摘要:

Rheology and microstructure of aqueous layered double hydroxide dispersions Louise Albiston," Kevin R. Franklin,"" Elizabeth Lee" and J. Bas A. F. Smeuldersb "Unilever Research Port Sunlight Laboratory, Quarry Road East, Bebington, Wirral, UK L63 3J W bPolymers and Colloids Group, Cavendish Laboratory, University of Cambridge, Madingley Road, Cambridge, UK CB3 OHE The rheology and microstructure of aqueous dispersions containing Mg-A1 layered double hydroxides (LDHs) are considered. The thickening and gelling of such dispersions can be controlled by the addition of electrolyte or, within a certain composition range, by the application of shear. In addition to electrolyte and LDH concentration, the degree of thickening is shown to depend on both the primary crystal size of the LDH and the crystallinity of the material.XRD, particle size measurements, and direct observation by optical and electron microscopy were used to probe the microstructure of unthickened and thickened dispersions. Thickening appears to result from the interaction of predominantly rod-shaped aggregates formed by the face-to-face association of the primary LDH crystals. No interlayer swelling or delamination of the LDH crystals was observed. In recent years interest has grown tremendously in the prep- aration, characterisation and properties of layered double hydroxides (LDHs),',~ and a wide range of technologically important applications have been identified. These include the use of LDHs as catalysts and catalyst precursors,374 as ant- acid~,~in the preparation of pigments,6 in the treatment of waste as sunscreen agentsg and as rheology modifiers for both aqueous and non-aqueous systems.l0,'' The LDHs are a large class of naturally occurring and synthetic clay-like materials with positively charged polymeric mixed-metal hydroxide layers separated by expandable inter- layer regions containing anions and water molecules.1.2.'2 The class includes the so-called hydrotalcite-like and manasseite- like material^,'^ as well as the compositionally atypical mate- rials known as mixed metal hydroxides (MMHS).'~,'~,'~ LDHS with a wide range of metal ions in the hydroxide layers are now known and materials with many different inorganic and organic anions have been prepared by ion ex~hangel~,'~ or similar procedure^.'^^'^ Despite extensive studies on LDH synthesis and on the physical and chemical properties of these materials, the rhe- ology and microstructure of LDH dispersions, and how this is affected by the morphological nature of the materials, has received little attention. This is surprising since the difficulties in filtering LDH suspensions will be apparent to anyone who has synthesised them.Furthermore, many of the applications of LDHs rely on their dispersion in aqueous or non-aqueous media. The thickening behaviour of MMH materials in aque- ous dispersions has been reported, in particular with respect to their action in combination with bentonite clays when used as drilling rn~ds'~*~~ and their sensitivity to electrolyte concen- tration." The MMHs, however, are not typical LDHs in that they have Mg :A1 ratios of ca.1.0 and, in our experience, upon prolonged storage as dispersions they recrystallise to form a more conventional hydrotalcite-like phase and gibbsite. In this paper we report studies on the rheology and associated micro- structure of aqueous dispersions containing nitrate anion form Mg-A1 LDHs with a hydrotalcite-like structure, and show how the LDH synthesis conditions can affect the rheology. Experimenta1 Materials Nitrate ion-form magnesium-aluminium LDHs were prepared by combining a solution containing 145.1 g (0.347 mol) of aluminium nitrate nonahydrate in 500ml water with a slurry containing 46.8 g (1.16 mol) magnesium oxide in 250 ml water.The mixture was shaken vigorously for 2 min and then either filtered and washed with distilled water immediately or aged at elevated temperature for up to 7 days. Ageing of materials at temperatures below 100°C was carried out by placing the reaction mixture in a 1000 ml polypropylene bottle and heating it in a thermostatted oven. Sample ageing at temperatures above 100°C was carried out by heating ca. 150g of the reaction mixture in a Berghof 250 ml capacity Teflon-lined stainless-steel autoclave. All ageing was carried out under static conditions. Products were filtered, washed thoroughly with water, freeze-dried, and finally stored in a desiccator over a saturated sodium chloride solution (water activity =0.75).Carbonate-, chloride- and sulfate-form Mg-A1 LDHs were prepared from nitrate-form materials by ion exchange. The nitrate-form material was contacted twice with a solution containing an excess of the sodium salt of the required anion for 2 h at 90°C. The products were then filtered, washed thoroughly with distilled water, and freeze-dried. Chemical compositions of the LDHs were determined by a combination of XRF (Philips XRFS PW1404), thermal analy- sis (Perkin Elmer TGA7 thermogravimetric analyser), and C,H,N microanalysis. Crystallinity and purity were established by X-ray powder diffraction (XRD; Siemens D5000 diffractometer), and crystal shape and size were determined using either a JEOL 200CX transmission electron microscope (TEM) or a Cambridge Instruments Stereoscan 360 scanning electron microscope (SEM).Sodium nitrate used in some of the dispersions was pur- chased from BDH and was AnalaR grade. Rheology measurements Stable (non-settling) LDH dispersions were prepared by disper- sing the LDH in distilled water using an ultrasonic probe or, for large samples, a Silverson high-shear mixer type L4R fitted with a square-hole high-shear screen. Dispersions prepared by the two methods had the same viscosities when measured under identical conditions. All comparative studies were, how- ever, carried out on samples prepared by the same method so as to avoid any effects from small differences in the degree of dispersion. Electrolyte was subsequently added where required.Low shear viscosity measurements were made with a Carri- Med CSLlOO controlled stress rheometer using a cone and plate geometry. All measurements were made at 25 "C and at J. Mater. Chem., 1996, 6(5),871-877 871 a shear rate of 20 s-'. More detailed rheological measurements were performed using a Rheometrics RDS-I1 controlled strain rheometer fitted with a 0.01 N m force rebalance transducer. A 25mm diameter cone and plate geometry was again employed. Measurements were made at 20°C. The term vis- cosity, as used throughout this paper, is defined quantitatively as shear stress divided by shear rate. Microstructure measurements LDH dispersions were prepared as described above. X-Ray powder diffraction studies were carried out on gelled LDH dispersions to probe any interlayer swelling or delamination.The gels were mounted in standard powder diffractometer slides. The sizes of the agglomerate particles in the dispersions were determined after dilution in like electrolyte or water using dynamic light scattering or by disc centrifuge measurements. Dispersion microstructures were visualised using optical microscopy, cryo-SEM and freeze-fracture TEM. No special sample preparation was employed for the optical microscopy. Cryo-SEM images were obtained on a Cambridge Instruments S120 SEM. Samples were prepared by rapidly freezing small amounts of the dispersion in a copper sample holder using a liquid nitrogen-cooled Emscope SP2000A Sputter cryo-system.The sample was then transferred to the liquid nitrogen-cooled stage of the SEM. The surface of the sample was carefully etched using controlled heating to remove water and reveal the dispersion microstructure. Samples for TEM were prepared as follows. A small droplet of the LDH dispersion was placed on a copper planchette, and the sample slammed against a copper mirror at -189°C in the Reichart-Jung MM80 Universal Cryofixation apparatus. The vitrified sample was transferred under liquid nitrogen into a Cressington CFE-50 freeze-fracture apparatus (again at -189 "C) where it was fractured. The sample was then etched at 90 "C for 60 s, and rotary shadowed with Pt/C at 45" and carbon at 90". The coating (replica) so produced was floated on water and picked up on a 700 mesh gold EM grid, which was subsequently left overnight in 1 mol dm-3 nitric acid to dissolve any LDH adhering to the replica.Replicas were finally observed in a JEOL 200CX TEM at 80 keV. At least four replicas were prepared of each LDH dispersion to ensure the observed structures were representative of the whole sample. In most cases stereomicrograph pairs were obtained to allow the three- dimensional nature of the structures to be investigated more readily. Results and Discussion LDH Characterisation The Mg-A1 LDHs prepared using an elevated temperature ageing step were all highly crystalline and were composed of well defined hexagonal plate-shaped crystals with an aspect ratio of ca. 10. The crystal plate diameter increased uniformly with ageing temperature from 60 nm at 75 "C to 770 nm at 180 "C (Table 1, Fig.1). All samples were devoid of amorphous material and no crystalline impurities were detected. XRD showed that all materials exhibited peaks consistent with a single basal spacing at 8.9k0.1 A, which is as expected for a nitrate ion-form material with an Mg :A1 ratio of ca. 2.2*21The determined chemical composition of the aged materials is largely in agreement with this; however, the Mg:A1 and NO3:A1 ratios were slightly lower for the higher-temperature reactions (Table 1). The unaged sample was composed of large aggregates of very small irregular plate crystals held within an amorphous mass. XRD showed the only crys!alline phase present was a LDH with a basal spacing of 8.9A.The peaks were, however, weak and broad, and chemical analysis indi- 872 J. Muter. Chem., 1996, 6(5),871-877 Table 1 Characterisation of the nitrate ion-form Mg-A1 LDHs ageing ageing basal crystal plate Mg A1 NO, A1 temp./"C time/days spacing/A diameter'/nm ratio ratio 75 7 8.8 60' 2.0 0.9 90 6 89 80' 20 10 120 5 8.9 170' 19 09 150 4 8.9 400b 18 08 180 3 89 770' 19 08 n a.' -89 < 50d 16 08 'Average diameter determined from TEM or SEM Hexagonal plate-shaped crystals with an aspect ratio of ca 10 n a.=not applicable. Sample contains aggregates of very small crystals held in a matrix of amorphous material Fig. 1 SEM images showing the variation in LDH crystal size with ageing temperature (a) 120, (b) 150, (c) 180"C (size bars = 1 pm) cated an Mg:Al ratio lower than for the aged materials (Table 1).The carbonate, chloride and sulfate LDHs prepared by ion exchange retained the crystal morphology of the parent mate- rial (Table 2). Although the anion contents of these materials were not chemically determined, the basal spacings obtained from XRD (Table 2) were consistent with those expected.2 Since the sulfate and nitrate forms have very similar basal spacings, IR analysis was used to assess the extent of exchange. The IR spectra showed the complete absence of the peak at ca. 1384cm-' which is indicative of nitrate-ion LDHs, and the presence of a peak at 11 10 cm-',characteristic of sulfate.22 Low shear viscosity measurements These measurements were made at a shear rate of 20 sdl and were carried out to establish how the morphology and com- position of the LDH and the electrolyte composition affected the viscosity of the aqueous LDH dispersions.Initial studies were carried out with the 80 nm crystal plate diameter sample of nitrate-form LDH. Dispersions in distilled water with mass fractions up to ca. 15% LDH had very low viscosity (ca.25 mPa s) and flowed freely like water. Mass fractions of between ca. 15 and 22% were also of low viscosity when prepared; however, subsequent shaking of these samples led to an irreversible increase in viscosity (Fig. 2) such that the sample would no longer pour from a beaker. The time required to achieve this thickening varied with the mass fraction and the shear applied, but the transition from the free-flowing state to the viscous state occurred very rapidly (over only a few seconds).This unusual behaviour was studied in greater detail and is discussed further in the next section. With mass fractions over ca. 22% the highly viscous state was formed during dispersion of the LDH (Fig. 2). Viscous gels were obtained at low mass fractions of LDH when electrolyte was added to the dispersion (Fig. 3). Gelling occurred instantly upon addition of the electrolyte. The gels so produced appeared far more viscous in the quiescent state than those obtained by shear-induced thickening of higher Table 2 Characterisation of ion-exchanged LDHs anion form basal spacing/A crystal plate diameter/nm nitratea 89 80 sulfate 89 85 chloride 77 80 carbonate 7.6 80 a Other anion-form LDHs were prepared by ion exchange of this material 100 10 v) Q!&g10 0.1 0.01 0 5 10 15 20 25 30 35 mass % LDH in dispersion Fig.2 Low shear viscosity of LDH dispersions in water as a function of dispersion composition and sample shaking. Dispersions made with the 80 nm crystal plate diameter LDH. x, as prepared; 0,after shaking. 10 % s v)8? 0.1 0.01 -6 -5 -4 -3 -2 -1 0 1 log [NaN03] Fig.3 Low shear viscosity of LDH dispersions as a function of electrolyte molar concentration x , 5% LDH; 0,10% LDH LDH mass fractions in the absence of electrolyte (see above), but were much more shear thinning.With respect to the electrolyte-containing dispersions, the crystal size of the LDH has a major effect. As seen in Fig. 4 the viscosity of the dispersions decreased sharply with crystal plate diameter up to ca. 100-200nm. Above this size little further decrease in viscosity was observed, but in this range the dispersions were essentially quite free flowing rather than gels. Dispersions prepared using the unaged LDH sample were free flowing under all conditions irrespective of electrolyte concentration and LDH mass fraction (up to at least 35%). It is believed that this is due to poor dispersion of the LDH aggregates; after post-synthesis drying of the LDH the amorph- ous components 'glue' the LDH crystals together preventing subsequent dispersion.More viscous gels are produced from unaged LDH samples if the material is not dried. The effect of LDH counter-ion on dispersion viscosity in distilled water was also investigated. Replacement of nitrate by carbonate ions had little effect on the dispersions produced. Replacement with chloride ions, however, reduced the tendency to gel; dispersions of up to at least 35% (m/m) chloride-form LDH could be prepared without any significant thickening. With sulfate ion forms some thickening occurred even at mass fractions as low as 5%. Shear induced thickening of LDH dispersions As mentioned above, electrolyte-free dispersions of between 15 and 22% (m/m) of nitrate-form Mg-A1 LDHs (crystal plate diameter, 80 nm) show unusual and complex rheological behav- iour.As prepared they are of relatively low viscosity and free flowing. However, after shaking the dispersions vigorously in a uniaxial manner for up to 1h (dependent on mass fraction of the LDH), irreversible thickening occurs. This process may crystal plate diameterhm Fig. 4 Low shear viscosity of 10% LDH dispersions in 1 mol dm-3 sodium nitrate solution as a function of LDH crystal plate diameter J. Muter. Chem., 1996, 6(5), 871-877 873 u) cdeg 0.1 00v) 5 0.01 0 lo00 2000 3000 4000 so00 t/s Fig. 5 Thickening of a 20% LDH dispersion in water during steady-state shear at 150 s-l The variability in samples from a single batch is clearly seen also be observed in a rheometer where at a constant rate of 150 s-l thickening takes well over 1 h to occur (Fig 5) Successive runs on samples from the same batch of dispersion show the same effect, however, the time required to achieve thickening varies between 3700 and 4400s This is probably caused by subtle variations in the rheometer loading process The increase in viscosity observed in all cases under these conditions was about 40-fold (Fig 5) If the sample is shaken for a few minutes in a uniaxial manner before the rheological study is made the time required to achieve thickening is much reduced and tends to be more reproducible This reduction in thickening time within the rheometer is illustrated in Fig 6, which also shows how increasing the shear rate in the rhe-ometer results in reduced thickening times It is not easy to determine whether the thickening time us shear rate obeys a power law or an exponential as their dependence is reasonably linear both on a log-log plot and on a lin-log plot (Fig 7) The measurements on these systems are further com-plicated by the fact that, after the sample has been shaken, the microstructural network responsible for thickening continues to form slowly upon standing The thickening time measured in the rheometer then decreases with standing time So far we have considered only the thickening of the dispersions with time at steady shear While, however, the dispersions do undergo thickening with shear, they are shear thinning with respect to increasing shear rate The combination of these effects is seen in Fig 8, where the dispersion viscosity as a function of shear rate is plotted for three samples measured pnor to the shear induced thickening, at a mid-point in the process and when the thickening process has been completed The shear thinning in all cases follows a power law with coefficients varying from 0 5 for the unthickened state to 02 for the fully thickened state Unlike the shear thickening 1 -.............................. ...--.;-..................... ---.,I.--i I v) lu !& g 0.1 -_........................... 8F --......-0.01 1 10 100 lo00 loo00 tls Fig. 6 Thickening of an 18% LDH dispersion in water as a function of shear rate Measured under steady shear conditions with applied shear rates from 50 to 300 s-in steps of 50 s-l 874 J Mater Chem , 1996, 6(5),871-877 10' lo3 lo2 i cj m 10' Id lo3 0 50 100 150 200 250 300 350 shear rate/$-1 Fig.7 Time at onset of thickening as a function of shear rate (taken from Fig 6) Linear plots in (a)and (b)would indicate power law and exponential dependence, respectively -L end ---\ ......I y; ...........e......i ; ...... "*..* v1 md +......... i .............. i _................... ................................................................. +......i,",".:p2 lo-l i ..... i .... v) *.%.*...,8 s> slart a,... i............ "'+........ .... ...../ '"'Q-10 I o..g Fig.8 Viscosity as a function of variable shear rate at various points on the thickening curve for an 18% LDH dispersion in water process, this shear thinning process is readily reversible, the dispersion viscosity prior to starting the shear thinning cycle is rapidly re-established when shearing ceases Dispersion microstructure To establish the dispersion microstructure responsible for the very interesting rheology of LDH dispersions as a function of both shear and electrolyte concentration, a number of non-imaging (XRD, particle size) and imaging techniques (SEM, TEM) were employed A major process in the gelling of smectite clay dispersions is the imbibition of water into the clay interlayer region and consequent swelling and delamination of the particles To our knowledge, no reports exist on such processes in LDHs, although some organo-anion forms are known to swell in organic solvents such as toluene 24 To investigate whether such a swelling process played a role in the thickening of LDH dispersions, samples gelled by either electrolyte or shear were subjected to X-ray diffraction (XRD) analysis With both mechanisms of gelling the LDHs in the dispersions gave sharp peaks corresponding to a basal spacing of 8.9 A, as was found for the LDH powder.This is demonstrated in Fig. 9. There is clearly, therefore, no evidence from this technique of the LDHs imbibing water and swelling and thus this can be ruled out as a thickening mechanism. Visualisation is probably the most convincing method for dispersion microstructure determination.It can, however, be extremely difficult to achieve good and artefact-free image^.^^-^^ The use of cryo-SEM to visualise LDH dispersions was found to be of limited use. Initial results suggested that, in the thickened state brought on by shear or electrolyte addition, the LDHs crystals associated to form a continuous ‘honey- comb’ structure with liquid-filled pores up to several pm in diameter (Fig. 10). This is, however, now believed to be an artefact of the sample freezing process where the LDH crystals are swept along by the solution freezing front. Despite many attempts to increase the rate of freezing to avoid this, such structures were always observed.These structures are clearly artefacts since the pore systems observed are large enough to be seen by optical microscopy where no sample freezing is necessary. In fact, no such structures are observed (Fig. 11). The only features clearly seen from the optical micrographs is that the thickened samples appear to contain more aggre-gated crystals. By employing slam freezing and freeze-fracture TEM, high resolution images can be obtained which we believe are free of freeze artefacts. Images of unthickened dispersions contain- ing 20% (m/m) LDH show the presence of small, largely non- connected aggregates (Fig. 12). Images of a shear-thickened sample of the same composition showed much larger rod- shaped aggregates, formed predominately by face-to-face con- tacts (Fig.13). Furthermore, many of the rod aggregates appear to interconnect to form much larger, randomly shaped struc- tures. Voids between aggregates appear to be of the order of 50 to 300 nm in diameter. Dynamic light scattering measure- ments of aggregate size in the unthickened and thickened states after suitable dilution in water give average aggregate diameters of cu. 300-400 nm and 800-1000 nm respectively, which is in reasonably good agreement with the TEM images; measure- ment of aggregate length from the TEM images suggested all aggregates in the unthickened sample were <250 nm in length, while those in the thickened sample were up to 700nm in length. It should, however, be appreciated that such measure- ments of aggregate length are somewhat subjective owing to the orientation of the aggregates and the difficulty in judging 28ldegrees Fig.9 XRD patterns for (a)the LDH powder and (b) the same material dispersed in water at a mass fraction of 20% Fig. 10 Cryo-SEM image of a 10% dispersion of LDH in 1 mol dm-3 sodium nitrate solution. The honeycomb structure is an artefact created by ice crystal growth. (size bar =5 pm.) where there are individual aggregates and where there are aligned associated aggregates. Both of these difficulties prob- ably lead to an underestimate in aggregate size. From these findings we suggest that the shearing process results in the growth of progressively larger rod-shaped aggre- gates from the face-to-face association of unswollen LDH crystals.The growth process is slow because of the strong repulsive forces between the positively charged LDH crystals; the probability of two crystals or crystal aggregates approach- ing close enough for the shorter-range attractive forces to become dominant is low. Uniaxial mixing is particularly effective in aiding the aggregation process, since it not only imposes driving forces for particles to collide and aggregate (also true for other modes of mixing), but also aids the alignment of the particles. Beyond a certain size limit (>600 nm in length) the aggregates associate to very extensive networks. Thinning of the dispersions with increasing shear rate probably results from the breaking of the network, but J.Muter. Chem., 1996, 6(5), 871-877 875 Fig. 11 Optical micrographs of (a) a 10% dispersion of LDH in water and (b)in 1 mol dm-3 sodium nitrate (size bars =12 pm) not the aggregates. Thus, once the shear force is removed the networks can readily reform. Images obtained for samples containing 10% (m/m) LDH, both unthickened and thickened by electrolyte (0.1 mol dmP3 sodium nitrate), gave much less clear images and it was only possible to say qualitatively that the aggregates in the thickened samples were larger. Again face-to-face association of the LDH crystals were observed to be predominant. The aggregates in these systems form more rapidly because the repulsive forces between the crystals are screened out by the electrolyte.Measured by disc centrifuge on 10% LDH samples in water and in 1 mol dmP3 sodium nitrate, made after suitable dilution in like solvent, gave average aggregate sizes of ca. 100 nm and 1700 nm respectively. Although no such large aggregates were observed for the electrolyte-thickened dispersions using freeze- fracture TEM, aggregates of this size are consistent with those observed by optical microscopy [Fig. 11(b)]. Conclusions The rheology of Mg-A1 LDH aqueous dispersions is complex. The thickening or gelling of such systems may be brought about by addition of electrolyte or, in certain composition regions, by applying prolonged shearing. In addition to the electrolyte and LDH level, the degree of thickening is also critically affected by the size of the primary LDH crystals and crystallinity of the material.Thickened dispersions are highly shear thinning, but rapidly regain their thickened state once the shear force is removed. The thickening of LDH dispersions appears to result from 876 J. Mater. Chem., 1996, 6(5),871-877 Fig. 12 Freeze fracture TEM image of an unthickened 20% dispersion of LDH in water (size bar= 138 nm) Fig.13 Freeze fracture TEM image of a shear-thickened 20% dispersion of LDH in water (size bar= 138 nm) the interaction of predominantly rod-shaped aggregates which form through the face-to-face association of the primary LDH crystals. No interlayer swelling or delamination of these crystals occurs during this process. The transition to the thickened state occurs very rapidly, but only once a critical number and 6 7 8 9 10 Y.Park, K. Kuroda and C. Kato, J. Chem. SOC., Dalton Trans., 1990,3071. J. Crosfield and Son, World Pat. Appl., WO 91/19850, 1991. Aluminium Company of America, US Put, US 4867882, 1989. Unilever PLC, Eur. Pat. Appl., EPA 0557089 Al, 1993. Dow Chemical Company, US Put., US 5015409,1991. size of aggregates is reached. 11 12 Giulini Chemie Gmbh, Ger. Pat. Appl., DE 3732265, 1989. R. Allman, Acta Crystallogr., Sect. B, 1968,24, 972. This work was carried out as part of the DTI Colloid Technology Link Project supported by Unilever, ICI, Zeneca and Schlumberger. The authors would like to thank Mr. I. Tucker and Miss S. Evans (XRD), Mrs. J. Munro-Brown (SEM) and Dr. I. Clarke (XRF) for their assistance in charac- terising the LDH powders and dispersions. All are employees of Unilever Research Port Sunlight Laboratory.The authors would also like to thank Dr. P. N. Segre (Dept. Physics and Astronomy, University of Edinburgh) for performing the 13 14 15 16 17 18 19 20 21 R. Allman, Chimia, 1970,24,99. Dow Chemical Company, US Put., US 4990268,1991. L. J. Fraser, in Proceedings of the 1992 SPE International Meeting on Petroleum Engineering, Texas, 1992, paper 22379. S. Miyata, Clays Clay Miner., 1983,31,305. M. Meyn, K. Beneke and G. Lagaly, Znorg. Chem., 1990,29,5201. D. L. Bish, Bull. Mineral., 1980, 103, 170. K. Chibwe and W. Jones, J. Chem. SOC., Chem. Commun., 1989,926. F. Lovoix and M. Lewis, Oil Gus J., 1992,September 28, p. 87. K. R. Franklin, E. Lee and C. C. Nunn, J. Muter. Chem., 1995, 5, 565. dynamic light scattering experiments. 22 M. J. Hernandz-Moreno, M. A. Ulibarri, J. L. Rendon and C. J. Serna, Phys. Chem. Minerals, 1985,22,34. 23 H. A. Barnes, J. F. Hutton and K. Walters, An Introduction to References Rheology, Elsevier, Amsterdam, 1989. 24 H. Hopka, K. Beneke and G. Lagaly, J. Colloid Interface Sci., 1988, 1 W. T. Reichle, Solid State Ionics, 1986,22, 135. 123,427. 2 F. Cavani, F. Trifiro and A. Vaccari, Cutal. Today, 1991,11,173. 25 B. Gu and H. E. Doner, Clays Clay Miner., 1992,40,246. 3 Henkel KGaA, Ger. Pat., DE 4010606 Al, 1991. 26 H. Vali and R. Hesse, Clays Clay Miner., 1992,40, 620. 4 C. Busetto, G. Del Piero, G. Manara, F. Trifiro and A. Vaccari, 27 B. Gu and H. E. Doner, Clays Clay Miner., 1993,41, 114. J. Catal., 1984,85, 260. 5 P. C. Schmidt and K. Beneke, Pharm. Acta Helv., 1988,63, 188. Paper 5/06567D; Received 5th October, 1995 J. Mater. Chern., 1996, 6(5), 871-877 877

ISSN:0959-9428    

DOI:10.1039/JM9960600871

出版商:RSC    

年代:1996

数据来源: RSC

摘要:

Preparation, characterization and surface structure of coprecipitated high-area Sr,Ti02 +J 0 <x <1) powders Josh Manuel Gallardo Amores," Vicente Sanchez Escribano," Marco Daturib and Guido Busca*b "Departamento de Quimica Inorganica, Faculdad de Quimicas, Universidad, Plaza de la Merced E-37008 Salamanca, Spain bIstituto di Chimica, Facolta di Ingegneria, Universita, P.le Kennedy, I-1 6129 Genova, Italy Powders with composition Sr,TiO, +, (x=0,0.02,0.11,0.42 and 1)have been prepared by coprecipitation from strontium hydroxide and titanium isopropoxide and have been characterized after drying at 393 K, and after calcination at 773 and 973 K, using XRD, TG-DTA, BET surface area and porosity measurements, FTIR-FTFIR and FT-Raman skeletal vibrational characterization, diffuse reflectance UV-VIS spectroscopy and FTIR spectra of the hydroxy groups and of adsorbed pyridine and CO,.Strontium was found to inhibit anatase crystallization, sintering and transformation to rutile. Strontium tended to be deposited at the surface of anatase and gave rise to a decrease of the surface acidity and an increase of surface basicity. Moreover, it shifted the absorption edge of anatase to higher energies. For x = 1, the truly cubic SrTiO, perovskite phase was produced at room temperature with high surface area. The surface of SrTiO, was definitely basic, with only strontium and oxygen ions exposed at the surface, and hydrogen impurities in the bulk. Perovskite-type metal titanates, e.g. SrTiO,, BaTiO, and PbTiO,, are widely applied in the ceramic and electronics industries, owing to their dielectric, ferroelectric, incipient- ferroelectric and piezoelectric properties.' These materials are also of interest for a number of other properties. SrTiO, is a semiconducting material, acting as a powerful resistivity sensor for ~xygen,~ is active in the photoelectrolysis of water4 and in other photocatalytic proce~ses.~ SrTiO, was also the first ceramic material found to become superconducting, at a very low temperature, below 0.3 K.6 Perovskite-like materials can also be applied in the field of heterogeneous catalysis as active phases7 or as potential catalyst supports.* Some of these properties appear to be peculiar to perovskite-type materials while others (like those related to photoconductivity and catalytic behaviour) appear to concern the nature of the Ti4+ -O2-bond and closely relate perovskite-like metal tita- nates to pure titanias and other titanates characterized by different crystal structures (e.g.ilmenites). Ti0,-anatase and rutile are mainly used in industry as white pigments' while anatase constitutes the active support of the industrial catalysts for o-xylene selective oxidation to phthalic anhydride' and for the selective catalytic reduction of nitrogen oxides by ammonia,'' both based on vanadium oxide. Anatase is also of interest because of the so-called strong metal support interaction (SMSI) phenomenon occurring when some metals are supported on it." Additives to Ti0,-anatase strongly influence several proper- ties, including its stability to sintering and to phase transform- ation to rutile.As a further development of our research on the surface behavior of titanias12?13 and perovskite-like fine powder^'^,'^ we investigated the preparation and the bulk and surface properties of high-area powders in the Sr,TiO,+, (0d xd 1) composition range. The aim of this study is twofold: (i) to test the surface properties and the thermal stability of high-area, nearly stoichiometric, SrTiO, powders; and (ii) to determine the properties of Sr,TiO,+, mixed oxides and the effect of Sr addition to high-area Ti0,-anatase powders. Experimental Preparation procedures Samples with composition Sr,TiO, +,(0dx d 1) were pre-pared via a conventional coprecipitation method, using Sr(OH), 8H,O (Strem Chemicals) and Ti[OCH(CH,),], (Aldrich) as precursors.Stoichiometric amounts of the stron- tium precursor dissolved in water, were added to the titanium precursor under continuous stirring. After water evaporation, the cake was dried at 393 K for 3 h in air. Finally, all samples were subjected to thermal treatments in air at different tempera- tures (773 and 973 K). Characterization techniques Curves were recorded by means of a Setaram TGA 92 12 apparatus, with a heating and cooling rate of 10 K min-'. XRD spectra were recorded on a Philips PW 2256/20 diffractometer (Co-Ka radiation, Ni filter; 40 kV, 20 mA). Cell parameters were calculated using dedicated least-squares software.The crystal size was evaluated using the Scherrer forrnula.l6 Microstructures were measured using N, adsorption at 77 K, with a conventional volumetric BET apparatus. TEM images were obtained using a high-resolution Zeiss instrument with a maximum magnification of 250000 x . The UV-VIS diffuse reflectance spectra were obtained using a Shimadzu spectrophotometer (model UV-240), previously calibrated with two analytical MgO samples. IR spectra were recorded with a Nicolet Magna 750 Fourier transform instrument. The skeletal spectra in the region above 400cm-' were recorded with KBr pressed disks and with a KBr beam splitter, while those in the far IR (FIR) region (400-50 cm -') were recorded using the powder deposited on polyethylene disks, and with a 'solid substrate' beam splitter.FT-Raman spectra were obtained through a Bruker RFSlOO instrument, with an Nd-YAG laser (1064 nm), using 30 mW laser power, 50 scans and 4cm-' resolution. The adsorption experiments were performed using pressed disks of the pure powders, activated by outgassing at 300-1070K into the IR cell. Results DTA and XRD characterization The DTA curves recorded in air for the Sr,TiO, +,(0d x 6 1) samples previously dried at 393 K are reported in Fig. 1. The pure titania sample shows a strong exothermic peak centred J. Muter. Chem., 1996, 6(5), 879-886 879 970 I R+P I 673 773 873 973 1073 1173 1273 TIK Fig. 1 DTA curves for Ti0, (a), Sr, ozTi02oz (b), Sr, llTi02 11 (c), Sr, 42T10242 (d), and SrTiO, (e) (Am =amorphous, A =anatase, R = rutile, P =perovskite at 970 K, with a tail at lower temperatures The XRD spectra of this matenal after drying already show the diffraction peaks of anatase (JCPDS file no 21-1272) with traces of those of brookite (JCPDS file no 29-1360), as reported in Table 1 The XRD patterns of this sample after calcination at 773 and at 973 K are shown in Fig 2(a) and 3(a) respectively At 773 K the material is still essentially composed of anatase with traces of brookite while at 973 K the material is almost totally converted to rutile According to these data and to our previous data12 13 17 the exothermic peak at 973 K relates to the anatase- rutile phase transition, while the previous signal is associated with anatase sintering The DTA curve for Sr, 02T102 02 is shown in Fig 1(b),while the XRD patterns of this sample are reported in Fig 2(b) and 3(b) The sample is still stable as anatase after calcination at 973 K and, correspondingly, the anatase-rutile phase transition is shifted to 1215 K in the DTA curve These data show that the addition of small amounts of Sr stabilizes the metastable form, anatase, towards its transformation to the thermo-dynamically stable form, rutile The data relative to the sample Sr,,,T10211 show a quite different picture In fact, the sample after drying appears amorphous to XRD analysis, while after calcination at both 773 and 973 K, virtually brookite-free and rutile-free anatase is observed [Fig 2(c) and 3(c)] The DTA trace shows a broad exothermic peak near 870 K, that is associated (according to XRD analyses) with the amorphous-anatase transition, occur- ring in the non-isothermal conditions of DTA experiments at higher temperatures with respect to isothermal calcination in the oven A weak peak observed near 1220K in the DTA curve is associated with a more complex phase transition from anatase to rutile +perovskite mixed phases, as deduced from the XRD analyses performed after the DTA run 880 J Mater Chem, 1996, 6(5),879-886 -I I11 73 65 55 45 35 25 2Bldegrees Fig.2 XRD of Ti0, (a), Sr, 02Ti02o2 (b), Sr, 11Ti0211 (c), Sr, 42Ti02 42 (d) and SrT103 (e), all calcined at 773 K Indexing in (c) is for anatase and in (e) for cubic perovskite 101 I I I211 'I0 r 1 7s 65 55 45 35 25 2Bldegrees Fig.3 XRD of Ti0, (a),Sr, ,,Ti02 02 (b),Sr, ,lTiOz 11 (c), Sr, 42T102 42 (d) and SrTiO, (e), all calcined at 973 K Indexing in (a) is for rutile, in (b) for anatase and in (e) for cubic perovskite Table 1 XRD and morphological data for Sr,TiO,+, samples cell parameters/,& composition crystal partic!e BET surface sample T/K phases SrTiO, 393 cubic perovskite +anatase 773 cubic perovskite +anatase 973 cubic perovskite +anatase 1243 cubic perovskite SrO 42Ti02 42 393 773 973 amorphous amorphous amorphous +cubic 1243 perovskite +anatase +rutile cubic perovskite Sr, ll-r-102 11 393 773 amorphous anatase 973 anatase 1243 cubic perovskite +rutile Sro 02T102 02 393 773 anatasef anatase +brookite 973 anatase 1273 cubic perovskite +rutile Ti0, 393 773 anatasef anatase +brookite 973 rutile +anatase (Yo)" a c sizeb/,& size'/A area/m2 g-' DMd/nm VTPe/ml g-' 97 3.9347 - 276 3 20 87 5.5 0.12 trace 98 3.9103 - 322 380 50 10.1 0.12 trace 98 3.9296 - 344 420 20 7.1 0.04 trace 100 3.9067 - 100 - - - - 339 5.1 0.40 100 - - - 60 118 7.4 0.22 41 7.0 0.07 3.9250 - 193 260 70 3.9079 380 100 - - - - 41 1 3.1 0.32 100 3.8012 9.4758 135 145 125 2.4 0.08 100 3.8036 9.47616 118 175 39 8.6 0.08 21 3.9109 380 79 4.5922 2.959 1 158 100 3.81222 9.36248 70 85 236 5.0 0.29 92 3.7889 9.43837 79 100 129 7.2 0.23 100 3.8051 9.5596 135 155 49 8.2 0.10 3.9212 365 92 4.5892 2.9591 241 - 3.8224 9.36817 49 120 240 5.6 0.33 92 3.7852 9.4905 147 160 75 10.9 0.20 94 4.5821 2.9531 538 590 5 9.9 0.01 "Evaluation based on the relative intensities of the XRD diffraction peaks.bFrom XRD data. 'From TEM data. dDM=average pore diameter. eVTp=total pore volume. 'Brookite is also present, according to Raman spectra. The DTA curve of the Sr0.42Ti02.42 sample shows a sharp split peak at 1025-1042 K [Fig. l(d)]. XRD analysis shows that this material is still completely amorphous after calci- nation at 773 K [Fig. 2(d)], while it consists of a mixture of amorphous material and of the perovskite SrTiO, (JCPDS file no.35-734) after calcination at 973 K. After the DTA run a poorly crystallized mixture of anatase, rutile (JCPDS file no. 21-1276) and perovskite is observed [Fig. 3(d)]. The DTA curve relative to the stoichiometric material SrTiO, does not show any definite peak. The XRD patterns show that the well crystallized cubic perovskite phase of SrTiO, (which is thermodynamically stable in air) is already present after drying [Fig. 4(u)], so that only a partial increase of crystal size by further calcination is observed [Fig. 4(b) and (c)]. These data show that the addition of Sr to TiO, at low Sr loadings inhibits the anatase-rutile phase transition while, at higher loadings, it also inhibits anatase crystallization from the amorphous state.In contrast, when Sr :Ti =1 : 1 the copre- cipitation produced the cubic perovskite structure, already well crystallized. This behaviour differs with respect to that of the Sr-Zr system which, at Sr :Zr z 1:1, gives rise to amorphous materials that crystallize to form the orthorhombic perovskite SrZrO, only after calcination at 973 K." This may be related to the larger size ofJhe Zr4+ cation with respect to the Ti4+ cation (0.80 US. 0.68 A, respectively, after Pauling18). Zr4+ does not enter readily into the BO, octahedra of the ABOJ perov- skites, thus inducing an orthorhombic distortion in the SrZrO, perovskite and inhibiting its crystallization. In contrast, owing to its smaller size, Ti4+ enters the perovskite structure more readily, and does not cause distortion when A =Sr2+.No phases other than 'pure' Ti02 and SrTiO, were found in our samples, in agreement with the lack of detection of any phase between TiO, (rutile) and SrTiO, (cubic perovskite) in the phase diagrams of the SrO-TiO, system." Analysis of the unit-cell parameters of rutile, anatase and perovskite phases in our powders (Table 1) does not provide any evidence of reciprocal solubility of TiO, and SrTiO,, according to phase diagrams." Morphology characterization The morphological properties of the materials after drying and after calcination at 773 and at 973 K are summarized in Table 1. The surface area measurements show that the addition of small amounts of Sr to anatase tends to increase its surface area, in agreement with the stabilization of the amorphous phase with respect to the crystallization of anatase and of the anatase phase with respect to its conversion to rutile. This agrees with the previous results which show that the dopants that inhibit the anatase-rutile phase transition also inhibit the loss of surface area of anatase, and with the mechanism we previously proposed for this.', The surface areas are smaller for SrTiO, in the crystalline perovskite form. However, the surface areas obtained for SrTiO, are definitely higher than those obtained with similar preparation procedures for other perovskites, including SrZrO,." This is certainly associated with the low temperature of crystallization of such a cubic phase, in contrast with SrZrO,, which cannot be produced in a crystalline form at low temperature.The parent compound BaTiO,, which can also be obtained in a crystalline, apparently cubic, form at very low temperat~res,'~ is also obtained with rather high surface areas. J. Muter. Chem., 1996, 6(5), 879-886 881 Comparison of the theoretical surface areas, measured from the XRD crystal size assuming cubic particles, with the exper- imental surface areas and those forecast on the basis of TEM images gives satisfactory agreement for SrTiO, According to TEM images, the SrTi0, powder consists, immediately after drying, of well defined globular particles This agreement indicates that the SrTiO, particles are composed essentially of single crystals In contrast, in Ti0,-anatase and in strontium- deficient materials, the 'theoretical' surface areas measured from the XRD crystal size are higher than those observed experimentally On the basis of TEM images, in these cases the particles are more irregular and are probably polycrystal- line, with possibly some amorphous material present also, even after calcination The porosity data for SrXTiO2+, (0<x< 1) samples are summarized in Table 1 The shapes of the experimental Nz adsorption isotherms of the dried samples are intermediate between type I and type IV in the IUPAC classification," corresponding to micro-mesoporous solids They covert to type IV isotherms for samples calcined at 773 K The increasing calcination temperature promotes isotherm transformation from type IV to type I1 and the disappearance of the hysteresis loop The variation of average pore diameter reaches a mini- mum value for the x=O 11 sample, which also shows the highest surface area after drying Bulk vibrational characterization In Fig 5 the FTIR skeletal spectra of the samples Ti02, Sr, 02T102 o2 and Sr, 42Ti02 42, all calcined at 773 K, are reported The spectrum of T102 corresponds with those reported in the literature and with those we discussed pre- viously12 The main absorptions near 700 (very broad shoulder), 425, 330 (shoulder), 260 and 175 cm-' (weak) are typically found in anatase samples with small particle sizes Weak additional components in the 550-650 cm-' region are probably associated with brookite impurities l2 The addition 100 75 65 55 45 35 2s 2Bldegrees Fig.4 XRD of SrTIO, dned at 393 K (a) and calcined at 773 K (b) and 973 K (c) Indexing is for cubic perovskite structure 882 J Muter Chem , 1996,6(5), 879-886 1200 1000 800 600 400 200 wavenumber/cm-l Fig. 5 Skeletal FTIR-FTFIR spectra of T102 (a), Sr, 02T102o2 (b), and Sr, 42T10242 (c),all calcined at 773 K of small amounts of Sr seems to cause mainly the decrease of the bands associated with brookite and of the lowest frequency band However, the general spectrum of anatase remains almost intact In contrast, the spectrum of Sr, ,,Ti02 42, which is X-ray amorphous, is definitely different, being dominated by a very broad band centred near 570 cm-' and a sharper band at 255 cm -The FTIR spectra of the SrTiO, samples (Fig 6) present a typical spectrum with the main bands at 550-555, 450 (shoulder), 405,250 and 150 cm-' The spectrum is similar but with significant band shifts with respect to those discussed previously for a commercial low-area SrTi0, sample,I2 and is also in approximate agreement with the spectra of different SrTi03 samples reported by Diaz-Guemes et aZZ1In our case, the calcination treatment does not change the spectrum of SrTi0, very much, thus showing that the crystal shape is not modified significantly The sharp band at 880-860cm-' for samples with x30 42 is due to traces of SrCO, (out-of-plane deformation of the carbonate ion) The FT-Raman spectra of Ti02 and of Sr, llTi02 11 calcined at 773 K are reported in Fig 7 The FT-Raman spectrum of the strontium-free sample (and of Sr, ,,T102 02, which is very similar) show the very intense peaks of Ti0,-anatase 639, 517, 398, 196 and 144cm-', which correspond to the six funda- mental modes of this structure, because of the superimposition of two of them However, weaker features already assigned to brookite can also be found at 453, 366, 323 and 246 cm-' l2 The FT-Raman spectrum of Sr, ,,TiO2 11 presents the anatase peaks, with nearly the same intensity as pure TiO,, but the brookite peaks decreased Moreover, the scattering baseline increases progressively from ca 1000 cm-' towards lower frequencies, and a broad component can be found in the region 950-700cm-' These features may be associated with the increased disorder corresponding to the decrease of crystal- linity Further increase of the Sr content causes the almost 1200 1000 800 600 400 200 wavenumberfcm-l Fig. 6 Skeletal FTIR-FTFIR spectra of SrTiO, calcined at 393 (a), 773 (6) and 973 K (c) y 0.25.1200 1000 800 600 400 200 wavenumberfcm-1 Fig. 7 Skeletal FT-Raman spectra of TiO, (a)and of Sr,,,,Ti02.1, (b), calcined at 773 K (two different scale expansions) complete disappearance of the anatase peaks (for samples calcined at 773 K). Very weak and broad Raman peaks are observed for Sro.42Ti02.42 calcined at 773 K, near 870 and 240 cm-I [Fig. 8(a)], which could be associated to scattering from amorphous materials. The broad peaks are also found in -~ k , , a 1 , , , , , , 1 1200 lo00 800 600 400 200 wavenumberkm-1 Fig.8 Skeletal FT-Raman spectra of Sro,42Ti02,42[(a),(c)] and SrTiO, [(b),(d)]calcined at 773 K [(a),(b)] and at 973 K [(c),(d)] the case of SrTiO, calcined at 773 K [Fig. 8(b)]with additional peaks at 549cm-' (rather broad) and at 180 and 149cm-1 (both sharp). Further heating at 973 K causes the almost complete disappearance of the broad peaks near 870 and 240 cm-' [Fig. 8(c) and (41,so that the Raman patterns are even weaker. However, in the case of SrTiO,, weak scattering peaks are still apparent at 550, 470, 178 and 148crn-'. The sharp peak near 1070 cm-' (out-of-plane deformation of the carbonate ion) confirms the presence of traces of SrCO, in the samples with x20.42.Note that the 'cubic perovskite structure' of SrTiO,, space group Pm3m =Ot,2 = 1, is first-order Raman-silent. In fact, the irreducible representation for the optical modes is: To,,=3 Flu(IR)+F,, (La.) This means that only three triply degenerate IR-active modes are expected, with no Raman-active fundamentals. However, Raman-active combinations and overtones can be found using low-temperature monocrystal mesurements.22 The peaks we observe, however, do not correspond to such a second-order spectrum. Nevertheless, the peaks we found correspond nicely to those assigned by Nilsen and Skinner22 at 551,450, 176 and 146cm-', for an 'impure' SrTiO, monocrystal, to the three IR-active translational fundamental optical modes (vq 55 1, v2 J.Muter. Chem., 1996, 6(5),879-886 883 176 and v1 146cm-l) and to the longitudinal component of v4 (450 cm-l) These authors attnbute the appearance of these peaks in the Raman spectra to a distorsion of the impure crystal at low temperature In our case, these modes may appear due to imperfect crystallinity and to the presence of defects The extremely weak Raman spectra show that the SrTi0, samples we prepared are 'truly' cubic perovskites, since they are essentially Raman-silent This differs strongly from the high-area sample of BaTiO, we investigated previo~sly,~~ that looks cubic according to XRD but gives rise to a very strong Raman pattern similar to that of the 'normal' tetragonal structure, whose tetragonal domains are probably very small and disordered, thus being averaged upon XRD analysis UV-VIS characterization In Fig 9 the UV-VIS diffuse reflectance spectra of TiO,, Sr, 02T102 02, Sr, llTi02 11, Sr, 42T10242, and SrTiO,, all cal- cined at 773 K, are reported All present an absorption edge in the region 350-450 nm, and possibly two main absorptions at lower wavelengths The electronic structure of T102polymorphs and of perovsk- ite-type titanates has been the object of previous experimen- ta15 23 24 and the~retical~~ investigations The valence band is generated by the 2p oxygen orbitals while the conduction band is essentially due to the 3d orbitals of titanium, thus the absorption edge is due to an 02-+Ti4+ charge-transfer transition Our spectra show a progressive shift to lower wavelengths of the edge by increasing Sr content for the samples TiO,, Sr, 02Ti0202, Sr, ,,T102 11 and Sr, ,,Ti02 42, from ca 390 to ca 320 nm Correspondingly, the onset shifts from ca 430 nm to 380nm (E, shifting from 288 to 3 26eV) Examination of the spectra seems to indicate that an higher wavelength component of the absorption shifts to lower wave-length or progressively disappears, while the lower wavelength component remains unaffected However, in the case of the crystalline SrTiO, samples a component grows again and gives rise to a nearly split edge at 320 and 370 nm The UV-VIS spectra of our SrTiO, samples do not change significantly with increased heating temperature and agree well with the spectra of bulk SrTiO, samples reported '24 According to Kutty and Avudaithai,' bulk SrTiO, has a bandgap of 3 2 eV (387 nm), evident as a pro- nounced shoulder similar to that found (at a slightly higher wavelength) in Fig 9, and also shows three absorption compo- nents whose positions depend on the preparation procedure The position of the higher wavelength component found by 200 400 600 BOO wavelengthfnm Fig.9 UV-VIS diffuse reflectance spectra of Ti0, (a), Sr, 02Ti02o2 (b), Sr, 11T10211 (c), Sr, 42T10242 (d) and SrTi03 (e),all calcined at 773 K 884 J Mater Chew, 1996, 6(5),879-886 these authors (300-350 nm) depends strongly on the crystal size, while the other two (265-270 and 210nm) are less sensitive According to the literat~re,~ 23 25 the position of the edge in titanates is also associated with the more or less pronounced deformation of the octahedron around titanium, and with the arrangement of the octahedra in the structure, which modifies the breadth of the lower energy t,, part of the Ti d band According to these data, the shift to lower wavelength of the edge with Sr addition in the SrxTi02+x samples (x<<l) seems to occur in parallel with the progressive decrease of the crystallinity of anatase to an amorphous material In effect, the dried precipitates, which are amorphous or poorly crystal- line, show their edges at even lower wavelengths The data reported above show that the addition of Sr causes a deep modification of the electronic state of anatase, which may be important in relation to the use of Ti0,-anatase as a support for vanadia catalysts These materials need mor- phology and structure stabilizers, in order to inhibit anatase sintering and transformation to rutile, both of which are favoured by vanadium oxide l7 On the other hand, basic dopants should also be added, to improve selectivity in the oxidation of hydrocarbons In both respects, Sr addition may be beneficial However, according to our interpretation, the role of anatase as an optimal support for these catalysts lies in the ability of this semiconductor phase to interact with V centres and exchange electrons with them The significantly different energy gaps in the two T10, polymorphs anatase and rutile (E, anatase >E, rutile) may be a reason for their different behaviour in this respect The addition of Sr to anatase, resulting in a further enhancement of E,, does not necessarily further improve the behaviour Surface characterization by FTIR of adsorbed probe molecules To gain information on the nature of the cationic centres exposed at the surfaces of these solids, we investigated the adsorption of the basic probe pyridine (Fig 10) As is well known, some bands of pyridine are sensitive to the strength of the coordinative interaction involving its own nitrogen lone pair The most sensitive bands are the so-called 8a, 19b, 12 and 1 modes, which are observed in liquid pyridine at 1580, 1438, 1029 and 991 cm-' 26 all of which tend towards higher wavenumbers the stronger the interaction, 1 e the stronger the Lewis acidity of the site 27 According to previous studies,28 the positions of the 8a and 19b modes for pyridine adsorbed on Ti0,-anatase are ca 1608 and 1445 cm-l, which manifest the medium Lewis acidity of Ti4+ cations The 12 and 1 modes fall close to the cut-off limit of T102 By increasing the Sr content, a new band arising from the 8a mode adsorbed on a different site grows progressively at 1595 cm-' and becomes predominant with respect to the band at 1608cm-I in the case of the sample Sr, ,,T102 11 Correspondingly, the cut-off limit shifts to lower wavenumbers and the 12 and 1 modes can be better observed They are both split and their higher frequency components at 1044 and 1012 cm-' decrease in intensity while their lower frequency components at 1034 and 1001 cm-' increase in intensity, caused by increasing the Sr content Meanwhile, the 19b mode shifts progressively to 1441 cm-' The evident splitting of the 8a, 12 and 1 modes in the Sr,TiO,+, samples with 0<x <1 demonstrates that two defi- nitely different sites are present on such surfaces, whose relative amounts progressively invert The 8a, 19b, 12 and 1 modes at 1608, 1446, 1044 and 1012cm-' are typical of pyridine mol- ecules coordinated on sites with medium Lewis acid strength, 1 e on coordinatively unsaturated Ti4+, very similar to those observed on anatase In contrast, the bands at 1596, 1441, 1034 and 1001 cm-' are associated with molecules coordinated I I 1&0 1400 12bo 1000 wavenumber/cm-1 Fig.10 FTIR spectra of the adsorbed pyridine species on pressed disks of T10, (a), Sr, ,,TiO, 02 (b), Sr, llTi02 l! (c), Sr, ,,Ti02 42 (d) and SrTiO, (e), all calcined at 773 K and activated at 773 K, and outgassed at 373 K after pyridine adsorption to sites with low Lewis acidity, which can be identified as Sr2 cations.+ The bands associated with pyridine interacting with Sr2+ ions are already intense for the sample Sr0.02Ti02.02, and are definitely predominant with respect to those of pyridine inter- acting with Ti cations for the sample Sro.llTi02 ll.This should indicate that Sr2+, although coprecipitated with TiO,, is located preferentially at the surface of anatase. We can calculate that for Sr, ,,T~O, ,,,assuming all Sr2+ cations are located at the surface, 88 A2 per Sr atom is available. The same calculation for the sample Sr, 11Ti02 11 gives rise to the a cation density of 1 Sr cation per 17.2A2. These data can be compared with the cation dEnsity for a perfect {OOl} plane of anatase, 1 Ti ion pFr 14.2 A,, and for the {1001 face of SrTiO,, 1 Sr ion per 15.2 A,. This means that, in the case of the sample Sr0.02Ti02.02, all Sr ions can be easily accommodated at the surface; for x= 0.11, Sr ions (if all are located at the surface) would already saturate the surface.For higher x, Sr is necessarily also in the bulk. Interestingly, on SrTiO, only the bands of pyridine adsorbed on Sr2+ are observed. Thus, it is concluded that at the surface of SrTiO, only Sr2+ cations are exposed while Ti4+ cations remain coordinatively saturated in a lower layer. The positions of the main bands of adsorbed pyridine on SrTiO, are similar here with respect to those observed pre- viously on SrZr03,15 on BaTi0,I4 and on lanthanum metal- 1ates.l' It can be concluded that the surfaces of ABO, perovskite-type structures, which are generated by the presence of very large A cations that cannot enter a close-packed array of oxygen ions where, instead, the B cation can enter, are dominated by the presence of A cations. These sites, due to their large size and relatively small charge, are very weak Lewis acids, even if they are coordinatively unsaturated.Correspondingly, studies of CO, adsorption on Sr,TiO, + show that, by increasing x, carbonate species in increasing amounts and adsorbed with increasing strength are formed. The desorption of such carbonate and bicarbonate species from TiO, can be obtained by outgassing at room or only slightly higher temperature, according to the literat~re.~~?~' In contrast, on SrTiO,, to desorb partially carbonated species outgassing at a high temperature such as 973 K must be accomplished. Similar data have been found over other perovskite-type corn pound^.'^^^^ This demonstrates that the very weak acidity of such surfaces goes together with a very strong basicity and/or nucleophilicity, reasonably due to highly uncoordinated exposed oxide ions.py,,,.,--7 400 i900 5600 3'00 3600 Ed0 3400 I300 3200 3130 wavenumberkm-1 Fig. 11 FTIR spectra of the OH groups of TiO, (a), Sro ,,T102 11(b), Sr, 42Ti0242 (c) and SrTiO, (d), all calcined at 773 K and outgassed at 773 K IR evidence of bulk hydrogen impurities in SrTiO, In the IR spectra of all samples (pure powder pressed disks) bands in the OH stretching region 3800-3000cm-1 are observed (Fig. 11). We can distinguish in all cases a weak complex band in the region 3730-3680cm-1. This is the typical region for the free surface OH groups on metal oxides, and, in particular, on anatase and rutile Ti0228.31 as well as on alkaline-earth-metal oxides.31 However, this absorption seems to be decreased in intensity by increasing the Sr content.This can be interpreted, assuming that Sr2+ exchanges H+ at the OH groups, according to its location on the surface, as deduced by pyridine adsorption experiments, and to its effect on sintering. At lower frequencies a strong sharp band is observed only in the case of SrTiO, at 3403cm-1 which may have a component on its lower frequency side. This band, which is definitely unusual for both binary and ternary metal oxide^,^"^^ corresponds to the band observed at 3477cm-1 in the IR spectrum of high-area BaTiO, ~0wders.l~ This band has also been observed by several authors in SrTiO, monocrystals, and was assigned to bulk hydrogen It is well known that hydrogen impurities can penetrate several perovskite structures36 as Hf bonded to a lattice oxygen in the form of an OH-group.These protons can compensate the cation charge defect due either to reduced centres like Ti3+ (or trivalent dopants) or to cation vacancies in non-stoichio- metric samples. Conclusions The conclusions from the present study are as follows. (i) The addtion of Sr to TiO, hinders both the crystallization of anatase from the amorphous state and the phase transform- ation to rutile. (ii) Sr addition also hinders the sintering of anatase, thus allowing higher surface areas to be obtained. (iii) Sr addition to anatase causes a significant shift of the absorption edge to higher energies, thus evidencing an import- ant electronic perturbation of the solid.(iv) Although coprecipi- tated with titanium, strontium cations tend to cover the TiO, surface, possibly by exchanging with H+ of the surface hydroxy groups, and this leads to a drastic modification of the surface chemistry of anatase. (v) When Sr :Ti =1:1 the sample crys- tallizes at room temperature into the cubic perovskite phase, with surface areas approaching 90 m2 g-' that progressively decrease by calcination. (vi) The Raman spectrum shows that this phase is truly cubic, since it is essentially Raman silent. Only extremely small distortions give rise to very weak Raman peaks. This behaviour differs from that of high-area BaTiO, which appears cubic to XRD but is 'microscopically' tetragonal J.Muter. Chem., 1996, 6(5),879-886 885 as clearly deduced by the very strong Raman pattern (viij The surface of SrTiO, exposes only Sr2+ cations (very weakly acidic) and very strongly basic oxide centres (viiij The bulk of SrTiO, contains hydrogen impurities in the form of internal OH groups This work was supported in part by MURST (Romaj J M G A acknowledges Iberdrola (Madrid, Spain) for a ‘Beca de Investigacion Cientifica y Desarrollo Tecnologico’ References 1 W Buchner, R Schliebs, G Winter and K H Buchel, in Industrial Inorganic Chemistry, VCH, Berlin, 1989, p 523 2 D M Smyth, Adv Ceram, 1987,23,339 3 T Bieger, J Maier and R Waser, Sensors Actuators, 1992,7,763 4 M Matsumura, M Hiramoto and H Tsubomura, J Electrochem SOC,1983,130,326 5 T R N Kutty and M Avulaithai, in Properties and Applications of Perovskite-type Oxides, ed L G Tejuca and J L G Fierro, Marcel Dekker, New York, 1993, p 307 6 I Amato, Ceram Acta, 1989,1, 15 7 L G Tejuca, J L G Fierro and J M DTascon, in Adv Catal, 1989,36,237 8 M F M Zwinkels, S G Jaras, P G Menon and T A Griffin, Catal Rev Sci Eng ,1993,35,319 9 M S Wainwnght and N R Foster, Catal Rev, 1979,19,211 10 H Bosch and F Janssen, Catal Today, 1988,2,369 11 G L Haller and D E Resasco, Adv Catal, 1989,36, 173 12 G Busca, G Ramis, J M Gallardo, V S Escribano and P Piaggio, J Chem SOC, Faraday Trans 1,1994,90,3181 13 J M Gallardo, V S Escribano and G Busca, J Muter Chem, 1995,5,1245 14 G Busca, V Buscaglia, M Leoni and P Nanni, Chem Muter, 1994,6,955 15 M Daturi, G Busca and R J Willey, Chem Muter, 1995,7,2115 16 A R West, Solid State Chemistry and its Applications, Wiley, New York, 1984, p 174 17 G Ohveri, G Busca, G Ramis and V S Escnbano, J Muter Chem, 1993,3,1239 18 L Pauling, The Nature of the Chemical Bond, Cornell University Press, Ithaca, NY, 3rd edn , 1960 19 A Cocco and F Massazza, Ann Chim Rome, 1963,53,892 20 S J Gregg and K S W Sing, Adsorption Surface Area and Porosity, Academic Press, NY, 2nd edn ,1982 21 M I Diaz Guemes, T Gonzalez Carreno and C J Serna, Spectrochim Acta, 1989,45, 589 22 K W G Nielsen and J G Skinner, J Chem Phys , 1968,48,2240 23 H Bevan, S V Dawes and R A Ford, Spectrochim Acta, 1958, 13,43 24 L G J DeHaart, A J DeVnes and G Blasse, J Solid State Chem , 1985,59,291 25 J K Burdett, T Hughbanks, G J Miller, J W Riochardson and J V Smith, J Am Chem SOC, 1987,109,3639 26 L Corrsin, B J Fax and R C Lord, J Chem Phys, 1953,21,1170 27 M Taillandier and E Taillandier, Spectrochim Acta, Part A, 1969, 25,1807 28 G Busca, H Saussey, 0 Saur, J C Lavalley and V Lorenzelh, Appl Catal, 1985,14,245 29 G Ramis, G Busca and V Lorenzelli, Muter Chem Phys, 1991, 29,425 30 C Morterra, A Chiorino, F Boccuzzi and E Fisicaro, 2 Phys Chem Neue Folge, 1981,124,211 31 H P Boehm and H Knozinger, in Catalysis Science and Technology, ed J R Anderson and M Boudart, Springer Verlag, Berlin, vol 4, 1983, p 39 32 G Busca, V Lorenzelli, G Ramis and R J Willey, Langmuzr, 1993, 9,1492 33 F G Wakim, J Chem Phys, 1968,49,3738 34 S Kapphan, J Koppitz and G Weber, Ferroelectrics, 1980,25,585 35 A Jovanovic, M Wohlecke, S Khappan, A Maillard and G Godefroy, J Phys Chem Solids, 1989,50,623 36 Yu M Baikov and E K Shalkova, J Solid State Chem, 1992, 97,224 Paper 5/07935G, Received 6th December, 1995 886 J Muter Chem , 1996, 6(5),879-886

ISSN:0959-9428    

DOI:10.1039/JM9960600879

出版商:RSC    

年代:1996

数据来源: RSC

摘要:

Photoluminescence study of (CaO), -x( ZnO), powder solids in air Loukia A. Loukatzikou*, Antonios T. Sdoukos and Philip J. Pomonis Department of Chemistry, University of Ioannina, Ioannina 45110, Greece Luminescence from (CaO), -,(ZnO), (x =0-1) powder solids, prepared by thermal treatment in air at 1423 K, was observed at room temperature by means of high-energy UV excitation (5.2-6.2 eV). In general, broad emission bands centred in the visible region of the spectrum were recorded, with features dependent upon the excitation wavelength as well as on the sample nature. The presence of small quantities of zinc within the CaO results in a dramatic enhancement of the photoluminescence intensity, reaching a maximum at 1YOaddition of Zn atoms/Ca +Zn atoms. In this case, a nearly twenty-fold increment in photoluminescence intensity was observed.The emission band showed a peak at 435 nm (2.85 eV) and was excited to a maximum by 210 nm (5.9 eV) light. Photoinduced electron transfer, followed by recombination processes involving oxygen anion vacancies as well as photogenerated F+ and F centres, are possibly the cause of the observed luminescence. In recent work based on collected literature data, Blasse' showed that complexes [M(d10)],[02-]y may show efficient luminescence at room temperature when excited with high- energy UV radiation. This report also drew attention to the fact that the spectral features of these complexes are similar to those of several oxides which do not contain any d10 ions. Also, Kunkely and Vogler2 reported highly efficient photo- luminescence of the molecular species Zn,O(CH, COO), in solution at room temperature.The reported spectra were similar to those of Zn,0(B02), crystal^.^ This similarity was attributed to the presence of a particular cluster of atoms, since both compounds were characterized by a tetrahedral arrangement of four Zn atoms around the central oxygen., Molecular cluster calculation^^-^ indicated the above arrange- ment to be a very good model of the oxygen chemical environment in ZnO. Photoluminescence from ZnO powders is well known.6-'' A narrow ultraviolet band and a broad green band are observed usually, with an intensity ratio dependent on the preparation as well as the excitation The ultraviolet emission is of an excitonic nat~re,~'~ while the green emission appears to be associated with the presence of oxygen anion vacancies near the surface of the ample.^.^?'' On the other hand, cathodoluminescence12 and X-ray l~minescence'~ have been observed from CaO powders in air although the well known photoluminescence of outgassed CaO powders in UUCUO~~-'~is reported to be 'quenched' by oxygen or air.Cathodolumin- escence has also been observed from CaO :Zn by Lehmann,I2 but no photoluminescence has been recorded so far, although Blasse reported that it might be possible.' In this paper we present for the first time the photoluminescence spectra of (CaO), -,(ZnO), polycrystalline powders in air. These mate- rials show an intense luminescence at room temperature, which was previously overlooked probably because of the high- energy UV radiation required for their excitation.Experimental Polycrystalline powders of (CaO), -,(ZnO), (x=0.000, 0.005, 0.010, 0.050, 0.100 0.300, 0.500, 0.700, 0.900, 0.950, 0.995 and 1.000) were prepared by firing intimate mixtures of CaO (Fluka, p.a) and ZnO (Merck, p.a). The trace element content of the starting materials was mainly Fe (~0.05%) and Pb (<0.01YO),with some Cu and Zn (<O.O05%) for the CaO and Pb (0.005Y0) for the ZnO. The starting materials were mixed and ground thoroughly before firing in air for 5 h at 1423 K. This procedure was repeated three times. Samples were quickly removed from the 673 K furnace after the high-temperature treatment, and were ground and stored in glass vials.These samples will be hereafter designated as CZ:y, where y is defined as the percentage of Zn atoms over the total metal (Zn +Ca) atoms in the sample [y =100Zn/(Ca+Zn); Table 1). We emphasise that neither Ca(OH), nor CaCO, were detect- able by thermal analysis examination of the solids treated as described above. X-Ray diffraction (XRD) patterns of the oxides were meas- ured at room temperature in a Philips system (PW 2253 lamp, PW 1050 goniometer, PW 1?65/50 analogue detector) by using Cu-Ka radiation (2. =1.542 A). Photoluminescence measurements were carried out in air at room temperature, using a Perkin Elmer LS-3 fluorescence spectrometer equipped with a 150 W xenon lamp and a side- window photomultiplier tube (EM1 9781 RA), with reflection grating monochromators with fixed slits of 10nm.Thus, the wavelength accuracy was +2 nm. The scanning rate was 120nm min-'. The powders were held in a front surface accessory and a 1% neutral density filter was placed in front of the emission monochromator in order to record the emitted light. All excitation spectra were corrected for the xenon lamp intensity as well as for the excitation monochromator efficiency. The emission spectra were not corrected for the photomultiplier response. The maxima of spectral bands were identified by point-to-point measurements of the luminescence intensity us. wavelength in the range of interest. Successive measurements of a maximum were found to lie within the instrument's repeatability (+ 1nm).Table 1 Designation of the prepared samples designation sample [y = 100 Zn/(Ca +Zn)] CaO cz:0.00 Ca0.995Zn0.0050 Ca0.99Zn0.010 Ca0.95Zn0.050 Ca0.90Zn0. loo CZ :0.5 cz:1 cz:5 cz:10 Cao.7oZno.3oO CZ: 30 Ca0.50Zn0.500 CZ: 50 Cao.3oZno.7oO CZ :70 Cao.1oZno.900 CZ : 90 OsZnO. 95 cz :95 Ca0.01Zn0.990 cz :99 Ca0.005Zn0.9950 ZnO cz : 99.5 cz:100 J. Muter. Chem., 1996, 6(5), 887-893 887 Results The XRD patterns of the CZ :y solids are shown in Fig. 1. The patterns are indicative of well crystallized solids and show the crystal phases of one or both of the parent oxides in varying ratios. CaO was the only phase detected in the case of CZ :0-1 solids, while the ZnO phase alone was detected in the CZ:99-100 solids.However, the patterns of both CaO and ZnO were detected in all the other solids, e.g. CZ: 5-95, indicating that they are mixtures of the parent oxides in varying ratios. The CaO phase dominates for the CZ: 5-30 samples, while that of ZnO is dominant in the CZ :50-95 ones. The emission spectra of (CaO), -x(ZnO), powders excited by 200, 215 and 240 nm light are shown in Fig. 2. Upon excitation in the 200-240nm range, an intense visible lumi- nescence visible to the naked eye was observed from all the samples and broad emission bands in the 300-650nm region were recorded. Thus, the Stokes' shifts of the various emission bands are very large (2.5-4 eV).The emission bands may consist of one, two or three components peaking in the visible range of the spectra, although ultraviolet emission is also recorded, either as a shoulder (A =360 nm for sample CZ :1) or as a tail of the main emission. A very weak red emission band, with a peak at ca. 726 nm, can also be distinguished for the samples CZ :0.5-5. A similar red emission is also apparent in the case of pure CaO, but not in that of pure ZnO. All the emission spectra recorded for excitation at 200nm consist of broad bands with at least three overlapping compo- nents [Fig. 2(a)]. Although the overall and the relative compo- nent intensities vary across the samples, the emission band maxima are similar. The main emission band has a maximum either in the blue (Lax=435-448 nm for CaO and CZ :0.5-95 solids) or in the blue-green (A,,, =473 nm for CZ :99-100) region of the spectrum.CZ:30 i I 0cz:10 I I0 0 cz:5 I 1 cz:1 i I 0 CaO I 1 1.5 2.0 2.5 Excitation of CZ: 0.5-70 solids at 215 nm [Fig. 2(b)] results in a broad emission band with a maximum in the blue region of the spectrum (Amax =435-448 nm). Nevertheless, the corre- sponding spectra of the other CZ :y solids are similar to that obtained from excitation at 200 nm. The most striking result of the Zn addition to CaO (0.5-90%) is a dramatic enhance- ment of the photoluminescence intensity, which is maximized for 1% Zn addition. In this case, a nearly twenty-fold increase of the luminescence intensity was observed with respect to that of CaO [Fig.2( b)]. The emission now reaches a maximum at 435 nm (2.85 eV) and is excited to its maximum by 210 nm (5.9 eV) light. The emission spectra of CZ :0.5-70 recorded upon excitation at 240 nm are similar to the corresponding ones obtained from 200 nm excitation (Fig. 2). However, excitation of CaO and CZ: 90-100 at 240 nm gives rise to some special spectral features [Fig. 2(c)]. In particular, the blue emission maximum of CaO (A=444-446 nm) is now replaced by a violet one at 409 nm. At the same time, blue-green (Amax =479 nm) and green (Lax=527 nm) emissions are recorded as shoulders on the main violet bond. The emission spectra of CZ:90-100 consist now of two partly overlapping bands with maxima at 508-523 nm and 406-416 nm.The green emission clearly dominates over the violet one. The photoexcitation spectra of some CZ :y solids are shown in Fig. 3. The corresponding excitation spectra of the other solids are quite similar to that of CZ:O.5. The CaO and CZ :99-100 excitation spectra recorded for an emission wave- length of 440nm show a broad band with a maximum at 200-208 nm, as well as a second weak band at 230-238 nm. However, the same emission (440 nm) is the result of only one excitation band peaking at 210-218nm in the case of CZ :0.5-95 solids. The above excitation bands also result in a green emission (520 nm), although such an emission arises also . , 1.....ll..l.t ' ..L .II CZ95 I I cz99 1.. ..11T r... I....I .. ~ 1.5 2.0 2.5 dlA Fig. 1 XRD patterns of CZ :y solids: 0,CaO; ,ZnO 888 J. Muter. Chem., 1996, 6(5),887-893 0.00 0.00 I I I I I 250 350 450 550 650 750 250 350 450 550 650 750 300 400 500 600 700 A/nm Fig. 2 Emission spectra of CZ :y solids ( y =0.00-100): (a),Iex200 nm, (b) ,Iex=215 nm, (c) I,,, =240 nm= from a new excitation band (380-390 nm) for CZ :70-100 (Fig. 3). Excitation of the CZ: 99-100 samples is also apparent in the whole A= 200-400 nm region. We emphasise that although the emission of CZ :0.5-70(4 solids peaks at nearly the double the excitation wavelength, the increased luminescence intensity is not a result of a ‘second- -100 order’ effect. In fact, the emission maximum shifts only slightly -99s with the excitation wavelength in the whole 200-240 nm range, -99 while the overall emission also shows an increased intensity. So, the increased luminescence intensity is an intrinsic property of the solids and is not enhanced by secondary effects.We have also noticed that the emission spectra were not corrected for the photomultiplier response. This could actually affect the spectral shapes, especially the relative band intensities recorded. However, all the spectra should be uniformly affected and the picture obtained would not be modified drastically. Discussion CaO The excitation spectrum of CaO powders in air (Fig. 3) is rather similar to the absorption spectrum of deformed CaO single crystals,” which exhibit maxima at 214 and 270 nm, although the main absorption peak is that at 270nm. A similarity has also been recognized between the absorption bands of deformed crystals and those of high-surface-area 0.5 alkaline-earth-metal oxide powders.20-22 The former was attri- -buted to vacancy clusters formed during deformation,” and0.00 the latter to transitions from low-coordination oxygen 200 300 200 300 400 anions.21,22 Excitation of the alkaline-earth-metal oxide powders inA/nm DUCUU, with energies in the region of their optical absorptions, Fig.3 Excitation spectra of CZ:y solids: (a) ,Ie,=440 nm, (b) ,Iem= produce lumines~ence~~-~~~~~ that has also been attributed to 520 nm transitions from low-coordinated surface ions.14-16 Such ions J.Muter. Chem., 1996, 6(5),887-893 889 have been assumed to be present in the outgassed alkaline- earth-metal oxide powders as a result of their high-temperature treatment zn ~UCUU,~'22 but their involvement in photoprocesses over powder oxides has been questioned24 26 However, the excitation energies of the used CaO powders in air are rather different from those of CaO powders zn z)ac~u,'~'' with the exception of an excitation near 240 nm The other excitation bands of outgassed CaO powders are all centred on smaller energies Luminescence from outgassed CaO powders is reported to be 'quenched' by molecular oxygenI4 So, the absence of luminescence when CaO powders were excited in air in the 260-380nm range may be attributed to the action of oxygen The emission spectra recorded for CaO powders are not identical to any spectrum of the oxide reported in the literature, although similar emissions have been observed Excitation of CaO at 200-215 nm causes a major blue emission with a maximum at 444-446 nm [Fig 2(a),(b)] A broad emission band which peaks at 426nm has also been observed from outgassed powders in uucuu by excitation in the 235-300nm region An emission peaking at 454nm has also been observed from CaO single crystals upon excitation at 270 nm 27 CaO spectra in air also show emission bands with maxima at 479-480 and 525-527 nm (Fig 2) The emission band from deformed single crystals has a maximum at 477 nm and a shoulder at cu 510nm is also clear," but in this case no distinct emission peak was observed at shorter wavelengths However, the emission spectra of deformed CaO crystals reported in the literature seem to be more similar to the present results Although broad emission bands centred at 495 or 477nm have also been recorded from outgassed CaO powders zn U~CUUby excitation at 330 nm,I4 a green emission has never been observed from the oxide zn U~CUU However, the green emission is the result of air-annealing of deformed single ~rystals,,~ although it has also been observed for undeformed crystals 28 Excitation of CaO at 240nm results mainly in a violet emission band peaking at 409 nm [Fig 2(c)] An emission band with a maximum at 405 nm has also been obtained from outgassed powders zn uucuu, by excitation at 282 or 310 nm Note that the violet emission is the major one observed from these powders at low temperatures for excitation over the whole 235-300 nm range l5 In conclusion, the photoluminescence spectra of CaO pow- ders in the present work are not quite identical to some spectra reported in the literature, although similar emissions have been observed However, the material treatment as well as the experimental procedure followed here are quite different from those reported in the literature Thus, the observed spectral differences may be attributable to the different preparation procedure and, in particullar the quick quenching of the solids Note that the luminescence from CaO powders reported here is not a result of grinding A bright blue luminescence quite similar to that of pressurized single crystals has been observed previously from thermally treated CaO powders after grinding,29 but a second firing step was found to remove this luminescence completely In contrast, the luminescence from CaO recorded here is permanent On the other hand, the observed luminescence is not due to the presence of small quantities of Ca(OH),, as a result of the CaO sensitivity to the ambient atmosphere Although no luminescence was observed previously from Ca(OH), after 254 or 365 nm excitation,'2 more re~ently'~ some spectral features of partially dehydroxylated CaO powders were ascribed to the presence of Ca(OH), These features consist of excitations at ca 270 and 300nm, far from the present CaO excitation region, and of emissions at 353, 405 and 443 nm l7 The recorded luminescence does not seem to be a result of the trace elements present in the starting material In fact, Fe is known to act as an efficient fluorescence trap The well 890 J Muter Chem , 1996, 6(5),887-893 known luminescence of Pb2+ impurity ions (<O 01) in CaO,,' 31 which shows vibrational structure at low tempera- tures, is centred on the ultraviolet region (A=348-368 nm), although 0 1% or more Pb in CaO was reported to emit almost exclusively in a broad distribution in the visible region The CaO Cu emission spectrum consists of two overlapping bands with maxima at 390 and 448 nm ZnO The luminescence spectra of the ZnO powders used, excited by high-energy UV light (A =200-215 nm) are different from those reported in the literature, owing to the different energy used for their excitation Thus, the spectra obtained here show very large Stokes' shifts, in contrast to the small Stokes' shift of the well known luminescence of ZnO obtained by excitation near its energy gap The excitation spectrum of ZnO for the green (520 nm) emission [Fig 3(b)] is quite similar to the absorption spectrum of a ZnO thin layer However, the blue (A,,, =435-445 nm) and blue-green (nmax=473-484 nm) emis- sion components recorded here [Fig 2(a)(b)] seem to be only a result of high-energy UV excitation, especially near 208 nm [Fig 3(a)] Excitation of ZnO near its energy gap (Aex= 320-380 nm) results exclusively in a broad green emission band, quite similar to that referred to in the literature67 No ultraviolet emission component was observed at the same time, although such an emission was obtained together with the main green one, after 240 nm excitation [Fig 2(c)] This is probably due to the sample treatment method The emission spectra of CZ 0 5-70 (iex=215 nm) exhibit close similarities to that of Zn,O(BO,), Excitation of these crystals at 250 nm results in a broad emission band with a maximum at 435 nm There is also good agreement between the exci- tation wavelengths recorded from CZ 0 5-90 solids (Aex= 210-218 nm) and the absorption of the Zn,O(CH,COO), solutions at 216nm,2 although the emission from the latter reaches a maximum at shorter wavelengths (A=372 nm) Thus, there is a clear similarity in the luminescence behaviour of Cz 0 5-70 and Zn,O(BO,), solids as well as Zn,O(CH,COO), solutions On the other hand, the lumi- nescence spectra of CZ 0 5-70 solids are quite similar to that of deformed MgO crystals which do not contain any Zn The latter absorb at 2175nm (57eV) and emit at 4275nm (2 9 eV) l9 The similarity between the Zn,O(CH,COO), and Zn40( BO,), luminescences has been recognized and has been related to the common tetrahedral arrangement of four Zn atoms around the central oxygen in these compounds A similarity has also been recognized between the luminescences of the above compounds with those of several oxides which do not contain any d'' ions In general, an analogy between the luminescences of 0x0-d" and 0x0-do complexes has been established,' 32 while the optical transitions involved seem to be of a very complicated nature ' 32 Thus, a question arises about the possible role of either Zn2+ ions (d") or a particular cluster of atoms, in the determinatim of the luminescence behaviour of CZ y solids, in particular CZ 0 5-70 No other crystal phases except for CaO and ZnO were detected in the solids, so the photo-effects obtained should be attributable to a synergistic action between the two oxides The luminescence spectra recorded for the CZ y solids show that such an effect exists This is especially clear in the emission spectra obtained by 240 nm excitation [Fig 2(c)] These spec- tra are characterized by a progressive 'red' shifting across the CZ y samples, as the Zn addition to CaO increases It can be anticipated that the above spectra arise from excitation in the range of a second, minor excitation state of pure CaO and ZnO oxides, which is absent in the case of CZ 05-95 solids However, the oxide’s XRD phase prevalence across the CZ :y samples does not affect their spectral features in an expected way.The influence of the ZnO phase prevalence becomes apparent only when the Zn concentration reaches 90%. Thus, although the luminescence behaviour of single-phase CZ:90-99.5 solids is quite similar to that of pure ZnO, the behaviour of CZ :50-70 differs. It rather resembles the behav- iour of CZ: 5-30 as well as that of single-phase CZ: 0.5-1 solids, which in turn is rather different to that of pure CaO.On the other hand, a synergistic action is also clear in the emission spectra obtained by 200 nm excitation [Fig. 2(a)]. A similarity between the spectral features of CZ:y solids as a whole is now clear. The Zn addition to CaO affects the overall luminescence intensity but leaves emission peak positions virtually unchanged. Although such a similarity is not clear in the emission spectra recorded by 215 nm excitation, it seems that the CZ :0.5-70 solids do not exhibit some special kind of luminescence, e.g. different from that of the other samples. The luminescence is now increased dramatically, probably owing to a synergistic action between the two oxides.The emission spectra seem to be rather structureless, possibly because of the relatively increased blue component emission (Lax= 435-448 nm), overlapping with components at longer wavelengths. In conclusion, the emission spectra of CZ :y solids recorded by excitation at 200-215 nm show that the percentage of Zn atoms present in CaO does not play any significant role in the determination of the emission energies, although it determines the overall and the relative intensities of the emission bands. Considerable differences are observed in the emission spectra of the solids recorded for excitation at 240 nm. These are related to the present Zn percentage and probably arise from the influence of a new excitation state of pure ZnO near 240 nm (Fig.3). Thus, a similarity really exists across the luminescence spectra of CZ:y solids as a whole, which also suggests a similarity in the nature of the photoprocesses involved. Proposed photoluminescence model In order to describe the photoluminescence processes for CZ :y solids, a simplified model is proposed based on the experimen- tal data. Indeed, a general energy diagram can be drawn as shown in Fig. 4. This is based on the following experimental observations: EemJ where Eex2denotes the energy corresponding to the weaker excitation band maximum at 230-238 nm (5.2-5.4 eV), recorded in the excitation spectra of some samples (CaO, CZ:99-loo), while Eeml, Eem2, Eem3 and Eem4denote the energies corresponding to the emission bands maxima for the blue-green (460-490 nm), green (500-530 nm) blue (415-450 nm) and violet (400-412 nm) components respectively.A general mechanism explaining the photoprocesses involved as well as the various energy levels shown in Fig. 4 could be as follows. High-energy UV light (Eexl=5.7-6.2 eV) can induce the formation of levels of electrons and holes in the solids, possibly via excitation of lattice 02-ions, Then, level 1 in Fig. 4 corresponds to the top of the valence band of the solid, while levels 2 and 3 correspond to the electron and hole levels generated, respectively. We propose that the elec- trons released from the valence band could arrive initially in the vacuum level.In fact, the vacuum level lies about 6eV above the top of the valence band of most semiconductor oxides,33e.g. near to the main excitation energy EeX1recorded for our solids. The main excitation energy of CaO is smaller 0 0 Fig. 4 General energy diagram describing the photoluminescence pro- cesses over CZ :y solids: 0,electron; 0,hole that its bandgap, but in this case the vacuum level lies inside the bandgap, as a result of its negative electron affinit~.~~.~~ Then, emission could be a result of recombination processes between photogenerated electrons and holes, possibly proceed- ing through oxygen vacancies present in the solids as a result of their treatment. The following general emission mechanism is possible: during the trapping of an electron from level 2 (Fig.4) by an oxygen vacancy, blue-green light (460-490 nm) is produced. This process results in the formation of an F+ centre. Next, a second electron from level 2 could be trapped by an Ff centre to form an F centre and produce the green emission (500-530 nm). The above processes are both accompanied by changes in the oxidation state of the ions of the electrons level. Then, the F+ and F centres electrons recombine with holes at level 3. The recombination of an electron coming from either an Ff or an F centre with a hole is accompanied by either the blue (415-450nm) or the violet (400-412 nm) emission respectiveiy. At the same time, an oxygen vacancy or an F+ centre is reformed. If so, levels 4 and 5 probably correspond to F+ and F centres levels, respectively.It must be stressed here that no distinct levels, but energy bands should be formed in the CZ:y solids during the irradiation, as the observation of broad excitation and emission bands suggests. However, although the problem of optical transitions involved in the above luminescence processes is probably of a very complicated nature, the proposed simplified model aims to provide a first approach to the matter. A complete answer to this problem requires at least a more detailed energy-level calculation. By applying the above general model in the cases of CaO, CZ :0.5 and ZnO solids, we obtain the energy diagrams shown in Fig. 5 (a), (b) and (c), respectively. These diagrams were designed by first fitting the experimental energies correspond- ing to the maxima of the emission bands recorded by 200 nm excitation, and then those corresponding to the maxima of the main excitation bands, recorded for the emission at 440 (CaO, CZ :0.5) or at 520 nm (ZnO).These energy values are indicated by bold characters in the Figures. The bandgaps of the pure oxides used are based on the literature data.34,35 In the case of CZ :0.5 we propose that E, >5.7 eV. The agreement of this model with the experimental data is very good in the case of CaO and satisfactory enough in the other cases. In fact, the estimated distance from Fig. 5(a) between electron and hole levels (5.38 eV) corresponds to the maximum of the second weak excitation band recorded from CaO (Fig.3) at 230 nm (5.4 eV). In addition, the estimated energy between F and hole levels (3.02 eV) corresponds to the violet emission at 409 nm (3.03 eV), recorded upon excitation of the oxide at 240 nm [Fig. 2(c)]. In contrast, neither a distinct emission peak near 3.08 eV (402 nm), nor a second excitation peak near 5.5 eV (225 nm) were recorded in the spectra of the CZ:O.5 solid. Finally, the emission at 406 nm (3.05 eV), J. Muter. Chem., 1996, 6(5),887-893 891 conduction band ,, level EL=? 5.47 cv, (5 I* m (473 nm) I ' . , 023 eV F+ Fig. 5 Energy diagrams describing the photoluminescence processes over CaO (a), CZ : 0.5 (b) and ZnO (c) recorded by excitation of ZnO at 240 nm, corresponds satisfac- torily to the F and holes levels distance, estimated to be 3.02 eV [Fig.5(c)], although the estimated distance between electron and hole levels (5.41 eV) does not correspond very well to the maximum of the second excitation band of the oxide at 240 nm (5.2 eV). According to the proposed energy diagrams, the F and F+ levels are estimated to be 3.16 and 3.39 eV below the conduc- tion band of CaO, as well as 0.03 eV and 0.26 eV below the conduction band of ZnO, respectively. The F-centre ground level has been found to be 3.11 eV below the conduction band of CaO single crystals,36 while the Ft centre absorbs at 3.6-3.7 eV,37-39 which corresponds to the F' distance from the valence band rather than the conduction band of CaO in Fig.5(a). In the case of ZnO, F and F' levels have been estimated to be 0.02-0.05 eV40-42 and 0.19 eV4' or 0.30-0.45 eV4' below the conduction band respectively, in satisfactory agreement to the corresponding energies estimated here from Fig. 5(c). Zn2+ ions, although containing closed d shells, facilitate the photoprocesses over CZ : y solids as described above, because they can really trap electrons and be converted to unstable Zn+ ions, which in turn can easily return electrons to the oxygen vacancies. Thus, an intense luminescence is produced. The blue emission (i,,,= 435-448 nm) dominates for the CZ : 0.5-70 solids, possibly because the recombination pro- cesses proceed mainly through the F+ centres. Bearing in mind that the energy gap of ZnO is 3.34 eV at room temperat~re,~' the F centres position (3.31 eV above the valence band) estimated from Fig.5(b) suggests that Zn2'-F pairs should be formed in the CZ : 0.5 solid after high-energy excitation. If so, the recombination processes should proceed niore easily cia Ff rather than via F centres, because F electrons could then easily be trapped by Zn2' ions. Similar effects are possible in the case of CZ: 1-70 solids. In contrast, the green emission (Amax = 508-523 nm) dominates for the CZ : 90-99.5 solids, as a result of the relatively high conductivity of ZnO. Now, although the F electrons are close to the conduction band [Fig. 5(c)], they prefer to recombine with holes in the hole level. The observation that their excitation by light of energy E < 5.2 eV (i240 nm) results exclusively in green (i,,,= 508-523 nm) and violet (i,,,= 406-416 nm) emission possibly indicates that F centres are formed exclusively under these conditions.The involvement of FC as well as F centres in the photo- luminescence processes over the alkaline-earth-metal oxide powders in mcuo has also been considered in the pa~t.'~,'*,~~ Charge-transfer transitions'6 and relaxation processes have also been con~idered.~~ The model developed here is in good agreement with the interpretation of the phosphorescence as well as the thermoluminescence of UV-irradiated CaO powders in cacu~.'~ The above luminescence, which also exhibits some similarity to the emission spectra of CaO in air.was ascribed to radiative tunnelling recombination processes between elec- tron and hole pairs in distant Ft and V- centres that were pho toformed.18 In conclusion, we have demonstrated in this work that an enhancement of the emission spectra of CaO-ZnO powders is observed by exciting them with i= 200-215 nm. The effect is particularly strong at 1% addition of ZnO into CaO where a 2000% increase of intensity is produced. A simplified photo- luminescence model is proposed, answering the question of the similarity observed across the CZ : y solids spectra as a whole and fitting satisfactory with the experimental data. We are grateful to Professor M. Marselos for the provision of the fluorescence spectrometer for the luminescence experi- ments, and to Dr.M. A. Demertzis for his help in the early stages of this study and for helpful discussions. References 1 G. Blasse, Chem. Phys. Lett., 1990,175,237. 2 H. Kunkely and A. Vogler, J. Cheni. Soc., Chem. Commun.: 1990, 1204. 3 A. Meijerink, G. Blasse and M. Glasbeek, J. Phys: Condens. Matter 2, 1990, 6303. 4 R. Bertoncello, M. Bettinelli, M. Casarin, A. Gulino, E. Tondello and A. Vittadini, Inorg. Chem., 1992,31,1558. 5 M. Casarin, E. Tondello, F. Calderazzo, A. Vittadini, M. Bettinelli and A. Gulino, J. Chem. Soc., Faradaj Trans., 199?,89,4363. 6 G. Heiland, E. Mollwo and F. Stockmann, Solid State Phys., 1959, 8, 191. 7 F. Van Craeynest, W. Maenhout-Van der Vorst and W. Dekeyser, Phys. Status Solidi, 1965,8, 841.8 W. Maenhout-Van der Vorst and F. V. Craeynest, Phys. Status Solidi, 1965, 9, 749. 9 M. Anpo and Y. Kubokawa, J. Phys. Chem., 1984: 88,5556. 10 M. Mienska, R. Leszczynski, S. Karolczak and S. Wysocki, Radiat. Phys. Chem., 1989,33,483. 11 H. Idriss and M. A. Barteau, J. Phys. Chem., 1992,96, 3382. 12 W. Lehmann, J. Luminescence, 1973,6,455. 13 G. Blasse and L. H. Brixner, Mater. Chem. Phys., 1991,28,275. 14 S. Coluccia, A. M. Deane and A. J. Tench, J. Chem. Soc., Faraday Trans. 1, 1978,74,2913. 892 J. Mater. Chem., 1996, 6(5), 887-893 15 16 17 18 19 20 21 22 23 24 25 26 27 28 S. G. MacLean and W. W. Duley, J. Phys. Chem. Solids, 1984, 45, 227. W. W. Duley, Philos. Mag. B., 1984,49, 159.W. W. Duley, High Temp. Sci., 1984, 17,409. Y. Yanagisawa, N. Inishi and A. Narumi, Phys. Rev. B, 1992, 46,11121. Y. Chen, M. M. Abraham, T. J. Turner and C. M. Nelson, Philos. Mag., 1975,32,99. R. L. Nelson and J. W. Hale, Discuss. Faraday Soc.,1971,52, 77. A. Zecchina, M. G. Lofthouse and F. S. Stone, J. Chem. SOC., Faraday Trans I, 1975,71,1476. E. Garrone, A. Zecchina and F. S. Stone, Philos. Mag., 1980, 42, 683. A. J. Tench and G. T. Pott, Chem. Phys. Lett., 1974,26,590. V. A. Shvets, A. V. Kuznetsov, V. A. Fenin and V. B. Kazansky, J. Chem. Soc., Faraday Trans 1,1985,81,2913. J. Cunningham, in Surface and Near-Surface Chemistry of Oxide Materials, ed. J. Nowotny and L. C. Dufour, Elsevier, Amsterdam, 1988,ch. 8.A. M. Stoneham and P. W. Tasker, in Surface and Near Surface Chemistry of Oxide Materials, ed. J. Nowotny and L. C. Dufour, Elsevier, Amsterdam, 1988,ch. 1. J. A. Carcia, A. Remon and J. Piqueras, Phys. Status Solidi. A, 1985,89,237. J. Llopis and J. Piqueras, J. Appl. Phys., 1983,54,4570. 29 H. Donker, W. M. A. Smit and G. Blasse, Phys. Status Solidi B, 1988,145,333. 30 A. F. Ellervee, Phys. Status Solidi B, 1977,82,91. 31 A. C. Van der Steen and L. T. F. Dijcks, Phys. Status Solidi B, 1981, 104, 283. 32 G. J. Dirksen, A. N. J. M. Hoffman, T. P. Vande Bout, M. P. G. Laudy and G. Blasse, J. Muter. Chem., 1991,1,1001. 33 P. A. Cox and A. A. Williams, Surf.Sci., 1986,175, L782. 34 0.V. Krylov, in Catalysis by Nonmetals, Academic Press, New York, 1970. 35 W. F. Wei, Phys. Rev. B, 1977,15,2250. 36 B. Henderson, S. E. Stokowski and T. C. Ensign, Phys. Rev., 1969, 183, 826. 37 A. E. Hughes and B. Henderson, in Point Defects in Crystals, ed. J. H. Grawford, Jr. and L. M. Slifkin, Plenum Press, New York, 1972, vol. 1. 38 L. S. Welch, A. E. Hughes and G. P. Summers, J. Phys. C: Solid State. Phys., 1980,13, 1791. 39 J. C. Kemp, W. M. Ziniker and E. B. Hensley, Phys. Lett. A, 1967, 25,43. 40 A. Poppl and G. Volkel, Phys. Status Solidi A, 1991, 125, 571. 41 W. Cope1 and U. Lampe, Phys. Rev. B, 1980,22,6447. 42 F. A. Kroger, in The Chemistry of Imperfect Crystals, North Holland, Amsterdam, 1964. Paper 5/06576C; Received 5th October, 1995 J. Muter. Chem., 1996,6(5), 887-893 893

ISSN:0959-9428    

DOI:10.1039/JM9960600887

出版商:RSC    

年代:1996

数据来源: RSC

摘要:

Study of the order-disorder transition in yttria-stabilised zirconia by neutron diffraction Iain R. Gibson" and John T. S. Irvineb "Departmentof Chemistry, University of Aberdeen, Meston Walk, Aberdeen, UK AB9 2UE bSchool of Chemistry, University of St. Andrews, St. Andrews, Fife, UK K Y16 9ST A comprehensive study of 8 mol% yttria-stabilised zirconia has been made between 150 and 1000 "C, using ac impedance spectroscopy and high-temperature neutron powder diffraction. It has been demonstrated that the conductivity anomaly, which occurs at ca. 650 "C, is structural in origin. A sharp decrease in the activation energy for conduction of ca. 0.2 eV was observed at ca. 650 "C.Additional broad, diffuse scattering peaks were observed below 600 "C in the neutron diffraction patterns; above 650 "C, the diffuse scattering peaks disappeared. A deviation from linearity was observed at a similar temperature in the plots of both Y/Zr and 0 isotropic temperature factors us.temperature. The low-temperature behaviour can be explained in terms of ordering of oxygen vacancy-(dopant) cation clusters to form microdomains, which are evidenced by the presence of diffuse scattering peaks. At high temperature, the association of vacancies with defects breaks down, or at least becomes randomised, allowing vacancies to move more freely as indicated by the decrease in activation energy for conduction. A discontinuity in thermal expansion coefficient (from neutron diffraction data) confirms the second-order nature of the transition. Cubic yttria-stabilised zirconia (YSZ) exhibits high oxide-ion conductivity at 1000°C and is used as the electrolyte in solid oxide fuel cell (SOFC) devices.It is generally agreed that the maximum conductivity occurs with an yttria content of around 8 mol%.',2 Increasing the amount of yttria dopant above 8 mol% produces a decrease in conductivity, which is thought to arise from an increase in oxygen vacancy-(dopant) cation associations, reducing the number of 'free' oxygen vacancies available to migrate.3 From conductivity measurements of 8 mol% YSZ between 600 and lOOO"C, a linear Arrhenius plot (log aTvs. 1/T) can be obtained and the activation energy for conduction calcu- lated, although curvature has been observed in Arrhenius plots at lower temperature^.^ In a more recent study, Badwal observed a change in the slope of the Arrhenius plot at approximately 550 OC,' indicating a higher activation energy at low temperatures.We generally observe similar behaviour, e.g. Fig. 1. An increase in activation energy at low temperatures is in accord with the behaviour expected for some form of vacancy- dopant cation association. A number of such models for oxide- ion conduction have been developed to account for results from studies of ceria and perovskite Kilner and Steele8 stated that the low-temperature activation energy com- prised an oxygen vacancy migration enthalpy (AmH),and an association enthalpy (AaH),due to the vacancy4opant com- 2 ---a 1 -. c .2 o--z .E -1 -2 *\ '*, plex.At higher temperatures, the vacancy-dopant complexes dissociate, allowing oxygen vacancies to migrate freely. The activation energy now contains only the migration enthalpy term, AmH.Atomistic calculations predict that the additional association enthalpy, equal to half the binding energy, is of the order of 0.1-0.2eV,' in accord with our recent studies, Table 1." As the amount of dopant increases, the additional association enthalpy at low temperatures also increases (Table 1) indicating an increase in the degree of association. The nature of the vacancy-dopant interaction in YSZ at low temperatures has been reviewed extensively." Although it is widely considered that oxygen vacancies are associated with the aliovalent dopant cation, Y3+, an EXAFS study of 10 mol% YSZ by Catlow et al.indicated that oxygen vacancies were preferentially sited adjacent to the Zr4 cation, and not + the dopant Y3+ cation.12 This results in seven-fold oxygen coordination of Zr4+, similar to its structural environment in monoclinic zirconia; Y3+ ions are in sites of eight-fold coordi- nation. Similar results were obtained in a series of EXAFS studies by Li et a1.13-1s Many structural studies have been made on cubic stabilised zirconias. Room-temperature neutron diffraction experiments on single crystals of YSZ by Steele and Fender,16 and CSZ (calcia-stabilised zirconia) by Moringa et and Cohen et ~1.'~concluded that oxygen atoms were displaced along the (100) direction.In contrast, Carter and Roth" and Horiuchi et studying CSZ and YSZ respectively, reported a dis- placement along the (1 11) direction. Diffuse scattering, or a modulated diffuse background, was a significant feature of the Table 1 Activation energies and association enthalpy of 3, 8-11 molo/o yttiria-stabilised zirconia (in eV) mol% yttria EA(300-500 "C) E, (600-800 "C) A,H 3" 0.94 0.89 0.05 8 1.08 0.92 0.16 9 1.11 0.93 0.18 10 1.14 0.95 0.19 11 1.18 0.98 0.20 " 3 mol% yttria-stabilised zirconia is single-phase tetragonal, the other compositions were single-phase cubic (4-7 mol% is a two-phase mixture). J. Muter. Chem., 1996, 6(S), 895-898 895 neutron diffraction pattern obtained by Steele and Fender,16 and possible short-range ordering by vacancy association was considered.Further work on YSZ and CSZ18*21-25 supports the concept of ordering of vacancy-cation complexes, or clusters, visible as diffuse scattering. A detailed study on the effect of high temperatures on the diffuse background scattering has not been made. High-temperature neutron diffraction experiments on YSZZ6 and CSZ27 have recently been reported. Both studies examined the change in temperature factors of the cations and anions with increasing temperature; however, only a few data points were obtained in the temperature region 300-10OO0C, where the deviation in the activation energy of conduction is observed. Martin et studying CSZ, observed a change in the increase in B,,,(oxygen) at ca.lOOO"C, which is close to the order-disorder tran~formation.'~ Proffen et again studying CSZ, observed a decrease in the diffuse scat- tering as this temperature was approached. Our conductivity data, and the neutron diffraction work on CSZ and YSZ previously discussed, suggest that a similar order-disorder transition may be observed in YSZ at ca. 650°C. Here, we make a detailed study of 8 mol% YSZ using time-of-flight neutron diffraction over the temperature region 25-1000 "C. The effect of increasing temperature on the change in diffuse scattering, isotropic temperature factor, and lattice parameter is discussed. Experimental The powder used in this study, TZ-8Y, was produced by Tosoh Corporation (Japan).For conductivity measurements, the powder was uniaxially pressed at 80 MPa in a 13 mm die. Samples were heated at 10°C min-', sintered at 1500°C for 2 h, and cooled at 10 "C min-l. Platinum electrodes were applied to both faces of the sample. Two terminal ac impedance measurements were performed using a HP4192A impedance analyser, over the frequency range 100 Hz-13 MHz. Measurements were made in air, over the temperature range 300-1000 "C. For neutron diffraction experiments, the powder was sintered at 1500 "C for 20 h. Powder neutron diffraction data were collected on the Polaris diffractometer at the UK spallation neutron source ISIS, Rutherford Appleton Laboratory. Samples were contained in vanadium cans, and a vanadium wound furnace was used between 25 and 1000°C.A thermo- couple was attached next to the sample can, ensuring accurate temperature control of f1 "C. The crystal structures were refined by the Rietveld method with the program TF14LS29730 using data collected over the time-of-flight range 2500-19500 ps. Refinements were carried out using a similar approach to that of Argyriou.26 No attempt was made to model any displacement of the anions in either the (1 11) or the (100) direction. All refinements were carried out in the space group Fm3m, with the occupancy and positions of Zr/Y and 0 fixed (Table 2). Background, lattice parameters and isotropic temperature factors for Zr/Y and 0 were refined for all temperatures. Table 2 Initial parameters for 8 mol% YSZ ZrV) 0 X 0.0 0.25 Y 0.0 0.25 Z 0.0 0.25 occupancy 0.8515 (0.1485) 0.9634 scattering lengths" 0.716 (0.775) 0.5805 " Ref.3 1. 896 J. Muter. Chem., 1996, 6(5), 895-898 Results and Discussion A clear change in the slope of the Arrhenius plot for conduction is observed at ca. 650°C (Fig. 1). Below this temperature the activation energy is higher (1.07 eV), and as the temperature is increased to above 650°C, the activation energy decreases (0.92eV). The transition between the two regions is quite sharp, certainly for an order-disorder process in an ionic conductor, although there is a limited region of slight curvature joining the two linear portions extending over less than 100 "C. Examples of observed and calculated neutron diffraction profiles and their difference curves, at room temperature and lOOO"C, are given in Fig.2. At room temperature, a heavily modulated diffuse background is visible in the difference curve. At 1000 "C, the background appears virtually flat, and the only features in the difference curve greater than 2 esd are associated with 'allowed reflections'. The diffuse scattering peaks observed at room temperature do not correspond to expected reflections from the space group Fm3m, although they can approximately be indexed on a primitive cubic unit cell with the same unit- cell edge. Previous reports indicate that this is due to some type of short-range ordering within the cubic lattice.16 As the temperature was increased, the diffuse background decreased, and above 600 "C it had effectively disappeared (Fig.3). The temperature region where the diffuse peaks are observed corresponds to the region where the activation energy for conduction is larger. This suggests that above 650"C, when the activation energy could be interpreted as arising solely from the enthalpy of migrati~n,~ the cubic lattice no longer shows signs of local ordering. A decrease in the weighted profile residual, R,,, with increasing temperature is another indication of the change occurring at ca. 650°C. It is normal to see a small increase in R,, with increasing temperature as the background noise increases; however, in Table 3, R,, is shown to decrease with increasing temperature, indicating an " 0.5 1 1.5 2 2.5 3 d-spacing/A $400 1 0.5 1 1.5 d-spacingA 2.52 3 4 -10 Fig. 2 Observed and calculated neutron diffraction profiles for 8 mol% yttria-stabilised zirconia at 20 "C (a) and 1000 "C(b) Table 3 Zr/Y and 0 istropic temperature factors (ITF) and R,, T/"C Zr/Y ITF 0 ITF RWP ~~ ~ 150 0.71(1) 2.19( 2) 3.83 225 0.76( 1) 2.28(2) 3.72 300 0.82( 1) 2.37( 2) 3.48 400 0.89( 1) 2.49(2) 3.23 500 0.96( 1) 2.60(2) 2.98 600 1.04( 1) 2.72(2) 2.81 700 1.12( 1) 2.84(2) 2.61 800 1.20(1) 2.98(2) 2.41 900 1.29(1) 3.14(2) 2.22 1000 1.38(1) 3.30(2) 2.24 707-----7 y60 t3850 Cg 40 30 20 11 I b.75 0.80 0.85 0.90 0.95 d-spacinglA 1:OO 1:05 Fig.3 Comparison of neutron diffraction profiles at 150, 300, 500, 700 and 900 "C, locations of low-temperature diffuse peaks are arrowed 3.50 3 00 2.50 2.00 1SO 1.oo LLk 1.40 1.20 1.oo 0.80 060 040 I 0 200 400 600 800 loo0 1200 1400 TIK Fig.4 Isotropic temperature factors (ITF) us. temperature in Kelvin for 0 (a) and Y/Zr (b).Guidelines show fitting to high temperature and low temperature regimes, intercepts with T=O K indicate static contributions to ITFs. improvement in the refinement due to the disappearance of the diffuse scattering. The change in isotropic temperature factor (ITF) of Zr/Y and 0 with temperature is shown in Fig. 4, and listed in Table3. The ITF for Zr/Y and 0 both appear to increase linearly between 150 and 600°C. Above 600°C a change in slope occurs, with a greater increase in ITF with temperature.The temperature at which this deviation is observed also 5.19 5.17 5 rp 5.15 5.13 I 0 200 400 600 800 lo00 T/"C Fig. 5 Unit cell parameter, a, us. temperature corresponds to the temperature at which the activation energy for conduction changes. The observed linear dependence of the temperature factor upon temperature is in accord with a simple Debye-type model, although strictly this model is only applicable to monatomic solids with the atoms centred on their ideal crystallographic sites. An approximation developed by Martin et for YSZ is to split the temperature factor into two components, giving a static term which is independent of temperature, and a temperature-dependent term which varies in accord with the Debye approximation [eqn.(l)]. BkexP =gkstatic + thermal (1) The plots in Fig. 4 show how the observed dependences of the temperature factor can be interpreted using this model. The behaviour of each atomic sublattice shows a low-tempera- ture and a high-temperature region; extrapolated values of the low- and high-temperature static values are presented in Fig. 4. For both atomic sublattices at low temperatures the static contribution is higher; however, the rates of change of the thermal contribution with temperature are lower at low tem- perature. This behaviour is consistent with microdomains containing ordered arrays of distorted subcells at low tempera- tures, with disordering of the microdomains at ca.650 "C. A plot of lattice parameter against temperature also shows a deviation at 600-700 "C (Fig. 5). Consequently, this deviation can be considered as a discontinuity in the thermal expansion coefficient, which is indicative of a second-order transition. Conclusions The change in activation energy of conduction with increasing temperature for 8 mol% yttria-stabilised zirconia was con-firmed as occurring at ca. 650 "C. Using high-temperature neutron diffraction, a modulated background was observed in all diffraction patterns below 650°C. Above 650°C, the diffuse scattering decreased significantly, indicating that short-range ordering only occurs at low temperatures. The deviation in both cation and anion isotropic temperature factors, again between 600 and 700 "C, suggested a larger static contribution to the ITFs at low temperatures, due to the displacement of ions in locally ordered domains.In the disordered regime at higher temperatures, the static contribution to the ITFs was much lower owing to randomisation or effective loss of local distortion. A deviation in thermal expansion coefficient between 600 and 700°C suggested a transition which was second order in nature. We gratefully acknowledge Tioxide Specialties plc for funding a research studentship and the CLRC for the award of neutron diffraction beam time at the Rutherford Appleton ISIS facility. J. Mater. Chem., 1996, 6(5),895-898 897 References 19 R E Carter and W L Roth, in Proc Electromotive Force Measurements in High Temperature Systems, ed C B Alcock, S P S Badwal, J Muter Sci, 1985,20,4593 S P S Badwal, Solid State Ionics, 1992,52,23 20 London Institution of Mining and Metallurgy, 1963, pp 125-144 H Horiuchi, A J Schultz, P C W Leung and J M Williams, D K Honke, Solid State Ionics, 1981,5, 531 J E Baurele and J Hrizo, J Phys Chem Solids, 1969,30, 565 21 Acta Crystallogr Sect B, 1984,40, 367 B Hudson and P T Moseley, J Solid State Chem , 1976,19,383 A S Nowick and D S Park, in Superionic Conductors, ed G Mahan and W Roth, Plenum Press, NY, 1976, pp 395-412 22 N H Andersen, K Claussen, M A Hackett, W Hayes, M T Hutchings, J E MacDonald and R Osborn, in Proc 6th 6 A S Nowick, D Y Wang, D S Park and J Griffith, in Fast Ion Ris0 Int Symp on Metallurgy and Materials Science, ed 7 Transport in Solids, ed P Vashishta, J N Mundy and G K Shenoy, North Holland, Amsterdam, 1979, pp 673-679 J A Kilner and C D Waters, Solid State Ionics, 1982,6,253 F W Poulsen, N Hessel Andersen, K Claussen, S Skaarup and 0 Toft Sarrensen, London Institution of Mining and Metallurgy, 1985, pp 279-284 8 J A Kilner and B C H Steele, in Non-stoichiometric Oxides, ed 23 R Osborn, N H Anderson, K Claussen, M A Hackett, 0 T Sarrensen, Academic Press, NY, 1981, pp 233-269 W Hayes, M T Hutchings and J E MacDonald, Muter Sci 9 C R A Catlow, Solid State Ionics, 1984, 12,67 Forum, 1985,7,55 10 I R Gibson and J T S Irvine, Solid State Ionzcs, submitted 24 S Hull, T W D Farley, M A Hackett, W Hayes, R Osborn, 11 B C H Steele, in High Conductivity Solid Ionic Conductors, ed N H Andersen, K Claussen, M T Hutchings and W G Stirling, T Takahashi, World Scientific, Singapore, 1989, pp 402-446 Solid State Ionics, 1988,28-30,488 12 C R A Catlow, A V Chadwick, C N Greaves and L M 25 R B Neder, F Frey and H Schulz, Acta Crystallogr, Sect A, Moroney, J Am Ceram Soc , 1986,69,272 1990,46,799 13 P Li, I-W Chen and J E Penner-Hahn, Phys Rev B, 1993, 26 D N Argyriou, J Appl Crystallogr , 1994,27, 155 48,10074 27 U Martin, H Boysen and F Frey, Acta Crystallogr , Sect B, 1993, 14 P Li, I -W Chen and J E Penner-Hahn, J Am Ceram SOC ,1994, 49,403 77,118 28 T Proffen, R B Neder, F Frey, D A Keen and C M E Zeyen, 15 P Li, I -W Chen and J E Penner-Hahn, J Am Ceram Soc ,1994, Acta Crystallogr , Sect B, 1993,49,605 77,1289 29 R I Smith and S Hull, Report RAL-94-115, Rutherford Appleton 16 17 D Steele and B E F Fender, J Phys C Solid State Phys, 1974, 7, 1 M Morinaga, J B Cohen and J Faber Jr, Acta Crystallogr, Sect A, 1979,35,789 30 31 Laboratory, 1994 W I F David, R M Ibberson and J C Mathewman, Report RAL-92-032, Rutherford Appleton Laboratory, 1992 V F Sears, Neutron News, 1992,3,26 18 J B Cohen, M Moringa and J Faber Jr, Solid State Ionics, 1981, 3/4,61 Paper 5/05805H, Received 1st September, 1995 898 J.Muter Chem., 1996, 6(5),895-898

ISSN:0959-9428    

DOI:10.1039/JM9960600895

出版商:RSC    

年代:1996

数据来源: RSC

摘要:

MATERIALS CHEMISTRY COMMUNICATION Investigation of the stability of the hexagonal-cubic boron nitride prism interface Jorg Widany," Thomas Frauenheim" and Walter R. L. Lambrechtb "Technische Universitat, Institut fur Theoretische Physik III, D-09107 Chemnitz, Germany bDepartment of Physics, Case Western Reserve University, Cleveland, OH 44106-7079, USA The stability of the hexagonal-cubic boron nitride prism plane interface has been investigated using a density-functional based tight-binding approach. The results for the interface energy and structure support recent experimental findings that hexagonal layered BN may play an important role in the homoepitaxial nucleation of cubic BN films. Boron nitride (BN) is the 111-V analogue of carbon and forms similar sp2- and sp3-bonded crystalline modifications: a layered hexagonal structure, h-BN,' with an AA' stacking sequence is the equilibrium stable low-pressure phase, while the cubic c- BN2 and hexagonal wurtzite w-BN3 structures are metastable.Cubic boron nitride has the most interesting properties, e.g. extreme hardness, high thermal conductivity and a wide bandgap. Kester and Messier4 have made much progress towards solving frequently reported problems (high residual compress- ive stress, cracking, loss of adhesion) in depositing thin films of high-purity c-BN at low temperature and pressure. They showed that the single most important parameter for c-BN formation is control of the momentum transfer of the incident ions onto the growing film.Analysing the deposited films by FTIR spectroscopy and cross-sectional high-resolution trans- mission electron microscopy (HRTEM), Kester ft ~1.~ reported homogeneous layered structures of 20 A thick amorphous a-BN followed by a layer of h-BN (20-60 A) and a top layer of polycrystalline c-BN at the silicon (100) substrate. The c axis of h-BN was oriented parallel to the Si( 100) surface and randomly around the surface normal. Medlin et ul. reported a layer of highly oriented turbostratic BN between the silicon substrate and a region dominated by c-BN crystal- lite~.~This suggests strongly that c-BN films can nucleate through the formation of a hexagonal buffer layer. The related epitaxial growth process and the evolving interface between h-and c-BN was proposed by McKenzie et to be the key for maximal stress release and the enabling factor fo; the growth of homogeneous films to thicknesses of ca.1000 A. Here we present a study of these interfaces using a density- functional based tight-binding (DFTB) molecular dynamics (MD) method which may assist further development of methods for growing highly oriented epitaxial films. The interatomic potentials and tight-binding Hamiltonian used have been shown recently to provide accurate results for a wide class of different-scale BN systems, including interactions with hydrogen;" they thus appear to be highly transferable. Similar to the graphite/diamond interface found in diamond nucleation on the prism planes of graphite" and studied theoretically using empirical potential modelling12 and the DFTB ~cherne,'~ there is a preferential epitaxial and simple geometric relationship between the two bulk crystalline BN phases in the case of h-BN(0001) on c-BN(111) as shown in Plate 1.Three layers of c-BN almost perfectly match two layers of h-BN. Furtbermore, the (111) projection of the BN bond length (1.485 A) is only 3% larger than the corresponding distance in the hexagonal phase (1.440 A) and the density ratio of the hexagonal (2.27gcmP3) to the cubic structure (3.48 g ~m-~) is close to 2: 3. Analogous to the discussion of the graphite/diamond interface by Lambrecht et a1.,12 this allows for low strain, well bonded interface. However, in contrast to graphite, which has AB stacking, h-BN exhibits AA' stacking such that the B atoms in layer A sit directly below the N atoms in layer A'.Moreover, owing to the heteronuclear nature of BN, two distinct interface models must be analysed in which the B and N atoms are interchanged. Conjugate gradient energy minimization and constant tem- perature molecular dynamics at 300K have been used to investigate the stability of both AA' and AB interface structures. The calculations were performed on thin slab models with 276 atoms and two-dimensional periodicity with the interface in the centre. To mimic the connection to the infinite bulk material in the third direction, the outermost layers of the cubic modifications parallel to the interface were held fixed during the simulation. In contrast, the outermost atoms of the hexagonal phase were only held fixed in the direction normal to the hexagonal planes, to allow free intralayer sliding.Both hexagonal and cubic phases in the non-periodic extension (normal to the interface) are saturated by hydrogen in order to remove the dangling orbitals. In cases where the H atoms come very close on the cubic side, we constructed pseudo- 9~ hydrogens, which do not interact with each other, but still serve the saturation effect on B and N. Because of the relatively small lattice mismatch of the cubic and hexagonal phases, discussed above, the bonds within the h-BN layers normal to the plane of the paper are slightly widened to 1.4.8 A, while the bonds normal to the interface adjust to 1.46 A under relaxation with the constraint of fixed interlayer spacing. As the most stable configurations the interfaces with the hexagonal (AA') layer stacking have been established (Plate 1).As seen from the bond distances, the model shown in Plate l(a) is expected to be the more stable, since sigqificant distortions only occupy a very small range (ca. 3.5 A) normal to the interface, which means that after one lattice translation in the xdirection the residual strain on the crystal becomes negligible, supporting high stability and strong adhesion. The configur- ation in Plate l(b) is found to be less stable :wing to the increased strain in the BN lpnds (1.71, 1.85 A) and bond angles at the interface.The h-BN interlayer space can be considered as a small free surface for the part of the c-BN prism plane. As shown in Plate l(a), we observe the usual trend that N anions move outwards, while the B cations move inwards, causing an graphitisation effect of the part of the c-BN prism plane between the h-BN layers. This is also supported by calculating the local electronic densities of states (LDOS) at these atoms. In Fig. 1 we present the results, in comparison with the LDOS behaviour of N and B atoms in bulk crystalline J. Muter. Chem., 1996, 6(5),899-901 899 Plate 1 Hexagonal-cubic-BN interfaces; side and top view. Bond lengths in the interface region are given in A. (a) Boron (blue atoms) termination; (b)nitrogen (orange atoms) termination c and h-BN.Note that for both N and B atoms of the prism plane c-BN interface region, pronounced n-n* states develop around the Fermi energy, similar to the situation for n-bonded atom pairs within hexagonal BN layers. The pronounced peak in the interface LDOS, however, also indicates significant unoccupied dangling bond or non-bonded p-orbital character which may be significant for the chemical reactivity of this interface. A similar feature has been shown to be present in the case of carbon.I4 For a quantitative analysis of the bonding we calculated the interface formation energies (n) per unit-cell area (A) as the additional energy required to generate the interface from 900 J. Muter. Chem., 1996, 6(5),899-901 20.02s*01 15.0 10.0 5.0 0.0 -1.0 -0.5 0.0 0.5 E lev Fig.1 Local electronic densities of states (LDOS): (a) N anion and (b)B cation in the c-BN prism plane part of the interface [solid lines, (i)] compared with the LDOS of corresponding atoms in bulk c-BN [dotted lines, (ii)] and h-BN [dashed lines, (iii)] corresponding equilibrium h-and c-BN supercells [eqn. (l)]: where Ehic,E, and Eh are the energies of supercells with the same number of atoms with half-hexagonal and half-cubic, all cubic and all hexagonal structures, respectively. To allow for comparison of energies, the saturation of the outer faces with hydrogen was chosen to be identical to the interface structures on both outer faces. The values obtained, 0.10 and 0.26 J m-2 for model (a)and (b),respectively, confirm the high stability of both configurations, with (a)being more stable.Note that the energies are quite low relative to the values reported for the graphite/diamond interfaces (ca. 5 J m-2),13 which suggests that the adhesion of the h-BN/c-BN interface is even stronger than that for graphite/diamond. However, is the procedure of allowing the h-BN layers to move freely in the c plane appropriate? As can be seen in the top views of Plate 1 (a) and (b),the layers have moved slightly away from the perfect AA’ stacking, to enable better bonding at the interface. This assumption of free sliding would be appropriate for the case of a finite model consisting of a few layers of free-standing h-BN.In the growth situation, this may be approximately the case if h-BN is rather loosely connected to the substrate. Experiments637 have revealed that an amorph- ous sublayer is often found between substrate and h-BN, which may provide an easily stretchable ‘buffer layer’ for each BN plane. However, if the h-BN is infinitely strongly bonded to the substrate, the layers cannot slide but must be stretched, which would lead to a strain energy proportional to the thickness of the h-BN layer and determined by the high in- plane elastic constants of h-BN. The variation between these two extremes is controlled by the adhesive force constant between h-BN and the substrate. For a thick single-crystalline h-BN substrate, we must also consider the energy associated with weak van der Waals interlayer bonding.Although the energy differences between different stackings of the h-BN basal planes are small per unit area, the accumulated energy for a thick h-BN layer may not be negligible. We ignore this possibility here because we are mainly interested in the case of a thin h-BN buffer layer. The strain effects on the interface formation energy of the discussed model structures is simulated by subsequently fixing the AA' stacking at the h-BN prism- plane surface at the average x,y coordinates of the terminating B and N atoms. After a final conjugate gradient relaxation the formation energy of interface (a) increases to 0.18 J mP2, but the effect on interface (b) is only marginal. Thus, there is at most 0.02 J m-2 strain energy per h-BN layer.We conclude that for the thicknesses of h-BN occurring in realistic growth situations (which is of the same order of magnitude as in our simulation) the accumulated strain energy will not change our conclusion of strong bonding between h-BN and c-BN. Graphite-like stacking (AB) and the related interface con- figurations, shown in Plate 2, give rise to much larger distor- tions of the interface region, in particular on the c-BN side. The development of five-membered rings in this case causes highly strained bond apgles and bonds between atoms of the same type (B-B 1.75 A). Calculated interface formation ener- gies are greater than the values for the AA' stacking models. The hexagonal-cubic BN prism plane interface has been shown to be at least mechanically stable and thermally stable up to 300 K, concluded from the strong interfacial bonding (or low interface formation energy) and the fact that no significant structural strain is found in the adjacent h-BN or c-BN regions.As in the case of carbon, c-BN nucleation may be favourably mediated through the formation of hexagonal buffer layers. The presence of interfacial empty dangling bond states may lead to chemical reactivity of the interface which may be involved in the conversion of h-BN into c-BN along the interface. Further extensions of the present models to the case of heteroepitaxial c-BN on diamond (and vice versa) growth are of interest. The planar graphitic layers may be in competition with direct heteroepitaxy maintaining the tetrahedral bonding throughout the interface.Nevertheless, some observations of graphitic nucleation layers have been reported for the growth of diamond on c-BNI5 and h-BN has been reported to occur also in c-BN growth on diamond.16 The role of hexagonal layers in c-BNldiamond heteroepitaxy is presently under investigation. We gratefully acknowledge support from the Deutsche Forschungsgemeinschaft and the trinational D-A-CH corpor- Plate 2 Graphite-like hexagonal-cubic BN interface; side and top view ation. W.L. acknowledges support by NSF (MRG grant DMR- 91 -21479). References 1 R. S. Pease, Acta Crystallogr., 1952,5, 536. 2 J. R. H. Wentorf, J. Chem. Phys., 1957,26,956.3 F. P. Bundy and R. H. Wentorf, J. Chem. Phys., 1963,38,1144. 4 D. J. Kester and R. Messier, J. Appl. Phys., 1992,72, 504. 5 D. J. Kester, K. S. Ailey and R. F. Davis, J. Mater. Res., 1993, 8, 1213. 6 D. J. Kester, K. S. Ailey and R. F. Davis, Diamond Relut. Mater., 1994,3, 332. 7 D. L. Medlin, T. A. Friedmann, P. B. Mirkarimi, P. Rez, M. J. Mills and K. F. McCarty, J. Appl. Phys., 1994,76,295. 8 D. R. McKenzie, J. Vac. Sci. Technol. B, 1993, 11, 1928. 9 D. R. McKenzie, W. D. McFall, W. G. Sainty, C. A. Davis and R. E. Collins, Diamond Relat. Muter., 1993,2,970. 10 J. Widany, Th. Frauenheim, Th. Kohler, G. Jungnickel,M. Sternberg, D. Porezag and G. Seifert, Phys. Rev. B, 1996, in press. 11 Z. Li, L. Wang, T. Suzuki, A. Argoitia, P. Pirouz and J. C. Angus, J. Appl. Phys., 1993,73, 711. 12 W. R. L. Lambrecht, Ch. H. Lee, B. Segall, J. C. Angus, Z. Li and M. Sunkara, Nature (London), 1993,364,607. 13 G. Jungnickel, D. Porezag, Th. Frauenheim, W. R. L. Lambrecht, B. Segall and J. C. Angus, Phys. Status Solidi, 1996, in press. 14 A. De Vita, G. Galli, A. Canning and R. Car, Nature (London), 1996, submitted. 15 S. Koizumi and T. Inuzuka, Jpn. J. Appl. Phys., 1993,32,3920. 16 R. F. Davis, K. S. Ailey, R. S. Kern, D. J. Kester, Z. Sitar, L. Smith, S. Tanaka and C. Wang, Mater. Res. SOC.Symp. Proc., 1994, 339, 351. Communication 5/08356G; Received 28th December, 1995 J. Muter. Chem., 1996, 6(5),899-901 901

ISSN:0959-9428    

DOI:10.1039/JM9960600899

出版商:RSC    

年代:1996

数据来源: RSC

摘要:

Electrochromism-Fundamentals and Applications. P. M. S. Monk, R. J. Mortimer and D. R. Rosseinsky. VCH Verlagsgesellschaft, 1995. Pp. xxvi +21 6. Price DM1 68.00. ISBN 3-527-29063-X. Electrochromism is a rapidly growing technology and under- scores the increasingly interdisciplinary nature of inorganic and organic materials. This book brings together three physical chemists who deliver an excellent up-to-date account of the subject. Part I of the review provides the reader with some funda- mental principles and importantly defines the terminology associated with this field of science such as colouration efficiency. Basic electrochemical theory is included in this section and covers such topics as cyclic voltammetry, charge transport, electron transport kinetics and semiconductor elec- trodes.Part I concludes with a section on the construction of electrochromic displays (ECDs), with the combinations of all components in solution, solution-solid and all-solid systems. Part I1 introduces the major thrust of the book where both inorganic and organic systems are discussed for use in elec- trochromic applications. The inorganic metal oxides and mixed metal oxides are covered in terms of structure, preparation, diffusion characteristics and spectroscopic and optical effects. An extended discussion on the operation of W03 ECDs is also included, together with an account of new rare-earth-metal lutetium bis( phthalocyanine), Prussian blue and related systems for potential electrochromes.A wide range of organic systems is covered. Much of the focus is related to device applications employing bipyridilium electrochromes and elec- troactive conducting polymers such as polyaniline, polypyrrole and polythiophene. Miscellaneous organic materials such as pyrazolines are also mentioned. Part 111 presents recent advances in current research. This section comprises polyelectrochromes and photoelectrochro- mism, including a discussion on electrochromic printing. This book is a forum for scientific issues central to electro- chromic display technology. It also provides a springboard for anyone wishing to grasp the basics of the field and is strongly recommended for specialist practioners in display technology. Raoul Cervini Received 10th January, 1996 ~ ~ ~ ~~~~~~~ One-Dimensional Metals.Siegmar Roth. VCH, Weinheim, 1995. Pp. xii +247. Price DM 148.00. ISBN 3-527-26875-8. One-dimensional metals continue to fascinate scientists from a broad cross-section of disciplines, with the quest for high- temperature superconductors acting as a major driving force. This book, written by a leading researcher in the field, provides an introduction to molecular and polymeric systems (mostly organic) with an emphasis on the underlying physical concepts and the solid-state properties of the materials. The book is divided into 10 chapters. Chapter 1 introduces the reader to the concepts of dimen- sionality in the solid state, with reference to compounds such as diamond, graphite, polyacetylene and fullerenes.Densities of states and Fermi surfaces are dealt with in a non-theoretical way. Chapter 2 is an overview of one-dimensional substances, e.g. Bechgaard salts (tetramethyltetraselenafulvalene-based superconductors) and other charge-transfer systems, stacked phthalocyanines and conducting polymers, with passing refer- ence to such compounds as nanotubes and polycarbenes. Chapter 3 deals with one-dimensional solid-state physics in a easy-to-read style for the non-specialist: doping, energy bands and reciprocal space are discussed lucidly. Chapter 4 considers electron-phonon coupling and Peierls transitions in some detail, again in a non-theoretical way. Chapters 5 and 6 concern conducting polymers, and they constitute a major part of the book.A historical perspective is presented and topics discussed include: charge defects, methods of measuring con- ductivity of polymer films, and mechanisms of charge transport in these systems. Chapter 7 is devoted to superconductivity. BCS theory is discussed and organic and metal oxide supercon- ductors are compared. This chapter would have benefited considerably from the inclusion of X-ray crystal structures of salts of BEDT-TTF to illustrate different packing modes. Charge-density waves are discussed from a more theoretical viewpoint in chapter 8. Chapter 9 is entitled Molecular Electronics, and it broadens the subject matter of the book to include ideas of soliton switching and molecular rectification.The final chapter presents an overview of the emerging appli- cations of one-dimensional metals in such fields as polymer batteries, solar cells, electrochemical sensors and light-emit- ting devices. An unusual and idiosyncratic feature of this book, which may not appeal to all readers, is the inclusion of several cartoon drawings. For example, the search for solitons in conducting polymers is illustrated with a half-page picture of butterfly hunters! Elsewhere there are drawings of camels in the Sahara, and a car trying to overtake a lorry on a winding road, the latter to illustrate an important aspect of one-dimensionality, namely, that obstacles cannot be circumvented. There are some errors in the chemical formulae, with double bonds missing (e.g. NMP on p. 33 and ATCNQ on p. 34). It is incorrect to classify pyridinium and related cations as donors (p. 33). Rather, it is their iodide counter-ions which act as the electron donor species in the reduction of TCNQ to the radical anion. There is a list of ca. 40-50 references at the end of each chapter, including some 1994 publications, and a good index. This book is interesting and informative and it is rec-ommended as a very readable introduction to the solid-state properties of one-dimensional systems. It is unfortunate that it is priced so high that very few graduate students will be able to afford it. Martin R. Bryce Received 9th February, 1996

ISSN:0959-9428    

DOI:10.1039/JM9960600903

出版商:RSC    

年代:1996

数据来源: RSC