Structure and properties of lanthanum galliogermanate glasses

Structure and properties of lanthanum galliogermanate glasses

Journal of Non-Crystalline Solids 240 (1998) 22±28 Structure and properties of lanthanum galliogermanate glasses Sungping Szu a a,* , Chanping Shu ...

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Journal of Non-Crystalline Solids 240 (1998) 22±28

Structure and properties of lanthanum galliogermanate glasses Sungping Szu a

a,* ,

Chanping Shu a, Luu-Gen Hwa

b

Department of Physics, National Chung Shing University, Taichung, Taiwan, ROC b Department of Physics, Fu-Jen Catholic University, Taipei, Taiwan, ROC Received 1 December 1997; received in revised form 20 May 1998

Abstract Lanthanum galliogermanate glasses were prepared. Raman spectra, molar volumes, glass transition temperatures and activation energies for glass transition and crystallization were obtained. For glasses having the same La2 O3 /GeO2 ratio, the molar volumes increase with the Ga2 O3 content, and the glass transition temperatures, activation energies for glass transition and crystallization, increase initially then decrease as the ratio of Ga2 O3 /GeO2 increases. For glasses having the same Ga2 O3 /GeO2 ratio, the molar volumes increase with La2 O3 content, and the glass transition temperatures increase as the ratio of La2 O3 /GeO2 increases. The change of glass structure and its properties with composition is correlated with the concentration of lanthanum ion. Ó 1998 Elsevier Science B.V. All rights reserved.

1. Introduction Glasses containing GeO2 and Ga2 O3 are interesting for structural change. The coordination of Ge and Ga will change from 4-fold to 6-fold as the compositions of glass changes [1±4]. The extended X-ray absorption ®ne structure (EXAFS) for measurements of sodium germanate glasses show no indication of coordination change for Ge [5]. Another 6-fold analysis of the Raman spectra of sodium germanate glasses has been interpreted in terms of changes in tetrahedral ring size, rather than as a change in coordination [6]. Gallium oxide is a conditional glass-former, which does not form glass by itself but does so when combined with other suitable network forming oxide. The structural role of Ga is assumed to be comparable to that of Al [7,8]. Ga will

* Corresponding author. Tel.: +886-4 284 0427; fax: +886-4 286 2534; e-mail: [email protected]

substitute for Ge in galliogermanate or Si in galliosilicate glasses with a nearby cation for charge balance. For compositions with Ga‡3 /M‡ or Ga‡3 / 1 M‡2 greater than 1, where M is an alkali or al2 kaline earth cation, some Ga are octahedral coordinated [9]. However, the EXAFS spectra of alkali galliosilicate and barium galliogermanate glasses show that gallium coordination is 4, regardless of glass composition [10±12]. The structural role of La‡3 in glass is also very interesting for its high ®eld strength compared to the ®eld strength of alkali or alkaline earth ion. Coon and Shelby [13], have presented evidence that lanthanum plays a dual role in alkali silicate glasses with concentrations <4 mol%. In lanthanum oxide substituted barium galliogermanate glasses, La‡3 enters the glasses as charge-compensator for GaOÿ 4 tetrahedron or as network modi®er for non-bridging oxygen [14]. In this paper lanthanum galliogermanate glasses in glass-forming region were prepared. The structure of these glasses was studied by Raman

0022-3093/98/$ ± see front matter Ó 1998 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 3 0 9 3 ( 9 8 ) 0 0 7 1 0 - 8

S. Szu et al. / Journal of Non-Crystalline Solids 240 (1998) 22±28

spectroscopy. The properties of these glasses, including densities, glass transition temperatures (Tg ), crystallization temperatures (Tc ), activation energies for glass transition (Eg ) and crystallization (Ec ) were measured. The change of properties with composition as well as structure are presented in this paper. 2. Experimental procedure Glass samples were prepared using, 99.99% La2 O3 , Ga2 O3 and GeO2 . Batches (5 or 15 g) were melted in a Pt crucible between 1550°C and 1650°C for 1 h. The weight loss of the melt due to volatilization was less than 1 wt%. Five gram batches were quenched in water; while 15 g batches were cooled in air to obtain bulk samples for Raman and density measurements. The amorphous state was con®rmed by powdered X-ray di€raction (XRD) (Siemens D5000). Sample compositions were chosen according to the formula xLa2 O3 ± yGa2 O3 ±GeO2 , where x ˆ La2 O3 mol%/GeO2 mol% and y ˆ Ga2 O3 mol%/GeO2 mol%. Di€erential thermal analysis (DTA) (TA instrument TA2000) was used to determined Tg and Tc . For every DAT measurement, the measurement was terminated as the ®rst crystallization peak (exothermic peak) was completed. The activation energies for glass transition, Eg , and crystallization, Ec , were determined by varying heating rate (5°C± 180°C/min) during DTA using the Kissinger formula [15]. The Kissinger equation states that

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the unpolarized 488 nm line of an argon ion laser with power 0.8 W. The intensity of an absorption peak is obtained by integrating the area under each peak. 3. Results The glass-forming region of lanthanum galliogermanate system is shown in Fig. 1. The ®lled circle symbol shows the glasses that can be made with 15 g batches, the un®lled circle shows the glasses that can be made with 5 g batches; while, compositions that cannot be melted at 1650°C or crystallized during cooling are indicated by triangle. Fig. 2 shows typical Raman spectra for glasses with constant y (Ga2 O3 mol%/GeO2 mol%) but with di€ering x (La2 O3 mol%/GeO2 mol%). There are three major absorption bands for each spectrum. The band positions are around 800, 540 and 300 cmÿ1 . As x increases, the band at 800 cmÿ1 shifts to smaller wavenumber, while the other two bands do not change their positions. The amplitudes of bands at 800 and 300 cmÿ1 increase with increasing x, but the amplitude of the band at 540 cmÿ1 decreases. Fig. 3 shows typical Raman spectra for glasses with constant x (x ˆ 1/3) but di€erent y. The

ln…Tg2 =a† ‡ constant ˆ Eg =RTg ; where a is the heating rate (K/min) and R is the universal gas constant. From the slope of ln(T2g /a) vs. 1/Tg , Eg is calculated. The equation is also applied to evaluate Ec . After DTA measurements, XRD was performed on the samples. The powder patterns were compared with the Joint Committee on Powder Di€raction Standards (JCPDS) to determine the crystalline phases. Densities were measured by the Archimedes technique using water as the medium and errors of measurement were within ‹0.01 g/cm3 . Raman spectra were obtained using 0.85 m double monochromater (Spex-1403). The light source was

Fig. 1. La2 O3 ±Ga2 O3 ±GeO2 glasses prepared for this study. The symbols d and s represent the glass that can be made with 15 g batch and 5 g batch respectively; while, compositions that cannot be melt at 1650°C or crystallized during the cooling are indicated by m. Dotted lines are drawn to enclose glass samples as guides for glass-forming region.

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Fig. 2. Raman Spectra for glasses of the series y ˆ 0.25, where y ˆ Ga2 O3 /GeO2 .

maximum amplitude at 800 and 300 cmÿ1 decreases with increasing y, but the amplitude of the peak at 540 cmÿ1 increases with y. The results of density measurements are listed in Table 1. The molar volumes are calculated according to their compositions. They are also listed in Table 1. The variation of molar volumes with composition are shown in Fig. 4(a) and (b). As shown in the ®gures, the molar volume increases linearly either with Ga2 O3 for samples having the same x, Fig. 4(a), or with La2 O3 mol% for samples having the same y, Fig. 4(b). The data are ®tted by root mean square method. The ®tting parameters are listed in Table 2. The intercepts represent the molar volumes of ®ctitious binary La2 O3 ±GeO2 glasses (Fig. 4(a)) and Ga2 O3 ±GeO2 glasses (Fig. 4(b)), which cannot be prepared in this study. As shown in Table 2, the variation of molar volumes with composition in binary systems has the same trend as that in the ternary system. The molar volume increases with La2 O3 or Ga2 O3 .

Fig. 3. Raman Spectra for glasses of the series x ˆ 1/3, where x ˆ La2 O3 /GeO2 .

Table 1 Density and molar volume of xLa2 O3 ±yGa2 O3 ±GeO2 glasses x

y

Density ‹ 0.01 (g/cm3 )

Molar volume ‹ 0.1 (cm3 /mol)

0.20 0.20 0.20 0.20 0.25 0.25 0.25 0.25 0.33 0.33 0.33 0.33 0.5 0.5 0.5 0.5

0.20 0.25 0.33 0.50 0.20 0.25 0.33 0.5 0.2 0.25 0.33 0.5 0.2 0.25 0.33 0.5

5.14 5.08 5.16 5.20 5.26 5.30 5.25 5.24 5.35 5.36 5.38 5.39 N.A 5.55 5.57 5.51

40.4 42.7 45.0 50.7 42.5 43.9 47.4 53.4 46.9 48.5 51.3 57.0 N.A 56.7 59.2 65.6

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Table 2 Root mean square ®tting parameters for molar volumes in Fig. 4(a) and (b) Fig. 4(a)

Slope

Intercept (cm3 /mol)

Correlation coecient

x 0.50 0.33 0.25 0.20

0.84 0.71 0.76 0.67

44.4 37.3 31.7 30.8

0.997 0.998 0.998 0.998

1.12 1.00 0.98 0.90

37.3 31.7 28.4 27.2

0.998 0.999 0.984 0.996

Fig. 4(b) y 0.5 0.33 0.25 0.2

Fig. 5. Glass transition temperatures, Tg , for glasses having constant x as a function of y. The inset is the Tg for glasses with x ˆ 1. Lines are drawn as guides for the eye.

Fig. 4. Molar volume vs. content of (a) Ga2 O3 mol% and (b) La2 O3 mol%. Solid lines are the best ®t by least square method.

Fig. 5 displays the glass transition temperatures of all samples. As shown, for glass series having the same x, Tg increases initially then decreases when y increases. The ®gure also shows that Tg increases with x for glasses having the same y, with the exception of y ˆ 0.2. Fig. 6(a) and (b) show Tg and Tc as a function of heating rate, a, for sample series with x ˆ 0.2. Eg and

Ec are obtained according to the procedure stated in Section 2, and the Eg and Ec for other samples are listed in Table 3. The correlation coecients for all ®ttings are P 0.96. These coecients indicate validation of the Kissinger formula in this study. Since the DTA measurements were terminated immediately after the ®rst crystallization peak was completed, all samples crystallized during DTA measurements. Fig. 7 shows the powder patterns for samples with x ˆ 0.2. The crystalline phase is Ga2 GeO5 (JCPDS No. 42-0048) and LaGaGe2 O7 (JCPDS No. 41-0969) in samples with y ˆ 0.2 and

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Fig. 6. (a) ln(T2g /a) vs. 1000/Tg , and (b) ln(T2c /a) vs. 1000/Tc for glasses with x ˆ 0.2 series. The functions ln(T2 /a) + constant ˆ E/RT are ®tted to the data by root mean square method.

0.25, respectively. For samples with y ˆ 0.33 and 0.5, the di€raction patterns correspond to a mixture of La2 Ge2 O7 and La4 GeO8 (JCPDS No. 420205, 40-1185). 4. Discussion According to the literature, the Raman spectrum between 700 and 900 cmÿ1 , shown in Fig. 2, is attributed to the antisymmetric stretching mode

of four coordinated Ge and Ga, that is the stretching of T±O±T and O±T±Oÿ , where T is Ge or Ga and Oÿ is non-bridging oxygen (NBO) [16± 20]. The response around 400±600 cmÿ1 is attributed to the symmetric bending and stretching of T±O±T [16,17,20±23]. Both bands are asymmetric which indicates many vibration modes coupled together [16,17]. Because of the similar mass and coordination of Ga and Ge, the stretching and bending vibrations between Ga±O and Ge±O are inseparable in Raman spectra. The band between 200 and 400 cmÿ1 is probably due to the interaction between La‡3 ion and its surrounding oxygens. For every GaOÿ 4 tetrahedron or NBO, a positive charge is needed for charge neutrality. La‡3 ions can provide these positive charges. Ruller and Jewell [23] have pointed out that the band around 350 cmÿ1 in PbO±Ga2 O3 ± SiO2 glass system is due to the vibration of Pb±O± Ga because of the ®eld strength of the Pb ion. The intensity of this Pb±O±Ga vibration mode increases with PbO content. The intensity of the band around 300 cmÿ1 in the spectra of our samples also increases with increasing La2 O3 . As shown in Fig. 2, the intensities of bands around 800 and 300 cmÿ1 increase with La2 O3 , while, the intensity of the band around 540 cmÿ1 decreases with increasing La2 O3 . The change of intensities of these three bands is related to the La‡3 content. The role of La‡3 is either as chargecompensator for GaOÿ 4 or as glass modi®er for NBO. As the content of La‡3 ion increases, the modi®ers will break the T±O±T bonding. Hence, the intensity for the band at 540 cmÿ1 , which represents the T±O±T bending, decreases. The relative intensity for band representing T±Oÿ , which is around 800 cmÿ1 , increases. The shift of that band to smaller wavenumber is due to the increase of the number of NBO [20]. As the content of La2 O3 increases, there are more La‡3 ±O interactions, hence, the intensity of band around 300 cmÿ1 increases. The change of relative intensities in Fig. 3 is also related to the La‡3 ion content. The decrease of band intensity at 300 cmÿ1 as y increases is due to the decreasing La2 O3 content. As y increases, the content of Ga2 O3 is increased, some of La‡3 ions need to act as charge-compensators for charge neutrality on GaOÿ 4 . This function will decrease the

S. Szu et al. / Journal of Non-Crystalline Solids 240 (1998) 22±28

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Table 3 Eg and Ec of xLa2 O3 ±yGa2 O3 ±GeO2 glasses x

y

Eg ‹ 80 (kcal/mol)

Correlation coecient

Ec ‹ 40 (kcal/mol)

Correlation coecient

0.20 0.20 0.20 0.20 0.25 0.25 0.25 0.25 0.33 0.33 0.33 0.33 0.5 0.5 0.5 0.5

0.20 0.25 0.33 0.50 0.20 0.25 0.33 0.5 0.2 0.25 0.33 0.5 0.2 0.25 0.33 0.5

940 1000 1100 1100 1000 1200 1000 800 1000 1300 1300 1100 1100 1200 1300 1100

0.993 0.987 0.992 0.994 0.994 0.978 0.980 0.960 0.996 0.986 0.996 0.990 0.972 0.991 0.976 0.973

430 480 420 410 390 480 430 430 500 530 480 480 430 470 570 530

0.984 0.991 0.997 0.998 0.993 0.960 0.968 0.972 0.993 0.987 0.976 0.963 0.961 0.960 0.995 0.960

number of NBO in GeO2 sub-network and increase the number of bridging oxygen in Ga2 O3 sub-network. Hence, the band intensity at 800 cmÿ1 decreases and the band intensity at 540 cmÿ1 increases. The molar volume increases linearly with increasing of Ga2 O3 or La2 O3 content along the constant La2 O3 /GeO2 or Ga2 O3 /GeO2 ratio. The increase of molar volume with increasing Ga2 O3 content, Fig. 4(a), is due to replacing larger GaO4 tetrahedra with smaller GeO4 tetrahedra. The study of EXAFS shows that the bond length of Ga±O in GaO4 tetrahedron is longer than that of Ge±O in GeO4 tetrahedron [11]. They are 0.184

Fig. 7. XRD patterns for glass series with x ˆ 0.2 and y varies from 0.5 to 0.2. The symbols on each peak represent the diffraction peaks for crystalline phase.

and 0.175 nm, respectively [11]. Therefore, the volume of a GaO4 tetrahedron is larger than that of a GeO4 tetrahedron. The molar volume of the glass increases as GaO4 replaces GeO4 . The increase of molar volume with increasing La2 O3 content is probably due to the large size of the La‡3 ion, and/or the increase of the number of the NBO on tetrahedron [24]. The variation of Tg with glass composition is shown in Fig. 5. For glasses having the same x, Tg increases initially then decreases as y increases. As more Ga2 O3 enters glass, there are more GaO4 tetrahedra in glass. The bond strength of Ga±O is weaker than that of Ge±O, 84.5 and 157.6 kcal/ mol, respectively [25]. Hence, Tg decreases as Ga2 O3 increases. Another obvious change of Tg with composition is the increase of Tg with increasing x when y > 2. The increase in Tg with rare earth additions has been observed for other rare earth containing glasses and were explained by the unique structural con®guration of three MO4 (M is a glass± former cation) groups around a single modi®er ion, La‡3 ion [26]. Jewell et al. have attributed the increase of Tg with lanthanum content to the increase of the number of bonding con®gurations in the structure [14]. Table 3 lists Eg and Ec and the correlation coecient of the ®tting. Among the samples having

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the same x, Eg and Ec increase initially then decrease as y increase. These variations of Eg and Ec are similar to that of Tg . We suggest that this similarity indicates that the change of Tg and Eg with composition is caused by the same process. The crystalline phases in these samples are Ga2 GeO5 and LaGaGe2 O7 for y ˆ 0.2 and 0.25, respectively, and are a mixture of La2 Ge2 O7 and La4 GeO8 for y ˆ 0.33 and 0.5. The similarity of the Ec s values for the last two samples agrees with the XRD result. These two samples have the same crystalline phases. Most of the samples in this study have Ga/13 La‡3 6 1. We assume that, for glasses with composition in this range, the gallium ion acts as glass-former, which is four-coordinated. A recent EXAFS study on ternary glass systems of alkali gallosilicate [10,12] and barium gallogermanate [11] showed no evidence of six-coordinated Ge and Ga in those glasses. The Raman and IR spectra in some Ga2 O3 or GeO2 containing ternary glasses show no direct evidence of existence of six-coordinated Ge or Ga [9,16,20,23]. The Raman band due to six-coordinated Ge is at 700 cmÿ1 [6]. From the Raman spectra shown in Figs. 2 and 3, the possible of existence of six-coordinated Ge and Ga cannot be ruled out. However, the data indicate that formation of GeO6 and GaO6 species in these glasses is not detectable in the Raman spectra. 5. Conclusion The Raman spectra indicate that the majority of Ge and Ga are four-coordinated in the glass-network. The role of La‡3 ion acts either as chargecompensator for GaOÿ 4 or as glass modi®er for NBO. Both Tg and molar volume increase with the content of La‡3 ion. The intensity of Raman band around 300 cmÿ1 , which is assigned to the La±O vibration mode, also increases with La2 O3 content. As GeO2 is replaced by Ga2 O3 , Tg decreases and the molar volume increases. The decrease in Tg is explained by the bond strength of Ge±O which is greater than that of Ga±O. The increase of molar volume is due to replacing the GeO4 tetrahedrons by the larger GaO4 .

Acknowledgements This work is supported by the National Science Council through grant No. NSC85±2112±M005± 008. References [1] M.K. Murthy, E.M. Kirby, Phys. Chem. Glasses 5 (1964) 144. [2] C. Yin, K. Lu, Y. Zhao, J. Non-Cryst. Solids 112 (1989) 96. [3] P.W. Angel, C.S. Ray, D.E. Day, J. Am. Ceram. Soc. 73 (1990) 2965. [4] J.L. Piguet, J.E. Shelby, J. Am. Ceram. Soc. 68 (1985) C232. [5] J.P. Itie, A. Polan, G. Calas, J. Petiau, A. Fontaine, H. Tolentino, Phys. Rev. Lett. 63 (1989) 398. [6] G.S. Henderson, M.E. Fleet, J. Non-Cryst. Solids 134 (1991) 259. [7] L.A. Balewick, J.E. Shelby, J. Am. Ceram. Soc. 72 (1989) 1751. [8] J.L.L. Piguet, J.C. Lapp, J.E. Shelby, J. Am. Ceram. Soc. 68 (1985) 326. [9] C.I. Merzbacher, Phys. Chem. Glasses 33 (1992) 233. [10] P.L. Higby, J.E. Shelby, J.C. Phillips, A.D. Legrand, J. Non-Cryst. Solids 105 (1988) 139. [11] C.I. Merzbacher, D.A. McKeown, J. Non-Cryst. Solids 162 (1993) 81. [12] N. Iwamoto, N. Umesaki, S. Goto, T. Hanada, N. Soga, Transactions of JWRI 12 (1983) 173. [13] J. Coon, J.E. Shelby, Phys. Chem. Glasses 35 (1994) 47. [14] J.M. Jewell, P.L. Higby, I.D. Aggrawal, J. Am. Ceram. Soc. 77 (1994) 697. [15] H.G. Kissinger, Anal. Chem. 29 (1957) 1702. [16] K.E. Lipinska-Kalita, J. Non-Cryst. Solids 119 (1990) 41. [17] K.E. Lipinska-Kalita, D.J. Mowbray, J. Non-Cryst. Solids 122 (1990) 1. [18] F. Miyaji, S. Saka, J. Non-Cryst. Solids 134 (1991) 77. [19] K. Fukumi, S. Saka, Phys. Chem. Glasses 29 (1988) 1. [20] D.A. McKeown, C.I. Merzbacher, J. Non-Cryst. Solids 183 (1995) 61. [21] K.E. Lipinska-Kalita, D.J. Mowbray, J. Molec. Struc. 219 (1990) 107. [22] K. Fukumi, T. Kokubo, K. Kamiya, S. Saka, J. NonCryst. Solids 84 (1986) 100. [23] J.A. Ruller, J.M. Jewell, J. Non-Cryst. Solids 175 (1994) 91. [24] P.L. Higby, I.D. Aggarwal, J. Non-Cryst. Solids 163 (1993) 303. [25] J.B. Pedley, F.M. Marshall, J. Phys. Chem. Ref. Data 12 (1984) 967. [26] F. Branda, A. Buri, D. Caferra, A. Marotta, Phys. Chem. Glasses 22 (1981) 68.