CRYSTAL GROWTH & DESIGN
Investigations of Single Crystal Growth of PbMoO4
2006 VOL. 6, NO. 1 58-62
S. C. Sabharwal,* Sangeeta, and D. G. Desai Crystal Technology Laboratory, TPPED, BARC, Trombay, Mumbai-400 085, India ReceiVed December 22, 2004; ReVised Manuscript ReceiVed June 11, 2005
ABSTRACT: The synthesis of polycrystalline PbMoO4 material suitable for use as the starting charge for crystal growth was investigated employing a solid-state sintering method. The solidification behavior of the material was studied by the differential thermal analysis (DTA) technique, with an aim to model the growth by the Czochralski technique. Actual crystal growth runs were carried out to validate the modeling and identify conditions that yield good quality crystals. The results show that large changes in melt stoichiometry, which may arise due to the application of an improperly synthesized starting charge, drive growth from the initial melt growth regime to a self-flux assisted growth regime. The application of a properly synthesized starting charge and the growth direction are found to play a crucial role in determining the quality of the grown crystal. Introduction
Experimental Methods
Lead molybdate (PbMoO4, LMO) single crystals find applications as acousto-optic modulators. It is one of those few crystals that has been developed to the extent that devices based on them have been commercialized, but not much has been reported in the scientific literature on different aspects of its single-crystal growth. The limited literature that is available on LMO shows that there are several aspects of crystal growth that warrant further investigations.1,2 For instance, the crystal transparency, coloration, cracking, etc. are some major problems with the material. While the exact origin of some of these disorders has not been investigated in sufficient detail, crystal transparency and coloration have been generally attributed to crystal nonstoichiometry. This is because MoO3 exhibits a higher vapor pressure than PbO at the elevated temperatures that are required either for LMO phase formation or for crystal growth.3 The crystal cracking has been attributed to the selection of growth direction; however, the results reported by different groups are found to be contradictory.1,2 In this article, we report the results of our investigations on the preparation and characterization of polycrystalline LMO material for crystal growth and the solidification behavior of LMO using the differential thermal analysis (DTA) technique. Guided by the results obtained, the single-crystal growth has been modeled. A proper selection of the growth direction was found to be the most important factor to avoid extensive cracking that occurs otherwise. For seeding with a Pt wire, the natural growth direction was observed to be a few degrees off the a-axis, and the grown crystals in this case suffered extensive cracking during cooling to room temperature. The growth axis lying between 20 and 30° off the c-axis was found to yield crackfree crystals. A slightly molybdenum-rich homogeneous mixture of the constituent oxides, rapidly raised to the melting point and held there for several hours, has been found to be best suited for crystal growth, as compared to that sintered at a temperature below the melting temperature. The application of a molybdenum deficit charge was observed to derive the growth from a case of pure melt growth to the self-flux assisted growth regime, as the melt composition underwent further changes during growth.
The constituent oxides used for the preparation of polycrystalline LMO were of 4N purity. The starting material used for crystal growth was prepared by mixing the constituent oxides taken in either stoichiometric ratio or with excess of MoO3. In some of the experiments, the mixture was sintered at different temperatures to determine the best conditions for LMO phase formation and its suitability for crystal growth. The effect of different sintering temperatures on the phase formation was investigated by recording of the powder X-ray diffraction (XRD) patterns of the resultant materials. The XRD patterns were recorded on a Rigaku-D/Max 2200 diffractometer using a Cu KR source. A thermal analyzer (SETRAM model 92) was used to investigate the thermal behavior of the materials. In all the measurements, a uniform heating and cooling rate of 10° min-1 was used, and the atmosphere provided was air. The single-crystal growth was carried out by the Czochralski technique under flowing air, using an automatic diameter controlled crystal puller (Cyberstar model Oxypuller). The transmission spectra of the crystals were recorded over the wavelength range of 200-1000 nm by a Techcomp spectrophotometer model 8500.
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Results (i) Materials Preparation and Characterization. The JCPDS data available in the literature on different phases in the PbO-MoO3 phase diagram, namely, PbMoO4, Pb2MoO5, Pb5MoO8, PbMo5O8, Pb0.77Mo4O6, and of the constituent oxides were analyzed to formulate criteria for the analysis of powder XRD patterns recorded for different LMO samples.4-10 The strongest reflections of Pb2MoO5 and MoO3 corresponding to d values of 3.46 and 3.445, respectively, are found to overlap. The Pb2MoO5 phase has a unique reflection corresponding to a d value of 2.16 and an intensity of 13%. Hence, this reflection together with its strongest reflection can be used to confirm the presence of this phase. For the confirmation of MoO3, the reflections having d values of 6.88, 3.84, and 3.75, which are unique to the species can be used. The strongest peaks of LMO and Pb5MoO8 phases have nearly the same d values viz. 3.245 and 3.264. However, the reflections having d values of 8.30, 7.66, 6.42, 5.81, and 2.74 are unique to the Pb5MoO8 phase, and these can be used to ascertain its presence. For LMO, the intensities of (004) and (200) reflections (corresponding to d values of 3.027 and 2.718, respectively) are found to be nearly the same. The strongest diffraction peak (111) of PbO has a very close d value to that of the (004) reflection of LMO phase. Hence, an increase in the intensity of the reflection (004), which can be inferred from an increase in the ratio of (004) and (200)
10.1021/cg0495678 CCC: $33.50 © 2006 American Chemical Society Published on Web 09/24/2005
Investigations of Single-Crystal Growth of PbMoO4
Figure 1. Powder XRD patterns for LMO material synthesized at 800 °C/48 h.
reflections, reveals the presence of PbO in the sample. The Pb2MoO5 phase can be detected by the presence of its two strong reflections having d values of 3.46 and 3.07. Since the strongest reflection of MoO3 also corresponds to d ) 3.46 and moreover it has a reflection with a d value equal to 3.07 and hence to ascertain the presence of Pb2MoO5 phase, the reflections corresponding to d values of 6.88, 3.84, and 3.75 for MoO3 should be monitored concurrently. The strong reflections of the PbMo5O8 phase having d values of 2.93, 2.97, and 3.11 overlap with the strong reflections of other phases, namely, Pb5MoO8, Pb2MoO5, PbO, Pb0.77Mo4O6, and hence this species can be detected by looking at its unique reflection corresponding to a d value of 4.61, which has a 40% intensity. For the samples sintered at 600, 800, and 1000 °C for 48 h, the powder XRD patterns were recorded to investigate the presence/formation of different chemical entities. The patterns obtained for the materials sintered at 600 and 800 °C were found to be similar. A typical powder XRD plot recorded for the 800 °C sample is shown in Figure 1. The 100% intensity reflection at d ) 3.24 together with the presence of its unique reflection corresponding to d ) 4.95 clearly establishes the formation of the LMO phase. All the other reflections characteristic of this phase are also observed here. The ratio of the intensities of (004) and (200) reflections is found to be 1.67. This ratio being much greater than one indicates the presence of unreacted PbO. The absence of any reflection corresponding to a d value of 3.71 rules out the presence of unreacted MoO3 in the sample. The presence of a reflection at d ) 3.46 establishes the formation of the Pb2MoO5 phase. The absence of any reflection with a d value of 4.61 rules out the formation of the PbMo5O8 phase. The formation of the Pb5MoO8 phase is also not observed as its unique and strong reflection corresponding to d ) 2.74 is absent. The absence of any reflection at d ) 4.81 also shows the absence of the Pb0.77Mo4O6 phase. From these observations, it follows that the phases rich in lead content are formed, possibly indicating some loss of molybdenum oxide during sintering. This observation is consistent with the fact that at the elevated temperatures MoO3 has a higher vapor pressure compared to PbO.3 Under the situation, the deviation from stoichiometry toward PbO-rich composition would favor the formation of the Pb5MoO8 phase together with LMO and Pb2MoO5. However, this has not been the case, possibly indicating the slow kinetics of Pb5MoO8 phase formation as compared to that of the Pb2MoO5 phase. The XRD pattern obtained for the material sintered at 1000 °C is reproduced in Figure 2. The ratio of (004) and (200) reflections (d ) 3.02 and 2.72) is found to be close to unity, implying the near absence of unreacted PbO. The absence of reflections due to Pb2MoO5 phase suggests that the amount of this impurity, if present, should be below the detection limits of XRD measurements. The pattern is found to match very well
Crystal Growth & Design, Vol. 6, No. 1, 2006 59
Figure 2. Powder XRD patterns for LMO material synthesized at 1000 °C/48 h.
Figure 3. DTA plots obtained for stoichiometric material prepared at 1000 °C/48 h during (a) heating and (b) cooling.
with that of the LMO phase.4 This result shows that on raising the temperature to 1000 °C, the transformation from the Pb2MoO5 to the LMO phase is facilitated. The Pb2MoO5 phase melts at 920 °C, while PbO melts at 888 °C, and hence these two phases act as flux in promoting the growth of the LMO phase. In conclusion, this investigation reveals that the best condition for the formation of polycrystalline LMO material is to raise a stoichiometric mixture of the constituent oxides directly to 1000 °C. (ii) Solidification Behavior. The polycrystalline materials prepared at 800 and 1000 °C were used to understand the solidification behavior of LMO and also the effect of minute amounts of the impurity phase on this behavior. The typical DTA plots obtained for the material sintered at 1000 °C are shown in Figure 3. Here, the melting and freezing occurred at 1060 and 1053 °C, respectively. For the material sintered at 800 °C, both the melting and the freezing points were found to have relatively lower values of 1051 and 1046 °C. This lowering of the two temperatures is attributed to the presence of a lower melting point phase, namely, Pb2MoO5 in the material, as revealed by the XRD plot of the material shown in Figure 1. To confirm this inference, DTA measurements were performed on a LMO sample containing a known amount of Pb2MoO5. The results obtained are shown in Figure 4, which reveal that for a sample containing 0.1% (by weight) of the impurity, the melting point gets lowered to 1051 °C, while the freezing point is shifted to 1028 °C. In conclusion, the DTA results show that the presence of even a very small amount of the impurity phase Pb2MoO5 has the
60 Crystal Growth & Design, Vol. 6, No. 1, 2006
Figure 4. DTA plots obtained for LMO material containing 0.1% Pb2MoO5 impurity and sintered at 800 °C/48 h, during (a) heating and (b) cooling.
effect of lowering the freezing point appreciably. The implications of this fact on crystal growth are investigated in the next section. (iii) Crystal Growth. From the results of Figure 3, it was concluded that LMO solidification required very small amounts (a few degrees only) of supercooling. Consequently, to effect the crystal growth a shallow longitudinal temperature gradient (few degrees per centimeter) at the solid melt interface was thought to be adequate. In practice, this was achieved by positioning the interface inside the crucible. To efficiently conduct away the heat of solidification, two conditions were applied: (i) the crucible position inside the rf coil was adjusted so that a longitudinal temperature gradient of ∼30 deg/cm was obtained above the crucible top and (ii) a platinum seed holder of sufficient length was used. Crystal growth was carried out using a platinum crucible. The presence of convection currents exhibiting a spokelike pattern was observed in the melt. Consequently, moderate crystal rotation rates were selected for growth. A crystal rotation rate of 20 rpm was applied for the growth of a 15-mm diameter crystal. Initially, a Pt wire attached to the seed holder was used to nucleate the melt. After bringing the wire in contact with the melt, we lowered the temperature slowly so that a small amount of the melt around the wire was solidified. Once the nucleation was over, the Pt wire was pulled up at the rate of 1 mm h-1 to promote bulk crystal growth. The crystal growth as modeled was indeed found to be very effective for the growth of LMO crystals. When seeded with a Pt wire, the most favored growth direction was not a well-defined crystallographic axis. Also, the ingots showed a total absence of the faceting effect. The surface of the grown crystal was found to be quite smooth, which implied the existence of a stable solid-melt interface. The grown crystals were found to crack on experiencing a thermal shock of ∼20 deg during the termination stage or during cooling to room temperature. The cracking was quite extensive extending throughout the crystal volume. Guided by the thermal expansion data available in the existing literature,1 the growth in the subsequent runs was effected using the seed crystals with different orientations. The best results were obtained when the growth direction was ∼20° off the c-axis. When the crystal growth was effected using the material synthesized at 1000 °C and applying a constant cooling rate, it was observed to progress smoothly until a certain percentage of the melt was solidified and thereafter it got tapered off. The subsequent growth runs always started at relatively lower temperatures (lower rf heater power levels), the growth setup remaining the same. This result suggested the occurrence of compositional changes in the melt during progress of growth.
Sabharwal et al.
Figure 5. DTA plots obtained for LMO crystals obtained from two successive growth runs carried out using the same charge, during (a) heating and (b) cooling. The -plots (s[s) and (sOs) correspond to crystals grown from fresh and leftover charges, respectively.
Figure 6. Optical transmission spectra recorded for samples prepared from the ingots grown from (a) fresh charge (-b-), (b) leftover charge (s9s), and (c) melt obtained by rapid heating of a stoichiometric mixture of the constituent oxides containing 1% excess MoO3 and soaked for 12 h (s1s). Note: Spectra have not been corrected for the reflection losses.
The DTA measurements were performed on the ingots obtained from two successive growth runs, and the results obtained are shown in Figure 5. These measurements showed very small changes in the melting or freezing temperatures, but the optical quality of the crystal grown from the leftover charge had deteriorated considerably. The optical transmission measured for the samples prepared from two such ingots are shown by plots (a) and (b) in Figure 6. The plot (b) shows the overall reduction in the optical transmission over a wide spectral range. The polished disk in this case was found to have a mosaic structure suggesting that the misorientation of individual crystallites was responsible for the enhanced losses whereby causing an overall reduction in crystal transmission. The origin of the mosaic texture of the crystal obtained from a leftover charge is understood as follows. The lowering of the melting point of charge suggests the presence of a chemical entity that has a lower melting point compared to that of LMO. The presence of such an entity would drive the growth process from the case of a pure melt growth to the self-flux assisted growth regime. A good self-flux assisted growth would require relatively much lower growth rates than the lowest pull rate of 1 mm h-1 used in the present investigation.11,12 Consequently, a high pulling rate would promote an uncontrolled twodimensional nucleation regime thereby resulting in the growth of the crystal with a mosaic structure. It has been mentioned earlier that starting with a charge of stoichiometric composition that was sintered at 1000 °C, only a small percentage of the melt could be crystallized into a good crystal. In light of the other results presented here this fact was attributed to some preferential loss of MoO3 from the charge
Investigations of Single-Crystal Growth of PbMoO4
Figure 7. Photograph of the LMO crystal grown under optimized conditions.
during its sintering. To substantiate this point, the experiments on growth were carried out by adding a known amount of MoO3 to the starting mixture of the constituent oxides, which was quickly raised to the melting point without holding it at an intermediate temperature for sintering and soaked for several hours. The excess amount of MoO3 was added essentially to take care of the losses during heating to the melting point, while the soaking of the melt ensured that there was no excess of MoO3 left in the melt. In our experiments, a charge containing 1% excess MoO3 and soaking of the melt for a period of about 12 h yielded the best results, as revealed by plot (c) of Figure 6 and the photograph of the ingot shown in Figure 7. It may be noted that contrary to the results of Bonner and Zydzik1 almost the entire charge could be pulled without any difficulty. This experiment brought out an important result that the evaporation from a stoichiometric melt should also be stoichiometric. To test this inference by an independent experiment, repeated DTA measurements were performed on the crystal sample for which the transmission spectra recorded are shown in Figure 6c. In these experiments, the sample was raised to 1150 °C and held there for 10 min prior to initiating the cooling. For four such cycles the plots recorded for melting are reproduced in Figure 8. Here, the change in the melting point observed was negligibly small, which supported the above inference. Discussion Bonner and Zydzik1 have investigated the single-crystal growth of LMO by the Czochralski technique using a stoichiometric charge and have reported the following results: (i) Less than 50% of the total melt could be crystallized without sacrificing crystal quality. (ii) The cracking toward the end of the boule occurs when a large percentage of the melt has been consumed. This has been attributed to an increase in the impurity level or an alteration in the stoichiometry. (iii) The evaporation from the melt consists of approximately equal amounts of PbO and MoO3. If the evaporation from the melt consists of equal amounts of PbO and MoO3 then the stoichiometry of the melt should not change as the growth progresses and hence results (i) and (ii) cannot be easily comprehended. However, all these
Crystal Growth & Design, Vol. 6, No. 1, 2006 61
Figure 8. DTA plots recorded during four successive heating cycles for the crystal shown in Figure 7. The plots (sOs), (sbs), (s]s) and (s4s) correspond to first, second, third, and fourth heating cycles, respectively.
results show that even very small stoichiometric variations are sufficient to effect the observed changes in crystal growth. In our experiments in which the mixture of the constituent oxides containing 1% excess MoO3 was quickly raised to the melting point we were able to pull almost the entire charge contrary to the first observation mentioned above. It may also be added that no measurable change in the transmission across the crystal length was observed. These results show that a melt of true stoichiometric composition can be fully crystallized, while a minute amount of nonstoichiometry has a pronounced effect on the growth process. The present study shows that the best recipe for obtaining a stoichiometric melt is to use a mixture of the constituent powders with 1% excess MoO3 and quickly raise it to the melting point. It may be mentioned that contradictory results regarding the suitability of seed orientation have been reported by Bonner and Zydzik1 and Zeng.2 While in the first case, the most suitable seed orientation has been identified to be 30° off the c-axis, in the other case the best results are reported for the seed orientation of [100]. The results obtained in the present investigation on the relationship among the growth direction and crystal cracking are indeed in agreement with those of Bonner and Zydzik.1 Our observations on the pulling speed have been found to be in agreement with those reported by others.1,2 A pull rate of 2.5 mm h-1 is found to be quite suitable. This is consistent with the DTA results, which show that a very small amount of supercooling is required for solidification. Zeng has reported the decrease of visible transmission in sequential order of crystal bulk along the axial direction,2 which has been attributed largely to degradation of the crystal quality and/or accumulation of crystal imperfections with growth time. However, an analysis of the present results attributes the degradation of crystal quality to the shift of the growth regime from a case of a pure melt to a self-flux assisted growth regime. This view is further supported by the observations made by Zeng on the crystal defects in LMO. He has reported accumulation of defects, which is regardless of sample location and growth time, implying the insignificant role of impurities. Almost constant dislocation density of ∼300 pits mm-2 is observed between the edge and center of the crystal, indicating that a
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time-independent phenomena should be responsible. The results of the present investigation show that such is the case when the charge employed for growth is not truly stoichiometric. Conclusions A stoichiometric mixture of the constituent oxides containing 1% excess MoO3 and quickly raised to the melting point together with soaking of the melt for several hours is found to be best suited for the single-crystal growth of LMO. The tolerable limit for nonstoichiometry in the present case is found to be very small. The results show that the degradation of crystal quality arising due to nonstoichiometry is a consequence of the shift in the growth regime from a case of pure melt to a self-flux assisted growth regime.
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References (1) Bonner, W. A.; Zydzik, G. J. J. Cryst. Growth 1970, 7, 65-68. (2) Zeng, H. C. J. Cryst. Growth 1997, 171, 136-145. (3) CRC Handbook of Chemistry and Physics, 65th ed.,; Weast, R. C., Astle, M. J., Beyer, W. H., Eds.; CRC Press: Boca Raton, FL, 1984; p D195. (4) Powder diffraction file No. 44-1486, JCPDS-ICDD, USA, 1997. (5) Powder diffraction file No. 24-0579, JCPDS-ICDD, USA, 1997. (6) Powder diffraction file No. 37-1086, JCPDS-ICDD, USA, 1997. (7) Powder diffraction file No. 47-1086, JCPDS-ICDD, USA, 1997. (8) Powder diffraction file No. 42-0314, JCPDS-ICDD, USA, 1997. (9) Powder diffraction file No. 35-1482, JCPDS-ICDD, USA, 1997. (10) Powder diffraction file No. 47-1320, JCPDS-ICDD, USA, 1997. (11) Sabharwal, S. C.; Sangeeta. J. Cryst. Growth 1998, 187, 253-258. (12) Sabharwal, S. C.; Babita, T.; Sangeeta. J. Cryst. Growth 2003, 249, 502-506.
CG0495678
