Stable and Metastable Phases within the GeO2-Rich Part of the Binary ...

Journal of Materials Synthesis and Processing, Vol. 9, No. 2, 2001

Stable and Metastable Phases within the GeO2 -Rich Part of the Binary PbO–GeO2 System Marco Scavini,1 Corrado Tomasi,2 Adolfo Speghini,3 and Marco Bettinelli3,4

Glasses of composition (1-x)PbO . xGeO2 with 0.50 ≤ x ≤ 1.00 were produced by melt quenching. Aliquots of each sample were thermally treated in air for various times at 6608 C and then characterized by X-ray powder diffraction (XRPD) and differential thermal analysis (DTA). The XRPD patterns of devitrified samples show the presence of one or more crystalline phases, depending on x. For 0.50 ≤ x ≤ 0.75, short treatments (2 h) resulted in the presence of monoclinic PbGeO3 , accompanied by phases of PbGe4 O9 stoichiometry. On the contrary, prolonged thermal treatments (360 h) produced a slight decrease in the intensity of the XRPD peaks of the PbGeO3 phase and the transformation of the remaining material into orthorhombic PbGe3 O7 . For 0.80 ≤ x ≤ 0.95, short treatments (2 h) resulted in the formation of hexagonal GeO2 , accompanied by orthorhombic PbGe3 O7 and by PbGe4 O9 phases. In this case, prolonged thermal treatments (360 h) do not affect strongly the XRPD patterns. KEY WORDS: Lead germanates; phase diagram, devitrification; metastable phases.

1. INTRODUCTION

potential of these materials, it is crucial to investigate, in detail, the thermal stability of the crystalline and noncrystalline phases. For this purpose, studies on the structure and the thermal evolution of phases belonging to this system were carried out by various authors, but the results appear to be contradictory. First of all, despite several detailed investigations [10], the structure of lead germanate glasses appears to be, as yet, not fully understood. In particular, different authors have different points of view about the presence of six-fold coordinated Ge atoms in lead germanate glasses. Umesaki et al. [11] have interpreted their neutron scattering data on the basis of the presence of a sizable proportion of six-fold coordinated Ge, while, in a previous study, Ribeiro et al. [12] did not confirm this type of coordination by means of EXAFS, IR, and Raman measurements. Moreover, the local structure around Pb2+ ions has not been definitely established. Second, the nature of the stable crystalline phases has not been fully established. Speranskaya [13] proposed that five lead germanate crystalline phases

Crystalline and noncrystalline materials belonging to the PbO–GeO2 system have been widely investigated due to their promising chemical and physical properties, which make them suitable for several technological applications. In fact, lead germanate glasses can be fabricated as optical fibers and photosensitive waveguides for use in the field of optoelectronics [1, 2] and are excellent hosts for lanthanide ions for the development of luminescent devices [3, 4]. On the other hand, crystalline phases in the PbO–GeO2 system show interesting ferroelectric [5, 6], pyroelectric [7], electrooptic and photorefractive [8, 9] properties. In order to exploit the technological 1 Dipartimento

di Chimica Fisica ed Elettrochimica, Universita` di Milano, Milano, Italy. 2 C.S.T.E.-C.N.R. and Dipartimento di Chimica Fisica, Universita ` di Pavia, Pavia, Italy. 3 Dipartimento Scientifico e Tecnologico, Universita ` di Verona, Verona, Italy. 4 To whom all correspondence should be addressed.

93 1064-7562/ 01/ 0300-0093$19.50/ 0  2001 Plenum Publishing Corporation

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(Pb6 GeO8 , Pb3 GeO5 , Pb5 Ge3 O11 , PbGeO3 , and PbGe3 O7 ) are stable besides pristine oxides. On the other hand, the phase diagram reported by Phillips and Scroger [14] looks significantly different. These authors reported the stability of four lead germanate compounds (Pb4 GeO6 , Pb3 Ge2 O7 , PbGeO3 , and PbGe4 O9 ) at room temperature, as well as a fifth one with composition PbGe2 O5 , existing only in a small temperature range (700–7408 C). A third different phase diagram was drawn by Gouju et al. [15], who reported the presence of four stable lead germanate phases of compositions Pb3 GeO5 , Pb3 Ge2 O7 , PbGeO3 , and PbGe4 O9 . The phase relations in the PbO-rich part of the system were studied in detail by Hasegawa et al. [16] by devitrification of glassy samples. The authors reported evidence of the existence of an hexagonal Pb5 Ge3 O11 compound, besides hexagonal Pb3 Ge2 O7 and orthorhombic Pb3 GeO5 phases. In a more recent work by Yamaguchi et al. [17], samples were produced by hydrolysis of alkoxides. For compositions up to 62.5 mol% PbO the authors detected the presence of the hexagonal PbGeO3 phase, as well as orthorhombic Pb5 Ge3 O11 , but did not confirm the presence of the Pb3 Ge2 O7 compound. The investigations on the GeO2 -rich side of the phase diagram were complemented by X-ray and neutron diffraction studies carried out on PbGeO3 , PbGe3 O7 , and PbGe4 O9 single crystals [7, 18–23]. In the case of PbGeO3 and PbGe4 O9 , the presence of several polymorphs was evidenced. An overview of the literature shows that the reports on the stability of the various phases in the whole compositional range significantly disagree. In particular, the GeO2 -rich part of the system seems to be more obscure. For this reason, we found it interesting to establish the phase relationships in the GeO2 -rich part of the PbO–GeO2 system, by means of differential thermal analysis (DTA) and X-ray powder diffraction (XRPD) measurements on samples obtained by devitrification of glassy materials of suitable compositions. The characterization and properties of relative glasses and the mechanisms of the devitrification processes will be considered elsewhere.

2. EXPERIMENTAL

Glasses of composition (1-x)PbO . xGeO2 , with x ranging from 0.50 to 1.00, were prepared by mixing appropriate amounts of analytical grade PbO (Aldrich, 99.9+ %) and GeO2 (Aldrich, 99.998%). The powders

Scavini, Tomasi, Speghini, and Bettinelli were melted in a platinum crucible in an electrically heated furnace in air. The batches were held at 11008 C (0.50 ≤ x ≤ 0.70), 12008 C (x c 0.75 and 0.80), 12258 C (x c 0.85), 12708 C (x c 0.90), 13008 C (x c 0.95), and 15008 C (x c 1) for 1 h, in order to pour melts having similar viscosity. The melts were quenched in air on a copper plate. The samples with x c 0.95 and 0.50 were “water quenched” by soaking the crucible containing the melt in cold water, making sure to avoid direct contact between water and melt. This procedure was adopted to prevent crystallization. According to preliminary differential scanning calorimetry (DSC) runs, samples were devitrified in air at 6608 C for different times in Pt trays. The temperature control was achieved by positioning a Chromel-Alumel thermocouple close to the Pt sample holder. After the firing, samples were rapidly withdrawn from the furnace. DTA measurements were carried out in alumina cups, at 108 C/ min in static air, by means of the 2910 DSC (TA Instruments), fitted out with a high-temperature DTA cell (16008 C). Preliminary DSC measurements were carried out using the same equipment fitted with a standard DSC cell. Room-temperature XRPD patterns were collected between 10 and 808 (2v range) with a Philips 1710 diffractometer operating with a Cu anode at 40 kV and 35 mA. The apparatus was equipped with a Philips PW 1050 vertical goniometer and a graphite crystal monochromator (CuKa radiation).

3. RESULTS AND DISCUSSION

Preliminary DSC runs (not shown) evidenced that heat treatments carried out at 6608 C for 2 h achieved devitrification for all the samples investigated. Figures 1 and 2 summarize the XRPD patterns collected on samples fired for 2 h at 6608 C. For the sake of clarity, the patterns pertinent to samples in the compositional range 0.50 ≤ x ≤ 0.75 are shown in Fig. 1, while Fig. 2 contains patterns relative to samples in the compositional range 0.80 ≤ x ≤ 1.00. In these figures, only small 2v intervals are shown in order to obtain information on the phases present in the samples as a function of the molar composition of the system. The XRPD pattern of the x c 0.50 sample in Fig. 1 shows only peaks relative to the monoclinic phase of the metagermanate PbGeO3 [21] (2v c 32.0, 32.78 ). On passing from x c 0.50 to x c 0.75, the intensity of the

Stable and Metastable Phases in the PbO–GeO2 System

Fig. 1. Indexed XRPD patterns of the samples 0.50 ≤ x ≤ 0.75, heat treated for 2 h at 6608 C. * is pertinent to the PbGe4 O9 phases (H, hexagonal; M, monoclinic) and O to monoclinic PbGeO3 .

peaks relative to the monoclinic PbGeO3 phase decreases and reflections assigned to both the monoclinic [20] and the hexagonal [7, 24] modifications of PbGe4 O9 appear (2v c 30.5 and 33.18 ). In the 0.80 ≤ x ≤ 1.00 compositional range, we note that by increasing x, the intensity of the XRPD peaks relative to hexagonal GeO2 [25, 26] increases (Fig. 2); moreover, peaks attributed to PbGe4 O9 and orthorhombic PbGe3 O7 [19] are present. This is particularly evident for the x c 0.80 sample, whose composition corresponds to the PbGe4 O9 stoichiometry. The XRPD patterns shown in Fig. 1 seem to evidence that the stable compounds, in this compositional range, are PbGeO3 and PbGe4 O9 . On the other hand, the patterns shown in Fig. 2 indicate that the situation is less simple than expected: besides PbGe4 O9 and GeO2 , peaks relative to PbGe3 O7 are also present. The DTA measurements carried out on samples heat treated for 2 h at 6608 C are summarized in Figs. 3 and 4. We point out that in the compositional range

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Fig. 2. Indexed XRPD patterns of the samples 0.80 ≤ x ≤ 1.00, heat treated for 2 h at 6608 C. * is pertinent to the PbGe4 O9 phases (H, hexagonal; M, monoclinic), × to orthorhombic PbGe3 O7 , and ˚ to hexagonal GeO2 .

0.50 ≤ x ≤ 0.75, no thermal effects were detected at temperatures above 10008 C for all heat treatments. Figure 3 shows the scans performed on samples in the range 0.50 ≤ x ≤ 0.75. The curve corresponding to x c 0.5 exhibits a peak of melting at ≈7908 C, while the curve for x c 0.55 shows a weak peak at ≈7508 C followed by a double endothermic effect. A double peak, in the range 750–8008 C, is also present for x c 0.6, while for x c 0.65, no peak at 7508 C is detected. In this curve, only a sharp peak at about 7708 C and a broader one at higher temperature are apparent. The thermal effect at ≈7508 C is present also for x c 0.70 and x c 0.75, always followed by a second endothermic effect. Figure 4 shows the scans performed on samples in the range 0.80 ≤ x ≤ 1.00: the sample with x c 0.80 presents an endothermic peak followed by a tail; this behavior suggests the presence of an incongruent melting compound. Samples with 0.85 ≤ x ≤ 0.95 all exhibit a double peak between 800 and 8508 C, followed

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Fig. 3. DTA curves of samples heat treated for 2 h at 6608 C for the compositional range 0.50 ≤ x ≤ 0.75.

by broad endothermic effects at higher temperatures. For pure GeO2 (x c 1), a single thermal effect at ≈11008 C is detected. The presence of shoulders on several thermal effects and the anomalies of some XRPD patterns suggest that in these conditions (firing for 2 h at 6608 C) equilibrium is not fully reached. In fact, as one can remark in Fig. 2, most of the XRD patterns show the simultaneous presence of four phases, including two polymorphs of PbGe4 O9 ; we point out that this behavior does not agree with the phase rule. On the basis of the present finding, we decided, at first, not to change the firing temperature, as this would introduce an additional degree of freedom, but to extend the heat treatment time up to 360 h. This choice was considered to be a good compromise between a reasonable duration of the firing and the time required to reach stable phase equilibria. Figures 5 and 6 show XRPD patterns collected on samples, which have undergone a prolonged firing (360 h) at 6608 C. In addition, in this case, only small 2v intervals are displayed for the sake of clarity. In Fig. 5, the XRPD patterns pertinent to samples in

Scavini, Tomasi, Speghini, and Bettinelli

Fig. 4. DTA curves of samples heat treated for 2 h at 6608 C for the compositional range 0.80 ≤ x ≤ 1.00. For the sake of comparison, DTA curves of both commercial GeO2 (Aldrich analytical grade) and GeO2 fired at 7508 C for 21 days are also included.

the compositional range 0.50 ≤ x ≤ 0.75 are collected. In this compositional range, the sample fired for 2 h showed peaks relative to monoclinic PbGeO3 and/ or both monoclinic and hexagonal 4 : 1 phases (see Fig. 1). All patterns shown in Fig. 5 still contain peaks relative to the monoclinic PbGeO3 phase, with the exception of sample with x c 0.75. We point out that there is no further evidence of peaks relative to both monoclinic and hexagonal 4 : 1 phases; on the other hand, peaks relative to the PbGe3 O7 orthorhombic phase appear, and their intensity increases with the GeO2 content. In particular, for the x c 0.50 sample, only the PbGeO3 phase appears to be present, while the x c 0.75 sample, of nominal composition PbO–3GeO2 , is only characterized by reflections characteristic of the PbGe3 O7 phase. In Fig. 6, the XRPD patterns pertinent to samples in the compositional range 0.80 ≤ x ≤ 1.00 fired for 360 h at 6608 C are displayed. It is apparent that: (1) by increasing x, the intensity of the peaks relative to hexagonal GeO2 increases; (2) peaks relative to both PbGe4 O9 and


Fig. 5. Indexed XRPD patterns of the samples 0.50 ≤ x ≤ 0.75, heat treated for 360 h at 6608 C. × is pertinent to orthorhombic PbGe3 O7 and O to monoclinic PbGeO3 .

orthorhombic PbGe3 O7 are still present; (3) in sample x c 0.80, orthorhombic PbGe3 O7 is the major phase but, by increasing x, the PbGe4 O9 / PbGe3 O7 peak intensity ratio increases. At variance with the 0.50 ≤ x ≤ 0.75 compositional range, the firing time does not seem to affect the phase composition and the main difference between Figs. 2 and 6 is the sharpening of peaks, probably due to an increase in crystallite size with heat treatment time. The DTA curves collected on samples heat treated for 360 h at 6608 C are shown in Figs. 7 and 8. In Fig. 7, both x c 0.50 and x c 0.60, are characterized by a single peak at ≈790 and ≈7508 C, respectively. The former is related to the melting of the PbGeO3 compound, whereas the latter can be ascribed to the formation of a eutectic, suggesting that this composition is fairly close to the eutectic. In fact, both the constant temperature and the change in intensity with composition of this latter peak are clear evidences of the formation of the eutectic. For x c 0.55, two endothermic effects are detected. The first peak falls at the same temperature as the one

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Fig. 6. Indexed XRPD patterns of the samples 0.80 ≤ x ≤ 1.00, heat treated for 360 h at 6608 C. * is pertinent to the PbGe4 O9 phases (H, hexagonal; M, monoclinic), × to orthorhombic PbGe3 O7 , and ˚ to hexagonal GeO2 .

shown in the curve for x c 0.60 and the second effect starts at temperatures just above it. Samples having x c 0.65 and 0.70 both display the endothermic peak at ≈7508 C, followed by a broader effect whose intensity and width increase with the GeO2 content. The sample with x c 0.75 exhibits a sharp peak at ≈8508 C with a tail extending up to ≈9208 C; this behavior is characteristic of melting of an incongruent melting compound. The weak effect at ≈7508 C is likely to be due to a small deviation from the stoichiometry. As stated above, no thermal effects were detected for this compositional range at temperatures above 10008 C. The DTA curves of samples richer in GeO2 are shown in Fig. 8. The sample with x c 0.80 presents a weak peak followed by a stronger one at temperatures above 8008 C, with a tail at higher temperatures. The curves for x c 0.85, 0.90, and 0.95 exhibit a double endothermic peak between 800 and 8508 C followed by thermal effects spread over a wide temperature range above

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Fig. 7. DTA curves of samples heat treated for 360 h at 6608 C for the compositional range 0.50 ≤ x ≤ 0.75.

9008 C. For pure GeO2 , both the devitrified sample and the commercial oxide (Aldrich) present a single sharp peak at ≈11008 C, corresponding to the melting. With regards to XRPD of devitrified GeO2 (see Fig. 6), it shows only reflections belonging to a-GeO2 , which is stable for T > 10508 C: the devitrification of pure glassy germanium oxide results in a metastable phase. This means that the present heat treatment (360 h at 6608 C) is not long enough to obtain tetragonal GeO2 (b), the low-temperature stable phase. On the other hand, even XRPD pattern of commercial GeO2 (Aldrich) shows that the material is in its high-temperature form; experimentally we observed that prolonged firing at higher temperatures (21 days at 7508 C) allows equilibrium to be reached. Besides XRPD, this fact is confirmed by the DTA measurements shown in Fig. 8, in which tetragonal GeO2 shows two thermal effects: the former, at 10508 C, corresponding to the b r a polymorphic transition [27] and the latter, at 11008 C, which is associated to the melting. We could also observe that the presence of a second phase acts as a “catalyst” in this transformation. In fact, (1) a 7-day firing at 7508 C of x c 0.90 sample results in


Fig. 8. DTA curves of samples heat treated for 360 h at 6608 C for the compositional range 0.80 ≤ x ≤ 1. For the sake of comparison, DTA curves of both commercial GeO2 (Aldrich analytical grade) and GeO2 fired at 7508 C for 21 days are also included.

an almost complete formation of the b phase, whereas pure hexagonal GeO2 remains mainly in the a form; and (2) similarly, after 1 h at 8408 C, a 4 : 1 stoichiometric mixture of a-GeO2 and PbO shows only GeO2 diffraction peaks pertinent to the b phase, whereas pristine aGeO2 fired for 24 h at the same temperature remains mostly unchanged. In the 0.80 ≤ x ≤ 1 range, DTA runs recorded on the sample heat treated both for 2 and 360 h at 6608 C (Figs. 4 and 8) exhibit, in the high-temperature region, a double endothermic effect. Although we could not confirm it by means of our x-ray diffraction measurements, similarities with the DTA curve of pure GeO2 fired at 7508 C induce us to tentatively assign the lower temperature peak to the b r a transition of GeO2 . The transformation of metastable hexagonal into tetragonal GeO2 is an exothermic process which should appear in DTA measurements. In fact, although our samples do not show the presence of b-GeO2 at room temperature, it could be formed during DTA scans. However, the kinetics of such a transformation is likely to be


Fig. 9. Indexed XRPD patterns of the sample with x c 0.75, heat treated for 2 min, 2 h, and 360 h at 6608 C. × is pertinent to orthorhombic PbGe3 O7 and * to the PbGe4 O9 phases (H, hexagonal; M, monoclinic).

slow (compared to the time scale of the measurements), resulting in a thermal effect spread over a wide temperature range and, therefore, difficult to observe. By analyzing the experimental data collected after the above firing procedures, for 0.50 ≤ x ≤ 0.75, we could evidence the presence of two stable compounds: PbGeO3 and PbGe3 O7 . In order to demonstrate the stability of PbGe3 O7 in this compositional range, Fig. 9 shows the XRPD patterns of x c 0.75 composition (corresponding to PbO–3GeO2 ) fired for 2 min, 2 h, and 360 h at 6608 C. Only the pattern relative to the sample fired for 2 min evidences peaks relative to the PbGe4 O9 phases, together with very weak peaks belonging to PbGeO3 ; no traces of the PbGe3 O7 phase appear. For the sample fired for 2 h, the situation is quite similar; however, this pattern also shows small traces of the strongest peaks of orthorhombic PbGe3 O7 . The sample treated for 360 h shows only the PbGe3 O7 phase; obviously the PbGeO3 minor phase also disappears as the x c 0.75 sample has the nominal composition PbO–3GeO2 .

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A similar phase evolution as a function of the heat treatment time, can be followed for each composition in the 0.50 ≤ x ≤ 0.75 range by direct comparison of Figs. 1 and 5. We stress that for the composition x c 0.75, this behavior can be easily observed because this nominal composition leads to a monophasic compound (PbGe3 O7 ). Experiments based on solid-state reactions have been performed in order to ascertain the stability of both PbGe3 O7 and PbGeO3 compounds. Stoichiometric amounts of powdered pristine oxides (PbO and GeO2 ) were ground and fired for a total time of 360 h at 7508 C. XRPD patterns showed that the equimolar mixture leads to monoclinic PbGeO3 whereas the 3 : 1 molar mixture results in monophasic orthorhombic PbGe3 O7 . A further confirmation of the phase equilibria in this region was obtained by treating the x c 0.70 sample for 24 h at 7708 C, i.e., above the eutectic temperature, followed by quenching to room temperature. As expected, XRPD measurements evidence the presence of an amorphous background and of reflections pertinent to orthorhombic PbGe3 O7 . No traces of other crystalline phases (such as PbGe4 O9 ) are present. These results indicate that the PbGe4 O9 phase is certainly metastable in the 0.50 ≤ x ≤ 0.75 range, where its formation is presumably due to its fast crystallization kinetics. This is in good agreement with the DTA results shown in Figs. 3 and 7, in which samples belonging to the same compositional range are run after firing for 2 and 360 h at 6608 C, respectively. We emphasize that, in the 0.50 ≤ x ≤ 0.75 range, the choice of the heat treatment of the samples seems to be crucial, in order to fully reach the equilibrium conditions. Besides, by comparing Figs. 3 and 7 the disagreement among phase diagram reported in literature can be easily understood. Concerning the 0.80 ≤ x ≤ 1 compositions, the situation is somewhat different. In fact, experimental results obtained after prolonged firing at 6608 C do not lead to a clear understanding of the phase equilibria, mainly for the following reasons: 1. In this compositional range (apart from pure GeO2 ) both PbGe3 O7 and PbGe4 O9 phases always coexist; 2. The phase composition of the crystalline mixture does not change remarkably with increasing firing time (i.e., passing from 2 to 360 h); 3. The PbGe3 O7 to PbGe4 O9 phase ratio decreases as the GeO2 content increases;


100 4. Following the adopted firing procedures, tetragonal GeO2 is never detected.

An attempt to better understand phase equilibria in this compositional range and, in particular, the stability of PbGe4 O9 , was made by further investigations on some selected samples: (I) An equimolar mixture of PbGe3 O7 and GeO2 was fired for 5 days at 8008 C, ground, and fired again for other 5 days. No evidence of either hexagonal or monoclinic PbGe4 O9 was found. (II) A sample with x c 0.85 was kept for 24 h at 9508 C, i.e., above the peritectic temperature and quenched to room temperature. As expected, XRPD showed the presence of both an amorphous background and of reflections pertinent to the tetragonal GeO2 . No other phases were detected. A further experiment was carried out in order to check the stability of the PbGe4 O9 phase. A stoichiometric molar 4 : 1 mixture of the pristine oxides GeO2 (hexagonal) and PbO (orthorhombic) was reacted at 7508 C for different times. Figure 10 shows the XRPD patterns of (1) the starting mixtures of the pristine oxides; (2) the mixture after heating for 2 h; and (3) after heating for 180 h. Pattern (1) is obviously composed of reflections of hexagonal GeO2 and orthorhombic PbO. After reacting for 2 h (pattern 2) peaks relative to hexagonal GeO2 are still present, while the reflections typical of orthorhombic PbO have disappeared. Moreover, peaks relative to at least one intermediate phase, possibly identified as monoclinic PbGeO3 , have appeared. The large width of these peaks, presumably due to the small size of the crystallites, does not allow a definitive assignment to the PbGeO3 phase. However, the presence of residual peaks of GeO2 , and not of PbO, indicates that the intermediate phase must be richer in lead than the starting mixture. The presence of PbGe4 O9 as a minority phase is possible, while the presence of PbGe3 O7 can be ruled out. Pattern (3) is composed of peaks typical of tetragonal GeO2 and orthorhombic PbGe3 O7 . Moreover, weak peaks relative to minority PbGeO3 are still present. We point out that, due to the uncertain identification of the phases present in the pattern (2), we have decided not to index the peaks shown in Fig. 10. In this case, the time evolution of the system also seems to rule out the stability of PbGe4 O9 . In the compositional range 0.50 ≤ x ≤ 1.00 the only stable crystalline phases appear to be PbGeO3 , PbGe3 O7 , and GeO2 . It has to be noted that the procedure described above (prolonged firing at 7508 C) does not allow obtaining a complete equilibrium, as further heat treatments do

Fig. 10. XRPD patterns of (a) a molar 4 : 1 mixture of the pristine oxides GeO2 (hexagonal) and PbO (orthorhombic); (b) the mixture after heating at 7508 C for 2 h, and (c) after heating at 7508 C for 180 h. ˚ is pertinent to hexagonal GeO2 , × to orthorhombic PbGe3 O7 , O to monoclinic PbGeO3 , # to orthorhombic PbO, and + to tetragonal GeO2 .

not lead to the disappearance of the PbGeO3 phase. The shape of the sample after each firing indicates the formation of liquid phases, probably due to the presence of eutectics between nonequilibrium phases present in the sample. A high crystallization rate of PbGeO3 from the liquid could explain its persistence. The present careful investigation of the PbO–GeO2 system shows that it is extremely difficult to establish an operative protocol for obtaining a phase equilibrium in the 0.75 < x < 1.00 compositional range. This behavior is probably due to the fact that at relatively low temperatures the transformation kinetics are slow, while at higher temperatures the formation of nonequilibrium liquid phases makes it hard the eliminate the metastable phases. Nevertheless, on the basis of the present experimental results, a tentative sketch of both metastable and stable phase diagrams of the GeO2 -rich portion of the


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are also based on the detection of the high-temperature form (hexagonal) of GeO2 , which is not supposed to be present below 10508 C. Figure 11B shows a tentative equilibrium phase diagram derived both from data collected on samples fired for 360 h at 6608 C and from some further specific experiments focused on establishing the stability of the phases (see above). Let us split the investigated part of the diagram into two regions: the first one with 0.5 ≥ x ≥ 0.75 (region I), and the second one with 0.75 > x ≥ 1 (region II). As far as region I is concerned, one can still remark the presence of the monoclinic PbGeO3 compound melting at ≈8008 C and, besides, of the orthorhombic compound with PbGe3 O7 stoichiometry (x c 0.75), characterized by incongruent melting at ≈8508 C. A eutectic is also formed for x ≈ 0.60 at ≈7508 C. In region II, the situation is not completely satisfactory, because of the above mentioned experimental problems; however, all the experiments we have performed indicate that the PbGe4 O9 phase is metastable and that in this compositional range only PbGe3 O7 and GeO2 are stable phases.

4. CONCLUSIONS

Fig. 11. Binary PbO–GeO2 phase diagram in the GeO2 -rich region based on the data obtained for the samples fired for (A) 2 h and (B) 360 h at 6608 C. 4 : 1, 3 : 1, and 1 : 1 indicate compounds with the specified GeO2 : PbO molar ratios.

binary system can be drawn (Fig. 11). Figure 11A shows the metastable phase diagram, which is obtained by relying on the data collected on the samples fired for 2 h at 6608 C. Besides pure GeO2 , two other compounds are present: a monoclinic PbGeO3 (x c 0.5) phase melting at ≈8008 C and a PbGe4 O9 phase (x c 0.80) characterized by incongruent melting at ≈8508 C. A eutectic is also formed for x ≈ 0.60 at ≈7508 C. For x > 0.80, the thermal effects lying just below the peritectic temperature could be related to metastable phase equilibria due to the presence of both PbGe3 O7 and PbGe4 O9 polymorphs. In fact, as discussed above, the comparison of X-ray with DTA data indicates that, in these conditions, phase equilibrium is not fully reached. These indications

The investigation of the phase diagram of the binary PbO–GeO2 system appears to be complicated by the formation of several metastable phases and by the presence of numerous polymorphs of the possible compounds. For these reasons, contradictory reports on this subject have appeared in the literature. In the present paper, we have tackled the problem of the stable phases in the portion of the phase diagram richer in germania, spanning the compositions from PbGeO3 to GeO2 . Whereas we have obtained a satisfactory picture of the situation for the range 0.50 ≤ x ≤ 0.75 (where x is the GeO2 molar fraction), the thermal behavior of the region with 0.75 < x ≤ 1.00 appears to be still partly unexplained, despite evidence that only GeO2 and PbGe3 O7 are stable phases in this compositional range. The experimental investigation of this region is complicated by the presence of eutectics between nonequilibrium phases and deserves to be further investigated. In particular, we have shown that thermal equilibrium for the compositions with 0.50 ≤ x ≤ 0.75 is reached only for relatively long firing times (360 h at 6608 C) and that the PbGe4 O9 phases appears to be metastable, while PbGe3 O7 is a stable compound. It has not been possible to establish a correlation between the structure of the various crystalline phases obtained

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by devitrification, and the structure of the corresponding glasses, as structural information on lead germanate glasses appears to be presently contradictory and incomplete. On the basis of our experimental results, we have traced two phase diagrams, a first one suitable for the description of the metastable situation after short firing times and a second one, which we propose to describe the equilibrium situation. It is interesting to note that a comparison of the above tentative equilibrium phase diagram with the data proposed in literature indicates a great similarity between the present results and the ones reported by Speranskaya [13]. On the contrary, the present nonequilibrium phase diagram is more similar to the ones proposed by Phillips and Scroger [14] and by Gouju et al. [15]. It is plausible that the experimental conditions of the latter studies have not allowed the achievement of thermal equilibrium.

ACKNOWLEDGMENT The authors gratefully thank Erica Viviani (Universita` di Verona) for expert technical assistance.

REFERENCES 1. D. Lezal, J. Pedl´ıkova´ , and J. Hora´ k, J. Non-Cryst. Solids 196, 178–182 (1996). 2. S. Mailis, A. A. Anderson, S. J. Barrington, W. S. Brocklesby, R. Greef, H. N. Rutt, and R. W. Eason, Opt. Lett. 23, 1751–1753 (1998). 3. J. R. Lincoln, C. J. Mackechnie, J. Wang, W. S. Brocklesby, R. S. Deol, A. Pearson, D. C. Hanna, and D. N. Payne, Electron. Lett. 28, 1021–1022 (1992).

Scavini, Tomasi, Speghini, and Bettinelli 4. M. Wachtler, A. Speghini, K. Gatterer, H. P. Fritzer, D. Ajo` , and M. Bettinelli, J. Amer. Ceram. Soc. 81, 2045–2052 (1998). 5. H. Iwasaki, K. Sugii, T. Yamada, and N. Nizeki, Appl. Phys. Lett. 18, 444–445 (1971). 6. G. Burns and B. A. Scott, Phys. Lett. 39A, 177–178 (1972). 7. A. Yu. Shashkov, V. A. Efremov, I. Matsichek, N. V. Rannev, Yu. N. Venevtsev, and V. K. Trunov, Zh. Neorg. Khim. 26, 583–587 (1981). 8. X. Yue, S. Mendricks, Y. Hu, H. Hesse, and D. Kip, J. Appl. Phys. 83, 3473–3479 (1998) and references therein. 9. D. Kip, S. Mendricks, and P. Moretti, Phys. Stat. Sol. A 166, R3–R4 (1998). 10. J. E. Canale, R. A. Condrate Sr., K. Nassau, and B. C. Cornilsen, J. Can. Ceram. Soc. 55, 50–56 (1986). 11. N. Umesaki, T. M. Brunier, A. C. Wright, A. C. Hannon, and R. N. Sinclair, Physica B 213 & 214, 490–492 (1995). 12. S. J. L. Ribeiro, J. Dexpert-Ghys, B. Piriou, and V. R. Mastelaro, J. Non-Cryst. Solids 159, 213–221 (1993). 13. E. I. Speranskaya, Izv. Akad. Nauk S.S.S.R. Otd. Khim. Nauk 1, 162–163 (1959). 14. B. Phillips, and M. G. Scroger J. Amer. Ceram. Soc., 48, 398–401 (1965). 15. D. Gouju, J. Fournier, and R. Kohlmuller, C. R. Acad. Sci. Ser. C 266, 1063–1065 (1968). 16. H. Hasegawa, M. Shimada, and M. Koizumi, J. Mater. Sci. 8, 1725–1730 (1973). 17. O. Yamaguchi, K. Sugiura, M. Muto, and K. Shimizu, Z. Anorg. Allg. Chem. 525, 230–236 (1985). 18. H. H. Otto, J. Appl. Crystallogr. 11, 157 (1978). 19. H. H. Otto, Z. Kristallogr. 149, 197–205 (1979). 20. H. H. Otto, and A. Dietrich, J. Appl. Crystallogr. 12, 421–422 (1979). 21. Yu. Z. Nozik, B. A. Maksimov, L. E. Fykin, V. Ya. Dudarev, L. S. Garashina, and V. T. Gabrielyan, Zh. Strukt. Khim. 19, 731–733 (1978). 22. H. H. Otto, Naturwissenschaften 63, 38 (1976). 23. L. Righi, G. Calestani, M. Bettinelli and A. Speghini, Atti XXVIII Congresso Nazionale Associazione Italiana di Cristallografia, Rimini, 1998, p. M2–20. 24. JCPDS Card n. 13–153. 25. G. S. Smith, and P. B. Isaacs, Acta Crystallogr. 17, 842–846 (1964). 26. JCPDS Card n. 36–1463. 27. E. Baiocchi, M. Bettinelli, A. Montenero, L. Di Sipio, and A. Sotgiu, Mater. Chem. Phys. 8, 379–386 (1983).