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Homogeneous–inhomogeneous models of Agx(Ge0.25Se0.75)100−x bulk glasses
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Physica B 389 (2007) 77–82 www.elsevier.com/locate/physb
Homogeneous–inhomogeneous models of Agx(Ge0.25Se0.75)100x bulk glasses B. Arcondoa,, M.A. Uren˜aa, A. Piarristeguya, A. Pradelb, M. Fontanaa a
LSA, Facultad de Ingenierı´a, Departamento de Fisica, Universidad de Buenos Aires, Paseo Colon 850, 1063 Buenos Aires, Argentina b LPMC (UMR 5617), Montpellier Cedex, France
Abstract Ge–Se system presents an extensive glass forming composition range even when different metals (Ag, Sb, Bi) are added. In spite that the addition of Ag (up to 30 at%) to Ge–Se does not affect substantially the glass forming tendency, it impacts significantly on the transport properties. (Ge0.25Se0.75)100xAgx is a fast ionic conductor with xX8 at% whereas it is a semiconductor for xo8 at%. This behavior is also reported in other chalcogenide systems containing Ag. One of the structural models proposed to explain the relation between transport and structure is based on an intrinsically inhomogeneous structure where zones rich in metals coexist with zones of the host material. Field effect scanning electron microscopy (FE-SEM) is performed on amorphous (Ge0.25Se0.75)100xAgx bulk samples. These results appear to sustain this model. However previous structural and thermal studies oppose it. Mo¨ssbauer spectrometry on samples (0pxp25) containing 0.5 at% of 57Fe is performed at Tp300 K. The main contribution to the glasses spectra correspond to low spin Fe2+ in octahedral coordination and high spin Fe2+ in distorted octahedral environments. The relative population of both sites changes continuously as Ag concentration varies denoting that the change in the transport behavior obeys to a percolation phenomenon. The low temperature results are discussed with the aim to throw light on the controversy about the homogeneity–inhomogeneity of the studied bulk glasses. r 2006 Elsevier B.V. All rights reserved. PACS: 71.23; 72.80; 82.45 Keywords: Chalcogenide glasses; Mo¨ssbauer spectroscopy; Fast ionic conductors
1. Introduction The physical and chemical properties of chalcogenide glasses are strongly dependent on their structure. The knowledge of the glass structure is necessary for obtaining properties with specific applications. The fundamental structural unit of pure chalcogen glasses is based on a single atom and as such the short range order is rather straightforward and Se coordination number is n ¼ 2. For binary glasses the situation is more complex [1]. In the family of chalcogenide glasses GeySe1y system is considered an excellent glass former for atomic fractions yp0.43 [2]. The structure of this network glass has been Corresponding author. Tel.: +54 11 43429184x279; fax: +54 11 43311852. E-mail address: barcond@fi.uba.ar (B. Arcondo).
0921-4526/$ - see front matter r 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.physb.2006.07.028
widely investigated either by means of X-rays and neutron diffraction [3–5] or employing local probes [6]. The basic polyhedral unit in the GeSe2 glass is a tetrahedron centered on the tathogen, e.g. GeSe4/2. Ge coordination number is n ¼ 4. In glasses with yo0.33 the Ge centered tetrahedra are connected by Se bridges whereas in glasses with y40.33 the tetrahedra are either corner sharing or edge sharing [7]. As a consequence, as y increases (0pyp0.43) /nS also increases (2p/nSp2.86) and the structure becomes increasingly rigid due to the reinforcement of the constraints, that is, the interatomic forces that control bondbending and bond stretching. Mechanical equilibrium would be attained when the number of constraints per atom equals the dimensionality of the network. Philips [8] and Thorpe [9] predicted that in three-dimensional covalent networks this condition would be fulfilled for /nS ¼ 2.4.
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Boolchand employing 119Sn and 125Te in transmission geometry and 129I in emission Mo¨ssbauer spectroscopy as well as Raman spectroscopy observed two inequivalent environments for the atom probe location. A tetrahedral site plus a non-tetrahedral site are reported for Sn bonded to Se [10] whereas two-fold coordinated Te parents either in Ge–Te–Ge bonds (A) or in Se–Te–Se or Se–Te–Ge bonds (B) originate two inequivalent 129I environments [11]. These inequivalent environments are attributed to intrinsic inhomogeneity in GeySe1y glasses where a nanophase formed by marginally rigid Ge2(Se1/2)6 units segregated from the Ge(Se1/2)4 bearing backbone with the consequent loss in the connectivity of the network and correspondingly a decrease in the glass temperature, Tg, for y40.31 [12]. When Ag or other metals are added to GeySe1y system the glass forming tendency is not substantially affected whereas they impact notably on the transport properties. GeySe1y glasses are semiconducting but when Ag is added above certain threshold concentration, Agx[GeySe1y]100x glasses behave as fast ionic conductors [13,14]. This behavior is also reported in other chalcogenide systems containing Ag. Ag–Ge–Se–Sb, [15] and Ag–Ge–S [15,16] glassy systems also present a huge variation of the Ag+ ionic conductivity with the concentration. Although the chemical nature of these systems is similar, notable differences are depicted in the s vs. x curves: the conductivity regime change is abrupt in Ag–Ge–Se glasses and it occurs at higher Ag concentration than in the abovementioned glasses [14]. Several models have been proposed to interpret this behavior, some of them attribute to the glasses an intrinsically inhomogeneous structure where zones rich in metals coexist with zones of the host material [17]. In this direction are addressed the works of Pradel and collaborators [18] and Mitkova and collaborators [19]. In the first one [18] glasses of the Ag2S–GeS2, Ag2S–As2S3 and Ag2S–GeS–GeS2 systems are studied. The evolution of the room temperature conductivity with Ag concentration is analyzed and an abrupt jump is reported for Ag2S–GeS2 and Ag2S–As2S3 glasses at 7 at% Ag. Correspondingly, scanning electron micrographs of samples with 7 at% Ag and below this concentration evidence two zones, zones rich in Ag (bright) and other where Ag is scarce (dark). The samples with higher concentration present dark clusters immersed in a continuous bright matrix whereas the samples with lower concentration present bright clusters in a dark matrix. On the other hand, the conductivity of Ag2S–GeS–GeS2 changes smoothly and scanning electron microscopy shows homogeneous glasses for any magnification. Indeed, when regions of the Ag-rich phase start to connect the change in the conductivity regime occurs. 2. Experimental procedure Bulk glasses of composition (Ge0.25Se0.75)100xAgx with x ¼ 0, 0.5, 1, 3, 5, 7, 8, 9, 9.3, 9.6, 9.9, 10, 15, 20 and
25 at%, were prepared from high-purity (4 N) elements by melt-quenching technique using an ice-water bath as described in Ref. [20]. Glass samples are named Agx according to their concentration in at%. The glass structure was characterized by X-ray diffractometry (XRD) with Cu Ka radiation. Calorimetric scans of powdered samples were registered under continuous heating regime (scan rates b ¼ 10, 20, 40 and 80 K min1) using a Perkin Elmer DSC-7 under dynamic Ar atmosphere as described elsewhere [21]. The electrical conductivity measurements were performed on bulk samples using the impedance spectroscopy technique in the frequency range from 5 Hz to 2 MHz in the temperature range 293–363 K as is described in Ref. [14]. Measurements of dc conductivity were performed at room temperature on samples with higher resistance using a Keithley 617 electrometer. FE-SEM measurements were carried out employing an acceleration voltage of 12 kV in a Cambridge S 360 instrument with an Oxford ISIS-300 microanalytical system using finish fracture of glass samples and also polished samples. In addition, [(Ge0.25Se0.75)100xAgx]0.995Fe0.005 glasses and Ag2Se, Ag8GeSe6 and Ge25Se75 crystalline standards with 0.5 at% Fe were prepared for Mo¨ssbauer experiments employing Fe isotopically enriched with 90% 57Fe. The measurements were performed at different temperatures, in transmission geometry, employing a 57Co(Rh) source. Measurements at room temperature (300 K), in a mixture of dry ice and acetone (195 K) and in liquid nitrogen (77 K) were performed. Mo¨ssbauer spectra were fitted employing sites or quadrupole splitting (D) distributions with Normos programs Dist and Site [22]. Isomer Shift (d) is reported relative to BCC-Fe. 3. Results and discussion Fig. 1 shows several representative calorimetric curves obtained with b ¼ 10 K min1. Samples Ag10, Ag15 and Ag20 have a similar calorimetric behavior showing a single glass transition and two exothermic peaks. Sample Ag25 is quite different. It presents a different shape of the glass transition shift, and a small exothermic peak (intermediate peak) between the two main peaks. The changes with composition of the glass transition temperature, Tg (defined as the temperature of the curve inflexion point in the glass transition), and the crystallization peak temperatures, Tp, are reported in Table 1. The results of a unique Tg in each glass transition do not agree with those reported in Ref. [19] where the presence of two Tg gives support to the model of inhomogeneous glasses. These differences can be attributed to the composition, to the samples preparation or to the use of different calorimetric techniques (they employ modulated DSC). In (Ge0.25Se0.75)100xAgx glasses, the glass temperature Tg is almost constant and is equal to Tg of the binary Ge0.25Se0.75 glass [23]. This fact indicates that the ternary
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Fig. 2. Conductivity of (Ge0.25Se0.75)100xAgx amorphous samples determined by impedance spectroscopy at room temperature.
Fig. 1. DSC scans with a constant heating rate of 10 K/m. From top to bottom samples Ag10, Ag15, Ag20 and Ag25 are shown.
Table 1 Glass temperature, Tg, first crystallization peak temperature, T 1p , second crystallization peak temperature, T 2p , and intermediate peak temperature, T int p , corresponding to the DSC curves depicted in Fig. 1. Sample
Tg73 K
T 1p 70.5 K
Ag10 Ag15 Ag20 Ag25
494 488 488 496
579.1 578.3 580.0 595.0
T int p 70.5 K
T 2p 70.5 K
615.2
672.8 665.0 657.5 652.4
glass thermal stability is controlled by the binary GeSe matrix and is almost independent of the Ag content. XRD performed on samples heated at a constant rate in the DSC up to different temperatures show that the first exothermic peak corresponds to a primary crystallization of Ag8GeSe6 and the second one to GeSe2 as was reported in Ref. [21].
The conductivity results obtained by means of impedance spectroscopy are presented in Fig. 2. Two conductivity regimes are observed: semiconducting behavior for xp7 at% and ionic conductivity for xX8 at%. Conductivity jumps abruptly from 1012 to 105 S cm1 in agreement with the results reported by Kawasaki et al. [13] for Ag–Ge–Se system and similarly to the behavior of Ag–Ge–S and Ag–As–Se reported by Pradel et al. [18]. DC measurements contribute to corroborate whether conduction is ionic (sample polarization occurs) or electronic. FE-SEM observations are reported in Fig. 3 for Ag15, Ag20 and Ag25 polished samples. The bright zones represent 36.8% and 63.2% of the total area for Ag15 and Ag20 and the size of the dark zones are in both samples 0.5odo0.8 mm. These results are in agreement to similar results presented in Ref. [18] for ionic conducting glasses with abrupt conductivity jump. Sample Ag25 is homogeneous in the observed magnification. FE-SEM images suggest that the huge jump in conductivity can be attributed to a percolative phenomenon at 7 at%oxo8 at%. That is, Ag rich bright zones isolated in a non-conducting matrix percolate for a threshold Ag concentration about 7.5 at% giving place to fast ionic conductivity. Bright zones may be similar either to Ag8GeSe6 or to Ag2Se ionic conductors, though these zones are amorphous. The fact that the first crystallization product observed [21] is Ag8GeSe6 let us to assume that Ag rich zones composition may be not far from Ag8GeSe6. Interpolating the overall density d of Ag7.5 (Ge0.25Se0.75)92.5 sample from the density values reported in [20] and assuming that all Ag atoms are in the bright zones which local density dB is lower than crystalline Ag8GeSe6 density (7.1 g cm3), the dark zones local density dD can be estimated and also the bright zones volume fraction xV(B). With d ¼ 4.8 g cm3, assuming dB ¼ 6 g cm3 (85%
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Ag2Se density (8.2 g cm3). It is concluded that the Scher–Zallen criteria is not satisfactory fulfilled in the last case. Room temperature Mo¨ssbauer spectra either for crystalline or for glass samples present two main contributions: high spin (HS) Fe2+ sites and low spin (LS) Fe2+ sites, as was discussed elsewhere [25]. The first sites are not present in Ag0 glasses but as Ag concentration increases the population of HS 2+ sites increases. The relative area of HS Fe2+ sites is above 60% in Ag25 spectrum. Room and low temperature Mo¨ssbauer spectra are depicted in Fig. 4 for sample Ag10. The spectra were fitted employing two D distributions, one for high spin Fe2+ sites in distorted octahedral environments and the other for low spin Fe2+ sites in octahedral coordination. The area of both distributions increase as temperature (T) decreases, but the presence of two different growth rates denote differences in the Lamb-Mo¨ssbauer factor (f) associated to each environment. From the dependence of f on T one can write [26] ln
3E 2g f A ¼ ln ¼ f0 A0 Mc2 KYD ( ) Z Z T 2 YD =T x T 20 YD =T 0 x dx 2 dx , ex 1 ex 1 Y2D 0 YD 0 ð1Þ
where A is the integrated area of each contribution at T and A0 the corresponding area at T0, Eg ¼ 14.4 keV, M the atomic mass of 57Fe, K the Boltzmann constant, YD the Debye temperature and T0 ¼ 300 K. As R Y =T TpYD I YD =T ¼ 0 D x=ðex 1Þ dx does not depend strongly on T. Therefore I(YD/T) can be replaced by its average value in the range ToYD. A linear dependence of ln(A/A0) on T2 is verified and, from the corresponding slope, m, YD is determined for HS and LS sites as "
3E 2g YD ¼ Mc2 Km
Fig. 3. Scanning electron micrographs of three glasses: from top to bottom Ag15, Ag20 and Ag25 polished samples.
crystal density), dD and xV (B) result 4.62 g cm3 and 0.13, respectively. This volume fraction is in good agreement with the well-known Scher–Zallen threshold for volume percolation near 0.15 [24]. Analogous calculations performed assuming bright zones in Ag7.5 (Ge0.25Se0.75)92.5 samples correspond to Ag2Se, let us to estimate dD ¼ 4.57 g cm3 and xV(B) ¼ 9.47 for dB ¼ 7 g cm3. The local density dB is assumed lower than crystalline
#1=3
1=3 I YD =T .
The Debye temperatures obtained are YD ¼ 290 K for HS Fe2+sites and YD ¼ 370 for LS Fe2+ sites for sample Ag10. The remarkable difference of YD, correspondingly f, for both environments is hard to justify in the case of a continuous random network and one would be tempted to attribute this difference to intrinsic segregation. Nevertheless, as Fe does not substitute any element of the network, this argument cannot be used straightforward. [25] In any case, if glasses were homogeneous, higher f environments can be attributed to Fe immersed in the GeSe backbone while lower f corresponds to Fe located in zones were the intermediate range order (IRO) has been affected by Ag [20]. If inhomogeneous glasses were assumed, according to the microscopy results, as Ag concentration increases the lower YD regions grow in correspondence to
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Fig. 4. Mo¨ssbauer spectra of sample Ag10 obtained at different T and fitted employing two D hyperfine distributions with linear dependence of d on D. On the left, the corresponding d and D distributions are depicted.
the increase of the bright zones in Fig. 3 whereas zones Ag depleted with a persistent IRO correspond to higher YD.
4. Conclusions XRD results [20] plus the determination of a single glass transition contribute to the homogeneous glass model. However, the abrupt jump in conductivity as well as SEM studies strongly support the heterogeneous glass model. That is, Ag rich bright zones coexist with Ag depleted dark regions. The crystallization of Ag8GeSe6 in the first crystallization peak at a relatively low temperature suggests that long distance diffusion is not needed in the nucleation and growth processes. This result encourages us to propose that bright zones approximately correspond to Ag8GeSe6. Under this assumption the bright zones volume fraction is estimated for Ag7.5 (Ge0.25Se0.75)92.5 in good agreement with the Scher–Zallen percolation criteria for inhomoge-
neous glasses. However, an accurate determination of the composition of both zones is still pending. 57 Fe Mo¨ssbauer results are not self sufficient in order to solve this controversy but are not contradictory to the previous results. HS Fe2+ plus LS Fe2+ are the main contributions to the spectra. The increase of HS Fe2+ environments with Ag concentration is observed. Debye temperatures corresponding to both environ ments are compared (YD(HS site)oYD(LS site)) and a remarkable difference, ascribable to intrinsic hetero geneity of the glasses, is observed. Therefore, HS Fe2+ sites in glasses can be attributed to Ag rich bright zones whereas LS Fe2+ sites can be attributed to Ag depleted dark zones. Further experiments are running. Neutron diffraction as a function of temperature in order to obtain detailed information of the crystallization products and a systematic Mo¨ssbauer study in a wider temperature range not only in the ionic conducting regime but also in semiconducting glasses.
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Acknowledgements The support of CNRS (France) and SECYT (Argentine) through ECOS Project A03E03 is greatly acknowledged. The argentine authors acknowledge the support of ANPCyT and Universidad de Buenos Aires. The friendly collaboration of Dr. M.F. Van Raap with low temperature Mo¨ssbauer experiments is also acknowledged. References [1] S.R. Elliott, in: J. Zarzycki (Ed.), Glasses and Amorphous Materials, Materials Science and Technology, vol.9, VCH, Weinheim, 1991 (Chapt. 7). [2] R. Azoulay, H. Thibierge, A. Brenac, J. Non Cryst. Solids 18 (1975) 33. [3] I.T. Penfold, P.S. Salmon, Phys. Rev. Lett. 67 (1991) 97. [4] I. Petri, P.S. Salmon, Phys. Chem. Glasses 43C (2002) 185. [5] P. Armand, A. Iban˜ez, H. Dexpert, D. Bittencourt, D. Raoux, E. Philippot, J. Phys. IV Coll. C2 (2) (1992) C2-189. [6] P. Boolchand, in: D. Adler, B.B. Schwartz, M.C. Steele, (Ed.), Physical Properties of Amorphous Materials, Plenum Press, New York, 1985, pp. 221–260. See also P. Boolchand, W. Bresser, M. Zhang, Y. Wu, J. Wells, R.N. Enzweiler, J. Non-Cryst. Solids 182 (1995) 143. [7] M. Micoulaut, Physica B 226 (1996) 268. [8] J.C. Philips, J. Non Cryst. Solids 34 (1979) 153; J.C. Philips, J. Non Cryst. Solids 43 (1981) 37. [9] M.F. Thorpe, J. Non Cryst. Solids 57 (1983) 355.
[10] P. Boolchand, J. Grothaus, W.J. Bresser, P. Suranyi, Phys. Rev. B 25 (1982) 2975. [11] W. Bresser, P. Boolchand, P. Suranyi, Phys. Rev. Lett. 56 (1986) 2493. [12] P. Boolchand, W.J. Bresser, Philos. Mag. B 80 (2000) 1757. [13] M. Kawasaki, J. Kawamura, Y. Nakamura, M. Aniya, Solid State Ion. 123 (1999) 259. [14] M.A. Uren˜a, A.A. Piarristeguy, M. Fontana, B. Arcondo, Solid State Ion. 176 (2005) 505. [15] E. Bychkov, V. Tsegelnik, Yu. Vlasov, A. Pradel, M. Ribes, J. Non Cryst. Solids 208 (1996) 1. [16] E. Bychkov, D.L. Price, C.J. Benmore, A.C. Hannon, Solid State Ion. 154&155 (2002) 349. [17] G.N. Greaves, Solid State Ion. 105 (1998) 243. [18] A. Pradel, N. Kuwata, M. Ribes, J. Phys. Condens. Matter 15 (2003) S1561. [19] Y. Wang, M. Mitkova, D.G. Georgiev, S. Mamedov, P. Boolchand, J. Phys. Condens. Matter 15 (2003) S1573. [20] A.A. Piarristeguy, M. Mirandou, M. Fontana, B. Arcondo, J. Non Cryst. Solids 273 (2000) 30. [21] M.A. Uren˜a, M. Fontana, B. Arcondo, M.T. Clavaguera-Mora, J. Non Cryst. Solids 320 (2003) 151. [22] R.A. Brand, NORMOS Programs, 1990. [23] X. Feng, W.J. Bresser, P. Boolchand, Phys. Rev. Lett. 78 (1997) 4422. [24] H. Scher, R. Zallen, J. Chem. Phys. 53 (1970) 3759 see also; R. Zallen, H. Scher, Phys. Rev. B 4 (1971) 4471–4479. [25] B. Arcondo, M. Uren˜a, A. Piarristeguy, M. Fontana, J. Metastable Nanocrystall. Mater. 22 (2004) 45. [26] P. Boolchand, W. Bresser, M. Zhang, Y. Wu, J. Wells, R.N. Enzweiler, J. Non Crystall. Solids 182 (1995) 143.
Physica B 389 (2007) 77–82 www.elsevier.com/locate/physb
Homogeneous–inhomogeneous models of Agx(Ge0.25Se0.75)100x bulk glasses B. Arcondoa,, M.A. Uren˜aa, A. Piarristeguya, A. Pradelb, M. Fontanaa a
LSA, Facultad de Ingenierı´a, Departamento de Fisica, Universidad de Buenos Aires, Paseo Colon 850, 1063 Buenos Aires, Argentina b LPMC (UMR 5617), Montpellier Cedex, France
Abstract Ge–Se system presents an extensive glass forming composition range even when different metals (Ag, Sb, Bi) are added. In spite that the addition of Ag (up to 30 at%) to Ge–Se does not affect substantially the glass forming tendency, it impacts significantly on the transport properties. (Ge0.25Se0.75)100xAgx is a fast ionic conductor with xX8 at% whereas it is a semiconductor for xo8 at%. This behavior is also reported in other chalcogenide systems containing Ag. One of the structural models proposed to explain the relation between transport and structure is based on an intrinsically inhomogeneous structure where zones rich in metals coexist with zones of the host material. Field effect scanning electron microscopy (FE-SEM) is performed on amorphous (Ge0.25Se0.75)100xAgx bulk samples. These results appear to sustain this model. However previous structural and thermal studies oppose it. Mo¨ssbauer spectrometry on samples (0pxp25) containing 0.5 at% of 57Fe is performed at Tp300 K. The main contribution to the glasses spectra correspond to low spin Fe2+ in octahedral coordination and high spin Fe2+ in distorted octahedral environments. The relative population of both sites changes continuously as Ag concentration varies denoting that the change in the transport behavior obeys to a percolation phenomenon. The low temperature results are discussed with the aim to throw light on the controversy about the homogeneity–inhomogeneity of the studied bulk glasses. r 2006 Elsevier B.V. All rights reserved. PACS: 71.23; 72.80; 82.45 Keywords: Chalcogenide glasses; Mo¨ssbauer spectroscopy; Fast ionic conductors
1. Introduction The physical and chemical properties of chalcogenide glasses are strongly dependent on their structure. The knowledge of the glass structure is necessary for obtaining properties with specific applications. The fundamental structural unit of pure chalcogen glasses is based on a single atom and as such the short range order is rather straightforward and Se coordination number is n ¼ 2. For binary glasses the situation is more complex [1]. In the family of chalcogenide glasses GeySe1y system is considered an excellent glass former for atomic fractions yp0.43 [2]. The structure of this network glass has been Corresponding author. Tel.: +54 11 43429184x279; fax: +54 11 43311852. E-mail address: barcond@fi.uba.ar (B. Arcondo).
0921-4526/$ - see front matter r 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.physb.2006.07.028
widely investigated either by means of X-rays and neutron diffraction [3–5] or employing local probes [6]. The basic polyhedral unit in the GeSe2 glass is a tetrahedron centered on the tathogen, e.g. GeSe4/2. Ge coordination number is n ¼ 4. In glasses with yo0.33 the Ge centered tetrahedra are connected by Se bridges whereas in glasses with y40.33 the tetrahedra are either corner sharing or edge sharing [7]. As a consequence, as y increases (0pyp0.43) /nS also increases (2p/nSp2.86) and the structure becomes increasingly rigid due to the reinforcement of the constraints, that is, the interatomic forces that control bondbending and bond stretching. Mechanical equilibrium would be attained when the number of constraints per atom equals the dimensionality of the network. Philips [8] and Thorpe [9] predicted that in three-dimensional covalent networks this condition would be fulfilled for /nS ¼ 2.4.
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Boolchand employing 119Sn and 125Te in transmission geometry and 129I in emission Mo¨ssbauer spectroscopy as well as Raman spectroscopy observed two inequivalent environments for the atom probe location. A tetrahedral site plus a non-tetrahedral site are reported for Sn bonded to Se [10] whereas two-fold coordinated Te parents either in Ge–Te–Ge bonds (A) or in Se–Te–Se or Se–Te–Ge bonds (B) originate two inequivalent 129I environments [11]. These inequivalent environments are attributed to intrinsic inhomogeneity in GeySe1y glasses where a nanophase formed by marginally rigid Ge2(Se1/2)6 units segregated from the Ge(Se1/2)4 bearing backbone with the consequent loss in the connectivity of the network and correspondingly a decrease in the glass temperature, Tg, for y40.31 [12]. When Ag or other metals are added to GeySe1y system the glass forming tendency is not substantially affected whereas they impact notably on the transport properties. GeySe1y glasses are semiconducting but when Ag is added above certain threshold concentration, Agx[GeySe1y]100x glasses behave as fast ionic conductors [13,14]. This behavior is also reported in other chalcogenide systems containing Ag. Ag–Ge–Se–Sb, [15] and Ag–Ge–S [15,16] glassy systems also present a huge variation of the Ag+ ionic conductivity with the concentration. Although the chemical nature of these systems is similar, notable differences are depicted in the s vs. x curves: the conductivity regime change is abrupt in Ag–Ge–Se glasses and it occurs at higher Ag concentration than in the abovementioned glasses [14]. Several models have been proposed to interpret this behavior, some of them attribute to the glasses an intrinsically inhomogeneous structure where zones rich in metals coexist with zones of the host material [17]. In this direction are addressed the works of Pradel and collaborators [18] and Mitkova and collaborators [19]. In the first one [18] glasses of the Ag2S–GeS2, Ag2S–As2S3 and Ag2S–GeS–GeS2 systems are studied. The evolution of the room temperature conductivity with Ag concentration is analyzed and an abrupt jump is reported for Ag2S–GeS2 and Ag2S–As2S3 glasses at 7 at% Ag. Correspondingly, scanning electron micrographs of samples with 7 at% Ag and below this concentration evidence two zones, zones rich in Ag (bright) and other where Ag is scarce (dark). The samples with higher concentration present dark clusters immersed in a continuous bright matrix whereas the samples with lower concentration present bright clusters in a dark matrix. On the other hand, the conductivity of Ag2S–GeS–GeS2 changes smoothly and scanning electron microscopy shows homogeneous glasses for any magnification. Indeed, when regions of the Ag-rich phase start to connect the change in the conductivity regime occurs. 2. Experimental procedure Bulk glasses of composition (Ge0.25Se0.75)100xAgx with x ¼ 0, 0.5, 1, 3, 5, 7, 8, 9, 9.3, 9.6, 9.9, 10, 15, 20 and
25 at%, were prepared from high-purity (4 N) elements by melt-quenching technique using an ice-water bath as described in Ref. [20]. Glass samples are named Agx according to their concentration in at%. The glass structure was characterized by X-ray diffractometry (XRD) with Cu Ka radiation. Calorimetric scans of powdered samples were registered under continuous heating regime (scan rates b ¼ 10, 20, 40 and 80 K min1) using a Perkin Elmer DSC-7 under dynamic Ar atmosphere as described elsewhere [21]. The electrical conductivity measurements were performed on bulk samples using the impedance spectroscopy technique in the frequency range from 5 Hz to 2 MHz in the temperature range 293–363 K as is described in Ref. [14]. Measurements of dc conductivity were performed at room temperature on samples with higher resistance using a Keithley 617 electrometer. FE-SEM measurements were carried out employing an acceleration voltage of 12 kV in a Cambridge S 360 instrument with an Oxford ISIS-300 microanalytical system using finish fracture of glass samples and also polished samples. In addition, [(Ge0.25Se0.75)100xAgx]0.995Fe0.005 glasses and Ag2Se, Ag8GeSe6 and Ge25Se75 crystalline standards with 0.5 at% Fe were prepared for Mo¨ssbauer experiments employing Fe isotopically enriched with 90% 57Fe. The measurements were performed at different temperatures, in transmission geometry, employing a 57Co(Rh) source. Measurements at room temperature (300 K), in a mixture of dry ice and acetone (195 K) and in liquid nitrogen (77 K) were performed. Mo¨ssbauer spectra were fitted employing sites or quadrupole splitting (D) distributions with Normos programs Dist and Site [22]. Isomer Shift (d) is reported relative to BCC-Fe. 3. Results and discussion Fig. 1 shows several representative calorimetric curves obtained with b ¼ 10 K min1. Samples Ag10, Ag15 and Ag20 have a similar calorimetric behavior showing a single glass transition and two exothermic peaks. Sample Ag25 is quite different. It presents a different shape of the glass transition shift, and a small exothermic peak (intermediate peak) between the two main peaks. The changes with composition of the glass transition temperature, Tg (defined as the temperature of the curve inflexion point in the glass transition), and the crystallization peak temperatures, Tp, are reported in Table 1. The results of a unique Tg in each glass transition do not agree with those reported in Ref. [19] where the presence of two Tg gives support to the model of inhomogeneous glasses. These differences can be attributed to the composition, to the samples preparation or to the use of different calorimetric techniques (they employ modulated DSC). In (Ge0.25Se0.75)100xAgx glasses, the glass temperature Tg is almost constant and is equal to Tg of the binary Ge0.25Se0.75 glass [23]. This fact indicates that the ternary
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Fig. 2. Conductivity of (Ge0.25Se0.75)100xAgx amorphous samples determined by impedance spectroscopy at room temperature.
Fig. 1. DSC scans with a constant heating rate of 10 K/m. From top to bottom samples Ag10, Ag15, Ag20 and Ag25 are shown.
Table 1 Glass temperature, Tg, first crystallization peak temperature, T 1p , second crystallization peak temperature, T 2p , and intermediate peak temperature, T int p , corresponding to the DSC curves depicted in Fig. 1. Sample
Tg73 K
T 1p 70.5 K
Ag10 Ag15 Ag20 Ag25
494 488 488 496
579.1 578.3 580.0 595.0
T int p 70.5 K
T 2p 70.5 K
615.2
672.8 665.0 657.5 652.4
glass thermal stability is controlled by the binary GeSe matrix and is almost independent of the Ag content. XRD performed on samples heated at a constant rate in the DSC up to different temperatures show that the first exothermic peak corresponds to a primary crystallization of Ag8GeSe6 and the second one to GeSe2 as was reported in Ref. [21].
The conductivity results obtained by means of impedance spectroscopy are presented in Fig. 2. Two conductivity regimes are observed: semiconducting behavior for xp7 at% and ionic conductivity for xX8 at%. Conductivity jumps abruptly from 1012 to 105 S cm1 in agreement with the results reported by Kawasaki et al. [13] for Ag–Ge–Se system and similarly to the behavior of Ag–Ge–S and Ag–As–Se reported by Pradel et al. [18]. DC measurements contribute to corroborate whether conduction is ionic (sample polarization occurs) or electronic. FE-SEM observations are reported in Fig. 3 for Ag15, Ag20 and Ag25 polished samples. The bright zones represent 36.8% and 63.2% of the total area for Ag15 and Ag20 and the size of the dark zones are in both samples 0.5odo0.8 mm. These results are in agreement to similar results presented in Ref. [18] for ionic conducting glasses with abrupt conductivity jump. Sample Ag25 is homogeneous in the observed magnification. FE-SEM images suggest that the huge jump in conductivity can be attributed to a percolative phenomenon at 7 at%oxo8 at%. That is, Ag rich bright zones isolated in a non-conducting matrix percolate for a threshold Ag concentration about 7.5 at% giving place to fast ionic conductivity. Bright zones may be similar either to Ag8GeSe6 or to Ag2Se ionic conductors, though these zones are amorphous. The fact that the first crystallization product observed [21] is Ag8GeSe6 let us to assume that Ag rich zones composition may be not far from Ag8GeSe6. Interpolating the overall density d of Ag7.5 (Ge0.25Se0.75)92.5 sample from the density values reported in [20] and assuming that all Ag atoms are in the bright zones which local density dB is lower than crystalline Ag8GeSe6 density (7.1 g cm3), the dark zones local density dD can be estimated and also the bright zones volume fraction xV(B). With d ¼ 4.8 g cm3, assuming dB ¼ 6 g cm3 (85%
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Ag2Se density (8.2 g cm3). It is concluded that the Scher–Zallen criteria is not satisfactory fulfilled in the last case. Room temperature Mo¨ssbauer spectra either for crystalline or for glass samples present two main contributions: high spin (HS) Fe2+ sites and low spin (LS) Fe2+ sites, as was discussed elsewhere [25]. The first sites are not present in Ag0 glasses but as Ag concentration increases the population of HS 2+ sites increases. The relative area of HS Fe2+ sites is above 60% in Ag25 spectrum. Room and low temperature Mo¨ssbauer spectra are depicted in Fig. 4 for sample Ag10. The spectra were fitted employing two D distributions, one for high spin Fe2+ sites in distorted octahedral environments and the other for low spin Fe2+ sites in octahedral coordination. The area of both distributions increase as temperature (T) decreases, but the presence of two different growth rates denote differences in the Lamb-Mo¨ssbauer factor (f) associated to each environment. From the dependence of f on T one can write [26] ln
3E 2g f A ¼ ln ¼ f0 A0 Mc2 KYD ( ) Z Z T 2 YD =T x T 20 YD =T 0 x dx 2 dx , ex 1 ex 1 Y2D 0 YD 0 ð1Þ
where A is the integrated area of each contribution at T and A0 the corresponding area at T0, Eg ¼ 14.4 keV, M the atomic mass of 57Fe, K the Boltzmann constant, YD the Debye temperature and T0 ¼ 300 K. As R Y =T TpYD I YD =T ¼ 0 D x=ðex 1Þ dx does not depend strongly on T. Therefore I(YD/T) can be replaced by its average value in the range ToYD. A linear dependence of ln(A/A0) on T2 is verified and, from the corresponding slope, m, YD is determined for HS and LS sites as "
3E 2g YD ¼ Mc2 Km
Fig. 3. Scanning electron micrographs of three glasses: from top to bottom Ag15, Ag20 and Ag25 polished samples.
crystal density), dD and xV (B) result 4.62 g cm3 and 0.13, respectively. This volume fraction is in good agreement with the well-known Scher–Zallen threshold for volume percolation near 0.15 [24]. Analogous calculations performed assuming bright zones in Ag7.5 (Ge0.25Se0.75)92.5 samples correspond to Ag2Se, let us to estimate dD ¼ 4.57 g cm3 and xV(B) ¼ 9.47 for dB ¼ 7 g cm3. The local density dB is assumed lower than crystalline
#1=3
1=3 I YD =T .
The Debye temperatures obtained are YD ¼ 290 K for HS Fe2+sites and YD ¼ 370 for LS Fe2+ sites for sample Ag10. The remarkable difference of YD, correspondingly f, for both environments is hard to justify in the case of a continuous random network and one would be tempted to attribute this difference to intrinsic segregation. Nevertheless, as Fe does not substitute any element of the network, this argument cannot be used straightforward. [25] In any case, if glasses were homogeneous, higher f environments can be attributed to Fe immersed in the GeSe backbone while lower f corresponds to Fe located in zones were the intermediate range order (IRO) has been affected by Ag [20]. If inhomogeneous glasses were assumed, according to the microscopy results, as Ag concentration increases the lower YD regions grow in correspondence to
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Fig. 4. Mo¨ssbauer spectra of sample Ag10 obtained at different T and fitted employing two D hyperfine distributions with linear dependence of d on D. On the left, the corresponding d and D distributions are depicted.
the increase of the bright zones in Fig. 3 whereas zones Ag depleted with a persistent IRO correspond to higher YD.
4. Conclusions XRD results [20] plus the determination of a single glass transition contribute to the homogeneous glass model. However, the abrupt jump in conductivity as well as SEM studies strongly support the heterogeneous glass model. That is, Ag rich bright zones coexist with Ag depleted dark regions. The crystallization of Ag8GeSe6 in the first crystallization peak at a relatively low temperature suggests that long distance diffusion is not needed in the nucleation and growth processes. This result encourages us to propose that bright zones approximately correspond to Ag8GeSe6. Under this assumption the bright zones volume fraction is estimated for Ag7.5 (Ge0.25Se0.75)92.5 in good agreement with the Scher–Zallen percolation criteria for inhomoge-
neous glasses. However, an accurate determination of the composition of both zones is still pending. 57 Fe Mo¨ssbauer results are not self sufficient in order to solve this controversy but are not contradictory to the previous results. HS Fe2+ plus LS Fe2+ are the main contributions to the spectra. The increase of HS Fe2+ environments with Ag concentration is observed. Debye temperatures corresponding to both environ ments are compared (YD(HS site)oYD(LS site)) and a remarkable difference, ascribable to intrinsic hetero geneity of the glasses, is observed. Therefore, HS Fe2+ sites in glasses can be attributed to Ag rich bright zones whereas LS Fe2+ sites can be attributed to Ag depleted dark zones. Further experiments are running. Neutron diffraction as a function of temperature in order to obtain detailed information of the crystallization products and a systematic Mo¨ssbauer study in a wider temperature range not only in the ionic conducting regime but also in semiconducting glasses.
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Acknowledgements The support of CNRS (France) and SECYT (Argentine) through ECOS Project A03E03 is greatly acknowledged. The argentine authors acknowledge the support of ANPCyT and Universidad de Buenos Aires. The friendly collaboration of Dr. M.F. Van Raap with low temperature Mo¨ssbauer experiments is also acknowledged. References [1] S.R. Elliott, in: J. Zarzycki (Ed.), Glasses and Amorphous Materials, Materials Science and Technology, vol.9, VCH, Weinheim, 1991 (Chapt. 7). [2] R. Azoulay, H. Thibierge, A. Brenac, J. Non Cryst. Solids 18 (1975) 33. [3] I.T. Penfold, P.S. Salmon, Phys. Rev. Lett. 67 (1991) 97. [4] I. Petri, P.S. Salmon, Phys. Chem. Glasses 43C (2002) 185. [5] P. Armand, A. Iban˜ez, H. Dexpert, D. Bittencourt, D. Raoux, E. Philippot, J. Phys. IV Coll. C2 (2) (1992) C2-189. [6] P. Boolchand, in: D. Adler, B.B. Schwartz, M.C. Steele, (Ed.), Physical Properties of Amorphous Materials, Plenum Press, New York, 1985, pp. 221–260. See also P. Boolchand, W. Bresser, M. Zhang, Y. Wu, J. Wells, R.N. Enzweiler, J. Non-Cryst. Solids 182 (1995) 143. [7] M. Micoulaut, Physica B 226 (1996) 268. [8] J.C. Philips, J. Non Cryst. Solids 34 (1979) 153; J.C. Philips, J. Non Cryst. Solids 43 (1981) 37. [9] M.F. Thorpe, J. Non Cryst. Solids 57 (1983) 355.
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