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Effect of CsI addition on the thermal properties, crystallization kinetics and short-wavelength absorption edges of GeSe2–Ga2Se3 glasses
Journal of Alloys and Compounds 474 (2009) 468–472
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Effect of CsI addition on the thermal properties, crystallization kinetics and short-wavelength absorption edges of GeSe2 –Ga2 Se3 glasses Cunming Liu a,b,∗ , Gao Tang a,b , Lan Luo a , Wei Chen a,∗ a b
Shanghai Institute of Ceramics, Chinese Academy of Sciences, 1295 Dingxi Road, Shanghai 200050, PR China Graduate School of the Chinese Academy of Sciences, Beijing 100039, PR China
a r t i c l e
i n f o
Article history: Received 21 May 2008 Received in revised form 23 June 2008 Accepted 25 June 2008 Available online 3 August 2008 Keywords: Chalcogenide glasses Thermal properties Crystallization kinetics Short-wavelength absorption edges
a b s t r a c t The glasses of (1 − x)(11/16GeSe2 –5/16Ga2 Se3 )–xCsI (x = 0.1, 0.2 and 0.3) were prepared by the melt quench technique. The thermal properties and crystallization kinetics of these glasses were investigated by the differential thermal analysis. It is found that the glass of x = 0.2 is the most stable and that the increase of CsI decreases the activity energy of glass transition Et and the activity energy of glass crystallization Ec . When evaluating the thermal stability of the studied glasses with Ec , the crystallization frequency factor K0 should be considered. The crystallization process of the x = 0.1 glass is determined by two-dimensional crystal growth, while that of the other two glasses is realized by three-dimensional crystal growth. Through analyzing the short-wavelength absorption edges of the glasses in detail, it can be concluded that the opt addition of CsI increases the Urbach energy Eu , optical energy Eg and Tauch energy Eg , and that all of the short-wavelength absorption edges are dominated by the indirect transition. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Selenium-based chalcogenide glasses with the transparence from 1 to 16 m are considered to be an alternative solution to replace the single crystalline germanium (Ge) usually used for thermal cameras and infrared (IR) detectors due to their much lower cost, easier fabrication of complex aspheric IR optical lenses and higher chemical stability [1,2]. Recently, Zhang et al. [3] revealed that selenium-based chalcogenide glasses containing alkali halides can generate nano-crystals to enhance their mechanical properties through heat treatment, and Yang et al. [4] reported that the transparent range of selenium-based chalcogenide glasses can be extended to the partial red light by alkali halides addition. These progresses make the selenium-based chalcogenide glasses much superior to the crystal of Ge; however, there are still some problems for their application, especially the poor thermo-mechanical properties and small transparent range in the visible region. An effective solution to these problems is designing the new selenium-based chalcohalide glasses with the larger visible transparent range and then crystallizing the glasses to produce nano-crystals, which can enhance the thermo-mechanical properties of glasses but cannot
∗ Corresponding author at: Shanghai Institute of Ceramics, Chinese Academy of Sciences, 1295 Dingxi Road, Shanghai 200050, PR China. Tel.: +86 21 52411024; fax: +86 21 52413903. E-mail addresses: [email protected] (C. Liu), [email protected] (W. Chen). 0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.06.116
change the optical properties of glasses significantly. In order to design the new glasses well and perform nano-crystallization on the glasses properly, it is necessary to understand the influence of alkali halides addition on the thermal properties, crystallization kinetics and short-wavelength absorption edges of the glasses. In this work, the glasses of (1 − x)(11/16GeSe2 –5/16Ga2 Se3 )–xCsI (x = 0.1, 0.2 and 0.3) were prepared. Then, their thermal properties, crystallization kinetics and short-wavelength absorption edges were investigated. 2. Experimental Glass samples of (1 − x)(11/16GeSe2 –5/16Ga2 Se3 )–xCsI (x = 0.1, 0.2 and 0.3) were synthesized by melting mixtures of the constituent elements (Ge, Ga, Se all of 99.999% purity, and CsI of 99.95% purity) in evacuated (10−2 Pa) and flame-sealed silica ampoule in a rocking furnace. The mixtures were melted at 1100–1200 K for 10 h. The temperature of the furnace was subsequently lowered to 1000 K and then held for another 2 h. After that, the ampoules were quenched in water, swiftly moved to a preheated furnace and annealed at a temperature lower than the corresponding glass transition temperature for 2 h. Glass rods were obtained by taking them out from the ampoules and finally cut into discs (Ø 9 mm × 2 mm) which were then polished for succeeding experiments. For the analysis of thermal properties and crystallization kinetics, the differential thermal analysis (DTA) measurements of bulk glass pieces (about 45 mg) with the same shape were carried out by a CDR-1P thermal Analyzer (SBIF, Shanghai, PR China) with an accuracy of ±1 K at continuous heating rates (ˇ = 2.5, 5, 7.5, 10, 15 and 20 K/min). The calorimeter was calibrated by the well-known melting temperatures and melting enthalpies of zinc and indium at each heating rate. Pure ␣-Al2 O3 powder was used as the reference material. The glass transition temperature Tg , the starting crystallization temperature Tx and the crystallization peak temperature Tp were obtained by the microprocessor of the thermal analyzer.
C. Liu et al. / Journal of Alloys and Compounds 474 (2009) 468–472
Fig. 1. DTA curves of the as-prepared glasses: (1 − x)(11/16GeSe2 –5/16Ga2 Se3 )–xCsI (x = 0.1, 0.2, 0.3).
In the interest of studying the short-wavelength absorption edges of these glasses, visible–near infrared (vis–NIR) absorption spectra were measured by a Cary 500 spectrophotometer (Varian, Palo Alto, USA). Otherwise, X-ray diffraction (XRD) data were collected using a Rigaku D/MAX-2550V diffractometer (Rigaku, Tokyo, Japan) with Cu K␣ radiation at 40 kV and 40 mA to identify the homogeneity and lack of crystallinity in each sample.
3. Results and discussions 3.1. Thermal properties Typical DTA curves of the three glasses recorded at a heating rate of 10 K/min are shown in Fig. 1. Founded on the values of Tg , Tx , Tp given in Table 1, the common stability criterion T = Tx − Tg [5], the weight stability criterion Hw = T/Tg and the S criterion S = (Tp − Tx )(Tx − Tg )/Tg [6] are used to evaluate the thermal stability of these glasses and their values are listed in Table 1. Generally, the more stable glass has the larger values of these parameters. From Fig. 1 and Table 1, it is found that: firstly, the glasses of x = 0.1 and 0.2 have two exothermic crystallization peaks while the x = 0.3 glass has one exothermic peak; secondly, the x = 0.2 sample is the most stable with a common stability criterion T larger than 160 K; thirdly, when the CsI increases, Tg decreases gradually, whereas Tx , Tp and the thermal stability criterion parameters (T, Hw and S) increase at first and then decrease. These results illustrate that the added CsI alters the glass structure. With regard to the influence of the CsI content on the glass structure, it may be similar to that of some other alkali halides like KBr or CsCl reported in some papers [7,8]. When lower amount of CsI (x ≤ 0.2) is introduced, the [GeSe4 ] units decrease, transforming into the [GeSe3 I] units, which loose the whole glass structure. On account of the strong complexforming ability of Ga3+ , the [GaSe3 ] triangles can easily transform into the [GaSe4 ] tetrahedral to make up network relaxation through capturing the lone pairs from Se [9]; however, [GaSe4 ] cannot sufficiently compensate the loss of [GeSe4 ] because of the appearance of the [GaSe3 I] units. Consequently, Tg somewhat declines. NevertheTable 1 Thermal parameters for the three glass samples
469
Fig. 2. Plots of Tg versus ln(ˇ) for three glass samples.
less, due to the higher space complication caused by the appearance of [GaSe4 ], [GaSe3 I] and [GeSe3 I] units, Tx , Tp and the thermal stability criterion parameters increase. As the content of CsI is higher (x > 0.2), larger amount of [GaSe3 I] and [GeSe3 I] units appear to further decrease [GeSe4 ] and [GaSe4 ] units. The space complication starts to become lower because of the much flabbier glass structure. As a result, Tg sequentially decreases, and Tx , Tp and the thermal stability criterion parameters begin to decrease. Meanwhile, more [GaSe3 I] and [GeSe3 I] units change the crystallization of glasses with the increase of CsI, which represents that two exothermic crystallization peaks are turned into a single exothermic crystallization peak on their DTA curves. 3.2. Analysis of crystallization kinetics The relationship between the glass transition temperature Tg and the heating rate ˇ can be discussed through the empirical equation: Tg = A + B ln(ˇ)
(1)
where A and B are constants for a given glass composition [10]. Plots of Tg versus ln(ˇ) for the prepared samples are shown in Fig. 2. From this figure, the values of A and B can be obtained by using the least square fit and listed in Table 2, and this equation holds good for the studied samples. To obtain the activity energy of glass transition Et , the data are fitted by the Kissinger method [11], which is most commonly used in analyzing crystallization data in DTA experiments. According to the Kissinger equation, the glass transition temperature Tg depends on the heating rate ˇ as follows:
ln
ˇ Tg2
=−
Et + const. RTg
(2)
based on the above relationship, straight lines between ln(ˇ/Tg2 ) and (1000/Tg ) can be given; slopes of these lines yield the values of Et /R where Et corresponds to the glass transition activation energy Table 2 Values of A, B and Et for the studied glasses
Composition
Tg (K)
Tx (K)
Tp (K)
T (K)
Hw
S (K)
Composition
A (K)
B (K)
Et (kJ/mol)
x = 0.1 x = 0.2 x = 0.3
647 619 606
727 782 732
754 799 751
80 163 126
0.124 0.263 0.208
3.34 4.48 3.95
x = 0.1 x = 0.2 x = 0.3
633.72 599.86 585.28
6.41 8.54 8.66
490.55 333.29 267.60
470
C. Liu et al. / Journal of Alloys and Compounds 474 (2009) 468–472
Fig. 3. Plots of ln(ˇ/Tg2 ) versus (1000/Tg ) for three glass samples. Fig. 4. Plots of ln(ˇ/Tp2 ) versus (1000/Tp ) for the studied glasses.
and R is the universal gas constant (R = 8.314 J/(mol K)). The calculated values of Et from the slopes of fitted lines shown in Fig. 3 are listed in Table 2. The result reveals that Et decreases with the increase of CsI like Tg . Concerning Et , it is absorbed by a group of atoms in the glass region when a jump from one metastable state to another is possible [12]. In other words, Et is involved in the molecular motions and rearrangements of the atoms around the glass transition temperature [13]. When the sample is reheated in the DTA furnace, the atoms undergo infrequent transitions between the local potential minima separated by different energy barriers in the configuration space where each local minimum represents a different structure. The more stable local minimum in glass region has lower internal energy. Accordingly, the atoms with minimum activation energy in a glass have higher probability to jump to the metastable (or local minimum) state of lower internal energy and hence a decreasing Tg is observed on the DTA curves when Et decreases. The Kissinger method [11] can be also used in the analysis of exothermic peaks to get the activation energy of glass crystallization Ec with the following relationship:
ln
ˇ
=−
Tp2
Ec + const. RTp
(3)
a new equation from the above equation was developed by Bansal et al. [14] as follows:
ln
ˇ Tp2
=−
Ec − ln RTp
E c
R
+ ln K0
(4)
where K0 is the crystallization frequency factor. From the linear dependence of ln(ˇ/Tp2 ) versus (1000/Tp ) shown in Fig. 4, Ec and K0 can be calculated from slopes of these fitted lines and their values are listed in Table 3. Table 3 reveals that firstly, the minimal value of K0 appears when the composition is x = 0.2; secondly, Ec decreases with the increasing content of CsI and its developing tendency is not consistent with that of the thermal stable criterion parameters in Table 1. Commonly, the glass is more stable when Ec is larger due to a larger
energy barrier represented by a larger Ec . In fact, the whole crystallization is complicated, and it is not exact that the glass thermal stability is judged only by Ec , so the crystallization frequency factor K0 should be considered [15,16]. To evaluate the glass thermal stability accurately, a stricter criterion, the crystallization rate constant K computed with Ec and K0 , was advanced by some authors [17,18]. The formula is listed below:
K = K0 exp
−Ec RTp
(5)
according to this equation and the values of Tp , Ec , K0 in Tables 1 and 3, the calculated values of K are 9.94 × 10−4 , 9.94 × 10−5 and 4.45 × 10−4 min−1 corresponding to the x = 0.1, 0.2 and 0.3 glasses, which illustrates that the x = 0.2 glass is the most stable. For the determination of the crystallization mechanism, the crystallization kinetics parameters n and m, depending on the nucleation process and growth morphology, can be obtained by the following modified Johnson–Mehl–Avrami equation [19]: ln[−ln (1 − )] = −n ln(ˇ) − 1.052
mEc + const. RT
(6)
where is the crystallized volume fraction at any temperature. When the quenched glasses contain no preexisting nuclei n is equal to m + 1, and when the reheated glasses contain a sufficiently large number of nuclei n is equal to m [19]. Also, m = 3, 2, and 1 are for three-, two- and one-dimensional crystal growth, respectively [19]. The value of n can be obtained by calculating the mean value of slopes of curves at different specific temperatures, as plotted in Fig. 5. Since all of glass samples were quenched in water and annealed at the temperatures lower than their own Tg , the value of m can be deduced from n = m + 1. As shown in Table 3, m = 2 and 3 are for the x = 0.1 glass and the other two glasses, respectively, which suggests that two-dimensional crystal growth works in the x = 0.1 glass and three-dimensional crystal growth is dominant in the other two glass matrices when the crystallization process occurs in the three glasses. 3.3. Analysis of short-wavelength absorption edges
Table 3 Crystallization kinetics parameters for glass samples Composition
Ec (kJ/mol)
K0 (min−1 )
m
n
x = 0.1 x = 0.2 x = 0.3
233.03 184.94 174.90
6.48 × 1015 4.02 × 1011 5.30 × 1011
2 3 3
3 4 4
Short-wavelength absorption edges of selenium-based chalcohalide glasses normally reside at the visible and near-infrared region. Based on absorption spectra of these glasses shown in Fig. 6 (plots of absorption coefficient, ˛ versus ), the cut-off wavelength 0 of short-wavelength absorption edges can be estimated
C. Liu et al. / Journal of Alloys and Compounds 474 (2009) 468–472
471
Fig. 5. Plots of ln[−ln(1 − )] versus ln(ˇ) for the x = 0.1 glass with the slope n.
by the empirical tangent method. The empirical optical energy opt Eg (named optical energy) can be calculated by the formula 0 =
Fig. 7. Plots of ln ˛ versus h and their Urbach fittings near short-wavelength absorption edges.
opt
1239.8/Eg [20] and their values are listed in Table 4. The relationship between absorption coefficient ˛ and energy h is given by an exponential function [21]: ˛() = ˛0 exp
h
(7)
Eu
where ˛0 is a constant, h is the Plank constant, is the photon frequency and Eu is the Urbach energy. This formula fits the absorption coefficient data very well at the short-wavelength absorption edge close to absorption tail in most of amorphous materials with absorption coefficient ˛ in the range 1 < ˛ < 104 [22]; however, because Eu and ˛0 are not directly observable from absorption spectra, the value E0 = −Eu ln ˛0 is introduced into Eq. (7), and a new equation is obtained as follows: ln ˛ =
1 (h − E0 ) Eu
according to this equation, E0 can be easily derived by extending the fitted line of ln ˛ versus h to where ln ˛ = 0 and Eu are obtained from the slope of the line. Fig. 7 shows the fitted curves of absorption coefficient data of the studied samples using Eq. (8). Only the lower energy part of the data deviates from the straight natural logarithm line, which may be caused by the absorption overlapping and multi-phonon absorption [23,24]. Table 4, in which the calculated values of Eu and E0 are listed, indicates that 0 moves to opt the short-wavelength region and Eu , E0 and Eg gradually increase when the CsI content becomes higher. In order to distinguish the transition character (direct or indirect) of the short-wavelength absorption edge, the relation between absorption coefficient ˛ and energy h can be written as [25,26]:
(8) (˛ × h)
Table 4 Derived optical characteristic parameters for the studied glasses opt
Composition
0 (nm)
Eg
x = 0.1 x = 0.2 x = 0.3
673.9 649.3 591.3
1.84 1.91 2.10
(eV)
E0 (eV)
Eu (eV)
Eg (eV)
1.65 1.70 1.86
0.088 0.093 0.103
1.77 1.82 1.99
2 2 2
Fig. 6. Vis–NIR absorption spectra of glass samples with the sample thickness 1.4 mm.
1/r
= A(h − Eg )
(9)
where A is a constant, Eg is the Tauc energy [27], and is an index that depends on the transition type of the absorption edge. = 1/2 represents the direct transition, and = 2 represents the indirect transition. The transition character can be decided by whether there is a linear relationship between (˛ × h)1/ ( = 1/2 or 2) and h,
Fig. 8. Relationships of (˛ × h)1/ ( = 1/2 and 2) versus h near the shotwavelength absorption edge of the x = 0.1 glass. Solid line is the guide to eye for linear regime.
472
C. Liu et al. / Journal of Alloys and Compounds 474 (2009) 468–472
Fig. 9. Evolution of the glass structure and energy band with the increase of CsI.
and Eg can be calculated with the slope of the fitted line. Fig. 8 shows relationships of (˛ × h)1/ ( = 1/2 and 2) versus h for the x = 0.1 glass as an example. The values of Eg and for all samples are listed in Table 4. The fitted results reveal that Eg increases with the added CsI and all glasses have the indirect transition in their short-wavelength absorption region. The change of short-wavelength absorption edge may be attributed to the evolution of the glass energy band structure. According to external electrons and electro-negativity of atoms in the studied glasses, it is known that the valence band (VB) and conduction band (CB) are mainly composed of Se 4p3 with lone pairs and Ge 4sp3 with empty orbits, respectively. Fig. 9(a), which represents the evolvement of the glass structure with the increasing content of CsI, shows that the content of Se decreases and lone pairs from Se are attracted not only by empty orbits of Ga3+ but also by the high electronegative I− like Cl− in the paper reported by Calvez et al. [9]. Consequently, the possibilities mentioned above lower the valence band but hardly affect the conduct band. In other words, opt the Eu , Eg and Eg become larger because of the increasing band gap, the sketch map of which is shown in Fig. 9(b). 4. Conclusion The infrared glass materials of (1 − x)(11/16GeSe2 –5/ 16Ga2 Se3 )–xCsI (x = 0.1, 0.2 and 0.3) were synthesized and studied. At x = 0.2, the glass is the most stable and its T is larger than 160 K. Owing to the cooperative effect of the [GeSe4 ], [GaSe4 ], [GeSe3 I] and [GaSe3 I] units with the increase of CsI, Tx , Tp and glass thermal stability criterion parameters of the studied glasses increase initially and then decrease although Tg is decreasing. When the dissolved amount of CsI changes from 10 to 30 mol%, Et and Ec decrease to 267.60 and 174.90 kJ/mol, respectively. The crystallization frequency factor K0 cannot be neglected as the thermal stability of the studied glasses is evaluated by using Ec . During the crystallization process of the three glasses, the x = 0.1
glass possesses of two-dimensional crystal growth and the other two glasses have three-dimensional crystal growth. The glass opt with a higher content of CsI has larger Eu , Eg and Eg , which is attributed to the lower valence band of the glass. In addition, the short-wavelength absorption edges of the three glasses are caused by the indirect transition. These studied results may offer valuable information for the design and crystallization of the new selenium-based chalcohalide glasses. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27]
X.H. Zhang, Y. Guimond, Y. Bellec, J. Non-Cryst. Solids 326–327 (2003) 519–523. A. Rogalski, Infrared Phys. Technol. 43 (2002) 187–210. X.H. Zhang, H.L. Ma, J. Lucas, J. Non-Cryst. Solids 337 (2004) 130–135. Z.Y. Yang, W. Chen, L. Luo, J. Mater. Res. 20 (2005) 2597–2602. A. Dietzel, Glasstech. Ber. 22 (1968) 41–45. M. Saad, M. Poulain, Mater. Sci. Forum 19–20 (1987) 11–18. G. Tang, Z.Y. Yang, L. Luo, W. Chen, J. Alloys Compd 459 (2008) 472–476. Z.Y. Yang, L. Luo, W. Chen, J. Am. Ceram. Soc. 90 (2007) 667–669. L. Calvez, H.L. Ma, J. Lucas, X.H. Zhang, Adv. Mater. 19 (2007) 129–132. M. Lasocka, Mater. Sci. Eng. 23 (1976) 173–177. H.E. Kissinger, J. Res. Nat. Bur. Stand. 57 (1956) 331–336. M.M.A. Imran, D. Bhandari, N.S. Saxena, Physica B 293 (2001) 394–401. P. Agarwal, S. Goel, J.S.P. Rai, A. Kumar, Phys. Stat. Sol. (a) 127 (1991) 363–367. N.P. Bansal, R.H. Soremus, A.J. Bruce, C.T. Moynihan, J. Am. Ceram. Soc. 66 (1983) 233–238. T.R. Dinger, R.S. Rai, G. Thomas, J. Am. Ceram. Soc. 71 (1988) 236–241. J.I. Chu, H.C. Park, C.C. Sorrell, Int. Ceram. Monogr. 1 (1994) 585–589. V.A. Shneidman, M.C. Weinberg, Ceram. Trans. 30 (1993) 225–228. S. Surinach, M.D. Barof, M.M.T. Clavaguera, N. Clavaguera, Mater. Sci. 19 (1984) 3005–3007. K. Matusita, S. Sakka, Phys. Chem. Glasses 20 (1984) 291–296. J.M. Carlsson, B. Hellsing, H.S. Domingos, P.D. Bristowe, Phys. Rev. B 65 (2002) 205122–205132. E. Urbach, Phys. Rev. 92 (1953) 1324–1327. J. Wong, C.A. Angell, Glass Structure by Spectroscopy, Marcel Dekker, New York, 1973. N.F. Mott, Philos. Mag. 22 (1970) 7–9. K. Tanaka, Phys. Rev. B 39 (1989) 1270–1279. J. Tauc, R. Grigorovici, A. Vancu, Phys. Stat. Solidi 15 (1966) 627–634. J. Tauc, Amorphous and Liquid Semiconductors, Plenum Press, New York, 1974. R.A. Smith, Philos. Mag. 2 (1953) 81–84.
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Effect of CsI addition on the thermal properties, crystallization kinetics and short-wavelength absorption edges of GeSe2 –Ga2 Se3 glasses Cunming Liu a,b,∗ , Gao Tang a,b , Lan Luo a , Wei Chen a,∗ a b
Shanghai Institute of Ceramics, Chinese Academy of Sciences, 1295 Dingxi Road, Shanghai 200050, PR China Graduate School of the Chinese Academy of Sciences, Beijing 100039, PR China
a r t i c l e
i n f o
Article history: Received 21 May 2008 Received in revised form 23 June 2008 Accepted 25 June 2008 Available online 3 August 2008 Keywords: Chalcogenide glasses Thermal properties Crystallization kinetics Short-wavelength absorption edges
a b s t r a c t The glasses of (1 − x)(11/16GeSe2 –5/16Ga2 Se3 )–xCsI (x = 0.1, 0.2 and 0.3) were prepared by the melt quench technique. The thermal properties and crystallization kinetics of these glasses were investigated by the differential thermal analysis. It is found that the glass of x = 0.2 is the most stable and that the increase of CsI decreases the activity energy of glass transition Et and the activity energy of glass crystallization Ec . When evaluating the thermal stability of the studied glasses with Ec , the crystallization frequency factor K0 should be considered. The crystallization process of the x = 0.1 glass is determined by two-dimensional crystal growth, while that of the other two glasses is realized by three-dimensional crystal growth. Through analyzing the short-wavelength absorption edges of the glasses in detail, it can be concluded that the opt addition of CsI increases the Urbach energy Eu , optical energy Eg and Tauch energy Eg , and that all of the short-wavelength absorption edges are dominated by the indirect transition. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Selenium-based chalcogenide glasses with the transparence from 1 to 16 m are considered to be an alternative solution to replace the single crystalline germanium (Ge) usually used for thermal cameras and infrared (IR) detectors due to their much lower cost, easier fabrication of complex aspheric IR optical lenses and higher chemical stability [1,2]. Recently, Zhang et al. [3] revealed that selenium-based chalcogenide glasses containing alkali halides can generate nano-crystals to enhance their mechanical properties through heat treatment, and Yang et al. [4] reported that the transparent range of selenium-based chalcogenide glasses can be extended to the partial red light by alkali halides addition. These progresses make the selenium-based chalcogenide glasses much superior to the crystal of Ge; however, there are still some problems for their application, especially the poor thermo-mechanical properties and small transparent range in the visible region. An effective solution to these problems is designing the new selenium-based chalcohalide glasses with the larger visible transparent range and then crystallizing the glasses to produce nano-crystals, which can enhance the thermo-mechanical properties of glasses but cannot
∗ Corresponding author at: Shanghai Institute of Ceramics, Chinese Academy of Sciences, 1295 Dingxi Road, Shanghai 200050, PR China. Tel.: +86 21 52411024; fax: +86 21 52413903. E-mail addresses: [email protected] (C. Liu), [email protected] (W. Chen). 0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.06.116
change the optical properties of glasses significantly. In order to design the new glasses well and perform nano-crystallization on the glasses properly, it is necessary to understand the influence of alkali halides addition on the thermal properties, crystallization kinetics and short-wavelength absorption edges of the glasses. In this work, the glasses of (1 − x)(11/16GeSe2 –5/16Ga2 Se3 )–xCsI (x = 0.1, 0.2 and 0.3) were prepared. Then, their thermal properties, crystallization kinetics and short-wavelength absorption edges were investigated. 2. Experimental Glass samples of (1 − x)(11/16GeSe2 –5/16Ga2 Se3 )–xCsI (x = 0.1, 0.2 and 0.3) were synthesized by melting mixtures of the constituent elements (Ge, Ga, Se all of 99.999% purity, and CsI of 99.95% purity) in evacuated (10−2 Pa) and flame-sealed silica ampoule in a rocking furnace. The mixtures were melted at 1100–1200 K for 10 h. The temperature of the furnace was subsequently lowered to 1000 K and then held for another 2 h. After that, the ampoules were quenched in water, swiftly moved to a preheated furnace and annealed at a temperature lower than the corresponding glass transition temperature for 2 h. Glass rods were obtained by taking them out from the ampoules and finally cut into discs (Ø 9 mm × 2 mm) which were then polished for succeeding experiments. For the analysis of thermal properties and crystallization kinetics, the differential thermal analysis (DTA) measurements of bulk glass pieces (about 45 mg) with the same shape were carried out by a CDR-1P thermal Analyzer (SBIF, Shanghai, PR China) with an accuracy of ±1 K at continuous heating rates (ˇ = 2.5, 5, 7.5, 10, 15 and 20 K/min). The calorimeter was calibrated by the well-known melting temperatures and melting enthalpies of zinc and indium at each heating rate. Pure ␣-Al2 O3 powder was used as the reference material. The glass transition temperature Tg , the starting crystallization temperature Tx and the crystallization peak temperature Tp were obtained by the microprocessor of the thermal analyzer.
C. Liu et al. / Journal of Alloys and Compounds 474 (2009) 468–472
Fig. 1. DTA curves of the as-prepared glasses: (1 − x)(11/16GeSe2 –5/16Ga2 Se3 )–xCsI (x = 0.1, 0.2, 0.3).
In the interest of studying the short-wavelength absorption edges of these glasses, visible–near infrared (vis–NIR) absorption spectra were measured by a Cary 500 spectrophotometer (Varian, Palo Alto, USA). Otherwise, X-ray diffraction (XRD) data were collected using a Rigaku D/MAX-2550V diffractometer (Rigaku, Tokyo, Japan) with Cu K␣ radiation at 40 kV and 40 mA to identify the homogeneity and lack of crystallinity in each sample.
3. Results and discussions 3.1. Thermal properties Typical DTA curves of the three glasses recorded at a heating rate of 10 K/min are shown in Fig. 1. Founded on the values of Tg , Tx , Tp given in Table 1, the common stability criterion T = Tx − Tg [5], the weight stability criterion Hw = T/Tg and the S criterion S = (Tp − Tx )(Tx − Tg )/Tg [6] are used to evaluate the thermal stability of these glasses and their values are listed in Table 1. Generally, the more stable glass has the larger values of these parameters. From Fig. 1 and Table 1, it is found that: firstly, the glasses of x = 0.1 and 0.2 have two exothermic crystallization peaks while the x = 0.3 glass has one exothermic peak; secondly, the x = 0.2 sample is the most stable with a common stability criterion T larger than 160 K; thirdly, when the CsI increases, Tg decreases gradually, whereas Tx , Tp and the thermal stability criterion parameters (T, Hw and S) increase at first and then decrease. These results illustrate that the added CsI alters the glass structure. With regard to the influence of the CsI content on the glass structure, it may be similar to that of some other alkali halides like KBr or CsCl reported in some papers [7,8]. When lower amount of CsI (x ≤ 0.2) is introduced, the [GeSe4 ] units decrease, transforming into the [GeSe3 I] units, which loose the whole glass structure. On account of the strong complexforming ability of Ga3+ , the [GaSe3 ] triangles can easily transform into the [GaSe4 ] tetrahedral to make up network relaxation through capturing the lone pairs from Se [9]; however, [GaSe4 ] cannot sufficiently compensate the loss of [GeSe4 ] because of the appearance of the [GaSe3 I] units. Consequently, Tg somewhat declines. NevertheTable 1 Thermal parameters for the three glass samples
469
Fig. 2. Plots of Tg versus ln(ˇ) for three glass samples.
less, due to the higher space complication caused by the appearance of [GaSe4 ], [GaSe3 I] and [GeSe3 I] units, Tx , Tp and the thermal stability criterion parameters increase. As the content of CsI is higher (x > 0.2), larger amount of [GaSe3 I] and [GeSe3 I] units appear to further decrease [GeSe4 ] and [GaSe4 ] units. The space complication starts to become lower because of the much flabbier glass structure. As a result, Tg sequentially decreases, and Tx , Tp and the thermal stability criterion parameters begin to decrease. Meanwhile, more [GaSe3 I] and [GeSe3 I] units change the crystallization of glasses with the increase of CsI, which represents that two exothermic crystallization peaks are turned into a single exothermic crystallization peak on their DTA curves. 3.2. Analysis of crystallization kinetics The relationship between the glass transition temperature Tg and the heating rate ˇ can be discussed through the empirical equation: Tg = A + B ln(ˇ)
(1)
where A and B are constants for a given glass composition [10]. Plots of Tg versus ln(ˇ) for the prepared samples are shown in Fig. 2. From this figure, the values of A and B can be obtained by using the least square fit and listed in Table 2, and this equation holds good for the studied samples. To obtain the activity energy of glass transition Et , the data are fitted by the Kissinger method [11], which is most commonly used in analyzing crystallization data in DTA experiments. According to the Kissinger equation, the glass transition temperature Tg depends on the heating rate ˇ as follows:
ln
ˇ Tg2
=−
Et + const. RTg
(2)
based on the above relationship, straight lines between ln(ˇ/Tg2 ) and (1000/Tg ) can be given; slopes of these lines yield the values of Et /R where Et corresponds to the glass transition activation energy Table 2 Values of A, B and Et for the studied glasses
Composition
Tg (K)
Tx (K)
Tp (K)
T (K)
Hw
S (K)
Composition
A (K)
B (K)
Et (kJ/mol)
x = 0.1 x = 0.2 x = 0.3
647 619 606
727 782 732
754 799 751
80 163 126
0.124 0.263 0.208
3.34 4.48 3.95
x = 0.1 x = 0.2 x = 0.3
633.72 599.86 585.28
6.41 8.54 8.66
490.55 333.29 267.60
470
C. Liu et al. / Journal of Alloys and Compounds 474 (2009) 468–472
Fig. 3. Plots of ln(ˇ/Tg2 ) versus (1000/Tg ) for three glass samples. Fig. 4. Plots of ln(ˇ/Tp2 ) versus (1000/Tp ) for the studied glasses.
and R is the universal gas constant (R = 8.314 J/(mol K)). The calculated values of Et from the slopes of fitted lines shown in Fig. 3 are listed in Table 2. The result reveals that Et decreases with the increase of CsI like Tg . Concerning Et , it is absorbed by a group of atoms in the glass region when a jump from one metastable state to another is possible [12]. In other words, Et is involved in the molecular motions and rearrangements of the atoms around the glass transition temperature [13]. When the sample is reheated in the DTA furnace, the atoms undergo infrequent transitions between the local potential minima separated by different energy barriers in the configuration space where each local minimum represents a different structure. The more stable local minimum in glass region has lower internal energy. Accordingly, the atoms with minimum activation energy in a glass have higher probability to jump to the metastable (or local minimum) state of lower internal energy and hence a decreasing Tg is observed on the DTA curves when Et decreases. The Kissinger method [11] can be also used in the analysis of exothermic peaks to get the activation energy of glass crystallization Ec with the following relationship:
ln
ˇ
=−
Tp2
Ec + const. RTp
(3)
a new equation from the above equation was developed by Bansal et al. [14] as follows:
ln
ˇ Tp2
=−
Ec − ln RTp
E c
R
+ ln K0
(4)
where K0 is the crystallization frequency factor. From the linear dependence of ln(ˇ/Tp2 ) versus (1000/Tp ) shown in Fig. 4, Ec and K0 can be calculated from slopes of these fitted lines and their values are listed in Table 3. Table 3 reveals that firstly, the minimal value of K0 appears when the composition is x = 0.2; secondly, Ec decreases with the increasing content of CsI and its developing tendency is not consistent with that of the thermal stable criterion parameters in Table 1. Commonly, the glass is more stable when Ec is larger due to a larger
energy barrier represented by a larger Ec . In fact, the whole crystallization is complicated, and it is not exact that the glass thermal stability is judged only by Ec , so the crystallization frequency factor K0 should be considered [15,16]. To evaluate the glass thermal stability accurately, a stricter criterion, the crystallization rate constant K computed with Ec and K0 , was advanced by some authors [17,18]. The formula is listed below:
K = K0 exp
−Ec RTp
(5)
according to this equation and the values of Tp , Ec , K0 in Tables 1 and 3, the calculated values of K are 9.94 × 10−4 , 9.94 × 10−5 and 4.45 × 10−4 min−1 corresponding to the x = 0.1, 0.2 and 0.3 glasses, which illustrates that the x = 0.2 glass is the most stable. For the determination of the crystallization mechanism, the crystallization kinetics parameters n and m, depending on the nucleation process and growth morphology, can be obtained by the following modified Johnson–Mehl–Avrami equation [19]: ln[−ln (1 − )] = −n ln(ˇ) − 1.052
mEc + const. RT
(6)
where is the crystallized volume fraction at any temperature. When the quenched glasses contain no preexisting nuclei n is equal to m + 1, and when the reheated glasses contain a sufficiently large number of nuclei n is equal to m [19]. Also, m = 3, 2, and 1 are for three-, two- and one-dimensional crystal growth, respectively [19]. The value of n can be obtained by calculating the mean value of slopes of curves at different specific temperatures, as plotted in Fig. 5. Since all of glass samples were quenched in water and annealed at the temperatures lower than their own Tg , the value of m can be deduced from n = m + 1. As shown in Table 3, m = 2 and 3 are for the x = 0.1 glass and the other two glasses, respectively, which suggests that two-dimensional crystal growth works in the x = 0.1 glass and three-dimensional crystal growth is dominant in the other two glass matrices when the crystallization process occurs in the three glasses. 3.3. Analysis of short-wavelength absorption edges
Table 3 Crystallization kinetics parameters for glass samples Composition
Ec (kJ/mol)
K0 (min−1 )
m
n
x = 0.1 x = 0.2 x = 0.3
233.03 184.94 174.90
6.48 × 1015 4.02 × 1011 5.30 × 1011
2 3 3
3 4 4
Short-wavelength absorption edges of selenium-based chalcohalide glasses normally reside at the visible and near-infrared region. Based on absorption spectra of these glasses shown in Fig. 6 (plots of absorption coefficient, ˛ versus ), the cut-off wavelength 0 of short-wavelength absorption edges can be estimated
C. Liu et al. / Journal of Alloys and Compounds 474 (2009) 468–472
471
Fig. 5. Plots of ln[−ln(1 − )] versus ln(ˇ) for the x = 0.1 glass with the slope n.
by the empirical tangent method. The empirical optical energy opt Eg (named optical energy) can be calculated by the formula 0 =
Fig. 7. Plots of ln ˛ versus h and their Urbach fittings near short-wavelength absorption edges.
opt
1239.8/Eg [20] and their values are listed in Table 4. The relationship between absorption coefficient ˛ and energy h is given by an exponential function [21]: ˛() = ˛0 exp
h
(7)
Eu
where ˛0 is a constant, h is the Plank constant, is the photon frequency and Eu is the Urbach energy. This formula fits the absorption coefficient data very well at the short-wavelength absorption edge close to absorption tail in most of amorphous materials with absorption coefficient ˛ in the range 1 < ˛ < 104 [22]; however, because Eu and ˛0 are not directly observable from absorption spectra, the value E0 = −Eu ln ˛0 is introduced into Eq. (7), and a new equation is obtained as follows: ln ˛ =
1 (h − E0 ) Eu
according to this equation, E0 can be easily derived by extending the fitted line of ln ˛ versus h to where ln ˛ = 0 and Eu are obtained from the slope of the line. Fig. 7 shows the fitted curves of absorption coefficient data of the studied samples using Eq. (8). Only the lower energy part of the data deviates from the straight natural logarithm line, which may be caused by the absorption overlapping and multi-phonon absorption [23,24]. Table 4, in which the calculated values of Eu and E0 are listed, indicates that 0 moves to opt the short-wavelength region and Eu , E0 and Eg gradually increase when the CsI content becomes higher. In order to distinguish the transition character (direct or indirect) of the short-wavelength absorption edge, the relation between absorption coefficient ˛ and energy h can be written as [25,26]:
(8) (˛ × h)
Table 4 Derived optical characteristic parameters for the studied glasses opt
Composition
0 (nm)
Eg
x = 0.1 x = 0.2 x = 0.3
673.9 649.3 591.3
1.84 1.91 2.10
(eV)
E0 (eV)
Eu (eV)
Eg (eV)
1.65 1.70 1.86
0.088 0.093 0.103
1.77 1.82 1.99
2 2 2
Fig. 6. Vis–NIR absorption spectra of glass samples with the sample thickness 1.4 mm.
1/r
= A(h − Eg )
(9)
where A is a constant, Eg is the Tauc energy [27], and is an index that depends on the transition type of the absorption edge. = 1/2 represents the direct transition, and = 2 represents the indirect transition. The transition character can be decided by whether there is a linear relationship between (˛ × h)1/ ( = 1/2 or 2) and h,
Fig. 8. Relationships of (˛ × h)1/ ( = 1/2 and 2) versus h near the shotwavelength absorption edge of the x = 0.1 glass. Solid line is the guide to eye for linear regime.
472
C. Liu et al. / Journal of Alloys and Compounds 474 (2009) 468–472
Fig. 9. Evolution of the glass structure and energy band with the increase of CsI.
and Eg can be calculated with the slope of the fitted line. Fig. 8 shows relationships of (˛ × h)1/ ( = 1/2 and 2) versus h for the x = 0.1 glass as an example. The values of Eg and for all samples are listed in Table 4. The fitted results reveal that Eg increases with the added CsI and all glasses have the indirect transition in their short-wavelength absorption region. The change of short-wavelength absorption edge may be attributed to the evolution of the glass energy band structure. According to external electrons and electro-negativity of atoms in the studied glasses, it is known that the valence band (VB) and conduction band (CB) are mainly composed of Se 4p3 with lone pairs and Ge 4sp3 with empty orbits, respectively. Fig. 9(a), which represents the evolvement of the glass structure with the increasing content of CsI, shows that the content of Se decreases and lone pairs from Se are attracted not only by empty orbits of Ga3+ but also by the high electronegative I− like Cl− in the paper reported by Calvez et al. [9]. Consequently, the possibilities mentioned above lower the valence band but hardly affect the conduct band. In other words, opt the Eu , Eg and Eg become larger because of the increasing band gap, the sketch map of which is shown in Fig. 9(b). 4. Conclusion The infrared glass materials of (1 − x)(11/16GeSe2 –5/ 16Ga2 Se3 )–xCsI (x = 0.1, 0.2 and 0.3) were synthesized and studied. At x = 0.2, the glass is the most stable and its T is larger than 160 K. Owing to the cooperative effect of the [GeSe4 ], [GaSe4 ], [GeSe3 I] and [GaSe3 I] units with the increase of CsI, Tx , Tp and glass thermal stability criterion parameters of the studied glasses increase initially and then decrease although Tg is decreasing. When the dissolved amount of CsI changes from 10 to 30 mol%, Et and Ec decrease to 267.60 and 174.90 kJ/mol, respectively. The crystallization frequency factor K0 cannot be neglected as the thermal stability of the studied glasses is evaluated by using Ec . During the crystallization process of the three glasses, the x = 0.1
glass possesses of two-dimensional crystal growth and the other two glasses have three-dimensional crystal growth. The glass opt with a higher content of CsI has larger Eu , Eg and Eg , which is attributed to the lower valence band of the glass. In addition, the short-wavelength absorption edges of the three glasses are caused by the indirect transition. These studied results may offer valuable information for the design and crystallization of the new selenium-based chalcohalide glasses. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27]
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