Some measured and calculated properties of RE3+:LiTaO3 precursor glass
Ferroelectric bulk crystals are widely used in optoelectronic devices because of their well combination of extraordinary optical and electronic properties. Their crystal structure is non-centrosymmetric and due to this structural anisotropy they exhibit many nonlinear optical properties, for example, the electro-optic effect (change in optical index with electric field), harmonic generation (changing frequency of light), and photorefraction (index change in response to light), to name a few. However, preparation of their defect-free optical quality transparent single crystal is very difficult, lengthy process, and requires sophisticated costly equipment. In recent past, to triumph over these difficulties, much attention has been paid for development of transparent ferroelectric glass-ceramics by the high speed glass technology route because of its low cost of fabrication, tailoring of properties and flexibility to give desired shapes. Lithium tantalate (LiTaO3, LT) and potassium niobate (KNbO3, KN) single crystals are the most important lead-free ferroelectric materials with the A1+B5+O3 type perovskite structure concerning the environmental friendliness. LT has the rhombohedral crystal structure with crystal symmetry class 3m (unit cell dimensions: a = 5.1530 Å and c = 13.755 Å), large nonlinear constant (d33 = 13.6 pm/V at 1064 nm), second harmonic generation (SHG) coefficient (= 40.0 with respect to KDP at 1060 nm) (Risk et al., 2003, JCPDS No. 29-0836, Moses, 1978) and Curie temperature (660°C). In contrast, KN has the orthorhombic crystal structure with crystal symmetry class mm2 (unit cell dimensions: a = 5.6896 Å, b = 3.9693 Å and c = 5.7256 Å), large nonlinear coefficient (d33 = 27.4 pm/V at 1064 nm) [Moses, 1978] and Curie temperature (435°C). Thus, they exhibit unique electro-optic, piezoelectric, acousto-optic, and nonlinear optical (NLO) properties when doped with rare-earth (RE) [4f1-13] elements combined with good mechanical and chemical stability (Abedin et al., 1997, Zhu et al., 1995, Mizuuchi et al., 1995, Zgonik et al., 1993, Xue et al., 1998). Very recently, potassium niobate ceramics were investigated with an aim to develop environmental friendly lead-free piezoelectric and nonlinear materials (Ringgaard & Wurlitzer, 2005).
The electronic structure of each trivalent RE element consists of partially filled 4f subshell, and outer 5s2 and 5p6 subshell. With increasing nuclear charge electrons enter into the underlying 4f subshell rather than the external 5d subshell. Since the filled 5s2 and 5p6 subshells screen the 4f electrons, the RE elements have very similar chemical properties. The screening of the partially filled 4f subshells, by the outer closed 5s2 and 5p6 subshell, also gives rise to sharp emission spectra independent of the host materials. The intra-subshell transitions of 4f electrons lead to narrow absorption peaks in the ultra-violet, visible, and near-infrared regions.
In this chapter, we report synthesis, structure, properties and application of transparent ferroelectric LiTaO3 (LT) and KNbO3 (KN) nanostructured glass-ceramics. They were prepared by controlled volume (bulk) crystallization of their precursor glasses with and without RE dopant. The crystallization processes were studied by differential thermal analysis (DTA), X-ray diffraction (XRD), field emission scanning electron microscopy (FESEM), transmission electron microscopy (TEM), Fourier transform infrared reflection spectra (FT-IRRS), fluorescence and excited state lifetime analyses and dielectric constant measurement. The X-ray diffraction (XRD) patterns, selected area electron diffraction (SAED) and transmission electron microscopic (TEM) images confirm crystallization of LiTaO3 and KNbO3 nanocrystals in the transparent glass-ceramics.
2. Experimental procedure
2.1. Precursor glass and glass-ceramics preparation
The LT precursor glasses having molar composition 25.53Li2O-21.53Ta2O5-35.29SiO2-17.65Al2O3 (LTSA) doped with RE ions (0.5 wt% oxides of Eu3+ and Nd3+ in excess) or undoped were prepared by the melt-quench technique. The melting of thoroughly-mixed batches was done at 1600°C. The quenched glass blocks were annealed at 600°C for 4 h to remove the internal stresses of the glass and then slowly cooled down (@ 1°C/min) to room temperature. The annealed glass blocks were cut into desired dimensions and optically polished for ceramization and to perform different measurements. The crystallization was carried out at 680°C in between 0-100 h duration.
The KN precursor glasses having composition (mol%) 25K2O-25Nb2O5-50SiO2 (KNS) doped with Er2O3 (0.5 wt% in excess) or undoped were prepared similarly as mentioned above by the melt-quench technique. The well-mixed raw materials were melted in a platinum crucible in an electric furnace at 1550°C and the quenched glasses were annealed at 600°C to remove the internal stresses of these precursor glasses. They were transformed into nanostructured transparent glass-ceramics by heat-treatment at 800°C in between 0-200 h duration.
The density of precursor glasses was measured using Archimedes principle using water as buoyancy liquid. The refractive indices of precursor glass and representative glass-ceramics (d) were measured either on a Pulfrich refractometer (Model PR2, CARL ZEISS, Jena, Germany) at wavelength (λe = 546.1 nm) or on a Metricon 2010/M Prism Coupler at different wavelength (λ = 473, 532, 633, 1064 and 1552 nm). Differential thermal analysis (DTA) of precursor glass powder was carried out up to 1000°C at the rate of 10°C/min with a SETARAM TG/DTA 92 or with a Netzsch STA 409 C/CD instrument from room temperature to 900°C at a heating rate of 10°C/min. to ascertain the glass transition temperature (Tg) and the crystallization peak temperature (TP). XRD data were recorded using a PANalytical X’Pert-PRO MPD diffractometer operating with CuKα = 1.5406 Å radiation to identify the developed crystalline phases. The data were collected in the 2θ range from 10° to 80° with a step size of 0.05°.
A high resolution FE-SEM (Model: Gemini Zeiss SupraTM 35 VP, Carl Zeiss) was used to observe the microstructure of freshly fractured surfaces of the heat-treated nano glass-ceramics after etching in 1% HF solution for 2 minutes and coated with a thin carbon film. The TEM images and selected area electron diffraction (SAED) of powdered glass-ceramic sample were obtained from FEI (Model: Tecnai G2 30ST, FEI Company) instrument. The FTIR reflectance spectra of all the glasses and glass-ceramics were recorded using a FTIR spectrometer (Model: 1615, Perkin-Elmer) in the wavenumber range 400-2000 cm-1 with a spectral resolution of ±2 cm-1 and at 15° angle of incidence. Optical absorption spectra were recorded on UV-Vis-NIR spectrophotometer (Model: Lambda 20, Perkin-Elmer) at room temperature. The UV-Vis fluorescence emission and excitation spectra of Eu3+ doped precursor glass and nano glass-ceramics were measured on a fluorimeter (Model:Fluorolog-II, SPEX) with 150 W Xe lamp as a source of excitation. The fluorescence decay curves were recorded on the same instrument attached with SPEX 1934D phosphorimeter using pulsed Xe lamp. On the other hand, the fluorescence emission and excitation spectra of rest of samples were measured on continuous bench top modular spectrofluorimeter (QuantaMaster, Photon Technology International) attached with gated Hamamatsu NIR PMT (P1.7R) as detector and Xe arc lamp as excitation source. The excited state lifetime was measured with the same instrument using a Xe flash lamp of 75 W. The dielectric constants of precursor glass and nano glass-ceramics were measured at room temperature using a Hioki LCR meter (Model: 3532-50 Hitester, Hioki) at 1 MHz frequency after coating the surfaces with a conductive silver paint followed by drying at 140°C for 1h. Second harmonic generations (SHG) at 532 nm in the undoped glass-ceramics have been realized under fundamental beam of Nd3+:YAG laser source (1064 nm). The input energy of Nd3+: YAG laser was fixed at 17 mJ. The input energy of laser was divided in two directions (50% energy in each direction) using reflecting neutral density filter. In one direction KDP was put for reference. The reference SHG signal was measured using photodiode. Second beam was passed through visible filter (which blocks all visible wavelengths but pass 1064 nm) and focused onto the test samples. The SHG generated from the sample was focused onto a second harmonic separator, which reflects 532 nm at 45° and transmit 1064 nm. The SHG signal reflected from SHG separator passed through IR filter was finally measured using PMT. The reference signals from photodiode and from PMT were measured simultaneously using Lecroy oscilloscope (bandwidth 1GHz).
3. Nanostructured LiTaO3 ferroelectric glass-ceramics
Lithium tantalate (LiTaO3, LT) single crystal is one of the most important lead-free ferroelectric materials in the A1+B5+O3 type perovskite family. The correlation of property alteration of LT single crystals, powders, thin films, glass-ceramics, etc. with processing parameters is an important area of exploration. In recent times researchers have demonstrated the property monitoring based on preparation of LiTaO3 powders (Zheng et al., 2009) and thin films (Cheng et al., 2005, Youssef et al., 2008) by different methods. Luminescence properties of Ho3+, Eu3+, Tb3+ etc. doped LiTaO3 crystals, an another important area of exploration, which have also been investigated by various researchers (Sokólska, 2002, Sokólska et al., 2001, Gasparotto et al., 2008, Gruber et. al., 2006). Rare-earth (RE) doped transparent LiTaO3 nanocrystallite containing glass ceramics, in which RE ions selectively incorporated into the LiTaO3 nanocrystals embedded in an oxide glassy matrix, can offer excellent luminescent properties due to the low phonon energy environment of LiTaO3 nanocrystallites for luminescent ions, and good mechanical and chemical properties of oxide glassy matrix. This ability, combined with inherent nonlinear optical (NLO) properties of ferroelectric crystals, could offer a possibility to design self frequency doubling laser sources. Hence, this new material has attracted great attention in the continuous research for the development of novel optoelectronic devices (Jain, 2004, Romanowski et al., 2000, Hase et al., 1996). Mukherjee and Varma have reported the crystallization and physical properties of LiTaO3 in a LiBO2-Ta2O5 reactive glass matrix, however, they have not explored RE doped LiTaO3 containing glass-ceramics (Mukherjee & Varma, 2004). As such, work performed on nanocrystalline LiTaO3 containing aluminosilicate glass-matrix materials is very rare due to the difficulties in preparation of transparent precursor glass in general and glass–ceramics in particular which involves high temperature (about 1600°C) for its precursor glass melting (Ito et al., 1978). For this reason, the structure, dielectric and fluorescence properties of Eu3+, Nd3+ and Er3+ ion doped transparent precursor glass and glass-ceramic composites of LiTaO3 with heat-treatment time have been studied and reported elaborately by Tarafder et al., 2009 & 2010, Tarafder et al., DOI:10.1111/j.1744-7402.2010.02494.x. Second harmonic generation (SHG) from bulk LiTaO3 glass-ceramics has also been studied (Tarafder et al., 2011). For better understanding, the structure, dielectric and fluorescence properties of Eu3+ and Nd3+ ion doped transparent precursor glass and glass-ceramic composites of LiTaO3 with heat-treatment time have been reported elaborately along with the second harmonic generation (SHG) from bulk LiTaO3 glass-ceramics.
4. Results and discussion
4.1. Differential thermal analysis (DTA)
The DTA curve of the Eu3+ doped precursor glass is shown in Fig. 1. This exhibits the inflection in the temperature range 680-715°C followed by the intense exothermic peak at 821°C (Tp) corresponding to the LiTaO3 crystallization. The glass transition temperature (Tg) has been estimated to be 696°C from the point of intersection of the tangents drawn at the slope change as is marked in Fig. 1.
4.2. Refractive index
The Eu3+ doped precursor LTSA glass samples were heat treated at 680°C near glass transition temperature for various heat-treatment durations (0, 1, 3, 5, 7, and 10 h) after nucleating at 650°C for 2 h. Similarly, the Nd3+ doped precursor LTSA glass samples were heat treated at 680°C for 0, 3, 5, 10, 20, 50 and 100 h and were labeled as a, b, c, d, e, f and g. The Nd3+ doped precursor glass and nano glass-ceramics are presented in Fig. 2. From the measured glass density (ρ) and refractive index (ne) at wavelength λe = 546.1 nm, other related optical properties of Eu3+ doped precursor glass have been determined using relevant expressions and the same is presented in Table 1. Fig. 3 present Cauchy fitting based on measured refractive indices at five different wavelengths (see experimental procedure) and shows the dependences of the refractive index on the wavelength for Nd3+ doped precursor glass (a) and representative heat-treated glass-ceramics samples. In general, refractive index decreases with increasing wavelength due to dispersion. In addition to this, the refractive index of the glass-ceramics samples has increased in comparison with precursor glass (a) that can be seen in Fig. 3. The refractive indices nF, nD and nC have been estimated at three standard wavelengths (λF = 486.1 nm, λD = 589.2 nm and λC = 656.3 nm respectively) from the dispersion curve (Figs. 3, curve a). Similarly, from the measured glass density (ρ) and refractive index (nD) at wavelength λD = 589.2 nm, other related optical properties of Nd3+ doped precursor glass have also been determined and the results are presented in Table 1. From Table 1, it is clear that the LTSA glass under study has high values of refractive index and density. The large refractive indices of this glass are due to high ionic refraction (23.4) of Ta5+ ions (Volf, 1984) having an empty or unfilled d-orbital (outer electronic configuration: 5d06s0) which contributes strongly to the linear and nonlinear polarizability (Yamane & Asahara, 2000). The high density of the glass has originated from the large packing effect of Ta in the glass matrix (Hirayama & Berg, 1963). For the same reason, this glass possesses a high value of molar refractivity and electronic polarizability. Due to formation of high refractive index LiTaO3 (RI = 2.1834 at 600 nm (Lynch, 1975)), the heat-treated sample exhibit higher refractive indices as shown in Fig. 3, curve-d.
4.3. X-ray diffraction analysis
The X-ray diffractogram of Eu3+ doped precursor LTSA glass and cerammed glass-ceramics are shown in Fig. 4. The XRD pattern of the precursor glass exhibits broad humps characterizing its amorphous structure. With progression of heat-treatment, several diffraction peaks have been appeared in the glass-ceramics. From the analysis of these peaks it has been concluded that these peaks are attributed to rhombohedral LiTaO3 (JCPDS Card
|Eu3+:LiTaO3 doped precursor glass||Nd3+:LiTaO3 doped precursor glass|
|Average molecular weight, Mav||142.47||142.37|
|Density, ρ (g.cm-3)||4.54||4.50|
|Refractive index||ne (at 546.1 nm)||nF (at 486.1 nm)|
nD (at 589.2 nm)
nC (at 656.3 nm)
|Molar refractivity, RM (cm3)||13.23||13.39|
|Electronic polarizability, α (cm3)||1.84×10-21||1.79 x 10-21|
|RE3+ ion concentration, NRE3+ (ions/cm3)||5.71×1019||5.66 x 1019|
|Glass transition temperature, Tg (°C)||696||702|
|Crystallization peak, Tp (°C)||821||820|
File No. 29-0836) except a few diffraction peak around 2θ = 23.0°, 25.5°, 44.5° and 47.0° which are due to the formation of β-spodumene (LiAlSi2O6) crystal phase (JCPDS Card File No. 35-0797) in minor quantity. It is clearly evidenced from the XRD analysis that the peak of LiAlSi2O6 (2θ = 25.5°) is more prominent in sample of 5 h heat-treatment and it got diminished with respect to LiTaO3 phase in longer heat-treated glass-ceramics, indicating the stabilization of LiTaO3 nanocrystallites with increase in heat-treatment duration. From the full width at half maximum (FWHM) of the most intense diffraction peak (012) of LiTaO3, the average crystallite size (diameter,
where λ is the wavelength of X-ray radiation (CuKα = 1.5406Å), β is the full width at half maximum (FWHM) of the peak at 2θ. The average crystallite size of each RE doped heat-treated glass-ceramics found to increase with heat-treatment duration.
4.4. FESEM and TEM image analyses
The morphology and LiTaO3 crystallite size of Eu3+ and Nd3+ doped nano glass-ceramics have been examined by FESEM and TEM image analyses. FESEM images of the fractured surface of Nd3+ doped nano glass-ceramics have been presented in Figs. 5(a) -(b). The Nd3+ doped glass-ceramics 5(a) is obtained by heat-treating the precursor glasses at 680°C for 5 h. Similarly, the glass-ceramics 5(b) is obtained by heat-treating the precursor glasses at 680°C for 20 h. From the FESEM micrographs, it is clearly observed that the glassy matrix of the heat-treated samples initially phase separated on nanometric scale followed by incipient precipitation of defined crystallites within the Li–Ta rich phase regions with increase in heat-treatment time. The droplets have irregular shapes and dispersed uniformly thought out the bulk glass matrix. The size of the droplets varies in the range 20-60 nm. TEM image of the Eu3+ doped 10 h heat-treated glass-ceramics (f) has been presented in Fig. 6(a). The SAED pattern of the observed crystalline phase is presented in Fig. 6(b). From this image, it is observed that many spheroidal LiTaO3 crystallites precipitated homogeneously from the glass matrix and remained homogeneously dispersed in the residual glass matrix. The crystallite size from this TEM image of sample f found to be around 18 nm. The presence of fine spherical rings around the central bright region in SAED pattern discloses the existence of LiTaO3 nanocrystallites in the glassy matrix.
4.5. Fourier transform infrared reflectance spectroscopy (FTIRRS)
The FTIR reflectance spectra of the Nd3+ doped precursor LTSA glass and heat-treated glass-ceramic samples in the wavenumber range 400-2000 cm-1 is shown in Fig. 7. It is seen from this figure that the precursor glass (curve-a) exhibits two broad reflection bands centered around 960 and 610 cm-1 as a result of wider distribution of silicon and tantalate structural units respectively. As alumina is one of the glass constituents, it prefers to enter into the silica rich phase and somewhere replace the Si4+ and the charge is compensated by Li+ ion. But, in order to maintain neutral charge condition, the later phase contains a higher amount of Li+ ions as the TaO6 octahedra are negatively charged (Fukumi & Sakka, 1988, Samuneva et al., 1991). Hence, from the rearrangement of the glassy matrix it can be indicated that both the phase separated compositions begin to crystallize producing a nanostructure with the prolonged heat-treatment time. The appearance of a low intensity band at 735 cm-1 upon heat-treatment related to the stretching mode of Al-O bond of AlO4 tetrahedra of β-spodumene (Burdick & Day, 1967). The prominent band occurred at 600 cm-1 corresponds to the stretching mode of O-Ta bond of TaO6 octahedral units of lithium tantalate (Ono et al., 2001, Zhang et al., 1999). The reflection band centered at 600 cm-1 is assigned as LiTaO3 crystal formation and the reflection band centered at 1000 cm-1 is assigned to Si-O stretching vibration of residual glass and β-spodumene crystal. The variation of Si-O (998 cm-1) and Ta-O (602 cm-1) stretching vibration bands intensities (here reflectivity) of Nd3+ doped samples with heat-treatment time is also recorded. It is seen that with progression of heat-treatment the band intensities increase rapidly initially and then become almost saturated after a certain time of heat-treatment (10 h). The gradual increase of relative intensity of band at 600 cm-1 clearly indicates formation of LiTaO3 crystal with the increase of heat treatment time. The results of the FT-IRRS are in good agreement with that of XRD, FE-SEM and TEM studies. A similar observation has also been reported by Ito et al., 1978.
4.6. Dielectric constant (εr)
The as prepared Eu3+ and Nd3+ doped LTSA precursor glasses exhibit relatively higher value (~20.0) of dielectric constant (εr) than the common vitreous silica (3.8) or soda-lime silicate (7.2) or borosilicate glasses (4.1-4.9) (Blech, 1986) due to high ionic refraction of Ta5+ ions (23.4) (Volf, 1984). This is due to its empty or unfilled d-orbital which contributes very strongly to its high polarizability (Yamane & Asahara, 2000, Risk et al., 2003). Its magnitudes show a sharp increase with increase in heat-treatment duration up to 5 h and thereafter it maintained saturation with a small decrease for any further heat treatment time as shown in Fig. 8. This suggests that, at the initial stages of heat treatment (1-3 h), separation of silica rich phase and Li–Ta enriched phases takes place and with the further heat-treatment, incipient precipitation of LiTaO3 crystalline phase of high dielectric constant (εr = 52) (Moses, 1978) and spontaneous polarization (Ps = 0.50 C/m2) (Risk et al., 2003) occurs gradually which becomes well defined at 5 h and attains the maximum volume fraction of the crystalline phase. Thus accumulation of Li+ ions in the phase-separated glass matrix initially could cause a slight increase of dielectric constant and with further heat treatment time due to formation of stable LiTaO3 ferroelectric crystals remarkably increase the dielectric constant reaching the highest value for 5 h heat treated sample and then maintain almost same on further course of heat-treatment. The variation in the dielectric constant (εr) values among the heat-treated nano glass-ceramics are mostly due to volume fraction of crystal phases contained and also the distribution of the LiTaO3 phase in the microstructure (Vernacotola, 1994).
4.7. UV-Visible-NIR absorption spectra
The room temperature measured absorption spectra of the Nd3+ doped precursor glass (a) and 100 h heat-treated glass-ceramic samples (g) in the visible-NIR range have been presented in Fig. 9. The spectra reveal absorption peaks due to the
(880 nm) are assigned in accordance with Carnall’s convention (Carnall et al., 1968, Chen et al., 2005). From this figure it is noticed that the base line of absorption spectra of heat-treated sample (g) has been elevated significantly with the diminished intensity of the absorption peak. This uplifting can be attributed to scattering of short wavelength light by the crystals (Beall & Duke, 1983 & 1969] or may be due to the difference in refractive index of crystalline phase (RI of LiTaO3 is 2.1834 at 600 nm (Lynch, 1975)) with that of residual glassy matrix.
4.8. Photoluminescence spectra
The photoluminescence emission spectra of precursor glass (a) and Eu3+:LiTaO3 containing nano glass-ceramics (c and f) are recorded with an excitation at 392 nm and depicted in Fig. 10(A). All the spectra exhibit emissions from 5D0 excited level to the ground state multiplets 7F0, 1, 2, 3, 4 levels of Eu3+ ions with overall dominance of electric dipole (ED) transition 5D0
The infrared photoluminescence spectra (λex = 809 nm) of the Nd3+ doped samples (a, c and g) around 1069 nm are shown in Fig. 10(B). The emission band intensity around 1069 nm decreases with progression of heat-treatment. This decrease in emission intensity is due to the clustering of Nd3+ ions which is extremely sensitive to concentration quenching (Riello et al., 2006). Dejneka, 1998, has demonstrated in fluoride glasses that clustering thereby quenching occurs when the Eu3+-Eu3+ ionic separation is around 40 Å. In the present case, the Nd3+-Nd3+ ionic separation (Ri) in the precursor glass is found to be about 26 Å which was calculated using the relation (Pátek, 1970):
where NNd3+ is the Nd3+ ion concentration. It is, therefore, seen that the Nd3+- Nd3+ ionic separations (Ri) are in the quenching region. Theoretically, the rate of relaxation due to concentration quenching varies as Ri-6 (Campbell & Suratwala, 2000, Kang et al., 2001, Zgonik et al., 1993). With the progression of heat-treatment, the LiTaO3 crystal phase has been formed and the Nd3+ ions partitioned into the residual glassy phase by reducing the inter-ionic separation less than 40 Å. This fact results in reduction in fluorescence intensity due to concentration quenching. The emission bands become sharper and take shapes as in crystalline host with progress of heat-treatment duration. All these observations indicate that the Nd3+ ions enter into the LiTaO3 crystalline phase and therefore, environment around Nd3+ ions is changed with progression of heat-treatment.
4.9. Second harmonic generation (SHG)
The variation of SHG output power (nJ) with time period in bulk LTSA glass and LT glass-ceramics has been shown in Figs. 11(a) and (b). It is seen from the Fig. 11(a) that the precursor LTSA glass does not exhibit any SHG output. This is due to inversion symmetry of the precursor glass. It is also observed that for a constant heat-treatment time (10 h), the SHG output power increase from 1.08 to 1.875 nJ when the temperature is increased from 680°C to 850°C. This is attributed due to the formation of ferroelectric (non-centrosymmetric) LiTaO3 crystals in the glassy matrix. Moreover, the increase in SHG output power with subsequent increase of heat-treatment temperature is due to the increase of LiTaO3 crystallites content and their sizes in the glassy matrix.
5. Nanostructured KNbO3 ferroelectric glass-ceramics
Ferroelectric potassium niobate (KNbO3, KN) has the A1+B5+O3 perovskite-type (orthorhombic) crystal structure (crystal symmetry class
In comparison to the other rare earth ions, Er3+ has been extensively used as the most suitable active ion in several hosts by normal fluorescence for optical amplification at 1.5 µm. Trivalent erbium ions characterized by the 4I13/2 → 4I15/2 intra-4f transition play the key role in near infrared (NIR) emission (Dai et al., 2005). Besides, Er3+ doped glasses are chosen for getting better lasing property that can act as an eye-safe laser sources near 1550 nm. The photoluminescent emission intensity of erbium doped glasses is mainly dependent on the chemical environment of Er3+ ion because in the excited state Er3+ ion can de-excite by either photon emission at 1.53 µm, or non-radiative relaxation through coupling to a quenching site. Also, if the symmetry of the local crystal field around the erbium ion in the host glass is not distorted, the emission of erbium ion is forbidden. Therefore, the erbium ions must be incorporated in a non-centrosymmetric host material for strong optical emission (Winburn, 1985, Hui et al., 2007). Similarly, under favorable conditions strong green emission at 551 nm is possible upon normal high energy excitation.
6. Results and discussion
6.1. Differential thermal analysis (DTA)
The DTA curve of the Er3+ doped precursor glass is shown in Fig. 12. It exhibit the inflection in the temperature range 647-689°C followed by the intense exothermic peak at 759°C (Tp) corresponding to the phase crystallization. The glass transition temperature (Tg) has been estimated to be 681°C from the point of intersection of the tangents drawn at the slope change as is marked in Fig. 12.
6.2. Refractive index
The precursor glass was heat-treated at 800°C near glass transition temperature for various heat-treatment durations after nucleating at 750°C for 2 h. The prepared glass was transparent with body color of pink. From the measured glass density (ρ) and refractive index (ne) at wavelength λe = 546.1 nm, other related optical properties of Er3+ doped precursor glass have been determined using relevant expressions and the same is presented in Table 2. The refractive indices of glass and heat-treated samples at five different wavelengths (473, 532, 633, 1064 and 1552 nm) are shown in Fig. 13. It is seen that the refractive index of 100 h heat-treated sample is higher as compared to the precursor glass. This is due to the formation of KNbO3 crystals having high refractive index (2.2912 at 600 nm (Palik, 1998)). The large refractive index of this glass is due to the presence of highly
polarizable Nb5+ ions with high ionic refraction, 24.5 (Volf, 1984). It is observed that ions with an empty or unfilled d-orbital such as Nb5+ ion (outer electronic configuration: 4d05s0) contributes very strongly to the linear and nonlinear polarizabilities (Yamane & Asahara, 2000). For the same reason, this glass is also possessing a high value of molar refractivity (RM = 14.95 cm3) and electronic polarizability (α = 5.592 x 10-24 cm3) (Vernacotola & Shelby, 1994).
|Er3+:KNbO3 doped precursor glass|
|Average molecular weight, Mav||120.05|
|Density, ρ (g.cm-3)||3.37|
|Refractive index||ne (at 546.1 nm)|
nF΄ (at 480.0 nm)
nC´ (at 643.8 nm)
|Molar refractivity, RM (cm3)||14.95|
|Electronic polarizability, α (cm3)||5.592×10-24|
|RE3+ ion concentration, NRE3+ (ions/cm3)||5.41×1019|
|Glass transition temperature, Tg (°C)||681|
|Crystallization peak, Tp (°C)||759|
6.3. X-ray diffraction analysis
Fig. 14 shows the X-ray diffractogram of precursor glass along with the glass-ceramic samples. The amorphous nature of the as-prepared glass is indicated by the XRD pattern consisting of only a broad and halo band at around 29° diffraction angle. The structuring of this halo band takes place in the XRD pattern of the heat-treated glass-ceramic samples of 1-100 h along with the appearance of other well defined peaks around 25°, 28.5°, 30°, 32.8° and
51.5° diffraction angles, which confirms the precipitation crystalline phase in the amorphous matrix. The crystalline phase resembles the JCPDS cards 32-821 and 32-822 of known potassium niobate. The calculated average crystallite sizes lie in the range 5-12 nm.
6.4. FESEM and TEM image analyses
The FESEM photomicrographs of the sample heat-treated at 800°C for 3 and 50 h duration are taken and from the FESEM micrographs, it is clearly observed that the glassy matrix of the heat-treated samples initially phase separated on nanometric scale followed by incipient precipitation of defined crystallites within the Nb-K rich phase regions on prolonged heat-treatments. The TEM image and SAED pattern of the sample heat treated for 50 h at 800°C have been presented in Figs. 15(a) and 15(b) respectively. The crystallite size from this TEM image of 50 h heat-treated sample found to be around 20 nm. The presence of fine spherical rings around the central bright region in SAED pattern discloses the existence of KNbO3 nanocrystallites in the glassy matrix.
6.5. Fourier transform infrared reflectance spectroscopy (FTIRRS)
Fig. 16 shows the comparative FTIR reflectance spectra (FTIRRS) of the precursor glass and samples heat-treated at 800ºC for 2 and 100 h duration in the wavenumber range 500-2000 cm-1. Its inset shows the reflectivity at 930 and 749 cm-1 of precursor glass and heat-treated glasses as a function of heat-treatment time. It is seen that the FTIRR spectrum of the precursor glass exhibits a broad reflection band centered at 930 cm-1 as a result of wider distribution of SiO4 structural units. The FTIR reflectance spectra of 2 and 100 h reveal narrowing of the main reflection band with additional features arising at 1128, 749 and 598 cm-1 in comparison to the precursor glass (Pernice et al., 1999). In the FTIRR spectra, the stretching modes of the Si-O-Si bonds of the SiO4 tetrahedra with nonbridging oxygen (NBO) atoms are active in 900-1000 cm-1 range and the stretching modes of the Nb-O bonds in the NbO6 octahedra occur in the 700-800 cm-1 range (Samuneva et al., 1991). The reflection band at 1128 cm-1 and 930 cm-1 wavenumber can be related to the asymmetric and symmetric stretching vibration modes of Si-O bonds in SiO4 tetrahedra respectively, while the band at 749 cm-1 is due to the Nb-O stretching modes of distorted NbO6 octahedra (de Andrade et al., 2000, Silva et al., 2006). The band observed around 598 cm-1 is assigned as ν2 bending vibrational modes of the Si-O bonds in the SiO4 tetrahedra.
6.6. Dielectric constant (εr)
Fig. 17 shows the magnitude of dielectric constant of precursor glass and heat-treated glass-ceramics. From this Fig., it is observed that the dielectric constant increase steeply from εr =17 (for precursor glass) to εr =31 (for 100 h heat-treated glass-ceramics) and thereafter it maintains almost saturation for any further heat treatment time. This suggests that on heat-treatment, at the initial stages, separation of silica rich phase and K–Nb enriched phases takes place and with the further prolonged heat-treatments incipient precipitation of KNbO3 having high dielectric constant (εr =137) (Simões et al., 2004) and spontaneous polarization, Ps = 0.41 C/m2 (Risk et al., 2003). The variation in crystallite size distributions and also the distribution of the KNbO3 phase in the microstructure are the causes for the differences in the dielectric constant values amongst the heat-treated samples (Vernacotola, 1994).
6.7. UV-visible-NIR absorption spectra
Fig. 18 shows the representative UV-Vis-NIR absorption spectra of Er3+ doped precursor glass and heat-treated samples for 100 h duration. The absorption occurs due to the
6.8. Photoluminescence spectra
Fig. 19 depicts the infrared emission spectra of precursor KNS glass and samples heat-treated for 2 and 100 h durations. The directly excited 980 nm emission spectra of the as prepared glass and the heat-treated samples exhibits emission from 4I13/2 excited level to the 4I15/2 ground level with Stark splitting at 1537 and 1566 nm. With increase in heat-treatment time the peak at 1537 nm for glass shifted to 1540 nm for 100 h heat-treated sample. The peak intensity ratio at 1566 nm to 1540 nm for precursor glass is ~0.93 which decreases down to ~0.79 for the heat-treated samples. It is generally seen that the glass-ceramics samples show more intense photoluminescence than the precursor glass. In this case it is found that the fluorescence intensity first decreases for the glass-ceramics heat-treated for shortest duration (2h) and then increases but with low intense than precursor glass. The reason behind this may be the fact that in the shorter durations of heat-treatment the samples are phase separated and stable KNbO3 phase grows with longer heat-treatment duration.
6.9. Second harmonic generation (SHG)
The SHG output power (nJ) of KNS glass and KN glass-ceramics heat-treated at 800°C for 10 h is shown in Figs. 20(a) and (b). It is seen from the Fig. 20(a) that the precursor glass does not exhibit any SHG output. This phenomenon once again explained on the basis of inversion symmetry of the precursor glass. With heat-treatment of the precursor glass at 800°C for 10 h duration, the SHG output power increase to 39.74 nJ. This is attributed due to the formation of ferroelectric (non-centrosymmetric) KNbO3 crystals in the glassy matrix.
The precursor glasses having molar composition 25.53Li2O-21.53Ta2O5-35.29SiO2-17.65Al2O3 doped with RE oxides (0.5 wt% of Eu2O3 and Nd2O3 in excess) and 25K2O-25Nb2O5-50SiO2 (KNS) doped with Er2O3 (0.5 wt% in excess) were prepared by melt-quench technique. The precursor glasses and the resulting LT and KN nano glass-ceramics were characterized by studying their thermal, structural, optical, dielectric properties. The results of XRD, FESEM, TEM and FT-IRRS confirmed the formation of nanocrystalline LT phases in the LTSA glass matrices and KN phase in the KNS glass matrix. The nanocrystallite size of LT and KN evaluated from TEM images found to vary in the range 14-36 nm and 5-12 nm respectively. The dielectric constants found to increase with heat-treatment time due to ferroelectric LT and KN formation. The photolminescence studies indicate the incorporation of RE ions into LT and KN crystal lattice. The derived LT and KN nano glass-ceramics also exhibit considerable SHG output which is very important for exploitation in self frequency doubling laser devices.