high Sb2O3 content, both Sb2O3 and B2O3 acts as network formers while K2O as network modifier. The incorporation of network modifiers increases the quantity of non-bridging oxygens (NBOs). Since NBOs are more easily excited than bridging oxygen, Eopt decreases with addition of K2O and removal of Sb2O3. Thus these glasses are wide gap materials. 4.2.6. Density, softening point, glass transition temperature and coefficient of thermal expansion The increase in density of the KBS samples is obviously due to the increase in proportion of the much higher mass molar weight of Sb2O3 than K2O or B2O3. The experimental density is less than the theoretical ones and their departure increases with increase in Sb2O3 content manifesting the volatilization of Sb2O3 during the melting process. The close correlation between the experimental and theoretically calculated (predicted) softening point and glass transition temperature clearly indicates that these properties are additive in nature with respect to the constituting chemical components (see Eq. (7)). It must be mentioned here that both Sb3+ and K+ have considerably less field strength (0.73 and 0.13, respectively) than B3+ (1.34) [7]. Consequently they play a major role of decreasing the softening point and glass transition temperatures. It is, thus, perceived in the synthesized samples. It is well-known that covalent bonds are less stronger than ionic bonds. Hence, generally, the compounds with a more covalent bonds or covalent character are weaker than those with a less covalent character. Thus, the softening points (Ts, Td, and Mg) and glass transition temperatures (Tg) decrease with a decrease in the more covalent Sb2O3 content However, one should also consider the influence of structural changes on these properties as well. Generally the thermal expansion increases as bond strength reduces. Since the coefficient of thermal expansion (CTE) decreases with decrease of Sb2O3 concentration, this observation indicates that Sb–O bonds are weaker than B–O bonds. In order to form a strong covalent bond, the participating orbitals of both the bond forming atoms should possess similar energy and size. Besides the orbitals should have same symmetry with respect to the bond axis such that the electron clouds overlap to a considerable extent.

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Generally network bond strengths and glass forming-regions of HMO glasses are small compared to conventional network formers. This is justified by the fact that heavy metal cation–oxygen bond strengths are relatively weaker, give rise to low fundamental vibration frequencies and thereby large IR transmission spectra. Another evidence is that in vitreous state B2O3 has interionic distance of 1.47 Å while in Sb2O3 it is about 2.0 Å [1,7]. Also in Sb2O3 glasses, the SbO3 polyhedral chains are interlinked by weak secondary Sb–O bonds having lengths greater than 2.6 Å. In fact, the bond strengths of Sb–O and B–O are 89 ± 20 and 192.7 ± 1.2 kcal mol1 [54]. The dielectric properties of materials are the inherent effect associated to the mechanism of polarization of the permanent and induced electrically charges by an external applied electric field. Various factors contribute to the dielectric constant of glasses. Among them are electronic, ionic, dipolar and space charge polarizations of the constituting ionic network. In the present study, the as-prepared KBS antimony glasses exhibits a relatively higher value of dielectric constant (e) than the vitreous silica (e = 3.8), soda-lime silicate (e = 7.2) and borosilicate glasses (e = 4.1–4.9) [55] and the magnitude of (e) shows a steady increase with an increase Sb2O3 content. This is essentially due to the spontaneous high ionic polarization of Sb3+ ions (aSb3þ ¼ 1:111 Å3) [43] under an applied electric field. 4.3. Theoretically evaluated physical properties The electronegativity and electronic polarizability of the constituting ions are the most important material properties that regulate their applicability in the fields of optics and electronics. Dimitrov and Komatshu [56] had established that the third-order non-linear susceptibility of glasses (v3) increases with increasing optical basicity and metallization of the glasses. This was correlated to the high electron donor capacity of the oxide ions and small optical band gap. In this paper, our data emphasizes Ts and Tg reduces while CTE and Eopt increase with increase in covalent character and optical basicity. This can also be associated to the high electron donor capacity of the oxide ions. The introduction of modifiers or glass formers like B2O3 generally leads to a small increase in the average oxide ion polarizability and refractive index. On the contrary, addition, addition of heavy metal oxides Sb2O3 causes large increase in oxide ion polarizability and refractive index. The oxide ion polarizability data obtained provide evidence for the additive nature of the oxide ion polarizabilities in glass systems. B3+ cation (1s2) possesses low polarizability (aB3þ ¼ 0:002) and high field strength value and strongly affects O2 electron cloud. Consequently O2 polarizability is low in B2O3 (1.345 Å3) [57] and glasses with high B2O3 content show P low n, Rm, am, act, aO2 ðnÞ values. Sb3+ possesses large ionic radii (0.89 Å), high aSb3þ and low field strength value and also possesses a highly polarizable lone pair in the valence shell (5s2). As a result the polarizing power of Sb3+ cation on the O2 ions is reduced and the O2 ions in polarizability is Sb2O3 is very large (3.429 Å3) [57]. Consequently glasses containing large amount of Sb2O3 possesses P high n, Rm, am, act, aO2 ðnÞ values. It is seen that optical basicity increases with increase in Sb2O3 content. This may be understood according to the relations (11) and (12). These equations show that the increase in polarizability, basicity also increases (Table 3). Secondly, the optical band gap (Eopt) in the present system is also P act found to increase with the increase in polarizabilities, am, and aO2 ðnÞ (see Fig. 9). Covalent character can be treated analogous to the optical basicity. The greater the polarizibility of the oxide ion, the higher is the covalency. Again, small value of interaction parameter, F, certifies weak interionic interaction resulting in large average unshared electron density available for donation that is, increased covalency and

optical basicity. It is well established that the third-order non-linear susceptibility (v3) value increases and M-value decreases. This is because when (1  Rm/Vm) becomes zero the transition to the metal states takes place. In other words, the small metallization criterion means that the width of both valence and conduction bands become large, resulting in narrow band gap and increased tendency for metallization of the glass. Therefore it can be predicted that KBS glasses and glass-ceramics possess high NLO properties. 5. Conclusions In summary, a series of new monolithic antimony glasses and nano glass-ceramics in the K2O–B2O3–Sb2O3 (KBS) system have been prepared by the simple melt-quench technique and characterized with respect to structure (short range order) and various properties. X-ray analysis shows that the size of the nano crystallites of the glass-ceramics varies in the range 7–16 nm. FESEM images reveal that a typical antimony nano glass-ceramic has granular, cell-wall like and card-in-house microstructures. Infrared spectra and X-ray analyses establish that their structure closely resembles valentinite form of Sb2O3. Infrared reflection spectra establish that these samples have low-phonon energy around 592–602 cm1. Low-phonon, high density (molecular weight) and high refractive index indicate that these monolithic glasses and glass-ceramics are potentially very good candidates for various photonic applications. These antimony based samples are found to have very low softening points (331–392 °C) and glass transition temperatures (234–264 °C), high dielectric constants (12.4–14.5) and remarkably high coefficient of thermal expansion values (201–222  107 K1). This shows the possibility of making them superior alternatives to lead based glasses for low melting applications like sealing or coating frits for electronic components and other opto-electronics applications. The optical band gap (Eopt) varies between 3.15 and 3.22 eV. This indicates that the samples function as wide gap materials. Further theoretical calculations of average oxide ion electronic polarizability, optical basicity, interaction parameter and metallization criterion (M) of the synthesized samples anticipated from the refractive index and optical band gap values endorse that K2O–B2O3–Sb2O3 glasses and glass-ceramics are potential candidates as non-linear optical devices as well. All the properties are found to be controlled by the fundamental properties such as covalent character and optical basicity of the glasses and glass-ceramics. Considering the various properties exhibited by the KBS glasses and nano glass-ceramics investigated in this study, it is reasonable to conclude that they are a new class of versatile materials. Acknowledgements TS would like to express her sincere gratitude for the financial support of the Council of Scientific and Industrial Research (CSIR), New Delhi in the form of NET-SRF under sanction number 31/ 015(0060)/2007-EMR-1. The authors thank Dr. H.S. Maiti, Director of the institute for his permission to publish this work. They also thankfully acknowledge the technical supports provided by the infrastructural facility (X-ray and Electron Microscopy Divisions) of this institute. References [1] W.H. Dumbaugh, J.C. Lapp, J. Am. Ceram. Soc. 75 (1992) 2315. [2] D. Lezal, J. Pedlikova, P. Kostka, J. Bludska, M. Poulain, J. Zavadil, J. Non-Cryst. Solids 284 (2001) 288. [3] J.A. Ruller, J.E. Shelby, Phys. Chem. Glasses 33 (1992) 177. [4] W.H. Heden, B.W. King, J. Am. Ceram. Soc. 3 (1967) 387.

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