摘要:

J O U R N A L O F C H E M I S T R Y Materials Solvent-free synthesis of binary inorganic oxides John N. Hay*† and Hema M. Raval Department of Chemistry, School of Physical Sciences, University of Surrey, Guildford, Surrey, UK GU2 5XH The non-hydrolytic sol–gel route has been used in the solvent-free synthesis of binary inorganic oxides based on silicon, aluminium and titanium. Where necessary, iron(III ) chloride was used as a catalyst.Clear evidence for the formation of true, amorphous binary systems was obtained only in the case of aluminosilicates. A crystalline aluminosilicate was obtained after calcination at 1000 °C. A titanium–silicon binary system gave only crystalline rutile and anatase following calcination, with no evidence for either crystalline silica or a binary oxide.Low temperature routes to inorganic oxides have attracted Mechanistically, the non-hydrolytic route should favour the considerable attention in recent years owing to their reduced formation of homogeneous binary oxides from diVerent ‘metal’ energy demand compared to traditional high temperature glass precursors since the reaction of M(OR)m with M¾Xn should forming processes and the potential opened up for combining form unsymmetrical MMOMM¾ linkages in the absence of the oxides with thermally labile organic compounds.The reversibility and ligand exchange reactions. Another way to oxides and their hybrids have a plethora of uses, including make non-hydrolytic sol–gels is to form the alkoxy groups in catalysis and catalyst supports, ceramics, sensor applications situ, by reacting the metal halide with an alcohol10 or other and active glasses.The most important synthetic route is the organic oxo-compounds, e.g. ethers, aldehydes and ketones.11 hydrolytic sol–gel method which allows inorganic oxides The non-hydrolytic route has been used to make non-metal and/or their immediate precursors to be produced from simple oxides, metal oxides, transition metal oxides and binary oxides, alkoxides or chelates via low temperature hydrolysis and with the outstanding contribution coming from the group of condensation reactions.This route has fairly widespread appli- Corriu.12 Non-metal oxides, e.g. silica gels, have been studied cability and has been the subject of a number of recent books most extensively.Stoichiometric reactions between silicon and reviews.1–7 For certain applications, however, this tetrahalides and various oxygen donors including alcohols, approach has a number of disadvantages, including the need benzaldehyde and tetrabenzyloxysilane, have been examined.13 to add water for the hydrolysis, the formation of condensation Overall, alcohols appear to be more eYcient oxygen donors by-products (e.g. water, alcohols) and the initial formation of than alkoxides, ethers and aldehydes. In the case of metal a solvent-swollen gel which undergoes substantial shrinkage oxides, alumina gels have been studied by Corriu et al., who on drying. In addition, for the production of uniform binary found solvents to be necessary.9,14 The main oxygen donors oxides, the large diVerences in hydrolysis rates of diVerent to aluminium halides were aluminium isopropoxide and diisoprecursors can lead to undesirable inhomogeneities in the propyl ether.The alumina obtained was amorphous up to product. This is true of binary oxides derived from the trans- 750 °C compared to those obtained via the hydrolytic route ition metal alkoxides, especially those of d0 transition metals, which usually crystallise below 500 °C.As the precursors to e.g. titanium and zirconium, which are widely used precursors the transition metal oxides are mostly liquids, these reactions for glasses and ceramics. The lower electronegativities of the often proceed without the need for solvents. Some of the transition metals compared to silicon make their alkoxides oxygen donors used with titanium tetrachloride include tetramore reactive towards nucleophilic reactions such as hydrolysis hydrofuran (THF), diisopropyl ether, dimethoxyethane and condensation. (DME), and titanium isopropoxide.9,15 Gel-times were longer One possible route to avoiding this complication is the non- than under hydrolytic conditions, which may be an advantage hydrolytic sol–gel process, which is based on an observation for titanium since this increases the potential for more control made by Gerrard et al.8 Whilst studying the interactions of the reaction.Other metal halides investigated include between organosilanes and alcohols, Gerrard et al. found silica niobium, molybdenum and tungsten chlorides, but these lead to be formed.Since then, this route has been adapted to to precipitates rather than gels. More recently, Guenther et al. involve the direct condensation of metal halides with metal reacted zirconium chloride with diisopropyl ether.16 Further alkoxides, e.g. silicon, titanium and zirconium (Scheme 1).9 work by our group on single inorganic oxides prepared by the The mechanism for this reaction involves the coordination of non-hydrolytic sol–gel route is reported separately.17 the oxygen donor, e.g.alkoxide, to the metal centre of the The rate of condensation in the non-hydrolytic process is metal halide. This usually evolves via three possible transition highly dependent on the nature of the oxygen donor. The states or intermediates and results in the formation of oxogeneral tendency of the non-hydrolytic process compared to bridges.The non-hydrolytic route involves the cleavage of the the hydrolytic route is to delay crystallisation18 and this carbon–oxygen bond instead of the metal–oxygen bond. The behaviour may find application in the preparation of heteromain limitation of this process is the insolubility of some metal geneous catalysts.DiVerent metals may also be used to form chlorides in non-aqueous solvents. It also has a general tendmonolithic binary mixed metal oxides via a cross-condensation ency to delay crystallisation of the metal oxide. This process reaction (Scheme 2).18–21 appears to be simpler to control than the conventional hydroly- The hydrolytic route to mixed metal oxides often leads to sis–condensation process found in the hydrolytic route.diYculties in controlling stoichiometry and homogeneity. Homogeneity depends on the rate of homocondensation (i.e. formation of MMOMM and M¾MOMM¾) versus the rate of † E-mail: j.hay@surrey.ac.uk J. Mater. Chem., 1998, 8(5), 1233–1239 1233M OR + M X M X M OR M O M R + + X– M O M X – + R+ M X M O R M O R X M + (1) (2) (3) Scheme 1 involved a reflux system, immersed in a temperature controlled mMXn + nM¢(OR)m MmM¢ nOnm + nmRX oil bath, connected to a series of nitrogen drying vessels.A Scheme 2 side arm was included above the reaction tube to allow the safe addition of air-sensitive liquid reactants from a syringe via a septum. heterocondensation (i.e. formation of MMOMM¾). One way to overcome this in the hydrolytic route is to subject the less reactive alkoxide precursor to partial hydrolysis before adding the more reactive one, or to use heterometallic alkoxides.In Instrumentation principle, the non-hydrolytic process potentially provides a A Shimadzu thermogravimetric analyser-50 was used to deterbetter route to mixed metal oxides than the conventional mine sample mass loss over a temperature range of 25–490 °C. hydrolytic sol–gel approach or mixed powder methods because A heating rate of 20 K min-1 was applied under a nitrogen of improved homogeneity during the reaction.22 This is in part atmosphere (flow rate, 50 ml min-1).For mass losses up to a consequence of the reaction mechanism, which is completely higher temperatures, a Perkin-Elmer TGA-7 interfaced with a diVerent from that of the classical sol–gel process.As in the Perkin-Elmer 7700 computer was used to measure sample hydrolytic route, the reactants react at low temperature in a mass loss over a temperature range of 40–900 °C. A heating homogeneous solution or liquid, thus avoiding phase separarate of 10 K min-1 was applied under a nitrogen atmosphere tion, crystallisation and chemical decomposition.DiVerent with a flow rate of 50 ml min-1. combinations of silicon, aluminium and titanium chloride have All calcinations were carried out in a Lenton UAF 16/21 been reacted with the corresponding isopropoxides.9,18,20 In furnace. An SMC 24127 calibrated digital K-type thermometer all cases gels were formed. Compositions can be controlled via was used to control temperatures to ±5 °C and the sample the composition of starting solutions, since the results show was contained in a mullite tube during the heating run.A that calcined gels are close in composition to initial solutions Siemens D500 X-ray diVractometer was used to determine and this agrees with the high yields achieved. Iwasaki et al. sample morphology. Samples were ground into fine powders have recently described the condensation of silicon tetraacetate before being scanned with Cu-Ka radiation at 40 kV, 40 mA with titanium acetylacetonatotriisopropoxide in THF.23 More from 15–70° 2h, with a step size of 0.02° and a count time per recently, it has been demonstrated that homogeneous zirstep of 10 seconds.Surface morphology and evidence of any conium titanate gels can be prepared by the non-hydrolytic microporosity in the samples were evaluated using a sol–gel route without the intermediate formation of zirconia CAMSCAM S4 field emission scanning electron microscope or titania.21 In this paper we describe the synthesis and with an accelerating voltage of 5 kV for the secondary electron characterisation of further binary inorganic oxides via the imaging (SEI) and 20 kV for the energy dispersive X-ray solvent-free non-hydrolytic sol–gel route.To date, there are (EDX) analyses. Prior to mounting, the specimens were carbon relatively few reports of the solvent-free synthesis of binary coated to prevent charging under the electron beam. EDX oxides by the sol–gel route, which is likely to be a prerequisite spectra were acquired from a number of areas to determine for the successful commercial development of products prethe elemental composition of the sample.pared by this method. Infrared (IR) spectra were recorded on a Perkin-Elmer 1750 FT-IR spectrophotometer interfaced with a Perkin-Elmer Experimental computer. Samples were ground and prepared for diVuse reflectance infrared Fourier transform spectroscopy (DRIFTS).Materials Abbreviations used to describe peaks are as follows: vs=very Silicon tetrachloride (Aldrich), tetraethylorthosilicate, TEOS strong; s=strong; m=medium; w=weak; vw=very weak; (Lancaster), aluminium trichloride (Aldrich), aluminium iso- b=broad; sp=sharp and sh=shoulder. propoxide (Aldrich), titanium tetrachloride (Aldrich), titanium Solid-state 29Si nuclear magnetic resonance (NMR) specisopropoxide (Lancaster), ethanol (Hayman Ltd.) and iron(III ) troscopy was undertaken at the University of Durham, at chloride (Aldrich) were all used as received.Propan-2-ol ambient temperature on a Varian UNITYplus spectrometer. (Fisons), diethyl ether (Fisons) and carbon tetrachloride (BDH) The spectrum was recorded against an external TMS standard were dried using 4 A ° molecular sieves before use.with magic angle spinning (MAS) at a spinning rate of 4300 Hz and an angle of 54.7°. The cross polarised (CP) spectrum was Method obtained as a single contact experiment with a contact time of 3 ms and a relaxation delay of 2.0 s (700 repetitions). An Many of the materials used to make non-hydrolytic sol–gels acquisition time of 9.6 ms was used.Silicon sites were labelled are very reactive and have strict handling conditions. As a with the conventional Qn notation where Q refers to tetrafunc- consequence, care should be taken when embarking on new tional SiO4 units and n to the number of bridging oxygen and untried experiments, particularly in the absence of a diluting solvent.The basic set-up used in these experiments atoms surrounding the central silicon atom. 1234 J. Mater. Chem., 1998, 8(5), 1233–1239Fig. 2 FTIR spectrum of aluminosilicate synthesised from aluminium Fig. 1 FTIR spectrum of aluminosilicate synthesised from silicon tetrachloride and aluminium isopropoxide (equimolar) trichloride and TEOS (453 molar ratio) Synthesis of binary oxides Preparation of aluminosilicate from silicon tetrachloride and aluminium isopropoxide (equimolar) with the aid of a solvent.An equimolar amount of silicon tetrachloride (4.00 g, 2.70 cm3, 0.024 mol) was added to aluminium isopropoxide (4.81 g, 0.024 mol) dissolved in a 451 (by mass) mixture of carbon tetrachloride and diethyl ether (4 cm3). A white solid was formed within approximately 40 minutes at 90 °C.This was washed in propan-2-ol and dried for 3 days at 90 °C followed by a further 2 days at 150 °C (2.13 g, 81% yield based on SiAlO3.5. Note: since the non-condensed sites of the network cannot be evaluated, yields are only included as information). The thermogravimetric analysis (TG) showed a mass loss of 14%. The IR spectrum (Fig. 1) showed absorptions at nmax#3400 cm-1 (m, b), 1630 cm-1 (m), 1200 cm-1 (vs, sp), 1150 cm-1 (sh) and 450 cm-1 (vs, sp). Preparation of aluminosilicate from aluminium trichloride and TEOS (equimolar) with the aid of a catalyst. Equimolar amounts of aluminium trichloride (2.06 g, 0.015 mol) and TEOS (3.12 g, 3.35 cm3, 0.015 mol) were reacted with iron(III) chloride (0.070 g, ca. 1% by mass).A solid product was obtained within 3.5 hours at 110 °C. The solid was washed with ethanol and subsequently dried at 90 °C for 3 days and at 150 °C for a further 2 days (1.50 g, 84% yield based on AlSiO4). The resulting monolith was dark brown in colour probably due to trapped catalyst. The TG showed 9% mass loss and the IR spectrum showed absorptions at nmax#3600 cm-1 (b), 1600 cm-1 (vs, sp), 1250 cm-1 (m, sh), 1000 cm-1 (s, b), 950 cm-1 (sh) and 450 cm-1 (m).Fig. 3 XRD spectrum of aluminosilicate prior to calcination (top) and after calcination (below) at 1000 °C for 18 hours accompanied by the Preparation of aluminosilicate from aluminium trichloride peak positions observed for Al6Si2O13 and TEOS (stoichiometric) with the aid of a catalyst. Aluminium trichloride (3.06 g, 0.023 mol) and TEOS (3.52 g, 3.76 cm3, 0.017 mol) were mixed together in the presence of iron(III) be highly porous, with the pores being approximately 200 nm in diameter.EDX analysis of the bulk material (Fig. 4) showed chloride (0.068 g, 1.03% by mass). A clear, dark brown monolith was formed after approximately 10 minutes at 90 °C. The the elements aluminium, silicon and oxygen.Iron was also present in some loose material covering the surface of the product was washed in diethyl ether ( left overnight), filtered and dried at 150 °C for 3 days (2.08 g, 94% yield based on sample. The IR spectrum (Fig. 5) showed absorptions at nmax#3600 cm-1 (vw), 2000 cm-1 (vw), 1880 cm-1 (w), Si3Al4O12). TG showed a 22% mass loss and the IR spectrum (Fig. 2) showed absorptions at nmax#3400 cm-1 (m, b), 1620 cm-1 (w), 1330 cm-1 (vs, sp), 1240 cm-1 (m), 1030 cm-1 (m), 930 cm-1 (m), 850 cm-1 (s), ~640 cm-1 (s), 529 cm-1 (s) 1611 cm-1 (m), 1250 cm-1 (s, sp), 1160 cm-1 (sh), 950 cm-1 (m), 800 cm-1 (m) and 500 cm-1 (m). and 420 cm-1 (m). The dark brown and black coloured product of this reaction was calcined in air at 1000 °C for 18 hours, resulting in Preparation of titanium–silicon binary oxide from silicon tetrachloride and titanium isopropoxide (equimolar).18–21 approximately 30% mass loss.The resulting specimen was heterogeneously coloured black, brown and oV-white. The X- Equimolar amounts of silicon tetrachloride (2.00 g, 1.35 cm3, 0.012 mol) and titanium isopropoxide (3.35 g, 3.5 cm3, ray diVraction (XRD) pattern of the oxide prior to calcination and that after calcination are shown in Fig. 3. The latter shows 0.012 mol), were reacted at 55 °C. A fine white precipitate was seen within 1 minute of stirring at room temperature and this a very good fit with an orthorhombic mullite structure, Al6Si2O13, with no crystalline peaks left unmatched. The disappeared at approximately 45 °C.A clear, pale brown, crazed solid resulted in less than 24 hours and was washed in scanning electron micrograph (SEM) showed this material to J. Mater. Chem., 1998, 8(5), 1233–1239 1235Fig. 7 XRD spectrum of titanium–silicon binary oxide calcined for 18 hours at 1000 °C accompanied by the peak positions observed for rutile and anatase Fig. 4 EDX spectrum of aluminosilicate calcined for 18 hours at 1000 °C (bulk) Fig. 5 FTIR spectrum of aluminosilicate calcined for 18 hours at 1000 °C Fig. 8 SEM of titanium–silicon binary oxide calcined for 18 hours at 1000 °C Fig. 6 29Si NMR spectrum of titanium–silicon binary oxide prior to calcination propan-2-ol and dried at 85 °C for 2 hours producing a creamy, shell-like product (1.70 g, 101% yield based on TiSiO4).TG showed a 20% mass loss. When subjected to further drying at Fig. 9 EDX spectrum of titanium–silicon binary oxide calcined for 18 hours at 1000 °C (bulk) 140 °C for 24 hours, a charcoal black solid was formed (1.36 g, 81% yield). The IR spectrum shows absorptions at nmax#3500 cm-1 (s, b), 1630 cm-1 (m), ~1100 cm-1 (s, b) and 600 cm-1 (s, b). The 29Si NMR spectrum is shown in Fig. 6. tetragonal rutile and anatase. There seemed to be more rutile than anatase present, but no silica phase was detected. A The black material was calcined in air at 1000 °C for 18 hours to leave a white solid (#28% mass loss). A thin black typical SEM is shown in Fig. 8. The EDX spectrum (Fig. 9) showed the bulk of the specimen to have a high titanium coating covered the surface of a few particles.Qualitative XRD results (Fig. 7) show a good fit with two titania phases; content in addition to the silicon and oxygen signals. 1236 J. Mater. Chem., 1998, 8(5), 1233–1239Preparation of titanium–silicon binary oxide from titanium tetrachloride and TEOS (equimolar). An equimolar amount of 4 AlCl3 + 3 Si(OC2H5)4 Al4Si3O12 + 12 C2H5Cl FeCl3 titanium tetrachloride (4.00 g, 2.32 cm3, 0.021 mol) was reacted Scheme 3 with TEOS (4.39 g, 4.70 cm3, 0.021 mol) for 5 minutes at 110 °C.The dark brown gel was washed in ethanol, filtered and dried at 150 °C for 4 days (3.66 g, 124% yield based on formation. At a 453 molar ratio of aluminium trichloride to TiSiO4). TG showed a 28% mass loss and the IR spectrum TEOS, a shorter gel-time was achieved at a lower temperature showed absorptions at nmax#3500 cm-1 (m, b), ~3000 cm-1 and the yield was the highest of the aluminosilicate sol–gel (m), 1850 cm-1 (vw, b), 1630 cm-1 (m, sp), 1430 cm-1 (m, sp), systems; however, the TG showed a higher percentage mass 1370 cm-1 (m, sp), 1250 cm-1 (vs, sp), 1160 cm-1 (s), 960 cm-1 loss, suggesting incomplete reaction.IR analysis (Fig. 2) was (m), 810 cm-1 (m) and 525 cm-1 (m).indicative of formation of a binary oxide, which might be expected to have the empirical formula Al4Si3O12 on the basis Preparation of titanium–aluminium binary oxide from alu- of the stoichiometry (Scheme 3). minium trichloride and titanium isopropoxide (453 molar ratio). The XRD pattern of the oxide prior to calcination (Fig. 3) An appropriate stoichiometric amount of aluminium trichlo- shows the sample to be amorphous with a broad peak at the ride (2.00 g, 0.015 mol) was reacted with titanium isopropoxide low angle end of the spectrum.Following calcination of the (3.20 g, 3.35 cm3, 0.011 mol) for 24 hours at 110 °C. The dark oxide at 1000 °C, XRD analysis reveals that the sample has brown monolith was washed in diethyl ether ( left overnight), crystallised, although an amorphous phase is still present.The filtered and dried for 3 days at 150 °C (1.88 g, 113% yield XRD trace shows a very good fit with an orthorhombic mullite based on Al4Ti3O12). TG showed a 26% mass loss. The IR structure, Al6Si2O13, with no peaks left unmatched. The stoichispectrum produced absorptions at nmax#3400 cm-1 (s, b), ometry of this structure is very diVerent to that expected from 1630 cm-1 (vs, sp), 1100 cm-1 (s), ~900 cm-1 (vs, b), 800 cm-1 the reaction stoichiometry and may result from the presence (s) and 600 cm-1 (m, b).of an amorphous silica phase in the calcined structure. The extent to which molecular or network rearrangement has Preparation of titanium–aluminium binary oxide from occurred during the calcination process is impossible to ascertitanium tetrachloride and aluminium isopropoxide (354 molar tain from these results.It should be noted that at 1000 °C, the ratio). Appropriate stoichiometric amounts of titanium tetra- thermodynamics control the crystallisation process of the chloride (4.00 g, 2.32 cm3, 0.021 mol) and aluminium isoprop- mullite. The possibility of nucleation initiated by the mullite oxide (5.74 g, 0.028 mol) were stirred together for tube used in the calcination can not be totally discounted.approximately 6 hours at 110 °C. The dark brown sample was Optimisation of the reaction stoichiometry might maximise washed with propan-2-ol and then dried for 4 days at 150 °C formation of a crystalline mullite after calcination, but this leaving a black shiny material (4.59 g, 147% yield based on was not carried out as part of this study.Longer heating times Ti3Al4O12). TG showed a 35% mass loss. The IR spectrum might also improve the conversion. showed absorptions at nmax#3500 cm-1 (m, b), 3000 cm-1 (m, The SEMs show the calcined product to be highly porous, sh), 1630 cm-1 (m), 1100 cm-1 (s), 1000 cm-1 (vs, sp), 800 cm-1 with pores of approximately 200 nm diameter.The sample (m, sh) and ~450 cm-1 (w). surface was covered with loose material which, using EDX analysis, was found to contain a high iron content (catalyst residue) in addition to the elements aluminium, silicon and Results and Discussion oxygen which were found in the EDX spectrum of the bulk material (Fig. 4). The aluminium signal is much weaker than Aluminium–silicon oxides the silicon signal, in contrast to the Al5Si ratio of 652 found Three experiments were undertaken on the synthesis of alumifor the crystalline phase using XRD analysis. This apparent nosilicates similar to those reported by Corriu et al.19 A solvent discrepancy may be explained by the fact that XRD analysis was employed for the initial reaction to moderate the potenprovides information only on the crystalline phase, while EDX tially vigorous reaction between silicon tetrachloride and aluanalysis averages the analysis of the bulk sample, including minium isopropoxide.A stoichiometric ratio of 151 was used, amorphous material (such as silica). An additional factor may compared to the halide to alkoxide ratio of 154/3 used by be the influence of the surface roughness of the sample on the Corriu et al.Use of a higher reaction temperature in our case consistency of the quantitative analyses obtained by EDX. led to rapid formation of the initial product, compared to a Following calcination, IR analysis showed that the siloxane literature gel-time of two days at 40 °C. The IR spectrum stretching vibration at #1180 cm-1 had disappeared and been (Fig. 1) of the white product confirmed the formation of replaced by a higher frequency band at 1330 cm-1 (Fig. 5), aluminosilicate.24 The presence of a band at #1200 cm-1 may which could result from crystallisation of the aluminosilicate. be indicative of aluminium substitution in the silica network. The peak at #1150 cm-1 appears as a shoulder and can be Titanium–silicon oxides assigned to the elongation vibration of SiMOMAl.Vibrations between 3500–3400 cm-1 and at 1630 cm-1 are due to the For the titanium–silicon binary oxide system prepared from silicon tetrachloride and titanium isopropoxide, the results stretching of hydroxyl groups from either hydrogen bonded water or propan-2-ol and surface silanol molecules.The latter tend to agree with published findings on related systems.19 IR spectroscopy provides limited evidence for the presence of absorption is due to the deformation mode of these hydroxyl groups. The TG indicated an incomplete reaction, probably TiMOMSi species in the initial amorphous product. Solidstate 29Si NMR spectroscopy was used to obtain data on the due to the presence of unreacted material and trapped solvent or by-product.extent of TiMOMSi bonding in the system (Fig. 6). The spectrum shows the expected SiMOMSi signals at -101 ppm No solvent was used for the reaction of aluminium trichloride with TEOS; however, based on studies by Corriu et al.,25 and -109 ppm. These were assigned to partially condensed, Q3 (ROSiO3) species and fully condensed, Q4 (SiO4) species.a catalyst was employed in the reactions to reduce reaction times and residual unreacted functionalities. Two ratios were No signals (expected at #-20 ppm) could be attributed to the presence of TiMOMSi species,26 suggesting the product studied. At a 151 ratio of aluminium halide to TEOS, a good yield was achieved after 3.5 hours at 110 °C, with a relatively was actually a phase separated mixture of silica and titania.The poor stability of TiMOMSi bonds may help to explain small TG mass loss due to residual volatiles, implying good oxide conversion. The IR analysis24 suggested aluminosilicate this result since any first-formed TiMOMSi species will J. Mater. Chem., 1998, 8(5), 1233–1239 1237rearrange on ageing to form SiMOMSi and TiMOMTi species.The XRD pattern (Fig. 7) verified that after calcination at 3 Ti(OPri)4 + 4 AlCl3 Ti3Al4O12 + 12 PriCl 1000 °C, the majority of the sample had crystallised. A good Scheme 4 fit was obtained with two titania phases, tetragonal rutile and anatase with more of the former present. No crystalline silica was detected, although other studies undertaken by us had shown that pure silica formed by the non-hydrolytic sol–gel 3 TiCl4 + 4 Al(OPri)3 Ti3Al4O12 + 12 PriCl Scheme 5 process could by crystallised under similar calcination conditions to form mixtures of quartz and cristobalite.17 Some possible reasons for this are: (a) the ‘silica’ is actually present Titanium–aluminium oxides as a genuine titanium–silicon hybrid oxide which does not crystallise under these conditions, or (b) the phase size of the Two complementary experiments were carried out to investisilica is too small for crystallisation to occur in this system— gate the formation of titanium–aluminium binary oxides.The a recent study has demonstrated that there is a strong size first is the reaction of aluminium trichloride with titanium dependence of crystallisation kinetics in inorganic nanocrys- isopropoxide (Scheme 4) and the second the reverse reaction tals.27 NMR evidence showing the absence of TiMOMSi of titanium tetrachloride with aluminium isopropoxide species in the original amorphous product suggests ration- (Scheme 5).In the first reaction between aluminium trichloride ale (b) may be the most reasonable explanation for the failure and titanium isopropoxide, the TG result confirms that the of the silica to crystallise.high yield obtained was most likely due to trapped solvent or Corriu et al.18–21 calcined a similar titanium–silicon oxide by-product. The IR spectrum was diYcult to interpret. In the at 500 °C for five hours. They also detected no silica via XRD absence of literature data on the assignment of titanium– analysis, only crystallites of anatase.No rutile was found. aluminium binary oxide peaks, all that can be deduced is Elemental analysis of the oxide showed it to be practically general oxide formation. The second reaction appeared to be carbon-free and have a metal content very close to the composi- even more incomplete. The yield achieved was very high and tion of the starting solution. The microchemical analytical the TG mass loss supported this result. The IR spectrum was study by means of an electron-probe analyser, i.e.energy- also diYcult to interpret. All that can be deduced is oxide dispersive spectroscopy, indicated a constant Ti5Si ratio (close formation, with AlMO absorptions24 at 1100 cm-1 and to unity), which is consistent with homogeneity at the 800 cm-1.A slight shift to a higher frequency in the tentatively micrometer level. Analysis of our sample via EDX (Fig. 9) assigned TiMO stretching band at #1000 cm-1 might indicate agreed qualitatively with this finding. Homogeneity on an AlMO substitution.30 Longer reaction times are undoubtedly atomic scale would lead to a true binary oxide on crystallisation required to force the titanium–aluminium reactions to cominstead of a mixture of individual oxides (TiO2 and SiO2); pletion and much more work would be needed to fully however, the empirical formula predicted for our product is characterise the products.This was not justified in the context TiSiO4 which is not a known compound.22 For pure titania of the present work.systems synthesised via the non-hydrolytic route, Corriu et al.15 concluded that compared to the conventional sol–gel route,28 Conclusions the crystallisation of titania is delayed as is the anatase to rutile transformation.29 In our titanium–silicon system, the The non-hydrolytic sol–gel route to inorganic oxides has been XRD results show that crystallisation of titania phases, in extended to the synthesis of a series of binary oxides based on particular rutile, is the primary crystallisation process, with no silicon, aluminium and titanium.In principle, one advantage evidence for formation of a crystalline binary oxide. of this route over the well established hydrolytic route is that The low magnification SEM shows the general, crazed reactivity diVerences between diVerent ‘metals’ are reduced, morphology of the surface of this oxide.Again, a lot of loose therefore oVering the promise of more homogeneous binary material is present. At a higher magnification, the faceted systems. Since many of the starting inorganic halides and nature of the crystallites is highlighted (Fig. 8). The EDX alkoxides are liquids, this oVers the prospect of carrying out spectrum (Fig. 9) shows the presence of both silicon and the reactions under solvent-free conditions, albeit paying due titanium as well as oxygen. respect to the high reactivity of some of these systems! Binary For the reaction of TEOS with titanium tetrachloride, no systems based on aluminium–silicon, titanium–silicon and catalyst was necessary owing to the high reactivity of the titanium–aluminium combinations were studied.transition metal halide. At 110 °C, this reaction reached gel- As expected, reactions were generally rapid in the absence ation rapidly. The TG result suggested that the high yield of a solvent and the presence, where necessary, of a catalyst obtained was due to the presence of unreacted precursors.The [iron(III ) chloride]. Formation of true aluminosilicates was IR spectrum suggested binary oxide formation, with a demonstrated, although in one case it was shown by XRD SiMOMTi stretching vibration at 960 cm-1 in addition to that the crystalline product resulting from calcination had a TiMOMTi elongation bands30 at 1250 cm-1 and 725 cm-1 structure (mullite) diVerent from that expected solely from the and SiMOMSi stretching absorption bands2,31,32 at reaction stoichiometry. The extent to which this resulted from 1220 cm-1, 810 cm-1 and 525 cm-1.Corriu et al.19 have network rearrangement processes during calcination was not performed a similar reaction but used silicon isopropoxide determined. At the calcination temperature of 1000 °C, thermo- instead of TEOS.They formed a gel in five days at 40 °C, with dynamics control the crystallisation of mullite. Amorphous a 100% oxide yield. A 10% mass loss was seen up to 1200 °C. material such as silica was also present. In the case of titanium– XRD, determined after five hours heat treatment of their silicon systems, there was no clear spectroscopic evidence for product at 500 °C, showed the crystalline phase to be anatase the initial formation of a binary oxide.The results after and the specific surface area, also after calcination at this calcination showed the presence of only pure crystalline titania temperature, was found to be 590 m2 g-1. In the reaction phases. The presence of silica nanophases could explain the between titanium tetrachloride and TEOS, the condensation lack of silica crystallisation under these conditions.In the case rate needs to be reduced in order to achieve greater reaction of the titanium tetrachloride/TEOS reaction, a lower reaction control, i.e. by reducing the reaction temperature a lower rate might improve product homogeneity. The titanium–alu- reaction rate would avoid premature gelation and improve minium reactions led to incomplete oxide formation and precursor mobility.This should improve the homogeneity of unequivocal evidence for binary oxide formation was not the two components and consequently allow the reaction to proceed further towards completion. obtained. Further optimisation of the reaction conditions and 1238 J. Mater. Chem., 1998, 8(5), 1233–123912 (a) R.J. P. Corriu and D. Leclercq, Angew. Chem., Int. Ed. Engl., more in-depth analysis of the products are required to demon- 1996, 35, 1420; (b) D. Leclercq and A. Vioux, Heterog. Chem. Rev., strate the potential for binary oxide formation in these systems. 1996, 3, 65. Overall, this study has shown that in the silicon, aluminium, 13 R. J. P. Corriu, D. Leclercq, P. Lefe`vre, P.H. Mutin and A. Vioux, titanium series, amorphous binary oxides can in some cases J. Non-Cryst. Solids, 1992, 146, 301. be formed directly via the non-hydrolytic sol–gel reaction, but 14 S. Acosta, R. J. P. Corriu, D. Leclercq, P. Lefe`vre, P. H. Mutin and A. Vioux, J. Non-Cryst. Solids, 1994, 170, 234. only the aluminosilicates appear to form true crystalline binary 15 (a) P. Arnal, R.J. P. Corriu, D. Leclercq, P. H. Mutin and A. Vioux, oxides after calcination at 1000 °C. Optimisation of the initial Mater. Res. Soc. Symp. Proc., 1994, 346, 339; (b) P. Arnal, reaction conditions and/or stoichiometry may lead to improved R. J. P. Corriu, D. Leclercq, P. H. Mutin and A. Vioux, J. Mater. homogeneity in the other systems. Chem., 1996, 6, 1925. 16 E. Guenther and M.Jansen, Chem.Mater., 1995, 7, 2110. 17 J. N. Hay and H. M. Raval, J. Sol-Gel Sci. T echnol., in press. The authors are grateful to the Engineering & Physical Sciences 18 (a) M. Andrianainarivelo, R. J. P. Corriu, D. Leclercq, P. H. Mutin Research Council (EPSRC) and the Defence, Evaluation & and A. Vioux, J. Mater. Chem., 1996, 6, 1665; Research Agency (DERA) for the award of a CASE studentship (b)M.Andrianainarivelo, R. J. P. Corriu, D. Leclercq, P. H. Mutin (to H. R.). We also thank DERA for use of their facilities and and A. Vioux, Chem. Mater., 1997, 9, 1098. for undertaking XRD and SEM analyses. Particular thanks 19 R. J. P. Corriu, D. Leclercq, P. Lefe`vre, P. H. Mutin and A. Vioux, Chem. Mater., 1992, 4, 961. are due to Dr M. Clegg and Dr D.Porter (both DERA) for 20 S. Acosta, R. J. P. Corriu, D. Leclercq, P. H. Mutin and A. Vioux, helpful discussions and suggestions. We also thank the Mater. Res. Soc. Symp. Proc., 1994, 345. University of Durham for provision of solid state 29Si NMR 21 M. Andrianainarivelo, R. J. P. Corriu, D. Leclercq, P. H. Mutin services. and A. Vioux, J. Mater. Chem., 1997, 7, 279. 22 H.Dislich, Angew. Chem., Int. Ed. Engl., 1971, 10, 363. 23 I. Iwasaki, S. Yasumroi, S. Shibata and M. J. Yamane, Sol-Gel Sci. T echnol., 1994, 2, 387. References 24 T. Lo�pez, M. Asomosa, L. Razo and R. Go�mez, J. Non-Cryst. 1 C. L. Bird and A. T. Kuhn, Chem. Soc. Rev., 1981, 10, 49. Solids, 1989, 108, 45. 2 C. J. Brinker and G. W. Scherer, T he Physics and Chemistry of 25 L. Bourget, R. J. P. Corriu, D. Leclercq, P. H. Mutin and A. Vioux, Sol–Gel Processing, Academic Press, London, 1990. in First European Workshop on Hybrid Organic-Inorganic 3 D. Avnir, Acc. Chem. Res., 1995, 28, 328. Materials, ed. C. Sanchez and F. Ribot, Paris, 1993, 305, 308. 4 B. M. Novak, Adv. Mater., 1993, 5, 422. 26 (a) F. Babonneau, New J. Chem., 1994, 18, 1065; (b) P. Prasad and 5 J. Livage et al., Prog. Solid State Chem., 1988, 18, 259. B. A. Reinhardt, Chem.Mater., 1990, 2, 660. 6 C. D. Chandler, C. Roger and M. J. Hampden-Smith, Chem. Rev., 27 C.-C. Chen, A. B. Herhold, C. S. Johnson and A. P. Alivisatos, 1993, 93, 1205. Science, 1997, 276, 398. 7 L. Hench and J. K. West, Chem. Rev., 1990, 90, 33. 28 (a) K. Kamiya, K. Tanimoto and T. Yoko, J.Mater. Sci. L ett., 1986, 8 W. Gerrard and P. F. GriVet, Chem. Ind., 1959, 55. 5, 402; (b) M. Aizawa, Y. Nakagawa, Y. Nosaka, N. Fujii and 9 (a) S. Acosta, P. Arnal, R. J. P. Corriu, D. Leclercq, P. H. Mutin H. Miyama, J. Non-Cryst. Solids, 1990, 124, 112. and A. Vioux, Mater. Res. Soc. Symp. Proc., 1994, 346, 43; 29 P. Arnal, R. J. P. Corriu, D. Leclercq, P. H. Mutin and A. Vioux, (b) R. J. P. Corriu, D. Leclercq, P. Lefe`vre, P. H. Mutin and Chem. Mater., 1997, 9, 694. A. Vioux, J. Mater. Chem., 1992, 2, 673; (c) R. J. P. Corriu, 30 M. Schraml-Marth, K. L. Walther and A. Wokaun, J. Non-Cryst. Solids, 1992, 143, 93. D. Leclercq, P. Lefe`vre, P. H. Mutin and A. Vioux, J. Sol-Gel Sci. 31 Analysis of Silicones, ed. A. L. Smith, Wiley-Interscience, New T echnol., 1997, 8, 89. York, 1974. 10 (a) W. Gerrard and J. V. Jones, J. Chem. Soc., 1952, 1690; 32 R. K. Iler, T he Chemistry of Silica, Wiley, New York, 1979. (b) W. Gerrard and K. D. Kilburn, J. Chem. Soc., 1956, 1536. 11 (a) A. Zappel, J. Am. Chem. Soc., 1955, 77, 4228; (b) R. Schwartz and W. Kucher, Chem. Ber., 1956, 89, 169. Paper 7/07549I; Received 20th October, 1997 J. Mater. Chem., 1998, 8(5), 1233–1239

ISSN:0959-9428    

DOI:10.1039/a707549i

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Preparation and thermal stability of La2O3–Al2O3 aerogels from chemically modified Al-alkoxide Hiroshi Kobayashi, Kiyoharu Tadanaga* and Tsutomu Minami Department of AppliedMaterials Science, College of Engineering, Osaka Prefecture University, Sakai, Osaka 599, Japan La-doped Al2O3 aerogels (containing 0.5–10 mol% La2O3) were prepared using Al-alkoxide chemically modified with ethyl acetoacetate.The specific surface area of non-doped Al2O3 aerogels heat-treated at 1000 to 1200 °C was greatly increased by doping a small amount of La2O3. X-Ray diVraction measurements showed that the crystalline phase of La-doped Al2O3 aerogels containing 1 mol% La2O3 heat-treated at 1200 °C for 2 h was h-Al2O3, while that of non-doped Al2O3 aerogels was a-Al2O3.The La2O3–Al2O3 aerogels heat-treated at 1000–1200 °C showed the largest specific surface area at the composition of 1 mol% La2O3 with the presence of the h-Al2O3 phase. Recently, catalyst supports with high thermal stability have PriOH for at least one week at 50 °C and the alcohol was renewed several times to wash out the remaining water in the been required because of the increase of high temperature reactions such as catalytic combustion at temperatures higher aging process.The aged wet gels were then supercritically dried in an autoclave at ca. 10 MPa at 270 °C, in which the than 1000 °C.1 However, the specific surface area of conventional c-Al2O3 supports is decreased by sintering at such high initial pressure of 4 MPa was introduced by nitrogen gas.Xerogels were also prepared as a reference by drying the wet temperatures, and thus catalytic activity of the supports is decreased. It has been reported that metal oxides such as gels at 50 °C for ca. 2 weeks for comparison. The aerogels and xerogels were heat-treated at several La2O3 were eVective in increasing the thermal stability of conventional c-Al2O3.2–7 temperatures for 2 h.The BET specific surface area and the pore size distribution of these heat-treated gels were measured Aerogels, which are obtained via the supercritical drying of wet gels prepared through, for example, the sol–gel method, by the nitrogen adsorption method (Micromeritics, Accusorb 2100). The precipitated crystalline phases of heat-treated gels are porous materials with porosity higher than 90%.8 They have many characteristic properties owing to the high porosity; were identified by X-ray diVraction (XRD) measurements (Rigaku Co., RINT 1000).for instance, large specific surface area, low refractive index, etc. Al2O3 aerogels have been investigated as catalysts and catalyst supports.9 We have already reported the microstruc- Results ture of Al2O3 xerogels and aerogels prepared from Al-alkoxides chemically modified with various stabilizing agents like b- Table 1 shows the appearance, bulk density, porosity and specific surface area of non-doped Al2O3 and La2O3–Al2O3 diketones and alkanolamines,10–12 and shown that Al2O3 aerogels have larger specific surface areas at high temperatures (containing 1 mol% La2O3) aerogels without heat treatment.Although some cracks and small shrinkage were observed in than Al2O3 xerogels. In this study, the addition of La2O3 to Al2O3 aerogels has non-doped Al2O3 aerogels, the appearance is monolithic and translucent. The bulk density of non-doped Al2O3 aerogels is been investigated in order to further increase the thermal stability. Non-doped Al2O3 and La2O3–Al2O3 aerogels were ca. 0.23 g cm-3, and the porosity calculated from the bulk density is ca. 94%. The specific surface area of non-doped prepared from Al-alkoxide chemically modified with ethyl acetoacetate. The specific surface areas, pore size distributions, Al2O3 aerogels is ca. 700 m2 g-1. In the case of La2O3–Al2O3 aerogels, the appearance is also monolithic and translucent, and crystalline phases were measured, and the influences of La2O3 content on the thermal stability of aerogels are and the bulk density and porosity are the same as those of non-doped Al2O3 aerogels.The specific surface area of discussed. La2O3–Al2O3 aerogels is ca. 650 m2 g-1. Fig. 1 shows the specific surface area of non-doped Al2O3 Experimental and La2O3–Al2O3 (containing 1 mol% La2O3) aerogels heat- Aluminium-tri-sec-butoxide (Kanto Chemical Co.), Al(OBus)3, treated at various temperatures for 2 h.The specific surface was used as a starting material. Al(OBus)3 (5.87 g) and isopro- area of the La2O3–Al2O3 aerogels with heat-treatment at pyl alcohol (7.16 g, Wako Pure Chemical Industries), PriOH, temperatures higher than 800 °C for 2 h is larger than that of were mixed and stirred at room temperature for 1 h.Ethyl non-doped Al2O3 aerogels, while that of the La2O3–Al2O3 acetoacetate (3.10 g, Kishida Chemicals), EAcAc, was added aerogels without heat-treatment is smaller than that of nonto the solution as a stabilizing agent, and the solution was stirred at room temperature for 3 h. Water (1.72 g) diluted Table 1 Physical properties of non-doped Al2O3 and La2O3–Al2O3 with PriOH (7.16 g) containing varying amounts of lanthanum (1 mol % La2O3) aerogels without heat treatment nitrate hexahydrate (Nakalai Tesque), La(NO3)3 6H2O, was added to the solution with stirring for hydrolysis.The sols bulk density porosity specific surface aerogel appearance /g cm-3 (%) area/m2 g-1 were kept in closed containers for gelation at 50 °C.The molar ratios of PriOH, EAcAc and H2O to the total of Al(OBus)3 Al2O3 translucent 0.23 94 700 and La(NO3)3 were set to 10, 1, and 4, respectively. The La2O3–Al2O3 translucent 0.23 94 650 amount of La2O3 was varied from 0 to 10 mol%. (1 mol% La2O3) For the preparation of aerogels, the wet gels were aged in J. Mater. Chem., 1998, 8(5), 1241–1244 1241Fig. 3 Specific surface areas of La2O3–Al2O3 aerogels ($) and xerogels Fig. 1 Specific surface areas of non-doped Al2O3 (#) and La2O3–Al2O3 (#) with various amounts of La2O3. All aerogels and xerogels were (containing 1 mol% La2O3) ($) aerogels heat-treated at various heat-treated at 1000 °C for 2 h. temperatures for 2 h doped Al2O3 aerogels (Table 1). The specific surface areas of both non-doped Al2O3 and La2O3–Al2O3 aerogels linearly decrease with an increase of the heat-treatment temperature.The extent of decrease for La2O3–Al2O3 aerogels is smaller than that for non-doped Al2O3 aerogels. Fig. 2 shows the pore size distribution of non-doped Al2O3 and La2O3–Al2O3 (containing 1 mol% La2O3) aerogels. The aerogels were heat-treated at 1000 °C for 2 h. Both non-doped Al2O3 and La2O3–Al2O3 aerogels have micropores with pore size of 1–40 nm in radius.The pore size distribution of La2O3–Al2O3 aerogels is very similar to that of non-doped Al2O3 aerogels, while the pore volume of La2O3–Al2O3 aerogels is larger than that of non-doped Al2O3 aerogels. In the range of pore size of 1–2 nm in radius, where pores with these sizes mainly aVect the specific surface area, the pore volume of La2O3–Al2O3 aerogels is clearly larger than that of non-doped Al2O3 aerogels.Fig. 3 shows the specific surface areas of La2O3–Al2O3 aerogels with various amounts of La2O3. The specific surface Fig. 4 Specific surface areas of La2O3–Al2O3 aerogels with various areas of the La2O3–Al2O3 xerogels are also shown for compariamounts of La2O3. All aerogels were heat-treated at various tempera- son.The aerogels and xerogels were heat-treated at 1000 °C tures for 2 h. for 2 h. The specific surface area of the aerogels are obviously larger than that of the xerogels at any composition. The La2O3–Al2O3 aerogels and xerogels show maximum specific surface areas at 1–3 mol% La2O3. At compositions larger than maximum specific surface area of ca. 235 m2 g-1 at 1–3 mol% 3 mol% La2O3, the specific surface area of both La2O3–Al2O3 La2O3 and the specific surface area of non-doped Al2O3 aeroaerogels and xerogels is decreased with an increase of the gel is ca. 176 m2 g-1. For heat-treatment at 1200 °C, the amount of La2O3. La2O3–Al2O3 (containing 1 mol% La2O3) aerogel shows a Fig. 4 shows the specific surface areas of non-doped Al2O3 maximum specific surface area of ca. 100 m2 g-1, and the and La2O3–Al2O3 aerogels heat-treated at 1000, 1200 and specific surface area of the non-doped Al2O3 aerogel is only 1300 °C for 2 h. As shown in Fig. 3, when the aerogels were ca. 9 m2 g-1. For heat-treatment at 1300 °C, the specific surface heat-treated at 1000 °C, the La2O3–Al2O3 aerogels show a areas of the La2O3–Al2O3 aerogels are ca. 25 m2 g-1, and show no changes with the amount of La2O3. Fig. 5 shows the XRD patterns of non-doped Al2O3 and La2O3–Al2O3 aerogels. The aerogels were heat-treated at (a) 1000 °C and (b) 1200 °C for 2 h. When the aerogels were heat-treated at 1000 °C, the XRD pattern of the Al2O3 aerogel shows the presence of c- and h-Al2O3 phases. The La2O3–Al2O3 aerogels are almost amorphous with the addition of more than 1 mol% La2O3.When the aerogels were heat-treated at 1200 °C, the a-Al2O3 single phase is observed for non-doped Al2O3 aerogels. At the composition of 1 mol% La2O3, only the h-Al2O3 phase is observed. At compositions of La2O3 larger than 3 mol%, the XRD patterns show that the La-b- Al2O3 phase, the composition of which is La2O3 11Al2O3, is present.In the aerogel containing 10 mol% La2O3, the XRD Fig. 2 Pore size distributions of non-doped Al2O3 (#) and pattern shows the presence of a lanthanum aluminate phase, La2O3–Al2O3 (containing 1 mol% La2O3) ($) aerogels heat-treated at 1000 °C for 2 h LaAlO3. 1242 J. Mater. Chem., 1998, 8(5), 1241–1244Discussion As shown in Table 1, the specific surface area of the La2O3–Al2O3 aerogels without heat-treatment is smaller than that of the non-doped Al2O3 aerogels.It is assumed that the specific surface area of aerogels as-prepared is aVected by the structure of wet gels. In the present study, the gelation time of the La-doped Al2O3 gels was longer than that of non-doped Al2O3 gels because of the addition of La(NO3)3. Therefore, it is suggested that the longer gelation time causes the diVerence in the microstructure between La-doped and non-doped Al2O3 gels; gels with longer gelation time have denser microstructure.11 The addition of La2O3 to conventional c-Al2O3 has been studied by several researchers.For example, Schaper et al.2–4 reported the thermal stability of commercial c-Al2O3 supports, to which La2O3 was introduced by the impregnation method.The maximum specific surface area was obtained with the addition of 1–2 mol% La2O3. Yamashita et al.6 investigated the addition of rare-earth metal oxides to Al2O3 through coprecipitation of aluminium nitrate and rare-earth metal nitrates, and reported that the specific surface area of Ladoped Al2O3 showed the largest value by the addition of 5 mol% La2O3.As shown in Figs. 1, 3 and 4, the thermal stability of Al2O3 aerogels and xerogels prepared by the sol–gel method is increased by the addition of La2O3; this is similar to conventional c-Al2O3 supports. As shown in Fig. 2, the pore volume of La2O3–Al2O3 aerogels is larger than that of Al2O3 aerogels. This also shows that the addition of La2O3 prevents Al2O3 aerogels from sintering. In the present study, the Al2O3 aerogels with 1 mol% La2O3 showed the maximum specific surface area.This result is similar to samples prepared by the impregnation method.3 In this method, it is expected that La ions exist on the surface of Al2O3 particles. In the sol–gel method, dopants are usually expected to be highly dispersed in the gel matrix. In the present study, Al-alkoxide was used as the Al source, while La(NO3)3 was used as the La source.Thus, it is assumed that La ions are not dispersed in the whole oxide network of Al2O3 particles but in the oxide network of the surface of Al2O3 particles, so that the addition of 1 mol% La2O3 causes a drastic change of thermal stability. As shown in Fig. 3, the eVect of La2O3 addition to aerogel and xerogel appeared in diVerent composition ranges.In aerogels, the increase of specific surface area was observed in the range of La2O3 content from 0.5 to 7 mol%. In the case of the xerogels, which had smaller specific surface areas than aerogels, an increase of specific surface area was observed by the addition of 0.5 to 3 mol% La2O3. As shown in Fig. 4, aerogels heat-treated at 1200 °C have almost the same specific surface areas and showed similar La2O3 content dependence to that of xerogels.These results show that the specific surface area is closely connected to the optimum dopant level, in other words there is an optimum dopant level per unit surface area. If La ions exist on the surface of Al2O3 particles, this composition dependence of specific surface area can be explained by the following reason: the La ion concentration per unit surface area shows only slight change with the amount of La2O3 for gels with large specific surface areas, whereas it is greatly changed by La2O3 content for gels with small specific surface areas.As shown in Fig. 4, the specific surface area of the La2O3–Al2O3 aerogels is about 1.3 times larger than that of Fig. 5 XRD patterns of La2O3–Al2O3 aerogels with various amounts non-doped Al2O3 aerogels when heat-treated at 1000 °C, about of La2O3. The aerogels were heat-treated at (a) 1000 °C and (b) 1200 °C 11.1 times at 1200 °C, and about 3.5 times at 1300 °C. This for 2 h (6, h-Al2O3; %, c-Al2O3; #, a-Al2O3; $, La-b-Al2O3; +, LaAlO3). indicates that the degree of increase of the thermal stability of the La2O3–Al2O3 aerogels is the largest when the aerogels were heat-treated at 1200 °C.Results of XRD measurement (Fig. 5) are compared with those of specific surface area (Fig. 4) for heat-treatment at 1200 °C. It is known that sintering of J. Mater. Chem., 1998, 8(5), 1241–1244 1243Al2O3 proceeded via surface diVusion,3 and when densification from crystallizing but also cause nucleation of La-containing crystals (La-b-Al2O3 and LaAlO3) by heat-treatment at ca.of Al2O3 proceeded with sintering, the specific surface area of 1200 °C. Therefore, the thermal stability of Al2O3 aerogels was Al2O3 was rapidly decreased as a result of nucleation and not increased by incorporation of La ions, and also the specific growth of a-Al2O3.4 From the results of Fig. 4 and 5, the surface areas of Al2O3 aerogels were decreased. specific surface area of non-doped Al2O3 aerogels was very small because of the crystallization of the a-Al2O3 phase after densification. When 0.5 mol% La2O3 was added to Al2O3 Conclusion aerogels, the La2O3–Al2O3 aerogels have larger specific surface We confirmed that the specific surface areas of Al2O3 aerogels areas than the non-doped Al2O3 aerogels. The XRD pattern were greatly aVected by the addition of La2O3.The addition shows the presence of a- and h-Al2O3 phases and this indicates of a small amount of La2O3 to Al2O3 aerogels prevented the that the addition of La2O3 prevents Al2O3 aerogel from aerogels from sintering and crystallizing to a-Al2O3, and led sintering and crystallizing to the a-Al2O3 phase.The largest to large specific surface areas at high temperatures. It was specific surface area is observed at 1 mol% La2O3 added to shown that the specific surface area of La2O3–Al2O3 aerogels Al2O3, and the XRD pattern shows only the formation of hbecame large when the h- and c-Al2O3 phases, low temperature Al2O3. When 3 mol% La2O3 was added to Al2O3, the XRD phases of Al2O3, were present.The specific surface area was pattern of the La2O3–Al2O3 aerogel shows that the La-b-Al2O3 decreased with an increase of the crystalline phases containing phase is present. Broad diVraction peaks of this crystalline lanthanum, like La-b-Al2O3 and LaAlO3. It was found that phase indicate that the gels were not well crystallized. Thus, the addition of 1 mol% La2O3 was most eVective to give large this low crystallinity caused the comparatively large specific specific surface areas at high temperatures.surface area. When more than 5 mol% La2O3 was added to Al2O3, La-containing crystalline phases (La-b-Al2O3 and The present study was supported by a Grant-in-Aid from the LaAlO3) were detected. Sharp diVraction peaks of these crystal- Ministry of Education, Science, Sports, and Culture of Japan. line phases show that the gels were well crystallized in these compositions. This shows that the high crystallinity causes the References small specific surface area.As a whole, it is shown that the specific surface area became the largest when the h-Al2O3 1 D. L. Trimm, Appl. Catal., 1993, 7, 249. phase, a low temperature phase of Al2O3, with low crystallinity 2 H.Schaper and L. L. van Reijen, Proc. 5th Int. Round Table was present. Conference on Sintering, Porotoz, 1981, p. 173. 3 H. Schaper, E. B. M. Doesburg and L. L. van Reijen, Appl. Catal., Kumar et al.7 reported the mechanism of the improved 1983, 7, 211. stability of La-doped Al2O3 as follows: the rare-earth cations 4 H.Schaper, E. B. M. Doesburg, P. H. M. de Corte, and L. L. van enter into the interstitial positions of the lattice of transition Reijen, Solid State Ionics, 1985, 16, 261. alumina (c-, h-, d-Al2O3, etc.) and decrease the oxygen vacancies 5 B.Be� guin, E. Garbowski and M. Primet, Appl. Catal., 1991, 75, formed in the neck regions which are the nucleation sites of 119. 6 H. Yamashita, A.Kato, N. Watanabe and S. Matsuda, J. Chem. a-Al2O3. In this study, it is expected that La ions exist in the Soc. Jpn. Chem. Ind. Chem., 1986, 1169. oxide network of the surface of Al2O3 particles as mentioned 7 K.-N. P. Kumar, J. Tranto, J. Kumar, and J. E. Engell, J. Mater. above, and it is expected that there are many defects in the Sci. L ett., 1996, 15, 266. oxide network of the surface of prepared Al2O3 gels. Thus, the 8 J. Fricke and A. Emmering, J. Am. Ceram. Soc., 1992, 75, 2027. results of the present study show that La ions should decrease 9 Y. Mizushima and M. Hori, Appl. Catal. A, 1992, 88, 137. the oxygen vacancy concentration. When the content of La2O3 10 K. Tadanaga, T. Iwami, N. Tohge and T. Minami, J. Sol–Gel Sci. T echnol., 1994, 3, 5. added to the Al2O3 aerogels is 0.5–3 mol%, La ions exist in 11 K. Tadanaga, T. Iwami, T. Minami and N. Tohge, J. Ceram. Soc. the oxide network of the surface of Al2O3 particles and inhibit Jpn., 1995, 103, 582. the sintering of Al2O3 via surface diVusion. Hence La ions 12 K. Tadanaga, S. Ito, T. Minami and N. Tohge, J. Non-Cryst. Solids, prevent Al2O3 aerogels from crystallizing to the a-Al2O3 phase. 1996, 201, 231. When the content of La2O3 added to Al2O3 aerogels was larger than 5 mol%, La ions not only prevent Al2O3 aerogels Paper 8/00778K; Received 28th January, 1998 1244 J. Mater. Chem., 1998, 8(5), 1241–12

ISSN:0959-9428    

DOI:10.1039/a800778k

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Azacrown-CH2-bipyridine receptors in silica xerogel. Optical and coordination properties† Andrzej M. K�onkowski,a Krzysztof Kledzik,a Tadeusz Ossowskia and Anna Jankowska-Frydelb aFaculty of Chemistry, University of Gdansk, ul. Sobieskiego 18, 80–952 Gdansk, Poland bInstitute of Experimental Physics, University of Gdansk, ul. Wita Stwosza 7, 80–952 Gdansk, Poland Samples of silica xerogel doped with encapsulated series of three receptors of the type (aza-3n-crown-n)-CH2-(2,23-bipyridine), where n=4, 5 and 6, were prepared by the sol–gel process.Fluorescence excitation spectra of the encapsulated receptors diVer from the absorption spectra. Emission of the samples can be quenched specifically owing to the coordination reaction with transition metal ions.Emission decays can be described by a double exponential, the lifetime values being short at room temperature. EPR spectra of CuII ions complexed on the surface of the receptor doped xerogel have fairly well resolved hyperfine structure and show an elongated rhombic octahedral environment for the receptors with 2,23-bipyridine groups. Owing to the utilizing of a low temperature preparation of (with two amine nitrogen atoms as donors) and aza-3n-crownn, i.e. 3n-membered macrocycle (with n-1 ethereal oxygen oxide materials, the so-called sol–gel process,1 the previous problem related to the inability to incorporate organic mol- atoms and an amine nitrogen atom as donors). The fluorescent sensing is studied with respect to Cu2+ as a model transition ecules into an inorganic oxide glass is resolved.Using this technique, a three-dimensional network of metal-oxide bonds metal ion and other transition metal ions for comparison. The aim of this study is to prepare material which could is formed at room temperature by the polymerisation–condensation reaction of metal alkoxides, followed by low temperature possibly be proposed as an optical recognition phase for a chemical optical sensor.The structure of metal complexes dehydration. The porous xerogel matrix thus obtained can trap receptor molecules.2 formed after reaction with the studied receptors in silica xerogel matrix is also reported. Inorganic oxides are superior to organic polymer matrices due to the fact that the excited state of the trapped organic molecule is capable of undergoing photochemical reactions Experimental with the surrounding organic matrix.This results in low photostability of both dopants and carrier.3 Chemicals Silica xerogels can play the role of host matrices doped with Tetramethoxysilane, copper(II ) perchlorate Cu(ClO4)2 6H2O encapsulated sensing organic indicators. The gels have the (both from Fluka A.G., Switzerland), NiII, CoII, MnII and CrIII obvious advantage of chemical inertness, mechanical stability nitrates (from Aldrich Co.) as well as vanadyl sulfate and optical transparency.What is more important, they are VOSO4 5H2O (from Merck A.G. Germany) of analytical physically able to encapsulate the indicator molecules in pores grade were used without further purification.in the gels, so that these molecules cannot be leached out in The series of (aza-3n-crown-n)-CH2-(2,23-bipyridine) ligands solution. At the same time the host material is suYciently (receptors) were synthesized by a preparation method similar porous to enable transport of metal ions, solvent and other to that in ref. 7 from azacrowns aza-3n-crown-n (abbr. A3nCn, small molecules into the interior.4 where n=4, 5 and 6) and 2,23-bipyridine (bpy), both from Several recent observations revealed the feasibility of making Aldrich Chemical Co.optical recognition phases with receptor molecules based on Methanol and ammonia were of analytical grade purity. The silica gel for chemical sensors.5 The exciting outcome of these water used was triply distilled from glass.observations is the feasibility of preparing gels boasting optical properties which change in the presence of target molecules. Sample preparation In particular, gels with receptors producing characteristic colour changes when exposed to the metal ions were prepared. The sols were prepared by a typical sol–gel procedure1 from It is also noted that many other organic molecules incorporated a starting mixture of: tetramethoxysilane (TMOS), methyl into the sol–gel optical materials exhibit luminescence through- alcohol as diluent, distilled water (TMOS5H2O=154), out the UV–VIS region.2,6 In this case metal ions penetrating NH3(aq) catalyst and receptor A3nCn-CH2-bpy.The mixture a porous gel matrix can be complexed by the organic molecules was vigorously mixed at room temperature.The sol was which results in quenching of the fluorescence. allowed to gel for 3 days and then dried. The xerogel obtained We are interested in investigating the spectroscopic properties and studying the behaviour of an optical material with supramolecular dopants which exhibits chemical sensing. The optical material is a silica xerogel prepared by the sol–gel method and the supramolecular dopants (Fig. 1) are receptors containing two coordinatively active subunits: 2,23-bipyridine † Presented in part at the 11th International Symposium on the Photochemistry and Photophysics of Coordination Compounds, Fig. 1 Molecular structure of A3nCn-CH2-bpy receptor, where n=4 Cracow, Poland, July 1995. J. Mater. Chem., 1998, 8(5), 1245–1249 1245was heated for 3 h at 120 °C to remove ammonia and methanol as well as some of the water from the pores.Concentration of the receptor in the xerogel was 2.5×10-5 mol g-1 SiO2. The dry xerogel was then ground in a mortar and passed through standard sieves. Particles 0.75–1.50 mm in size were immersed in aqueous solutions containing such metal ions as: Cu2+, Ni2+, Co 2 +, Mn 2 +, Cr 3 + or VO2+ 200 mg (±0.1 mg) of the doped material was agitated in 0.01 M aqueous solution of the appropiate metal salt, maintaining the molar ratio [metal ion]5[receptor]=1051.The immersed xerogel was filtered oV after 24 h, rinsed with distilled water and then dried. By this chemisorption method complexes with transition metal ions were formed. The rate of uptake of the metal ions on the surface of the receptor doped xerogel was measured by cyclic voltammetry on a hanging mercury drop electrode, shaking 200 mg of the material with the metal ion solution.Complexes of CuII with A3nCn-CH2-bpy ligands (in a molar ratio 151) were also produced in the sol just before gelation. After 3 days the complexes were encapsulated by the sol–gel process in wet silica gel which was then dried.The procedure in the case of the CuII complex with bpy diVered. This complex, which was used for comparison, was Fig. 2 Optical absorption spectra of A3nCn-CH2-bpy receptors in methanol solution: (a) n=4; (b) n=5; (c) n=6 formed (in molar ratio Cu5bpy=152) separately and was dissolved in the sol before gelation. The silica xerogel with the trapped complex was obtained after gelation and drying.Apparatus Optical absorption measurements in the UV–VIS region were recorded on a Beckman DU 650 spectrophotometer. Spectra of the crushed xerogel samples were obtained in a silicon oil mull and were collected between 250 and 900 nm. Fluorescence emission (lexc=315 nm) and excitation spectra (lem=360 nm) were measured with a Perkin-Elmer LS 50B spectrofluorometer with reflection spectra attachment.None of the excitation spectra were corrected for the lamp and photomultiplier response. Fluorescence decays were measured using an Edinburgh Analytical Instruments CD 900 fluorometer. EPR spectra were obtained on a SE/X spectrometer (Radiopan, Poznan). Sample holders were sealed quartz capillaries (1 mmdiameter).A magnetic field modulation of 100 kHz was applied. Standard deviations of the EPR spectra param- Fig. 3 Optical absorption spectra of (a) bpy, and of A3nCn-CH2-bpy eters were estimated as follows: g||±0.003, g)±0.005 and receptors in silica xerogel: (b) n=4; (c) n=5; (d) n=6 A)±4×10-4 cm-1. Moreover, in this case, the higher wavelength wing shows comparativelyigher intensity.Results Silica gels prepared under basic conditions (pH>7) and high Coordination process water to silane ratios produce highly branched clusters which The rate of the coordination process (see Fig. 4) indicates that behave as discrete species. Gelation occurs by linking clusters complex formation on the xerogel surface is rather high together.1 This procedure makes the xerogels porous, these consequently being able to encapsulate and attach large supramolecular receptors of the type shown in Fig. 1. Probably owing to the hydrogen bond between oxygen atoms in the ether crown group and the silanol group, the receptors are practically non-leachable in aqueous solution. Absorption spectroscopy The receptors in methanol solution show strong absorption spectra with the characteristic band for free bpy at 290 nm (Fig. 2), whereas Fig. 3 shows room temperature absorption spectra of the receptors (and bpy for comparison) encapsulated in the xerogel. The bands for the receptors with n=4, 5 and 6 are centered at 301, 299 and 298 nm, respectively. The band for the first receptor shows the greatest intensity of UV absorption among the bands compared.Compared with the bpy derivatives, the absorption spectrum of bpy is much Fig. 4 Copper(II) uptake as a function of time for an A12C4-CH2-bpy receptor encapsulated in silica xerogel broader but has almost the same position of the main band. 1246 J. Mater. Chem., 1998, 8(5), 1245–1249(measured in seconds). However, later, due to the metal enrichment within pores controlled by diVusion, a slow complexation rate is observed. In any case the maximum duration time of the process is about 60 min.Luminescence spectroscopy Very weak emission was observed from the xerogel samples with the bpy derivatives on excitation with light of 290 nm (absorption lmax). The fluorescence excitation (lem=360 nm) and emission (lexc=315 nm) spectra of the encapsulated ligands are given in Fig. 5(A) and (B). The bands of A12C4-CH2-bpy exhibit much higher intensity than those of the ligands with larger crown groups (n=5 and 6), but the band positions are the same. The excitation spectra of the bpy derivatives encapsulated in the xerogel diVer from the absorption spectra [the excitation peaks are near zero at 310 nm where the species reach the maximum absorption, cf.Fig. 3 and 5(A)]. The Fig. 6 Characteristic quenching sequence for A18C6-CH2-bpy in silica emission originates from the higher wavelength wing of the xerogel, if the receptor is (a) uncomplexed; and complexed with absorption spectrum. (b) Cr3+; (c) VO2+; (d)Mn2+; (e) Co2+; (f ) Ni2+ and (g) Cu2+ A reasonable fit of the decay curves can be achieved using the fitting functions with more than one exponential component.The best results are achieved for two exponential functions. The adequacy of this exponential decay fitting was judged by an inspection of the plots of the standard deviation and by the statistical parameters x2 (Table 1). In general, the two components have very short lifetimes t. Participation of the longer lifetime is smaller and decreases with increasing n.The sequence of the fluorescence quenching of the receptor with n=6 complexed with transition metal ions such as CuII, NiII, CoII, MnII, CrIII and VOII is shown in Fig. 6. This sequence Fig. 7 Fluorescence quenching eVect due to the coordination of transition metal ions in a silica xerogel doped with A3nCn-CH2-bpy: (a) n= 4; (b) n=5; (c) n=6. Receptors uncomplexed (A), and complexed with (B) VO2+ and (C) Cu2+.is similar for the complexed species with n=4 and 5 and is characteristic of the transition metal ions used. The room temperature fluorescence spectra of the free A3nCn-CH2-bpy ligands in silica and complexed ones owing to coordination of the representative Cu2+ and VO2+ ions are shown in Fig. 7. Quenching of fluorescence for the complexed receptors is observed.The quenching eVect is especially great in the case of the copper(II ) complex with the A12C4-CH2-bpy ligand as compared with the fluorescence intensity of the respective free receptor immobilized in silica [cf. in Fig. 7(a)]. The receptors complexed with CuII cation in each case exhibit near zero emission intensity, whereas vanadyl cation in this Fig. 5 Excitation (A) and emission (B) spectra of A3nCn-CH2-bpy in situation possesses an intermediate position among the trans- silica xerogel: (a) n=4; (b) n=5; (c) n=6.The excitation spectra were ition ions studied. obtained by monitoring at 360 nm and the excitation wavelength for the emission was 315 nm. Recorded at 295 K. EPR spectroscopy Table 1 Photophysical properties of A3nCn-CH2-bpy receptors encap- EPR spectroscopy is a powerful tool with which to identify sulated in silica gel (measured at 295 K) changes in the coordination environment of CuII complexed with the supramolecular entities.EPR spectra of CuII comreceptor, n x2 t/ns contribution (%) plexed with the receptors and then encapsulated in silica xerogel are shown in Fig. 8. The CuII species give rise to typical 4 1.41 0.49 76.9 2.1 23.1 axial spectra. However, in the case of the complexes with bpy 5 1.08 0.11 99.2 and bpy derivatives with n=5 and 6, the spectra exhibit two 1.9 0.8 components. The hyperfine structure of the intense perpendicu- 6 1.40 0.097 99.9 lar signal on the high-field side is not resolved. 2.6 0.1 The EPR spectra of the samples with complexes formed by J.Mater. Chem., 1998, 8(5), 1245–1249 1247Table 2 EPR spectral parameters of CuII complexes with bpy and A3nCn-CH2-bpy trapped in silica xerogel but formed before gelation (spectra recorded at 295 K) ligand,n g||a A||b/10-4 cm-1 g) c bpy 2.289 154 2.055 4 2.265 154 2.041 5 2.259 158 2.033 6 2.258 156 2.027 a±0.003. b±4×10-4 cm-1. c±0.005. Table 3 EPR spectral parameters of CuII complexes formed with bpy and A3nCn-CH2-bpy ligands in silica gel by coordination from aqueous solution (spectra recorded at 295 K) ligand,n g||a A||b/10-4 cm-1 g) c bpy 2.280 154 2.044 4 2.257 155 2.035 5 2.258 157 2.031 6 2.258 156 2.027 a±0.003. b±4×10-4 cm-1.c±0.005. CuII ions on the surface and in the pores of the silica xerogel [see Fig. 8(b)] are similar to each other.Only in the spectrum of the sample with the n=6 receptor are two components clearly visible. In view of the fact that the CuII ion concentration in the gel samples is low (about 10-5 mol g-1 SiO2) we can neglect the possibility of interactions between copper(II) ions in the matrix. Fig. 8 X-Band EPR spectra of silica xerogel doped with: (A) CuII The lineshape of the EPR spectrum conforms satisfactorily complexed in the reaction mixture before gelation by (a) bpy; and by with the theoretical one which is obtained assuming the spin A3nCn-CH2-bpy ligands with (b) n=4; (c) n=5; (d) n=6.Recorded at 295 K Hamiltonian,8 provided that the current lineshape theory developed by Kneubuehl9 is corrected in respect of peak width. Sets of the g||, A|| and g) values for the encapsulated Cu2+ complexes and owing to coordination of CuII ions are presented in Tables 2 and 3.Owing to poor resolution of the spectra, it is impossible to determine the spectral parameters for the second component in the cases mentioned above. g|| and g) diVer distinctly in the CuII complexes with bpy and with the bpy derivatives on the other side, regardless of the preparation method.For all the CuII complexes, a g||>g)>ge= 2.0023 parameter sequence is observed. Discussion 2,2¾-Bipyridine (bpy) and its derivatives are renowned for their ability to form coordination compounds with metal ions. The description, based on MO theory, of the coordinative bonds in these complexes requires that the central ion and the ligands be able to form s- and p-bonds.It is no wonder, therefore, that bpy is a molecular block par excellence for a wide variety of types of supramolecular devices.10 Bpy, similarly to many of the optical active organic molecules tested, tends to dimerize and aggregate at moderate concentrations in aqueous solution.11 This tendency reduces the fluorescence quantum yield significantly.12 It is important to notice that dimerisation is greatly reduced by the trapping process in the silica xerogel, even though concentrations could be quite high (up to 10-2 M).The de-aggregation is a general phenomenon which indicates the lower polarity of the oxide cage11 than water and confirms the matrix isolation of the trapped molecules.13 The eVect is that maxima of absorption of organic, optically active molecules, are slightly red-shifted (ca. 5 nm) in silica xerogels, as compared with their aqueous solutions. These redshifts confirm the slightly less polar nature of the silica cage which is composed of SiMOH (silanol ) and SiMOMSi Fig. 9 X-Band EPR spectra of silica xerogel doped with (a) bpy, and with A3nCn-CH2-bpy: (b) n=4; (c) n=5; (d) n=6, complexed with groups.13 CuII in aqueous solution.Recorded at 295 K. It is known from previous studies14 that bpy in alcohol at a 1248 J. Mater. Chem., 1998, 8(5), 1245–1249concentration of 10-5 M does not aggregate, but the aggregates Conclusions exist at higher concentration (10-3–10-1 M) in water. In the The sol–gel process appears in this study to be a straightfor- latter case lmax of emission is at 430 nm.15 Thus, in the present ward and versatile fabrication method for the preparation of studies (ligand concentration 10-5 mol g-1 SiO2 and lmax of recognition phases for optical chemical sensors.emission at 345 nm) aggregation of the bpy derivatives cannot With the help of intensity quenching experiments we investi- be expected. It seems that the emission is probably from the gated three diVerent luminescent A3nCn-CH2-bpy receptors azacrown derivatives of bpy bonded to the silica network by (where n=4, 5 and 6) encapsulated in porous silica xerogel.hydrogen bonds between the silanol groups and oxygen atoms Among the samples the system with A12C4-CH2-bpy is the in the crown group. most promising as a component of the recognition phase in The considerable intensity of fluorescence emission for silica the optical chemical sensor for Cu2+ ions.xerogel doped with A12C4-CH2-bpy [Fig. 5(a)] is one of the The elongated rhombic octahedral environment of CuII ion promising features for the recognition phase based on this complexed with ligands of the A3nCn-CH2-bpy type (where receptor. It seems that the phase studied is suitable for selective n=4, 5 and 6) consists of two bpy ligand groups in the analysis of CuII ions in solutions owing to substantial quenchequatorial plane.ing of the emission by these ions [Fig. 7(a)]. The order of the emission intensity MnII>CoII>NiII>CuII is the reverse of the Irving–Williams order, i.e. order of the The financial support of this work by the Polish Scientific stabilities of corresponding complexes of the bivalent ions of Research Council (grant 7 T08A 028 10) is gratefully the first transition series, irrespective of the particular ligand acknowledged.involved.16 The bi-exponential decay observed means that two species take part in the excitation process. Probably one of the species References is the receptor supramolecule anchored with the xerogel network by hydrogen bonding and the second species is the non- 1 C.J. Brinker and G. W. Scherer, Sol–Gel Science. T he Physics and Chemistry of Sol–Gel Processing, Academic Press, London, 1990. bonded supramolecule. If the content of the non-bonded 2 D. Avnir, V. R. Kaufman and R. Reisfeld, J. Non-Cryst. Solids, species decreases with n then one of the decay components 1985, 74, 395.decreases (cf. in Table 1). 3 N. S. Allen and J. F. McKeller, Photochemistry of Dyed and The CuII ion can give numerous bis-bipyridine complexes Pigmented Polymers, Applied Science, Publishers, London, 1980. with the general formula Cu(bpy)2X2 nH2O. Complexes with 4 B. Dunn and J. I. Zink, J.Mater. Chem., 1991, 1, 903. anions which are generally reluctant to bond to the metal such 5 M.A. Arnold, Anal. Chem., 1992, 64, 1015. 6 Y. Kobayashi, Y. Imai and Y. Kurakawa, J.Mater. Sci. L ett., 1988, as X=ClO4- are of some interest. For these complexes the 7, 1148. axial CuMO bonds are considerably longer than the equatorial 7 J. C. Rodriguez-Ubis, B. Alpha, P. Plancherd and J.-M. Lehn, CuMN bonds. It is therefore uncertain whether, in the strict Helv.Chim. Acta, 1984, 67, 2264. sense, the oxygen atoms are bonded to the copper ion or not. 8 A. H. Maki and B. R. McGarvey, J. Chem. Phys., 1958, 29, 31. Hathaway et al.17–19 concluded that the compound should be 9 F. K. Kneubuehl, J. Chem. Phys., 1960, 33, 1074. formulated as [Cu(bpy)2](ClO4)2 and that the complex cation 10 F. Voegtle, Supramolecular Chemistry. An Introduction, Wiley, has an essentially planar structure, although the two bpy Chichester, 1991. 11 F. P. Schafer, T opics in Applied Physics, vol. 1: Dye L asers, 2nd ligands are twisted mutually by 10–30° towards a tetrahedral edn., Springer, Berlin, 1977. coordination. This is not entirely in agreement with results 12 K. Kemnitz, T. Murao, I. Yamazaki, N. Nakashima and obtained by Nakai20 regarding the aspect of the coordination K.Yoshihara, Chem. Phys. L ett., 1983, 101, 337. of the ClO4- ion. This author presented the conclusion that 13 D. Avnir, D. Levy and R. Reisfeld, J. Phys. Chem., 1984, 88, 5956. in the [Cu(bpy)2](ClO4)2 crystal the oxygen atoms complete 14 E. Castellucci, L. Angeloni, G. Marconi, E. Venuti and I. Baraldi, a distorted octahedron in the axial direction, though one of J.Phys. Chem., 1990, 94, 1740. 15 S. Dyanya and Bhattacharyya, J. Photochem. Photobiol. A: Chem., the CuMO distances is considerably longer than the other. 1992, 63, 179. Thus, the complex could be regarded as six-coordinate. 16 N. N. Greenwood and A. Earnshaw, Chemistry of the Elements, Since the EPR spectral parameters obey the order Pergamon Press, Oxford, 1984, p. 1065. g||>g)>ge for all of the studied coordination species (see 17 J. M. Procter, B. J. Hathaway and P. Nicholls, J. Chem. Soc. A, Tables 2 and 3), it could be proposed that there exists a 1968, 1678. tetragonal coordination environment in the samples with CuII 18 D. S. Brown, J. D. Lee, B. G. A. Melson, B. J. Hathaway, complexes prepared by encapsulation and chemisorption. J. M. Procter and A. A. G. Tomlinson, Chem. Commun., 1967, 369. 19 B. J. Hathaway, J. M. Procter, R. C. Slade and A. A. G. Tomlinson, Assuming that complexes of the type [Cu(bpy)2]2+ are unable J. Chem. Soc. A, 1969, 2219. to adopt a square-planar configuration because of steric inter- 20 H. Nakai, Bull. Chem. Soc. Jpn., 1971, 44, 2412. action between hydrogen atoms to the nitrogen,21 the coordi- 21 P. J. Burke, D. R. McMillin and W. R. Robinson, Inorg. Chem., nation environment of CuII should be elongated rhombic 1980, 19, 1121. octahedral.22 22 J. Foley, D. Kennefick, D. Phelan, S. Tyagi and B. Hathaway, An increase of g|| indicates decreasing tetragonality of the J. Chem. Soc., Dalton T rans., 1983, 2333. 23 B. J. Hathaway, in Comprehensive Coordination Chemistry, ed. coordination sphere of CuII.23 It suggests that the tetragonal G. Wilkinson, Pergamon Press, Oxford, 1987, vol. 5, pp. 667 et distortion of the complexes increases as the ligands change seq., 730. from bpy to bpy derivatives. In addition, this eVect is greater for complexes created by the chemisorption method than in the reaction mixture before gelation. Paper 7/08131F; Received 11th November, 1997 J. Mater. Chem., 1998, 8(5), 1245–1249 1249

ISSN:0959-9428    

DOI:10.1039/a708131f

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials An oxalate–peroxide complex used in the preparation of doped barium titanate Sven van der Gijp,* Louis Winnubst and Henk Verweij L aboratory for Inorganic Materials Science, Department of Chemical T echnology, University of T wente, P.O.Box 217, 7500 AE Enschede, T he Netherlands A method is described for the preparation of homogeneously doped barium titanate, which can be applied in non-linear dielectric elements.Ba and Ti salts are dissolved, mixed with hydrogen peroxide and added to a solution of ammonium oxalate, resulting in the formation of an insoluble peroxo–oxalate complex. The presence of oxalate ions and a high pH are necessary for the formation after calcination of a stoichiometric and sinterable powder.To characterise the structure of the precipitating complex, the thermal decomposition of the complex is studied by means of XRD, d-TGA and FTIR. It is found that the precipitating complex is BaTi0.91Zr0.09O2(C2O4) 2H2O. The calcined powder prepared with the peroxo–oxalate method contains no second phase, in contrast to powders prepared with the oxalate method and the peroxide method.of oxalate, which changes the complexation properties of the Introduction titanium peroxo complex. The diVerence between the oxalate Doped barium titanate is used in a broad range of electro- process and the peroxo process is that the oxalate process is ceramic devices. One example is a pulse-generating device, conducted at high temperatures, low pH and in the absence which can be applied in lamp starters in a fluorescent lamp.1 of hydrogen peroxide.Hydrogen peroxide changes the solu- For this type of application non-linear dielectric behaviour is bility of titanium peroxo complex drastically.7 However, needed. The required material properties for this type of neither the peroxide nor the oxalate method have been used application are: a high relative permittivity, a steep gradient up till now for Zr-doped BaTiO3.of the polarisation versus electric field hysteresis curve and In order to obtain dielectric properties next to a singlestable non-linear characteristics. These specifications can be phase polycrystalline material, a dense ceramic is required. met when the ceramic devices are made with a sinterable, Therefore, no aggregates in the powder may be present. The homogeneous Zr-doped barium titanate powder.1–3 grain size of the final ceramic must be large, because large Nowadays, doped barium titanate powders are mainly pro- grains reduce internal stress and low values of internal stress duced commercially by the mixed oxide process.The powders result in a higher value for the dielectric maximum.12 made by this route have certain disadvantages, like a large In this paper the eVect of oxalate on the powder properties and non-uniform particle size and a bad chemical homogeneity.is described. The precipitating complex in the peroxo–oxalate However, the mixed oxide process is easy to perform and when process and the thermal decomposition of the precipitating no evaporation occurs, stoichiometric powders can easily peroxo–oxalate complex are characterised by means of thermobe obtained.gravimetric analysis (d-TGA) and infrared spectroscopy However, an improvement in homogeneity and morphology (FTIR). The peroxo–oxalate method should result in a sinof the powder can be obtained when wet-chemical routes are terable, homogenous and single-phase polycrystalline powder.utilised. This development will lead to better dielectric behaviour. Unfortunately, during wet-chemical preparation all kinds of undesired side-reactions may occur, which makes control of Experimental procedure the stoichiometry more complicated. Powder and ceramic preparation An already existing commercial wet-chemical preparation route is the oxalate process.4–6 In this process pure, undoped Peroxo–oxalate method.Titanium oxychloride (0.15 mol; barium titanate is produced stoichiometrically by the formation 0.15 M) and zirconium chloride (0.015 mol; 0.015 M) were added and precipitation at low pH and 60 °C of a barium titanium to concentrated nitric acid (25 ml ) and water. The total volume oxalate complex. Another process for the preparation of barium of the solution was 1 l.This solution was mixed with an titanate is the so-called peroxo process based on the formation aqueous solution of barium nitrate, prepared from barium and precipitation of titanium peroxide complexes at high pH carbonate (0.17 mol; 0.085 M) and nitric acid (50 ml; 0.25 M). and room temperature.7 This process is described not only for The total volume of the mixed solution was 3 l.Finally, to this BaTiO3, but also for the production of titanium-rich materials mixture hydrogen peroxide (30%, 100 ml) was added. This so- (BaTi2O5) as well as barium zirconates and calcium and called ‘precursor’ solution was added dropwise to a solution magnesium titanates.7–11 of aqueous ammonia (0.5 mol) and diVerent amounts of oxalic The peroxo–oxalate process described in this paper is a acid with a total volume of 5 l.During the addition of the combination of the oxalate process and the peroxide process. precursor solution the pH was kept constant by the addition This powder preparation method is suitable for the production of small amounts of aqueous ammonia. of Zr-doped barium titanate (BaTi0.91Zr0.09O3). Doping is In this paper the ratio of oxalate ion concentration over the possible due to the simultaneous presence of hydrogen peroxide concentration of all the metal ions (Ba+Ti+Zr) present in and ammonium oxalate at high pH.solution is called x (0x0.5); note that when x=0 this Note that the only diVerence between the peroxide process process is the peroxide process. All experiments were performed and the peroxo–oxalate process described here is the presence at room temperature. After two hours stirring, the formed precipitate was filtered, washed with ethyl acetate, dried at 150 °C and calcined at †E-mail: s.vandergijp@ct.utwente.nl J.Mater. Chem., 1998, 8(5), 1251–1254 1251900 °C. Powders were first pre-pressed uniaxially at 80 MPa and subsequently isostatically pressed at 400 MPa.All compacts were sintered at 1400 °C for 5 h; heating rate 2 °C min-1, cooling rate 4 °C min-1. Oxalate method. Barium carbonate (0.1 mol) was carefully added to an aqueous solution of nitric acid (50 ml ). To this solution titanium oxychloride (0.091 mol) and zirconium chloride (0.009 mol) were added. This solution had a volume of 2 l. This solution was added dropwise to a solution of oxalate (0.25 mol) in water (1 l ) at a temperature of 60 °C.4 After 2 h stirring the precipitated complex was separated by filtration and calcined at 900 °C.Characterisation The decomposition of the dried precipitate was studied by Fig. 1 XRD of calcined powder prepared with increasing amounts of using TGA (Stanton Redcraft STA 625, heating rate 5 °C min-1 oxalate, x, as indicated on the right; *denotes the (cubic) perovskite to 1000 °C), and Fourier-transform infrared spectroscopy structure (FTIR).FTIR measurements were performed in situ at temperatures from 200–800 °C at temperature intervals of 20 °C Apparently, the presence of oxalate during complex forma- (with a holding [analysis] time at each temperature interval of tion results in a diVerent structure of the final precipitated a few minutes).complex. Doping with Zr seems possible due to the simul- XRD measurements were performed using a Philips PW taneous presence of oxalate and hydrogen peroxide. 1710 diVractometer with filtered Cu-Ka1-radiation, l= When a Zr-doped BaTiO3 calcined powder is prepared 1.4508 A ° . The chemical composition was measured with X-ray according to the oxalate method (in the absence of hydrogen fluorescence spectroscopy (XRF, X-ray spectrometer, Philips peroxide and in acidic environment), secondary phases are PW 1480/10).detected too, mainly BaZrO3, indicating that the oxalate Particle size distributions were measured with a Microtrac method is not suitable for the preparation of BaTi0.91Zr0.09O3.X-500 (Leeds and Northrup). The morphology of the powder and the microstructure of the ceramic were studied with Identification of the precipitating complex scanning electron microscopy (JEOL, JSM 35CF at 15 kV). The microstructure of the ceramics was revealed by etching at The weight loss found from TGA, after calcination to a a temperature 30 °C below the sintering temperature in a temperature of 1000 °C, for the dried precipitates prepared nitrogen atmosphere.To study the influence of oxalate on the with the peroxide method (x=0) is 20%. This corresponds to ligand structure of the complex, the diVerence in absorption the theoretical weight loss found for the thermal decomposition maximum was measured for a solution at pH=4, containing of BaTiO2(O2) 2H2O, the complex formed with the peroxide titanium and hydrogen peroxide, before and after the addition method,7 to BaTiO3.Note that both oxygen ligands in of oxalic acid. UV–VIS measurements were carried out with a BaTi(O2)O2 3H2O are diVerent. Philips PU 8740 spectrophotometer. When oxalate is introduced in the process, the weight loss Non-isothermal densification was studied on a Netzsch 410 increases to 31% for x=0.5.Also the thermal decomposition dilatometer; heating rate 2 °C min-1, cooling rate 4 °C min-1, behaviour changes when more oxalate is added in the process, holding time 3 h. Density measurements were performed with as can be seen in Fig. 3 ( later). In this figure the derivatives of the Archimedes technique using mercury.the TGA (d-TGA) results are given for three precipitates prepared with x=0.125, x=0.25 and x=0.50. It can be seen Results and Discussion that with an increasing amount of oxalate in the reaction mixture the d-TGA signals change, especially at approximately The complexation starts with the addition of a red-coloured 250 and 700 °C. Therefore, the addition of oxalate in the aqueous solution of titanium, zirconium, barium and hydrogen process results in the formation of a diVerent precipitating peroxide to a solution of ammonium oxalate.This results in complex as compared to the precipitating complex in the case the direct formation of a yellow precipitate. During the reaction of the peroxide method (x=0). some gas formation is observed, probably oxygen formed by Confirmation that such a change in ligand structure does the partial decomposition of hydrogen peroxide in the basic occur in the presence of oxalate ions (x>0), is provided by aqueous environment.the shift of the absorption maximum of the complex. This shift takes place from 356 nm in the absence of oxalate to 392 nm Influence of oxalate on powder properties in the presence of oxalate, as measured with UV–VIS spectroscopy in acidic aqueous environment.Fig. 1 shows XRD results of the calcined powders prepared with increasing oxalate content. It is clear that the amount of To characterise the structure of the precipitating complex in the case of x=0.5 FTIR measurements and TGA were per- second phase in the calcined powder as determined by XRD depends on the oxalate concentration. The powder prepared formed in situ as a function of temperature in the range 1000–4000 cm-1.The spectra for x=0.5 are given in Fig. 2. according to the peroxide method (x=0) contains secondary phases. The secondary phases present are mainly BaTi2O5 and IR signal 1 (at 3500 cm-1) corresponds to water. It is clear that water is still present at temperatures up to 600 °C.Most BaCO3. For increasing amounts of oxalate, the amount of these secondary phases decreases. At 0.50 molar equivalent of the water is removed at temperatures between 400 and 500 °C. Signal 2 (1700 cm-1, a characteristic signal13 at oxalate no secondary phases are present and only the cubic perovskite phase is present. 1300 cm-1 is present but is however not visible in this plot) correspond to CO vibrations of an oxalate ligand.14 These In spite of the change in phase composition of the powders, XRF measurements indicate that there is no diVerence in vibrations have disappeared at 500 °C.This means that this initial complex is no longer present above 500 °C. Peaks composition between the powders prepared according to the peroxide method (x=0) and the oxalate method (x=0.5).marked 3 (2500, 1750, 1500 and 1050 cm-1) are attributed to 1252 J. Mater. Chem., 1998, 8(5), 1251–1254Fig. 3 d-TGA of dried precipitate prepared with increasing amounts of oxalate, x (0.25, 0.5 and 1) theoretical weight loss for the proposed complex corresponds to the weight loss found with TGA. The temperatures of the above reaction mechanism steps correspond to the three maxima in the d-TGA spectrum for x=0.5 in Fig. 3. Considering the FTIR and XRD data these temperatures can be regarded as onset temperatures for these reactions. The maxima in Fig. 3 found for lower concentrations of oxalate are a mixture of the maxima found for the decomposition of the peroxo–oxalate structure and those maxima found in the thermal decomposition of the peroxo complex as pro- Fig. 2 Infrared spectra of the decomposition of the peroxo–oxalate complex (temperature intervals 20 °C), x=0.5. 1: water, 2: complex, 3: duced with the peroxide method (x=0) and described below.7 BaCO3, 4: Ti(OH)x BaTiO2(O2) 3H2O CA 300 °C BaTiO2(O2) H2O+2H2O CO vibrations in (barium) carbonate.13,14 The intensity of these signals first increases and reaches a maximum at 600 °C, BaTiO2(O2) H2O CA 500 °C BaTiO2(O2) H2O after which the intensity decreases again.This indicates that the intermediate formed after decomposition of the initial complex consist at least partially of BaCO3. The vibrations BaTiO2(O2) CA 750 °C BaO2+TiO2 (rutile)�BaTiO3+1/2O2 corresponding to BaCO3 are still present at 800 °C. Finally, signal 4 (3550 cm-1) corresponds to OH vibrations of Powder morphology and densification Ti(OH)x.15 This signal arises at temperatures above 700 °C. XRD analyses indicate the onset of crystalline perovskite The presence of ammonium oxalate also has an influence on formation at 800 °C, therefore the intermediate containing the the powder morphology.An increase in the oxalate concen- BaCO3 formed from the initial complex starts to decompose tration leds to a decrease in particle size (see Table 1).The at this temperature. smaller agglomerate size as determined from light scattering FTIR analysis at 800 °C still shows the presence of BaCO3. studies after ultrasonic treatment results in an increase of the This was confirmed by room temperature XRD measurements, green density as shown in Table 1.where perovskite as well as BaCO3 signals were found after In Fig. 4 a micrograph is given of a calcined powder prepared heating to 800 °C. XRD analysis of a powder calcined at 900 °C with the peroxo–oxalate method (x=0.5). It is clear that the showed a 100% perovskite crystal structure. Note that the powder consists of agglomerates, which in their turn consist absence of barium carbonate at 900 °C could not be confirmed of aggregates, with a size comparable to the size determined by FTIR measurements because the maximum temperature from light scattering measurements (1 mm).Within these aggrefor the high-temperature FTIR equipment used is 800 °C. gates smaller particles are visible. The average aggregate sizes Using the combined data of TGA, FTIR and XRD measure- of the various powders are given in Table 1.ments the following structure is proposed for the precipitating TEM-EDX measurements are used to study the homogeneity complex (x=0.5): BaTi0.91Zr0.09O2(C2O4) 3H2O. The thermal of the material. TEM-EDX revealed that the composition of decomposition mechanism of this complex is described below, the particles remained constant for 10 selected particles, which a mechanism which is closely related to the decomposition is an important indication that the powder is homogeneous. mechanism for the complex formed with the oxalate method.5 Dilatometer experiments on an isostatically pressed sample revealed a dense (95% rel.density) sample at a temperature of BaTiO2(C2O4) 3H2O+O2 CA 250 °C [BaCO3 TiO2 H2O] 10 °C.The maximum densification rate was at approximately 1180 °C. A sample heated for 10 h sintered at 1400 °C had a +CO2+2H2O grain size of 62 mm. High sintering temperatures are necessary [BaCO3 TiO2 H2O] CA 500 °C [BaCO3 TiO2]+H2O Table 1 EVect of the oxalate concentration on the aggregate size and green density [BaCO3 TiO2] CA 800 °C BaTiO3+CO2 oxalate average aggregate ratio, x size/mm green density (%) In these reactions Zr is left out for reasons of simplicity, it is 0 7 55 assumed that in this case Zr reacts in the same manner as Ti. 0.125 5 58 It can be seen that the complex formed in the peroxo–oxalate 0.25 4 59 complex (x=0.5) is very similar to the complex formed with 0.375 3 60 the peroxide method (x=0), namely BaTi0.91Zr0.09O2- 0.5 1 67 (C2O4) 3H2O instead of BaTi0.91Zr0.09O2(O2) 3H2O.The J. Mater. Chem., 1998, 8(5), 1251–1254 1253amount of oxalate in the process is x=0.5. The peroxide method (x=0) and the oxalate method result in the formation of a second phase. The presence of oxalate during the process ensures complete stoichiometric precipitation and an increasing amount of oxalate results in an increase of the green density. The precipitating complex in the peroxo–oxalate process is BaTi0.91Zr0.09O2(C2O4) 3H2O. The thermal decomposition of this complex is described.Finally, a calcined powder prepared using the peroxo–oxalate method (x=0.5) is homogenous and sinterable. The authors are indebted to Philips Forschungslaboratorien, Aachen, Germany for financial support.Special thanks are due to Dr D. Hennings and Dr O. Steigelmann. References 1 S. Iwaya, H. Masumura, Y. Midori, Y. Oikawa and H. Abe, US Patent, 4,404,029, 1983. 2 D. Hennings and A. Schnel, J. Am. Ceram. Soc., 1982, 65, 539. 3 S. M. Neirman, J.Mater. Sci., 1988, 23, 3973. 4 W. S. Clabaugh, E. M. Swiggard and R. Gilchrist, J. Res. Natl. Bur. Stand., 1956, 56, 289. 5 M. Stockenhuber, H. Mayer and J. A. Lercher, J. Am. Ceram. Soc., 1993, 76, 1185. 6 H. Yamamura, A. Watanabe, S. Shirasaki, Y. Moriyoshi and M. Tananda, Ceram. Int., 1985, 11, 17. 7 G. PfaV, Z. Chem., 1988, 28, 76. Fig. 4 SEM of calcined powder, magnification 25.000, x=0.5 8 G. PfaV, J.Mater. Sci. L ett., 1990, 8, 1145. 9 G. PfaV,Mater. Sci. Eng. B, 1995, 33, 156. 10 G. PfaV,Mater. L ett., 1995, 24, 393. to obtain the large grain sizes. Dielectric measurements show 11 G. PfaV, T hermochim. Acta, 1994, 237, 83. an er value of 27000 at a temperature of 90 °C. At 70 and 12 D. Hennigs, Int. J. High T echnol. Ceram., 1987, 3, 91. 110 °C the relative permittivity has 20% of its maximum value. 13 G. Busca, V. Buscaglia, M. Leoni and P. Nanni, Chem. Mater., 1994, 6, 955. 14 The Sadtler Standard spectra, Sadtler Research Laboratories, Conclusions Philadelphia, USA. The peroxo–oxalate method with suYcient oxalate results in single-phase perovskite, BaTi0.91Zr0.09O3, the optimum Paper 8/01466C; Received 20th February, 1998 1254 J. Mater. Chem., 1998, 8(5), 1251–1254

ISSN:0959-9428    

DOI:10.1039/a801466c

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Hydrothermal synthesis and characterization of a new aluminium vanadium oxide hydroxide Al2(OH)3(VO4) Brigitte Pecquenard, Peter Y. Zavalij and M. Stanley Whittingham*† Chemistry Department and Material Research Center, State University of New York at Binghamton, Binghamton, NY 13902-6000, USA A new aluminium vanadium oxide hydroxide analog of the mineral augelite has been prepared by hydrothermal treatment of a 15152.65 aqueous solution of V2O5, Al(NO3)3·9H2O, and TMAOH at 200 °C for 5 days.X-Ray powder diVraction data show that this phase crystallizes with monoclinic symmetry with space group C2/m. The structure was solved by using direct methods and then full profile Rietveld refinement was carried out; it’s cell parameters are: a=13.5634(2) A ° , b=8.2267(2) A ° , c=5.31232(9) A ° and b=112.741(1)° in space group C2/m.The structure contains clusters of edge sharing AlO6 octahedra and AlO5 trigonal bipyramids. These clusters are joined together by VO4 tetrahedra. steps and 15 s per step. The TGA was obtained on a Perkin- Introduction Elmer model TGA 7, the FTIR on a Perkin-Elmer 1600 series, In a constant search for new compounds which could be used and the electron microprobe on a JEOL 8900.as cathodic materials for advanced lithium batteries, we focussed our research on new metastable vanadium oxides. Results and Discussion The structure of the oxide lattice and a possible addition of a second metal are critical to obtain lithium insertion and The resulting powder has a light brown color.The electron reversible reaction. Indeed, a recent study on a Fe0.11V2O5.16 microprobe picture (Fig. 1) shows that the morphology of this compound obtained by the sol–gel process has shown that the new compound is well formed spherical particles, with an presence of iron(III ) ions in the orthorhombic host lattice average diameter close to 70 mm. It seems that these particles induced a real improvement of its electrochemical properties result from a germination/growth process, with a nucleation compared to V2O5.1 The use of ‘chimie douce’ allows us to point at the center of each particle.The X-ray diVraction obtain new materials which are sometimes impossible to be pattern indicates the presence of sharp lines. Thermal graviformed by using solid state reactions at high temperature.metric analysis of the compound under nitrogen (Fig. 2) shows Among these, the hydrothermal approach under mild con- a total weight loss of about 13.1% appearing in one step ditions, using organic templates such as tetramethylammonium essentially due to loss of OH groups. From TGA and X-ray ion was found to be particularly adept at forming new data, the general formula of the compound was found to be structures.2–17 It is noticeable that the template is often Al2(OH)3(VO4), the tetramethylammonium not being incorretained in the structure, as for TMA0.17V2O5, TMAV3O7,18 porated in the structure. The infrared spectrum (Fig. 3) con- TMAV4O105,9 or TMAV8O20.19 Recently, a study achieved by firmed the absence of tetramethylammonium ion which is Zhang et al.17 indicate the possible synthesis of iron and zinc usually characterized by three bands at 945, 1386 and double vanadium oxides using TMA.Here we report the synthesis of a new aluminium vanadium oxide analog of the mineral augelite Al2(PO4)(OH)320 by using tetramethylammonium ion as a template. Experimental The title compound was prepared by hydrothermal treatment by mixing V2O5 and Al(NO3)3·9H2O powder from Johnson and Matthey with 25% tetramethylammonium hydroxide (TMAOH) aqueous solution from Alfa in a 15152.65 molar ratio.Typically, 5 g of V2O5, 10.3 g of Al(NO3)3·9H2O and 26.5 g of TMAOH were mixed together and the initial pH of the resulting solution was 3.21. Then, the solution was transferred to a 125 ml Teflon-lined autoclave (Parr bomb), sealed and reacted hydrothermally for 5 days at 200 °C.At shorter and longer reaction times second phases were formed, the compositions of which have not yet been determined. The resulting light brown powder was filtered, washed with distilled water and dried in air. The pH of the solution after reaction was 6.32, higher than that of the initial mixture.X-Ray powder diVraction was performed using Cu-Ka radiation (l=1.5418 A° ) on a Scintag XDS2000 h–h diVractometer. The data were collected from 17° 2h to 90° 2h with 0.03° 2h Fig. 1 Electron microprobe picture of Al2(OH)3(VO4) †E-mail: stanwhit@binghamton.edu J. Mater. Chem., 1998, 8(5), 1255–1258 1255Fig. 2 Thermal gravimetric analysis of Al2(OH)3(VO4) under nitrogen Fig. 4 Calculated X-ray diVraction pattern resulting from final Rietveld refinement (thin line), experimental data (dotted line) and a diVerence plot (on the bottom) Table 1 Crystallographic data for Al2(OH)3(VO4) compound Al2(OH)3(VO4) crystal system monoclinic space group C2/m (no. 12) a/A ° 13.5634(2) b/A ° 8.2267(2) c/A ° 5.31232(9) b/° 112.741(1) cell volume, V/A° 3 546.68(3) calculated density, Dc/g cm-3 2.672 absorption coeYcient, m/cm-1 191.7 radiation, l(Cu-Ka)/A ° 1.54178 diVractometer Scintag XDS2000 Fig. 3 Infrared spectrum of Al2(OH)3(VO4) indexing method Ito software CSD mode of refinement full profile 1492 cm-1.21 All the bands are observed in the range 2h (max) 90° 450–1050 cm-1 which correspond to the MMO bonds domain number of refined parameters 29 (21 atomic) (M being a metallic atom).The precise assignment of IR peaks number of reflections 87 R(I) 0.062 is however not obvious. Only characteristic frequency ranges R(prof ) 0.086 are known for ‘condensed’ and ‘isolated’ AlMO octahedra and Rw (prof ) 0.105 tetrahedra.22 Peak positions are really dependent on the structure. After thermal treatment under air, the compound is decomposed into V2O5 and AlVO4.The powder diVraction Table 2 Atomic coordinates and thermal parameters for pattern was indexed with a monoclinic symmetry and the Al2(OH)3(VO4) space group C2/m using the Ito method from the CSD Software.23 The cell parameters of this new compound are: atom x/a y/b z/c Uiso/A° N a=13.5634(2) A ° , b=8.2267(2) A ° , c=5.31232(9) A ° and b= Al(1) 0 0.1911(5) 0.5 0.78(1) 4 112.741(1)°.Integrated intensities of 87 peaks were used in Al(2) 0.1824(3) 0 0.4637(8) 1.01(1) 4 direct methods to solve the structure. The powder X-ray V 0.3510(2) 0 0.1210(5) 0.88(6) 4 diVraction patterns of the observed and calculated data after O(1) 0.2463(7) 0 0.229(2) 1.438 4 Rietveld refinement are shown in Fig. 4. The final refinement O(2) 0.2962(6) 0 -0.242(2) 1.438 4 was achieved by using the CSD program and the results are O(3) 0.4234(5) 0.1672(7) 0.233(1) 1.438 8 O(4)* 0.9171(6) 0 0.295(1) 1.438 4 collected in Table 1.This gave R(F2)=0.062 and R(profile)= O(5)* 0.1065(5) 0.1788(6) 0.358(1) 1.438 8 0.086. The calculated density of this new compound is 2.672 g cm-3. The atomic positions, the selected interatomic aO(4) and O(5) are bonded to hydrogen.distances and the bond angles are given in Tables 2 and 3. The structure of this new aluminium vanadium hydroxide contains two types of polyhedra around aluminium (Fig. 5). of a continuous network. The projection of the structure along the b-axis (Fig. 7) shows that the groups of aluminium Al(1) belongs to an octahedron with four OH groups and 2 O atoms.The cation–anion distances in the polyhedron vary polyhedra are concentrated in layers parallel to the 00l plane, and separated by d(001). It also highlights the existence of from 1.821 A ° to 1.995 A ° with an average value of 1.896 A ° . The other aluminium, Al(2), is surrounded by five neighbors: two tunnels which contain hydrogen from the OH groups. The structure is very close to that described for the mineral O atoms and 1 OH group located in the basal plane and the other two OH groups quasi-perpendicular to this plane.The augelite Al2(OH)3(PO4). The cell parameters for the augelite structure are a=13.124 A ° , b=7.988 A ° , c=5.066 A ° and cation–anion distances in this trigonal bipyramid vary from 1.721 A ° to 2.193 A ° with an average value of 1.84 A ° . 2 Al(1) b=112.25°.20 The substitution of phosphorus by vanadium induces an increase of the cell parameters along the three and 2 Al(2) are bound together, sharing OH–OH edges. These aluminium clusters are connected to vanadium tetrahedra by directions. The array of the diVerent polyhedra is similar, but the anion–cation distances in some polyhedra are slightly sharing an oxygen apex (Fig. 6), which induces the existence 1256 J. Mater. Chem., 1998, 8(5), 1255–1258Table 3 Selected interatomic distances and bond angles for Al2(OH)3(VO4) VO4 tetrahedron VMO(1) 1.73(1) VMO(2) 1.778(8) VMO(3) 1.661(6) (×2) O(1)MVMO(2) 108.0(4) O(1)MVMO(3) 109.3(3) (×2) O(2)MVMO(3) 109.2(3) (×2) O(3)MVMO(3) 111.8(3) Al(1) octahedron Al(1)MO(3) 1.821(6) (×2) Al(1)MO(4) 1.995(6) (×2) Al(1)MO(5) 1.873(7) (×2) O(3)MAl(1)MO(3) 100.4(3) O(3)MAl(1)MO(4) 92.7(3) (×2) O(3)MAl(1)MO(5) 91.0(3) (×2) O(3)MAl(1)MO(4) 164.3(3) (×2) O(3)MAl(1)MO(5) 93.0(3) (×2) O(4)MAl(1)MO(5) 95.5(3) (×2) Fig. 6 Projection of the structure of Al2(OH)3(VO4) along the c-axis. O(5)MAl(1)MO(5) 173.8(3) Al(1), Al(2) and vanadium respectively occupy dark grey octahedra, medium grey trigonal pyramids and light grey tetrahedra.Al(2) trigonal bipyramid Al(2)MO(1) 1.770(1) Al(2)MO(2) 1.721(9) Al(2)MO(4) 2.193(9) Al(2)MO(5) 1.758(6) (×2) O(1)MAl(2)MO(2) 97.3(4) O(1)MAl(2)MO(5) 99.3(4) (×2) O(2)MAl(2)MO(4) 90.4(3) O(4)MAl(2)MO(5) 76.7(3) (×2) O(5)MAl(2)MO(5) 113.6(3) O(1)MAl(2)MO(4) 172.3(4) O(2)MAl(2)MO(5) 119.9(4) (×2) Fig. 7 Projection of the structure of Al2(OH)3(VO4) along the b-axis kinetic control which may exhibit new structures and properties.A new aluminium vanadium oxide hydroxide Al2(OH)3(VO4) analog of the mineral augelite has been synthesized hydrothermally and structurally characterized. This was Fig. 5 Representation of an aluminium polyhedral group connected achieved by mixing vanadium pentoxide, aluminium nitrate to a vanadium tetrahedron with the atom-labelling scheme.The and tetramethylammonium hydroxide. The last, used as a aluminium polyhedral group is composed of two Al(1) octahedra and template, is not retained in the structure after reaction. This two Al(2) trigonal pyramids bound together by sharing OH–OH edges. three dimensional structure is built up of clusters of aluminium octahedra and trigonal pyramids sharing edges.These clusters are connected together by vanadium tetrahedra. The structure diVerent. The Al(1) octahedron is quite similar with cation–anion distances which range from 1.826 A ° to 1.983 A ° of the mineral augelite Al2(OH)3(PO4), which is very similar, exhibits smaller cell parameters because of the shorter tetra- and average 1.891 A ° . The Al(2) trigonal bipyramid is less distorted in the augelite structure.In this case, the cation–anion hedral PMO bonds. Furthermore, this new compound as well as the mineral augelite is one of the few compounds to contain distances range from 1.750 A ° to 2.054 A ° and average 1.833 A ° which correspond to four short bonds, longer than in our aluminium in a trigonal pyramidal environment. A study by 27Al fast MAS NMR spectroscopy is now in progress.compound, and a long bond, which is shorter than in our compound. The average distance in the phosphate tetrahedron There are numerous aluminium phosphate hydroxide structures, but Al2(OH)3(VO4) is the first related vanadate. There is 1.519 A ° compared with 1.706 A ° for the vanadate, which induces a bigger unit cell, 546.7 A ° 3 for the vanadate and is just one other report of an aluminium vanadate hydroxide, but it is hydrated with the formula Al6V10(OH)12O28·29H2O, 491.5 A ° 3 for the phosphate.which indicates a decavanadate cluster.24 Conclusion This work was partially supported by the National Science Foundation through grant DMR-9422667. The authors are Hydrothermal reactions occur naturally in the crust of the earth, aiding in the formation of minerals.They therefore indebted to Bill Blackburn for his contribution to the use of the electron microprobe. might provide a viable route to synthesize materials under J. Mater. Chem., 1998, 8(5), 1255–1258 125712 D. Riou and G. Ferey, J. Solid State Chem., 1996, 124, 151. References 13 L. F. Nazar, B. E. Koene and J.F. Britten, Chem. Mater., 1996, 1 S. Maingot, R. Baddour, J. P. Pereira-Ramos, N. BaYer and 8, 327. P. Willmann, J. Electrochem. Soc., 1993, 140, L158. 14 Y. Zhang, J. R. D. DeBord, C. J. O’Connor, R. C. Haushalter, 2 M. S. Whittingham, J. Li, J. Guo and P. Zavalij, ‘Hydrothermal A. Clearfield and J. Zubieta, Angew. Chem., Int. Ed. Engl., 1996, Synthesis of New OxideMaterials using the T etramethyl Ammonium 35, 989.Ion’, in Soft Chemistry Routes to NewMaterials,Mater. Sci. Forum, 15 A. Muller, H. Reuter and S. Dillinger, Angew. Chem., Int. Ed. Engl., Trans Tech Publications Ltd., Nantes, France, 1994, 152–153, 99. 1995, 34, 2328 . 3 P. Zavalij, J. Guo, M. S. Whittingham, R. A. Jacobson, 16 Y. Zhang, R. C. Hausalter and A. Clearfield, Inorg. Chem., 1996, V.Pecharsky, C. Bucher and S. J. Hwu, J. Solid State Chem., 1996, 35, 4950. 123, 83. 17 F. Zhang, P. Zavalij and M. S. Whittingham, Mater. Res. Bull., 4 J. Guo, P. Zavalij and M. S. Whittingham, Chem. Mater., 1994, 1997, 32, 701. 6, 357. 18 T. Chirayil, E. A. Boylan, M. Mamak, P. Y. Zavalij and 5 M. S. Whittingham, J. Guo, R. Chen, T. Chirayil, G. Janauer and M. S. Whittingham, Chem.Commun., 1997, 33. 19 T. Chirayil, P. Y. Zavalij and M. S. Whittingham, J.Mater. Chem., P. Zavalij, Solid State Ionics, 1995, 75, 257. 1997, 7, 2193. 6 P. Zavalij, M. S. Whittingham, E. A. Boylan, V. K. Pecharsky and 20 T. Araki, J. J. Finney and T. Zoltai, Am. Mineral., 1968, 53, 1096. R. A. Jacobson, Z. Kristallogr., 1996, 211, 464. 21 G. L. Bottger and A. L. Geddes, Spectrochim. Acta, 1965, 21, 1701. 7 T. Chirayil, P. Zavalij and M. S. Whittingham, Solid State Ionics, 22 P. Tarte, Spectrochim. Acta, Part A, 1967, 23, 2127. 1996, 84, 163. 23 L. G. Akselrud, P. Y. Zavalij, Yu. N. Grin, V. K. Pecharsky, 8 T. Chirayil, P. Zavalij and M. S. Whittingham, J. Electrochem. B. Baumgartner and E.Wolfel, Mater. Sci. Forum, 1993, 133, 335. Soc., 1996, 143, L143. 24 A. A. Ivakin, M. V. Kruchinina, L. V. Chashchina, T. A. Denisova 9 E. A. Boylan, T. Chirayil, J. Hinz, P. Y. Zavalij and and O. V. Koryakova, Russ. J. Inorg. Chem., 1991, 36, 12. M. S. Whittingham, Solid State Ionics, 1996, 90, 1. 10 D. Riou and G. Ferey, J. Solid State Chem., 1995, 120, 173. 11 D. Riou and G. Ferey, Inorg. Chem., 1995, 34, 6250. Paper 7/09151F; Received 22nd December, 1997 1258 J. Mater. Chem., 1998, 8(5), 1255–1258

ISSN:0959-9428    

DOI:10.1039/a709151f

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials High temperature phase segregation of a new host for Er3+ upconversion: Cs3Tl2Cl9 Lukas Kamber, Philipp Egger, Bernhard Trusch, Rudolf Giovanoli and Ju�rg Hulliger* Department of Chemistry and Biochemistry, University of Berne, Freiestrasse 3, CH-3012 Berne, Switzerland Cs3Tl2Cl9 is one of the few air stable low-phonon host lattices of interest to Er3+ upconversion.Solution growth at ambient temperature demonstrated that water and a number of other molecular liquids do not yield Er3+ doped crystals. Optical and diVerential scanning calorimetry measurements revealed a phase segregation at 310 °C which is reversible in the presence of 1 atm Cl2. Growth from molten salt solutions or gas phase deposition techniques is hence restricted to a deposition temperature of less than 300 °C.Preliminary results show that molten ZnCl2 may be used as a flux to obtain lanthanide doped single crystals or epitaxial layers. the site symmetry for the dopant is 3m, a non-centrosymmetric Introduction site being a requirement for a high absorption cross section. High optical storage density devices and displaying appli- Single crystals of Cs3Tl2Cl9 have been obtained from aquecations demand compact and eYcient laser sources in the ous solution.7 However, lanthanide ions strongly coordinate visible spectral range.One implementation may be resonant to the oxygen of water molecules. It was therefore not possible conversion of infrared to visible light by solid state devices, to grow Er3+ doped Cs3Tl2Cl9 crystals from H2O solutions so-called upconversion (UC).1 For UC to take place, appro- near 300 K.The ultimate alternative is high temperature priate host lattices are doped with diVerent lanthanide ions, growth either from a flux or from the vapour. Both routes mainly Er, Tm, Pr or Ho. Recent investigations mainly used need knowledge on the thermal stability field of Cs3Tl2Cl9.fluoride and oxide crystals as well as heavy metal glasses as Solid state UC devices demonstrated so far are single crystals host materials.2 As the reduction of the phonon energies of the or waveguiding glass fibers grown at high temperatures.2 In host significantly slows down non-radiative multiphonon relax- particular, single crystalline waveguides allow the combination ation3,4 and thus increases UC eYciencies, there is considerable of the advantages of bulk crystals (narrow bands leading to interest in low-phonon UC materials.Because lowering phonon high gain coeYcients) with those of the fiber geometry (high energies can be achieved by introducing heavier ligands, chlor- pump intensities along the whole absorption length). High ide and bromide compounds have been investigated, showing quality waveguiding layers are most eVectively grown by liquid pronounced UC eYciencies.5,6 Due to the high sensitivity of phase epitaxy (LPE),8 pulsed laser deposition9 or high energy most of these materials to air and moisture, applications such ion implantation.10 A first fluoride LPE-grown UC waveguide as those mentioned before are chemically demanding.We have has been reported recently.11 A waveguide of Cs3Tl2Cl9 may found Cs3Tl2Cl9 to be one of the only low-phonon UC hosts be achieved by growing a lanthanide doped epitaxial layer on which is really stable in air and can be doped with, for example, a Cs3 Tl2Cl9 substrate. Here, too, knowledge of the thermal Er3+.6 Cs3Tl2Cl9 crystallizes in the rhombohedral space group behaviour of this potential host or substrate material is basic R39c featuring two face-sharing distorted octahedra (Fig. 1).7 to any further improvement. Tl3+ ions are situated in the centre of the octahedra, leading The thermal behaviour of powdered Cs3Tl2Cl9 samples in to short Tl3+–Tl3+ distances within a dimer (d#4 A° ) and open and closed systems has been investigated by Richter.12 facilitating energy transfer in the case of erbium.Furthermore, Following this analysis, Cs3Tl2Cl9 undergoes phase segregation at 301 °C, i.e. it decomposes into Cs2TlITlIIICl6+CsCl+Cl2 with an intermediate product identified as Cs11Tl6Cl29 (see ref. 12, p. 59). A weight loss of 6.2% has been found by thermogravimetry on unsealed samples. Compared to Cs3Tl2Cl9, a significantly simpler powder diVraction pattern was measured at temperatures above 301 °C.The Fm3m symmetry was attributed to Cs2TlITlIIICl6. Raman spectroscopy demonstrated a reversible phase segregation. In this work we present a thermal analysis relying on single crystals of Cs3Tl2Cl9 being kept under Cl2 (ca. 1 atm) in closed glass capillaries. Experimental conditions for the phase segregation study have been chosen as close as possible to parameters needed for high temperature single crystal and epitaxial growth of Cs3Tl2Cl9 to follow.Experimental Due to the toxicity of thallium compounds, care had to be Fig. 1 Dimeric structure of [Tl2Cl9]3- (ref. 7) taken. When handling powders or samples at elevated temp- J. Mater. Chem., 1998, 8(5), 1259–1262 1259eratures, they were kept in a special dry box (used for Tl The ampoules were glued to a thin glass plate with silver paste in order to improve the thermal contact.compounds only) or in sealed ampoules, respectively. Synthesis of Cs3Tl2Cl9 (ref. 13) Results and Discussion Wet TlCl3 (4 g; hygroscopic) and 5 g CsCl were dissolved While heating up to 330 °C DSC showed a strong endothermic separately in 10 ml distilled H2O, poured together and stirred.peak around 310 °C (10.6±0.65 kJ mol-1) accompanied by A white, finely spread precipitate was immediately obtained. several weaker endothermic signals in the range of 180–250 °C CsCl was added in excess to bring all TlCl3 to reaction. The (Fig. 3). Crystals tempered above 200 °C under N2 (48 h) did precipitate was recrystallised by addition of 70 ml distilled not show such spurious signals. In any case, there were no H2O, heating to 60 °C and slow cooling in a previously heated endothermic peaks around 100 °C, indicating that the crystals dewar. Transparent, needle shaped crystals were finally do not represent a hydrate or show water bubble inclusions.obtained which could be filtered oV and dried at 80 °C under Upon cooling an exothermic peak around 280 °C was observed.vacuum for several hours. Interesting thermal features were observed by orthoscopy: sealed crystals were positioned under crossed polars to yield Single crystal growth maximum extinction. Several bright spots indicated defects and strained areas [Fig. 4(a)]. On heating up to 180–250 °C Single crystals (Fig. 2) were grown using the DT method.14 An these areas dimmed and some of them came to complete excess of Cs3Tl2Cl9(s) and a saturated aqueous Cs3Tl2Cl9 extinction. Upon further heating, the entire crystal volume lit solution were filled into a specially designed tube. The tube up around 310 °C [Fig. 4(b) and (c)]. This phenomenon started was inclined at an angle of 30° to the horizontal.A DT of 1.5 °C was suYcient to induce constant convection, transporting saturated solution to the upper part of the tube where cooling led to supersaturation and subsequent nucleation. Nucleation could be promoted, i.e. controlled by use of a Peltier cooling element which was placed in the upper part of the tube.15 dT pulses allowed the supersaturation to be increased for a short time of 20–30 s.By X-ray diVraction we determined the growth direction to be the c-axis of the R39c system. The lattice constants are c=18.27 A ° and a=12.82 A ° .7 DiVerential scanning calorimetry DSC measurements were carried out using a Mettler Toledo DSC 25. The heating/cooling rates were 2 °C min-1. Crystals of ca. 10 mg were placed into sealed aluminium crucibles under N2.Fig. 3 DSC measurement of a Cs3Tl2Cl9 single crystal from room temperature to 330 °C Gandolfi X-ray photographs These were taken using an UB/71 OYcine Elettrotecnica Di Tenno instrument. A typical exposure time for single crystals was 4 h at 900W Cu-Ka1 radiation. Photographs at elevated temperatures were obtained by heating crystals sealed in Lindemann type capillaries (normal or quartzlass) under ca. 800 mbar of Cl2 by means of a heated air flow. The exposure time for rotated samples was typically 9 h at 900W Cu-Ka1. Additional photographs of single or polycrystalline samples were taken with the Gandolfi camera without rotating the sample, with sample exposure for 2 h. Orthoscopic investigation Orthoscopic investigations were performed using a Linkham THMS 600 heatable microscope table and a Leitz Orthoplan microscope with tenfold magnification.The samples were mounted in small Pyrex ampoules under N2 or Cl2 atmosphere. Fig. 4 Orthoscopy of the phase transition close to 310 °C (dashed lines Fig. 2 Single crystals of Cs3Tl2Cl9 grown from aqueous solution by indicate crystal edges). (a) T<310 °C; (b) start of the phase segregation around 310 °C; (c) end of the phase segregation.the DT method14 1260 J. Mater. Chem., 1998, 8(5), 1259–1262from crystal edges and defects, which did not come to complete present pattern could not be attributed to either Cs2TlITlIIICl6, Cs11Tl6Cl29 or CsCl, proposed by Richter (the products of extinction in the first phase of heating. From there, the illuminated sections spread over the entire crystal.No diVer- phase segregation of powder samples).12 The release of Cl2 at the phase change complicates the DSC ence was noticed whether working under N2 or Cl2 atmosphere. Upon cooling and about 30 °C below the transformation measurements on cooling, because Cl2 may react with the aluminium pan. However, at 280 °C an exothermic peak temperature (heating cycle), crystals showed some increase of local extinction, albeit not complete.supports reversibility of the transition under Cl2. Thermodynamical calculations16 applied to TlCl(s) and Optical inspection carried out in addition to orthoscopy monitored an increasing absorption (yellow colour) TlCl3(s) show the preferred formation of Tl+ at high temperatures, thus indicating the presence of at least one Tl+ containing above 200 °C.Around 300 °C, the crystal sharply turned yellow–brown. compound after the phase segregation. The sharp colour change to yellow around 310 °C may therefore be attributed Preliminary Gandolfi measurements were acquired without rotation of the crystal, showing a typical single crystalline to the formation of a mixed valence compound containing Tl3+ and Tl+.pattern at T<310 °C replaced by a powder pattern above 310 °C. DiVraction of rotated samples taken above the trans- Finally, strain as a source of defects mentioned in the paragraph on orthoscopy can be excluded, as relaxation of ition temperature produced photographs showing only a few lines (Table 1). Samples cooled to room temperature recovered strain would lead to exothermic peaks in the DSC.The observed relaxation temperatures varied with the crystals lines typical of Cs3Tl2Cl9. Taking into account all the observations mentioned above, examined. We assume that these areas act as nucleation centres for the phase segregation behaviour observed by optical means. the endothermic peak and the loss of birefringence around 310 °C can be attributed to a segregation reaction, transforming Compared to the transition temperature reported for powdered samples of Cs3Tl2Cl9,12 the single crystal transition single crystalline Cs3Tl2Cl9 into a composite of one or more new crystalline phases, at least one of them being yellow.On temperature is shifted by about 10 °C. This can be explained by the presence of fewer defects and a smaller surface.cooling, Cs3Tl2Cl9 is recovered in form of a polycrystalline, opaque material. However, this phase segregation was found to be reversible only if carried out in a closed system of either a small volume (under N2) or under Cl2 (ca. 1 atm). It may be Conclusions and preliminary results on solvent that Cl2 gas is released in the course of this reaction.Note systems that the recovery of the Cs3Tl2Cl9 phase required non-reactive vessels, such as Pyrex or quartz glass. In the case of normal Single crystals of Cs3Tl2Cl9 could be grown to optical quality and dimensions that allowed phase transition phenomena to Lindemann capillaries a completely diVerent and also unknown diVraction pattern was obtained. be tracked, indicating significant deviations from results gained on powder data.12 As the high temperature diVraction data revealed only a few lines, we conclude that predominantly one solid phase is Investigations on single crystals are indicative of phase segregation at 310 °C. At this temperature, an as yet unknown formed, other minor contributions being below the detection limit.The small number of lines in the high temperature new phase is formed.However, there is strong evidence that Cl2 and Tl+ are involved in the reaction. Within experimental diVraction pattern indicate a new phase of high symmetry. The Table 1 d-Values (A ° ) of the new and so far unknown high temperature phase, measured by a Gandolfi camera (further columns contain powder data of products reported by Richter;12 there is no agreement between the data of the first column with d-values reported by Richter) new high-temp. Cs3Tl2Cl9 (Ref. 18) Cs11Tl6Cl29 a-Cs2TlTlCl6 (Ref. 12) b-Cs2TlTlCl6 (Ref. 12) CsCl (Ref. 20) phase (300 K) (Ref. 19) (low-temp.)/(high-temp.) (300 K) 7.0562 (m) 6.4105 (m) 6.3551 (w) 4.4192 (m) 4.2322 (w) 4.120 (m) 4.2305 (m) 3.9893 (s) 3.9414 (w) 3.8951 (s) 3.9468 (s) 3.7944 (m) 3.8659 (m) 3.7024 (s) 3.7522 (m) 3.5308 (w) 3.5364 (w) 2.917 (s) 2.8934 (m) 2.8372 (s) 2.8158 (m) 2.7544 (s) 2.7925 (m) 2.7741 (m) 2.7914 (m) 2.67 (m) 2.7066 (m) 2.7331 (m) 2.59 (s) 2.5562 (w) 2.5695 (w) 2.5226 (m) 2.29 (m) 2.3564 (m) 2.380 (w) 2.24 (s) 2.2525 (w) 2.2181 (m) 2.2534 (w) 2.2793 (m) 2.2096 (s) 2.2092 (m) 2.2367 (m) 2.1375 (m) 2.1402 (w) 2.1489 (w) 2.1205 (m) 2.1045 (w) 1.9952 (m) 1.9465 (w) 1.9738 (m) 2.062 (w) 1.8972 (m) 1.8739 (m) 1.8510 (s) 1.8576 (m) 1.65 (w) 1.8442 (m) 1.844 (w) 1.59 (s) 1.7171 (w) 1.7145 (w) 1.7514 (w) 1.7655 (w) 1.683 (m) 1.6772 (w) 1.6643 (w) 1.7307 (w) 1.457 (w) 1.34 (w) 1.6469 (w) 1.6718 (w) 1.32 (s) 1.5358 (w) 1.5496 (m) 1.26 (w) 1.4190 (w) 1.5465 (w) 1.5412 (w) 1.4920 (w) 1.374 (w) (s), (m), (w) denote strong, medium and weak, respectively.J. Mater. Chem., 1998, 8(5), 1259–1262 1261error, the transition was reversible at ca. 1 atm Cl2 pressure. We thank Prof. T. Peters for the use of the Linkham microscope, Prof. T. Armbruster for the use of a Gandolfi camera, These important new results imply that bulk and epitaxial and Dr. C. Ba�rlocher for the use of the program TREOR.growth of Cs3Tl2Cl9 will have to be carried out at temperatures This work was supported by the Swiss National Science below T#300 °C and under Cl2 atmosphere. Foundation (project 20–43116.95) and the Priority Program The evident approach to bulk and epitaxial growth would ‘Optics II’ (project no. 232) of the Swiss Board of the Federal then be the use of a suitable low temperature solvent or high Institutes of Technology.temperature flux under Cl2 atmosphere. In the course of this work several solvents (DMSO, acetonitrile, liquid NH3, etc.) were tested. All of them converted the Cs3Tl2Cl9 single crystals References into a white mush, probably due to the soft Tl3+ ion binding 1 M. C. Brierley, J. F. Massicott, T. J. Whitley, C. A. Millar, soft Lewis bases.Recrystallisation from liquid NH3 yielded an R. Wyatt, S. T. Davey and D. Szebesta, BT T echnol. J., 1993, unknown powder pattern. Cs3Tl2Cl9 was insoluble in ethanol 11, 128. and methanol. Commonly used inorganic acids, HCl( l ) and 2 R. M. Macfarlane, J. Phys. IV (suppl.), 1994, 4, C4–289. some other solvents (Cl3PO) either yielded no significant 3 L. A. Riseberg and H.W. Moos, Phys. Rev., 1968, 174, 429. solubility or led to redox reactions. We therefore turned to an 4 J. M. F. van Dijk and M. F. H. Schuurmans, J. Chem. Phys., 1983, exploration of several relatively low melting fluxes: rhodanides 78, 5317. 5 M. P. Hehlen, K. Kra�mer, H. U. Gu� del, R. A. McFarlane and and nitrates were not suitable due to coordination to Tl3+. R. N. Schwartz, Phys.Rev. B, 1994, 49, 12475. SnCl2 systems were oxidised to SnCl4 under Cl2 here. 6 P. Egger, P. Rogin, T. Riedener, H. U. Gu� del, M. S. Wickleder and A mixture of Cs3Tl2Cl9 and ZnCl2 melting at 230 °C turned J. Hulliger, Adv. Mater., 1996, 8, 668. out to be a suitable system, showing suYcient solubility for 7 J. L. Hoard and L. Goldstein, J. Chem. Phys., 1935, 3, 199.both Cs3Tl2Cl9 and ErCl3. However, care has to be taken in 8 P. Rogin and J. Hulliger, J. Cryst. Growth, 1997, 179, 551. 9 D. S. Gill, A. A. Anderson, R. W. Eason, T. J. Warburton and order to avoid precipitation of ZnCl2·2CsCl.17 Further D. Shepherd, Appl. Phys. L ett., 1996, 69, 10. investigations on the use of the ZnCl2 flux in crystal growth 10 S. J. Field, D. C. Hanna, A. C.Large, D. B. Shepherd, experiments are in progress. A. C. Tropper, P. J. Chandler, P. D. Townsend and L. Zhang, An alternative approach would be physical vapour depos- Electron. L ett., 1991, 27, 2375. ition. In preliminary experiments, CsCl, ErCl3 and Cs3Tl2Cl9 11 P. Rogin and J. Hulliger, Opt. L ett., 1997, 22, 1701. 12 R. Richter, PhD thesis, Albert-Ludwigs-Universita�t, Freiburg, were evaporated separately in a nitrogen flow, carried over a 1993.substrate crystal of Cs3Tl2Cl9 held below 300 °C. Powder 13 J. H. Pratt, Z. Anorg. Chem., 1895, 9, 23. diVraction patterns of the materials deposited showed lines 14 J. Hulliger, Angew. Chem., Int. Ed. Engl., 1994, 33, 143. which were diVerent from those of Cs3Tl2Cl9, with the yellow 15 O. Ko� nig, P. Rechsteiner, B. Trusch, C. Andreoli and J. Hulliger, colour indicating the presence of Tl+. This could well be due J. Appl. Cryst., 1997, 30, 507. to Cl2 loss. In order to avoid reduction, the Cs3Tl2Cl9 source 16 I. Bahrin, T hermodynamical Data of Pure Substances, VCH Verlagsgesellschaft mbH, Weinheim, 1989. should therefore be kept under a Cl2 flow. Thermodynamical 17 B. F. Markov, I. D. Panchenko and T. G. Kostenko, Ukr. Khim. calculations using data given in ref. 16 and applied to TlCl3 Zh., 1956, 22, 290. and TlCl show that the reduction of TlCl3 to TlCl is exergonic 18 JCPDS database, card no. 44–716. at a sublimation temperature of about 400 °C in a N2 atmos- 19 G. Thiele and R. Richter, Z. Kristallogr., 1993, 207, 142. phere. In contrast, TlCl3 is thermodynamically favoured under 20 JCPDS database, card no. 5–607. a Cl2 atmosphere of about 5 atm to promote the formation of the Cs3Tl2Cl9 phase. Paper 7/09190G; Received 23rd December, 1997 1262 J. Mater. Chem., 1998, 8(5), 1259&ndas

ISSN:0959-9428    

DOI:10.1039/a709190g

出版商:RSC    

年代:1998

数据来源: RSC

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J O U R N A L O F C H E M I S T R Y Materials Synthesis of carbon nanotube–Fe-Al2O3 nanocomposite powders by selective reduction of diVerent Al1.8Fe0.2O3 solid solutions Ch. Laurent, A. Peigney and A. Rousset L aboratoire de Chimie desMate�riaux Inorganiques, ESA CNRS 5070, Universite� Paul-Sabatier, F 31062 T oulouse cedex 4, France Al1.8Fe0.2O3 solid solutions have been prepared as amorphous, g (cubic) and a (corundum) phases.The oxides have been reduced in a H2–CH4 gas mixture at 900 or 1000 °C, giving rise to composite powders containing alumina, a- and c-Fe, Fe3C and diVerent forms of carbon including nanotubes, thick tubes and spheroidal particles. The powders have been investigated using a combination of chemical analysis, X-ray diVraction, Mo�ssbauer spectroscopy, scanning and transmission electron microscopy, thermogravimetric analysis and specific surface area measurements.Using the stable form (corundum) of Al1.8Fe0.2O3 as starting material favours the formation of carbon nanotubes compared to the other forms of carbon. This could partly result from the fact that the metal nanoparticles formed upon reduction of the a solid solution, which act as a catalyst for CH4 decomposition and possibly nanotube nucleation, are smaller than when using amorphous or g solid solutions.Moreover, the crystallization of these latter compounds during the reduction in some way provokes the entrapment of carbon within the oxide grains. The nanotubes, most of which are less than 10 nm in diameter, are arranged in bundles several tens of micrometers long.shown19 that when using a hydrogen–hydrocarbon gas mixture Introduction instead of pure hydrogen for the reduction of an a-Al1.9Fe0.1O3 The discovery of carbon nanotubes1 triggered a worldwide solid solution, the pristine Fe nanoparticles formed in situ research eVort devoted to improving their synthesis,2–19 to upon reduction of the very homogeneously dispersed surface determining their structure20–24 and to calculating and measur- Fe3+ ions are active at a size adequate for the catalytic growth ing their physical properties.25–33 Particularly, regarding the of nanotubes.The resulting carbon nanotube–Fe-Al2O3 commechanical properties, the Young’s modulus of multi-walled posite powder contains a huge amount of single- and multicarbon nanotubes has been calculated to be up to 1.4 times walled carbon nanotubes, which have diameters in the that of a graphite whisker, i.e.about 1 TPa,34 and values 1.5–15 nm range. The nanotubes are arranged in bundles derived from thermal vibration experiments performed on smaller than 100 nm in diameter which may be up to 100 mm several multi-walled carbon nanotubes in a transmission elec- long.It has been calculated that the total bundle length in 1 g tron microscope are in the 0.4–3.7 TPa range.35 Moreover, of this powder is equal to about 100 000 km. Indeed, the metaltheir flexibility is remarkable36 and the bending may be fully oxide grains are uniformly covered by a web-like network of reversible up to a critical angle value as large as 110° for bundles and the powder is so densely agglomerated that it single-walled nanotubes.37 retains the shape of the reduction vessel when transferred to a Carbon nanotubes are most often prepared by arc-discharge storage box.between carbon electrodes in an inert gas atmosphere.1,2 Since the only metal particles active for the catalytic Catalytic species such as iron or cobalt can be used during the decomposition of the hydrocarbon are those located at the arc-discharge to improve both the quantity and the length of surface of the matrix grains, the active-to-inactive particle ratio tubes and to favour the formation of single-shell nanotubes.3–8 could be increased using powders with a high specific surface However, with this process, carbon nanotubes are obtained as area such as poorly crystallized or uncrystallized oxide solid a mixture with several other carbon forms, including much solutions. The aim of this work is to determine the influence amorphous carbon.Thus, purification has to be carried out of the crystallization level of an Al1.8Fe0.2O3 solid solution on and the yield of nanotubes is no more than 2%.38 Recently, the formation of carbon nanotubes, in order to increase the ‘ropes’ of single-walled carbon nanotubes were obtained by an tube yield, to maximize their length and to minimize their adapted laser-ablation technique.9,10 An alternative method is diameter, with the ultimate objective of obtaining a composite the catalytic decomposition of hydrocarbons on small metal powder suitable for the preparation of dense composite mateparticles (Fe, Co, Ni, Cu), which produces carbon filaments rials that may benefit from the exceptional properties of the including some authentic nanotubes.11–19 The minimal tube carbon nanotubes. diameter that can be achieved in this way is that of the catalytic particles. In order to maximize the nanotube yield, several Experimental authors investigated the influence of the temperature and of the nature of both the catalyst and the treatment atmos- The appropriate amounts of (NH4)2(C2O4) 2H2O, phere.13,39,40 Using a zeolite-supported Co catalyst, Ivanov Al(NO3)3 9H2O, and Fe(NO3)3 9H2O were mixed in an et al.14 and Hernadi et al.41 reported carbon tubes only 4 nm aqueous solution heated at 60 °C.The obtained clear solution in diameter and tubes of length 60 mm, but they point out that was cooled to room temperature and rapidly added to an the longest tubes are also the thickest. alcoholic medium, in which precipitation of the mixed In previous works, we have prepared metal-oxide nanocom- ammonium oxalate (NH4)3[Al0.9Fe0.1(C2O4)3] nH2O occurrposite powders by selective reduction in hydrogen of oxide ed immediately. After filtering, washing and oven drying, the solid solutions.42–47 In these materials, the metal particles (Cr, oxalate was finely ground and decomposed at 400 °C for 2 h.Fe, Co, Ni and their alloys) are generally smaller than 10 nm This powder will be denoted as C400 hereafter. Samples of the in diameter and are located both inside and at the surface of obtained C400 batch were calcined in air for 2 h, either at 850 °C (specimen C850) or at 1100 °C (powder C1100).The the matrix grains (Al2O3, Cr2 O3, MgO, MgAl2O4). We have J. Mater. Chem., 1998, 8(5), 1263–1272 1263three diVerent oxide powders were reduced in a H2–CH4 gas [Fig. 1(c)]. These results are in agreement with an earlier work48 showing that the thermal decomposition at 400 °C of mixture (6 mol% CH4) during 4 h at 900 °C or at 1000 °C (specimens denoted C400/R900, C400/R1000 etc.), giving rise a mixed oxalate (NH4)3[Al0.9Fe0.1(C2O4)3] gives rise to an Xray amorphous Al1.8Fe0.2O3 solid solution, which upon calci- to the carbon–metal-oxide composite powders.The materials were studied using scanning and transmission nation at the appropriate temperature crystallizes into the g (cubic) or a (corundum) structure.electron microscopy (SEM and TEM), X-ray diVraction (XRD) using Co-Ka radiation (l=0.17902 nm) and 57Fe Mo� ssbauer SEM observations show that the powders are made up of 10–20 mm agglomerates of submicronic or nanometric primary spectroscopy. The Mo�ssbauer spectra were recorded at room temperature with a constant acceleration spectrometer using a grains. The specific surface area measured for each oxide powder (Sss; Table 1) is consistent with the crystallization level 50 mCi 57Co (Rh) source; the spectrometer was calibrated by collecting at room temperature the spectrum of a standard Fe revealed by XRD analysis: 80 m2 g-1 for the amorphous Al1.8Fe0.2O3 solid solution (C400), 30 m2 g-1 for the poorly foil and the center shift (CS) values quoted hereafter are with reference to this standard.The composite powders were oxid- crystallized g-Al1.8Fe0.2O3 powder (C850) and 3.1 m2 g-1 for the a-Al1.8Fe0.2O3 powder (C1100). ized in air at 850 °C in order to eliminate all or part of the carbon, as required for the specific surface area study. The specific surfacethe oxide powders (Sss), of the nano- Nanocomposite powders composite powders obtained after reduction (Sn) and of the X-Ray diVraction. The XRD patterns of the nanocomposite powders oxidized at 850 °C (Son) were measured by the BET powders prepared by reduction of the solid solutions at 900 method using N2 adsorption at liquid N2 temperature.We or 1000 °C are shown in Fig. 2. In C400/R900, we observe used a Micromeritics FlowSorb II 2300 apparatus that gives wide peaks accounting for the presence of diVerent forms of a specific surface area value from one point (i.e. one adsorbant transition alumina (g, c, h) and only traces of a-Al2O3. In pressure) and requires calibration. For a given powder, we addition, a-Fe and Fe3C (cementite) are detected. A wide peak successively measured Sss, Sn and Son using the same caliwhich could correspond to the distance between graphene bration.The reproducibility of the results is in the±5% range, layers (d002=0.34 nm) is also detected. Since neither the (hk0) corresponding to ±10% for DS=Sn-Son, which accounts for nor the other (hkl) reflections, which would have much smaller the quantity of carbon nanotubes as detailed later in this intensities for nanotubes as well as for graphite,33 are detected, paper.Furthermore, we checked that the specific surface area it is not possible from this XRD pattern to discriminate measurements performed by the present method are in good between graphite and other graphenic forms of carbon such agreement with those performed using a multipoint as nanotubes.Thus we have labelled this peak Cg (Cgraphene). Micromeritics Accusorb 2100 E apparatus. The carbon content The same phases are present in C400/R1000, but the intensities was determined by flash combustion. The oxidation of the of the peaks representing a-Al2O3, a-Fe and carbon are much nanocomposite powders was investigated by thermogravihigher. The XRD patterns of C850/R900 and C850/R1000 are metric analysis (TGA) in flowing air (heating rate 1 °C min-1).similar to those of C400/R900 and C400/R1000, respectively. The main diVerence is the variation in the proportions of the Results and Discussion diVerent forms of alumina on the one hand and in the proportions of a-Fe and Fe3C on the other hand. Also, the Oxide powders (110) peak of a-Fe is distinctly wider in the C850-derived The XRD patterns of the C400, C850 and C1100 powders are composites, denoting that the decrease in specific surface area shown in Fig. 1. The pattern of specimen C400 is blank whereas of the oxide solid solution favours the formation of smaller that of the C850 powder presents wide peaks characteristics metal particles, as shown previously.42 Analysis of the XRD of an g-Al1.8Fe0.2O3 solid solution as well as traces of the patterns of C1100/R900 and C1100/R1000 clearly reveals the (104), (113) and (116) peaks of an a-Al2O3-type (corundum) presence of a-Al2O3, a-Fe and Fe3C.The proportions of phase [Fig. 1(a), (b)]. The g-phase peaks are very wide (more carbide and, notably, carbon are much lower than in the other than 10° 2h), reflecting a small crystallite size and a poor powders.c-Fe may be present in all or some powders, but is crystallization level. For the C1100 powder, all peaks are in extremely diYcult to detect on the XRD patterns because the good agreement with an a-Al1.8Fe0.2O3 solid solution c-Fe (111) diVraction peak (d111=0.208 nm) is probably masked by the base of the corundum (113) peak (d113= 0.2085 nm), and more so if Fe3C (d210=0.206 nm) is present as well.A FeAl2O4 spinel phase, which is known to form upon reduction of Al1.8Fe0.2O3 in pure H2 at temperatures lower than 1000 °C,44–47 is not detected in the R900 powders. However, one can not rule out its presence in small quantities, particularly in the C400/R900 and C850/R900 specimens, because the FeAl2O4 main diVraction peaks (d113=0.245 nm, d004=0.202 nm and d044=0.143 nm) may be masked by transition- alumina and a-Fe peaks.Carbon content. The carbon content measured in the nanocomposite powders (Cn; Table 1) is in the 1–20 wt.% range. Cn is much higher in the composites derived from the amorphous Al1.8Fe0.2O3 and g-Al1.8Fe0.2O3 phases than in those derived from the stable a-Al1.8Fe0.2O3 phase.This could result from the higher specific surface area of the former powders, which would allow more metal particles to nucleate and grow on the surface of the oxide grains and therefore to be active as catalysts for the decomposition of CH4. We also note that for a given starting powder, the carbon content is higher after reduction at 1000 °C than after reduction at 900 °C.According Fig. 1 XRD patterns of the oxide powders prepared at diVerent temperatures: (a) 400 °C; (b) 850 °C; (c) 1100 °C to Wagman et al.,49 the CH4 equilibrium content in H2–CH4 1264 J. Mater. Chem., 1998, 8(5), 1263–1272Table 1 Specific surface area (Sss: oxide solid solutions; Sn: nanocomposites; Son: nanocomposites oxidized at 850 °C; DS=Sn-Son) and carbon contents (Cn: nanocomposites; Con: nanocomposites oxidized at 850 °C) of the powders Sss Sn Son Cn Con DS DS/Cn oxide /m2 g-1 composite /m2 g-1 /m2 g-1 (wt.%) (wt.%) /m2 g-1 /m2 g-1 C400 79.5 C400/R900 23.5 20.0 13.4 0.65 3.5 26 C400/R1000 16.0 11.0 20.1 0.58 5.0 25 C850 30.5 C850/R900 44.1 39.3 8.9 0.30 4.8 54 C850/R1000 31.0 25.1 18.3 0.21 5.9 32 C1100 3.1 C1100/R900 6.7 4.1 1.65 0.00 2.6 155 C1100/R1000 8.4 3.5 6.2 0.00 4.9 79 that the reducibility of the Fe3+ ions substituting in the alumina lattice, and thus the formation of metal nanoparticles, are hampered by a decreasing specific surface area of the oxide powder44 and that consequently a 100 °C increase in reduction temperature has more impact with low specific surface area oxides. Mo�ssbauer spectroscopy.All composite powders except C400/R1000 were studied by Mo�ssbauer spectroscopy. Each spectrum was fitted assuming it is the sum of diVerent subspectra: a sextet accounting for ferromagnetic a-Fe, a sextet representing Fe3C and a singlet characteristic of non-ferromagnetic Fe (Fig. 3 and Table 2). It is noteworthy that neither Fe3+ nor Fe2+ ions were detected on the Mo�ssbauer spectra, even in the composites prepared from the C1100 oxides, i.e.the less easily reducible phase (a-Al1.8Fe0.2O3). This shows that the presence of CH4 in the reducing gas mixture tremendously favours the reduction of the iron(III ) ions substituting for aluminium in the corundum lattice: indeed, when using pure H2, Fe3+ contents of 35 and 65% are measured in C1000/R1000 and C1200/R1000 composite powders, respectively,44 from which one may estimate that the Fe3+ proportion would be of the order of 50% in a C1100/R1000 specimen. It has been shown50 that treatment at 1300 °C is necessary to fully reduce the Fe3+ ions substituting for aluminium in the corundum lattice to the metallic state.Fe2+ ions could account for the presence of a phase of the FeAl2O4 spinel type.44–47 As the detection limit in Mo�ssbauer spectroscopy is about 4%, a very small amount of spinel phase could indeed be present in the R900 composite powders while no corresponding pattern is seen in the spectra.44 However, considering the above observation that CH4 strongly favours the formation of metallic Fe, it is reasonable to assume that no residual Fe3+ and Fe2+ ions are present in the composite powders.Since Fe3C has two inequivalent crystallographic sites,51 the Mo�ssbauer parameters of the sextet accounting for Fe3C correspond to the average of the two Fe-site parameters one may obtain using two sextets for the fit. Our average values are in good agreement with those reported by Le Cae�r et al.52 for bulk Fe3C and by Bi et al.53 for Fe3C nanoparticles. The non-ferromagnetic phase of metallic Fe corresponding to the singlet could be either antiferromagnetic, paramagnetic or superparamagnetic a-Fe, or c-Fe.Paramagnetic a-Fe has Fig. 2 XRD patterns of the nanocomposite powders prepared by been observed in Fe/Ni-MgAl2O4 nanocomposite powders54 calcination (C) and reduction (R) treatments at diVerent temperatures: and superparamagnetic a-Fe was fresent in some (a) C400/R900; (b) C400/R1000; (c) C850/R900; (d) C850/R1000; Fe-Al2O3 specimens.44,45,47 However, the negative value of the (e) C1100/R900; (f ) C1100/R1000; a: a-alumina; O: transition alumina; +: Fe3C; *: a-Fe; Cg: corresponding to d002 in multi-walled nanotubes CS points towards the face-centered-cubic c-Fe phase.This and/or in graphite could reflect the formation of a c-Fe–C alloy, rather than pure Fe. It is interesting to note that Baker et al.55 have shown that gas mixtures is lower at 1000 °C (about 1%) than at 900 °C metallic Fe, and not cementite, is the active phase responsible (about 2%). Thus, the CH4 supersaturation level in the gas for the formation of carbon nanofilaments.In the present mixture used in the present study (6 mol% CH4) is much work, the proportions of the diVerent Fe species vary from higher at 1000 °C than at 900 °C and consequently more one powder to another, but analysis of the data does not allow carbon is deposited during the reaction at 1000 °C. Moreover, one to draw precise conclusions from their evolution. The very the increase in carbon content when the reduction temperature low value (2%) observed for Fe3C in the C400/R900 powder is changed from 900 to 1000 °C is higher when using solid solutions calcined at a higher temperature. This could reflect is probably a fit artefact and the relatively high proportion of J.Mater. Chem., 1998, 8(5), 1263–1272 1265in the composite powders reduced from the stable a solid solutions (C1100/R900 and R1000) and the latter one in the specimens prepared from the high specific surface area oxides, because this favours the formation of surface metal particles.Electron microscopy. Low-magnification SEM observations of each of the six composite powders show that the matrix grains (about 20 mm long) are covered by a web-like network of carbon filaments (Fig. 4), in line with earlier results.19 However, higher magnification SEM images [Fig. 5(a)–(f )] clearly reveal some diVerences depending on the calcination and/or the reduction temperature used for the preparation of the powders. In the C400/R900 powder [Fig. 5(a)], we observe long, tight bundles with a diameter smaller than 50 nm, as well as curved filaments both thicker (up to 100 nm in diameter) and shorter (about 1 mm or less), and some clusters of spheroidal grains in the 20–50 nm diameter range.The C400/R1000 powder [Fig. 5(b)] presents less carbon bundles and thick filaments but more spheroidal grains which seem to be larger than in the C400/R900 powder. Both the C850/R900 and C850/R1000 powders [Fig. 5(c), (d)] contain tight bundles and small spheroidal clusters in various quantities from one powder grain to another.Thus, the diVerence in spheroidal cluster quantities between the two images does not reflects a diVerence between the two powders. Images of the C1100/R900 and the C1100/R1000 powders [Fig. 5(e)–(f )] show more bundles but interestingly neither thick filaments nor spheroidal clusters have been detected on any observed matrix grain.We also observe homogeneously dispersed nanometric particles which could be Fe or Fe3C particles. TEM [Fig. 6(a)–(c)] and HREM [Fig. 6(d)–(f )] images of nanocomposite powders show details of the diVerent forms of carbon previously revealed by SEM observations. A thick tube,12 irregularly shaped, is shown in the C400/R900 powder Fig. 3 Mo�ssbauer spectra of the nanocomposite powders prepared by calcination (C) and reduction (R) treatments at diVerent temperatures: (a) C400/R900; (b) C850/R900; (c) C850/R1000; (d) C1100/R900; (e) C1100/R1000; I: ferromagnetic a-Fe; II: non-ferromagnetic c-Fe; III: Fe3C the same phase in the C850/R1000 powder is in line with the XRD results.The sextet accounting for a-Fe is the major component in all specimens and the proportion of the singlet representing c-Fe is in the 20–35% range.The intragranular Fe particles, which number extremely high in such materials, 42,45 are probably protected from alloying with carbon and crystallize in the stable a form. Moreover, surface particles that form an Fe–C alloy but grow to a large enough size undergo the c–a transformation upon cooling from the Fig. 4 SEM image showing the network of filaments on a grain of the C400/R1000 powder reduction temperature. The former cause could be preeminent Table 2 Room temperature Mo�ssbauer parameters of some nanocomposite powders [ferro: ferromagnetic; non-ferro: non-ferromagnetic; H: hyperfine field/kG; CS: center shift/mm s-1; C: half-line width/mm s-1; 2eQ: quadrupole shift/mm s-1, P: proportion (%) ferro a-Fe non-ferro Fe Fe3C specimen CS H C P CS C P CS H 2eQ C P C400/R900 0.00 328 0.16 73 -0.10 0.16 25 0.10 187 0.00 0.10 2 C850/R900 -0.01 333 0.20 49 -0.09 0.20 35 0.20 205 0.06 0.27 16 C850/R1000 -0.01 332 0.18 38 -0.09 0.19 20 0.19 205 0.03 0.21 42 C1100/R900 -0.02 328 0.22 42 -0.11 0.22 35 0.21 204 0.06 0.20 23 C1100/R1000 -0.01 329 0.21 55 -0.11 0.18 32 0.19 201 0.06 0.25 13 1266 J.Mater. Chem., 1998, 8(5), 1263–1272Fig. 5 High magnification SEM images of the nanocomposite powders prepared by calcination (C) and reduction (R) treatments at diVerent temperatures: (a) C400/R900; (b) C400/R1000; (c) C850/R900; (d) C850/R1000; (e) C1100/R900; (f ) C1100/R1000 [Fig. 6(a)]. Its inner and outer diameter are about 10 nm and observations of the same powder show several Iijima-type1 nanotubes [Fig. 6(d)], between 3 and 4 nm in diameter, made 50 nm respectively and the inside cavity is partly filled with elongated Fe and/or Fe3C particles. The short, thick filament up of two or three concentric layers (d002=0.34 nm) as well as Fe and/or Fe-carbide particles covered with three well crys- in the C400/R1000 powder [Fig. 6(b)] appears to be made up of several hollow forms, one of them containing an Fe and/or tallized graphene sheets [Fig. 6(e)]. An image of the C1100/R1000 powder [Fig. 6(f )] reveals a two shell carbon Fe3C particle. This filament is superimposed with a large Fe and/or Fe-carbide particle, and is connected to a matrix grain nanotube, 2.8 nm in diameter, superimposed with an alumina grain and Fe and/or Fe3C particles, the smaller ones (4–8 nm) by what seems to be a carbon thin film.In the C1100/R900 powder [Fig. 6(c)], we observe several bundles composed of being probably dispersed inside the oxide grain. A previous study42 on the selective reduction in H2 of Al2-2xFe2xO3 filaments less than 10 nm in diameter, a few nanometric irregularly shaped carbon species, and a lot of nanometric Fe (0<x0.2) has shown that the size and size distribution of the Fe particles formed upon reduction mainly depend on the and/or Fe-carbide particles, appearing black on the image.We have also observed that the spheroidal clusters (20–50 nm in mono- or biphasic nature (0<x0.1 or x>0.1, respectively) of the starting oxide and also from their specific surface area.diameter) detected by SEM, particularly in the C400/R1000 powder [Fig. 5(b)], are graphitic carbon nanoparticles,56 most Starting with a monophasic oxide such as the present Al1.8Fe0.2O3 compound favours a more monomodal size of them containing Fe and/or Fe-carbide particles. HREM J. Mater. Chem., 1998, 8(5), 1263–1272 1267Fig. 6 TEM (a–c) and HREM (d–f ) images showing diVerent forms of carbon revealed in the nanocomposite powders: thick tube in the C400/900 powder (a); thick, short filament in the C400/R1000 powder (b); bundles of nanotubes in the C400/900 powder (c); four nanotubes pointed by arrows in the C1100/R900 powder (d); Fe and/or Fe-carbide particles covered with three well crystallized graphene layers in the C1100/R900 powder (e); a nanotube partially superimposed with an alumina grain containing nanometric Fe and/or Fe3C particles in the C1100/R1000 powder (f ) distribution and a low specific surface area favours a smaller Al1.8Fe0.2O3 (C850) and a-Al1.8Fe0.2O3 (C1100) powders.The variations of S could result from the superimposition of two average size of the metal particles.phenomena: on the one hand, crystallization of the oxide and sintering of the primarains inside the agglomerates lead to Specific surface area measurements. Comparison of the specific surface areas (S) of the solid solutions (Sss; Table 1) a decrease of S when the reduction temperature is higher than the calcination temperature, i.e. when C400 and C850 powders with those of the nanocomposite powders (Sn; Table 1) shows that the reduction treatment induces a decrease of S when are reduced.On the other hand, the formation of carbon species, particularly carbon nanotubes and, to a much lesser amorphous starting powders (C400) are used, but that, in contrast, it induces an increase of S when starting from the g- degree, the formation of metal nanoparticles at the surface of 1268 J.Mater. Chem., 1998, 8(5), 1263–1272the matrix grains lead to an increase of S upon reduction of Ni-MgAl2O4 nanocomposite powders,43 it is inferred that the low-temperature weight gain (DTG peak in the 225–380 °C all the powders (C400, C850 or C1100). The former eVect outweighs the latter for the C400/R900 and C400/R1000 range) corresponds to the oxidation of the Fe and Fe-carbide particles located at the surface and in the open porosity of the powders; this is reversed for the C850/R900, C1100/R900 and C1100/R1000 composites.The two processes balance each oxide matrix. For the C1100/R900 and C1100/R1000 powders [Fig. 7(e), (f )], the high-temperature weight gain occurs in two other in the case of the C850/R1000 powder.This makes comparisons between the diVerent specimens steps (DTG peaks at 695 and 1165 °C, and 775 and 1180 °C respectively) as in Fe/Cr-Al2O357 nanocomposite powders. It diYcult and does not allow the derivation of precise data on the amount of nanotube formed upon reduction. Therefore, could similarly correspond to the oxidation of the intragranular metal particles. In contrast, for the other powders we oxidized the composite powders in air at 850 °C, in order to eliminate all carbon species, and we again measured the [Fig. 7(a)–(d)], the high-temperature weight gains are much smaller and only the first step is apparent (DTG peak in the specific surface area (Son; Table 1). In fact, as we will discuss below, TGA (Fig. 7) and carbon analyses of specimens oxidized 630–720 °C range).Two assertions can explain this point. Firstly, in the composite powders derived from the reduction at 850 °C (Con; Table 1) show that some carbon remains in some of these samples (C400/R900 or R1000, and C850/R900 of an amorphous Al1.8Fe0.2O3 solid solution (C400) or a poorly crystallized g-Al1.8Fe0.2O3 powder (C850), the Fe and Fe3C or R1000), but in an amount smaller than 10% of the initial content, which can be neglected.Since the oxidation tempera- particles are mainly located at the surface of the matrix grains and are thus oxidized at a low temperature. Secondly, a high- ture (850 °C) is lower than the reduction temperature (900 or 1000 °C), this treatment should not aVect the matrix and thus temperature weight loss takes place and probably masks the end of the first high-temperature weight gain.However it is Son is a good approximation of its specific surface area. We assume that the Fe2O3 particles formed upon oxidation of the obvious that above 1100 °C, the weight is constant [Fig. 7(a), (c)], which indicates that the second high-temperature weight surface Fe and Fe3C nanoparticles57 will be of a similar nanometric size and thus that the contribution of the oxide gain does not take place for these powders.The weight losses, which are supposed to represent the particles to Son will be roughly similar to that of the Fe and Fe3C nanoparticles to Sn. oxidation of carbon (free or combined with Fe), mainly occur between 350 and 650 °C and also between 850 and 1000 °C for As proposed elsewhere,19 the increase in specific surface area per gram of powder, DS=Sn-Son, essentially represents the some powders (but in a much smaller proportion) [Fig. 7(a)–(f )]. The DTG curves [Fig. 7( b), (d), (f )] show that quantity of nanotubes (more precisely of nanotube bundles) and the increase in specific surface area per gram of carbon, the first weight loss is generally made up of two or three steps which could represent the oxidation of diVerent forms of DS/Cn, can be considered as ‘quality’ data, a higher figure for DS/Cn denoting a smaller average tube diameter and/or more carbon. However, the steps are not well enough separated on the TGA and DTG curves to draw any conclusions.Further carbon in tubular form. Walker et al.58,59 have also reported specific surface areas in the 35–170 m2 g-1 range for carbon investigations are in progress to clarify this point.Comparisons can be made between the total weight loss filaments, which they correlate to a high degree of internal porosity because the geometrical surface area of their filaments measured in TGA and the carbon content (Table 1) on the one hand and between the total weight gain measured in TGA (100 nm in diameter and 1 mm long) does not exceed 15 m2 g-1.In contrast, the present electron microscopy observations have and the theoretical weight gain corresponding to the oxidation of all the Fe-containing species as Fe2O3 on the other hand. revealed important quantities of very long tubes of nanometric diameter, the geometrical surface area of which is of the order Weight losses are always smaller than the corresponding carbon contents, except for the C1100/R900 powder, because of several hundred m2 g-1.Other researchers60,61 have reported increases in specific surface area upon the catalytic gains and losses are generally superimposed. For the same reason, weight gains are always smaller than the theoretical formation of carbon nanofibers, which are in qualitative agreement with the present results. values, which have been calculated assuming that no residual Fe2+ or Fe3+ were present in the powders as indicated by On the one hand, whatever the initial solid solution (C400, C850 or C1100), DS is higher when the reduction is performed Mo�ssbauer spectroscopy.Note that for the C1100/R900 powder, a weight loss higher than the measured carbon content at 1000 °C than when it is performed at 900 °C, indicating that a higher CH4 supersaturation level enhances both the quantity could be considered as being a consequence of the cumulative errors on the two measurements.Thus, it appears that TGA of nanotubes and the carbon content (Cn; Table 1). On the other hand, for a given reduction temperature, there is no of the present composite powders can not give quantitative results for either the determination of the content of diVerent great diVerence in DS for the diVerent solid solutions (Table 1).Consequently, the ratio DS/Cn is much higher for C1100/R900 carbon forms or the reduction yield of the solid solutions. Interestingly, the high-temperature weight loss is detected or R1000 powders for which Cn is low; the C1100/R900 powder thus presents the higher quality figure (DS/Cn=155 m2 g-1).only for the C400- and C850-derived powders [Fig. 7(a)–(d)]. For the C1100-derived specimens [Fig. 7(e)–(f )], this weight Comparisons of DS and DS/Cn values between C850/R900 and R1000 or between C1100/R900 and R1000 powders (Table 1) loss either does not occur or is masked by the large weight gain arising from the oxidation of intragranular Fe particles. show that the increase in quantity obtained when the reduction temperature is increased from 900 to 1000 °C is detrimental to The carbon content was determined for each of the six powders after oxidation in air at 850 °C (Con; Table 1), the temperature the quality.at which the high-temperature weight loss phenomenon begins.A small quantity of carbon is measured for each of the four Thermogravimetric analyses. The oxidation of the nanocomposite powders was investigated by thermogravimetric analysis C400- and C850-derived powders, but it is much smaller than the observed high-temperature weight loss [Fig. 7(a), (c)]. For (TGA) in flowing air up to 1400 °C.The TGA curves [Fig. 7(a), (c), (e)] and the corresponding diVerential thermogravimetric the C1100/R900 and C11/R1000 powders, carbon was not detected, and therefore we infer that the weight loss is not (DTG) curves [Fig. 7( b), (d), (f )] show that the oxidation of the composites occurs in several steps, with both weight gains masked but rather does not occur. For the sake of comparison, TGA was performed on an and weight losses.For all powders, the first weight gain occurs at temperatures alumina powder previously calcined at 400 °C and then heattreated at 900 °C in the H2–CH4 gas mixture, as the C400/R900 lower than 500 °C and other gains at temperatures exceeding 600 °C. From the conclusions of previous works on the oxi- composite powder.Three steps are detected on the TGA and DGA curves (Fig. 8). At low temperature (DTG peak at dation of Fe/Cr-Al2O3 and Fe/Cr-Cr2O357 as well as Co- and J. Mater. Chem., 1998, 8(5), 1263–1272 1269Fig. 7 Thermogravimetry (TGA and DTG) curves measured in flowing air of the nanocomposite powders: C400/R900 and C850/R900 (a, b); C400/R1000 and C850/R1000 (c, d); C1100/R900 and C1100/R1000 (e, f ) 505 °C), we observe a small weight loss which probably of some carbon entrapped inside the alumina grains.It is noteworthy that the third step is more important for this accounts for the oxidation of crystallized form(s) of carbon located at the surface of the alumina grains. The second step powder than for the Fe-containing specimens, because no Fe particles can act as catalyst for CH4 decomposition and (DTG peak at 900 °C) is a weight gain, probably corresponding to an oxygen gain due to the establishment of alumina therefore preferentially interact with the resulting deposited carbon. In fact, the second and third phenomena may overlap stoichiometry during the beginning of the crystallization in the a-form.Indeed, before this oxidation, the thermal treatment in and thus the actual weight loss may be higher than determined by TGA.In spite of this, the total weight loss (5.93%) is higher the H2–CH4 gas mixture at 900 °C produced transition alumina, which is probably oxygen-deficient. The third step (DTG than the measured carbon content (5.49%). As pointed out above for the composite powders, the high-temperature weight peak at 1065 °C) is a weight loss, probably due to the oxidation 1270 J. Mater.Chem., 1998, 8(5), 1263–1272(C1100) is used, carbon seems to be obtained only as bundles of nanotubes, thin graphitic films at the surface of Fe and/or Fe3C particles, and combined with Fe in cementite. The consequences are a low carbon content (Cn; Table 1) and a surface of carbon which is rather high (DS; Table 1) and thus a rather high quality parameter (DS/Cn; Table 1). This could partly result from the fact that the metal nanoparticles formed upon reduction of the a solid solution, which act as catalyst for CH4 decomposition and possibly nanotube nucleation, are much smaller than when using amorphous or g solid solutions42 and therefore appear to strongly favour the catalytic deposition of carbon in the form of nanotubes, smaller than 10 nm in diameter and arranged in bundles several tens of micrometers long, corresponding to aspect ratios as high as 1000–10000.19 It should be noted that the macroscopic characteristics of the powders are repeatable provided the specific surface area of the starting solid solution is the same.Reducing a solid Fig. 8 Thermogravimetry (TGA and DTG) curves measured in flowing solution of a given structure but with a lower specific surface air of an alumina powder calcined at 400 °C and then heat-treated at area yields less carbon and thus lower values of Sn and DS. 900 °C in H2–CH4 gas mixture Influence of the reduction temperature. As shown by the above Mo�ssbauer spectroscopy results, 900 °C is a high enough loss is also higher than the carbon content measured, before this step occurs, in powders oxidized at 850 °C.A possible temperature, in the H2–CH4 atmosphere, to fully reduce to the metallic state the Fe3+ ions substituting in the diVerent explanation for these observations is that the flash combustion technique used to determine the carbon contents does not solid solutions.Thus the increase in both carbon content Cn and DS (Table 1) when the reduction temperature is raised allow the release of all the carbon entrapped inside the alumina grains, thereby decreasing its amount. from 900 to 1000 °C are not a consequence of a more complete reduction producing a higher number of catalytic particles, Similarly, in the composite powders in which the alumina matrix is not completely crystallized in the a-form (C400/R900 but result from the higher CH4 sursaturation level in the reducing atmosphere.However, a lower quality parameter and R1000, C850/C900 and R1000), some carbon is probably entrapped inside the transition-alumina grains and it is oxid- (DS/Cn; Table 1) is observed in the R1000 powders, notably when starting with an a-solid solution (C1100).ized only at high temperature. Thus a compromise has to be found between a high value of DS denoting a high quantity (C1100/R1000) and a high Influence of the form of the initial Al1.8Fe0.2O3 solid solution. These characterization results have shown that, depending on value of DS/Cn denoting a high quality (C1100/R900). Since a huge number of carbon nanotubes in the composite powder is the crystallization level of the initial Al1.8Fe0.2O3 solid solution, the reduction process in a H2–CH4 gas mixture can lead to desirable to reach a suYciently high volume fraction, suitable for a possible enhancement of the properties of materials made the formation of several carbon species in the composite powder. The microscopic observations can be directly corre- from these powders, a high reduction temperature (1000 °C or more) is to be preferred.We have yet to study the influence of lated with macroscopic results, particularly with specific surface area measurements. other synthesis parameters to enhance the quality of the obtained nanotubes. When amorphous (C400) or transition (C850) solid solutions are used, carbon species other than bundles of carbon nanotubes appear, including thick tubes, hollow carbon forms Conclusions and/or clusters of graphitic carbon nanoparticles, most of them containing Fe and/or Fe3C particles. The consequence is a An amorphous Al1.8Fe0.2O3 solid solution has been prepared from the precipitation and thermal decomposition of a mixed high carbon content (Cn; Table 1), but a moderate carbon surface (DS; Table 1) and thus a low quality parameter (DS/Cn; ammonium oxalate.Calcination in air of the amorphous oxide at the appropriate temperatures produced the g (cubic) and Table 1). Since these solid solutions have a relatively high specific surface area (Sss; Table 1), most of the Fe3+ ions are the stable a (corundum) forms of Al1.8Fe0.2O3. The three oxides have been reduced in a H2–CH4 gas mixture at 900 or at located at or near the surface of the grains and are easily reduced to the metallic state, giving rise to a very high number 1000 °C, giving rise to composite powders containing alumina, a- and c-Fe, Fe3C, and several forms of carbon, including of metal particles which in turn can easily coalesce at the surface of the oxide grains to form much larger particles.42 carbon nanotubes, thick tubes and clusters of graphitic carbon nanoparticles.Thus, the diameter of most of these metal particles may be too high for the catalytic formation of nanotubes, in agreement The powders have been studied using a combination of chemical analysis, X-ray diVraction, Mo�ssbauer spectroscopy, with Baker and Rodriguez12 who claim that only suYciently small (less than 20 nm in diameter) metal particles can lead to scanning and transmission electron microscopy, thermogravimatric analysis and specific surface area measurements.In nanotubes. Moreover, we have shown that upon heating in a H2–CH4 atmosphere, the crystallization of the amorphous or particular, we have made use of the specific surface area of carbon (DS, see text) as a representation of the quantity of g-Al1.8Fe0.2O3 solid solutions in some way provokes the entrapment ofbon within the oxide grains.This entrapped carbon nanotubes in a powder and we have considered the ratio of this value to the carbon content (DS/Cn) as a quality carbon, as well as other non-tubular forms of carbon, would probably deteriorate the mechanical properties of massive value, a higher figure for DS/Cn denoting a smaller average tube diameter and/or more carbon in tubular form.composites prepared from these powders. Consequently, amorphous (C400) or transition (C850) solid solutions can not The presence of CH4 in the reducing gas mixture has been found to markedly favour the reduction of the iron(III ) ions be retained as precursors of composite materials including carbon nanotubes.substituting for aluminium in the corundum lattice. The composite powders prepared from the amorphous or g solid In contrast, when a stable a-Al1.8Fe0.2O3 solid solution J. Mater. Chem., 1998, 8(5), 1263–1272 127123 D. Ugarte,Microsc.Microanal.Microstruct., 1993, 4, 505.solutions contain important quantities of non-tubular carbon, 24 L. A. Bursill, J. L. Peng and X. D. Fan, Philos. Mag. A, 1995, resulting in a poor value of the quality parameter DS/Cn. 71, 1161. Moreover, some carbon is entrapped within the oxide grains 25 S. J. Tans, M. H. Devoret, H. Dai, A. Thess, R. E. Smalley, upon the crystallization of alumina during the reduction step.L. J. Geerligs and C. Dekker, Nature (L ondon), 1997, 386, 474. 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ISSN:0959-9428    

DOI:10.1039/a706726g

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Crystal chemistry and physical properties of complex lithium spinels Li2MM¾3O8 (M=Mg, Co, Ni, Zn; M¾=Ti, Ge) Hiroo Kawai,a Mitsuharu Tabuchi,b Mikito Nagata,c Hisashi Tukamotoc and Anthony R. Westa aDepartment of Chemistry, University of Aberdeen, Meston Walk, Aberdeen, Scotland, UK AB24 3UE bOsaka National Research Institute, 1-8-31 Midorigaoka, Ikeda, Osaka 563, Japan cCorporate R&D Centre, Japan Storage Battery Company L imited, Nishinosho, Kisshoin, Minami-ku, Kyoto 601, Japan The spinels Li2MM¾3O8 (MM¾=MgTi, CoTi, CoGe, NiGe and ZnGe) are cubic with space group P4332.Simple crystal field theory qualitatively explains the distribution ofMover tetrahedral and octahedral sites: Ni occupies only octahedral sites, whereas Zn, Mg and Co show strong preference for tetrahedral sites. 153 cation ordering of Li/Mand M¾ occurs on the octahedral sites. The titanates undergo an order–disorder phase transition involving the octahedral cations at high temperatures, whereas the ordered phase is maintained until melting for the germanates. Solid solutions Li2-2XM1+3XM¾3-XO8 form at both sides of the Li2MM¾3O8 stoichiometry for the titanates; but there is no substantial range of solid solution for Li2ZnGe3O8 and Li2NiGe3O8.The occurrence of order–disorder phenomena and solid solutions in the titanates is attributed to the similarity in size of Li,Mand Ti, whereas the smaller Ge is less able to disorder with Li/M. Mis shown to be divalent from magnetic susceptibility measurements (for Co and Ni) with the support of conductivity data.The samples containing Co and Ni are paramagnetic down to 5 K. From impedance measurements on pellets with blocking electrodes, the main conductive species is deduced to be Li+: the activation energies for conduction are high, 0.55<DH/eV<2.14. Cyclic voltammograms show a set of reversible peaks at ca. 1.5 V vs. Li/Li+ for the titanates, attributed to the Ti3+/4+ couple, but no Li could be electrochemically extracted from either titanates or germanates up to 5 V vs.Li/Li+. The ideal spinel structure consists of a cubic close-packed reported on the line Li4Ti5O12–Zn2TiO4: these show 153 cation order for a small range of compositions close to Li2ZnTi3O8, array of anions, with one eighth of the tetrahedral and one half of the octahedral interstices occupied by cations, having but are disordered to either side.13 In order to better understand the crystal chemistry and the general formula A[B2]X4, where A is a tetrahedrally surrounded cation, B an octahedrally surrounded one, and X properties of this series of spinels, it is necessary to examine an anion.The crystal chemistry of binary spinel oxides such the distribution, valence state, and ordering of cations more asM2+M3+2O4 andM4+M2+2O4 has been studied in detail.1–4 accurately and to survey the formation of solid solutions for The cation distributions were only partially explained on the lithium spinels Li2MM¾3O8.In this paper, we focus on the basis of ionic bonding and were aVected as well by the compositions Li2MM¾3O8 (M=Mg, Co, Ni, Zn; M¾=Ti, Ge), individual site preference of cations.1–4 In contrast to binary and discuss (1) the cation distributions, (2) valence states of spinel oxides, the crystal chemistry of ternary spinel oxides is transition-metal ions, (3) order–disorder phase transitions, (4) less well understood.formation of solid solutions, and (5) magnetic, electrical and A considerable number of ternary lithium spinels with the electrochemical properties.composition Li2MM¾3O8 has been synthesised.5–12 Blasse9 determined approximate cation distributions for several of these: no systematic correlation with the individual site prefer- Experimental ence of cations was obtained, but the cation distributions were Starting materials were Li2CO3, MgO, CoO, NiO, ZnO, TiO2 partially accounted for by introducing an anion polarisation and GeO2, all reagent grade.Stoichiometric amounts of the eVect, in addition to the electrostatic energy term. dried reagents were mixed in an agate mortar, using acetone Complex spinel phases allow either tetrahedral or octahedral to form a paste, dried and fired in air, initially at 650 °C for a sites to be occupied by more than one kind of cation. The few hours to decarbonate and then at 800 °C for 12 h.The distributions over both sets of interstices tend to be random calcined powders were reground, pelleted and reacted between at high temperatures, while they are often ordered at low 900 and 950 °C for 12 h to 3 days, depending upon starting temperatures with lowering of the symmetry. 153 cation material and composition. The products were quenched in air ordering in octahedral sites is commonly observed at the from the reaction temperatures. composition Li2MM¾3O8.5–11 In Li2ZnTi3O8, for instance, Li Samples were initially examined for completeness of reaction and Ti show 153 cation ordering over two sets of octahedral and phase purity by powder X-ray diVraction (XRD) with a sites below 1150 °C at which temperature an order–disorder Philips Hagg Guinier focusing camera, Cu-Ka1 radiation. For phase transition occurs.13 lattice parameter determination and Rietveld refinement, data The composition Li2MM¾3O8 may be considered as a spinel were collected with a Stoe Stadi/P diVractometer in trans- composition on the line between the two binary compositions mission mode using a small linear position sensitive detector Li4M4+5O12 and M2+2M4+O4; the cation to anion ratio of resolution 0.02°, Cu-Ka1 radiation.A scan range of remains at 354 throughout, but 153 cation order is possible at 8<2h/°<113 in steps of 0.2° was used for Rietveld refinement. this composition. A complete range of spinel-like solid solution, Li2-2XZn1+3XTi3-XO8: -1/3X1 has been recently The data analysis was carried out with the Stoe software J.Mater. Chem., 1998, 8(5), 1273–1280 1273package: lattice parameter refinement was performed using the appearance of samples after each stage of a stepwise heating cycle. The possible formation of solid solutions LATREF, and Rietveld refinement using the pattern fitting structure refinement (PFSR) program.Powder neutron Li2-2XM1+3XM¾3-XO8 was checked at either side of the Li2MM¾3O8 stoichiometry (X=0), for compositions X=0.11 diVraction data were collected for Li2ZnGe3O8 on the Polaris medium resolution diVractometer at the UK spallation neutron and -0.067. source ISIS, Rutherford Appleton Laboratory. The crystal data, collected over the time-of-flight range 2.300–19.570 ms, Results and Discussion in the highest resolution, backscattering detectors were refined by the Rietveld method with the program TF14LS.Phase-pure spinels were obtained at the following composi- For conductivity measurements, pellets (8 mm diameter; tions: Li2MgTi3O8, Li2 CoTi3O8, Li2 CoGe3O8, Li2 NiGe3O8 2–3 mm thickness) were cold-pressed uniaxially at 150 MPa and Li2ZnGe3O8, and had the colours shown in Table 1.Their and sintered at 950 °C for 24 h, in order to increase their powder XRD patterns were indexed in the cubic space group mechanical strength. Gold paste electrodes were coated onto P4332 with lattice parameters, #8.18<a/A ° <#8.38, Table 2. opposite sides of the sintered pellets which were gradually The titanates had larger lattice parameters than the correheated to 800 °C to decompose the paste and harden the Au sponding germanates, consistent with the larger size of Ti4+ residue. Impedance measurements between 30 mHz and compared with Ge4+.15 Li2MgGe3O8 and Li2NiTi3O8 showed 1 MHz, used combined Solartron 1250/1286 and Hewlett- complicated XRD patterns which could not be indexed as Packard 4192 instrumentation. The data were analysed in single phase spinels; further work on them is in progress.three formalisms, the complex impedance (Z*), admittance Li2ZnTi3O8 was not prepared; data for it13 are listed in Tables 1 (Y *), and modulus (M*),14 using in-house software. and 2. Thermal analysis in air, heating/cooling rates of 10 °C min-1, used a Stanton Redcroft STA1500 simultaneous TG/DTA Structure refinement instrument with alumina as an inert reference for diVerential thermal analysis, DTA.Magnetic measurements were per- Structure refinement using the Rietveld method has been recently carried out by Ohtsuka’s group16,17 for the spinel formed under He from room temperature to 5 K, using a MB- 3 Shimazu Faraday balance and Quantum Design MPMS2 phases obtained in this study. The refinement was repeated on our diVraction data, using more extensive structure models SQUID magnetometer.The standard used for magnetic susceptibility calibration was Tutton’s salt, (NH4)2Mn(SO4)2 6H2O. within the space group P4332. The starting atomic positions were those given in ref. 16 and 17. The value of 0.05 was used Cyclic voltammograms were made over the potential range 1–5 V vs.Li/Li+ with scan speed 0.1 mV s-1, using a Hokuto- as the initial isotropic thermal parameters (Uiso) of all the atomic positions. The total occupancy of 8c, 4b and 12d cation Denco HA151 potentiostat galvanostat linked to a Hokuto- Denco HB111 function generator. A three-electrode cell was sites was fixed at unity; oxygens were situated at 8c and 24e sites with full occupancy. The eVect of varying the cation used.The working electrode consisted of Li2MM¾3O8, mixed with 5 mass% acetylene black (active material ) and 8 mass% distributions on 8c, 4b and 12d sites was examined initially by generating theoretical powder XRD patterns, using the starting binder made of poly(vinylidene fluoride) and N-methyl-2- pyrrolidone.The counter and reference electrodes consisted of atomic coordinates and thermal parameters, with the program THEO in the Stoe software package. Rietveld refinement was a strip of Li metal foil. The electrolyte was 1 M LiPF6 dissolved in either propylene carbonate or 50/50 vol.% mixture of then carried out for several likely models, to examine detailed occupancy of cation sites.ethylene carbonate and diethyl carbonate. Melting temperatures were determined approximately from Profile parameters were initially allowed to refine. After Table 1 Physical data for spinel phases Li2MM¾3O8 melting points DTA peaksa,b formation of s.s.c formation of s.s.c composition colour /°C /°C X=0.11 X=-0.067 Li2ZnTi3O8 white 1410d 1165d,e yesd yesd Li2ZnGe3O8 white 1000 none no no Li2MgTi3O8 white 1410 990e yes yes Li2CoTi3O8 green 1250 907e yes yes Li2CoGe3O8 purplish blue 980 none yes no Li2NiGe3O8 greenish sky 970 none no no aDTA peak temperatures on the heating cycle.Peak temperatures on cooling were 15–35 °C lower. bObserved from room temperature to ca. 50 °C below the melting points for the germanates, and at least up to ca. 1200 °C for the titanates. cs.s. denotes the solid solution Li2-2XM1+3XM¾3-XO8. dAfter Hernandez et al., ref. 13. eSharp peak. Table 2 Structural data for spinel phases Li2MM¾3O8 symmetry lattice parameter octahedral-site preference composition (space group) a/A ° cation distributionsa Y b energy for M2+c/kcal mol-1 Li2ZnTi3O8 cubicd (P4332d) 8.3710(2)d Li0.5Zn0.5[(Li0.5)Ti1.5]d 0.5 0 (Zn2+) Li2ZnGe3O8 cubic (P4332) 8.1961(5) Li0.5Zn0.5[(Li0.5)Ge1.5] 0.5 0 (Zn2+) Li2MgTi3O8 cubic (P4332) 8.3774(9) Li0.55Mg0.45[(Li0.45Mg0.05)Ti1.5] 0.45 Li2CoTi3O8 cubic (P4332) 8.3766(12) Li0.55Co0.45[(Li0.45Co0.05)Ti1.5] 0.45 2.1 (Co2+) Li2CoGe3O8 cubic (P4332) 8.2098(4) Li0.55Co0.45[(Li0.45Co0.05)Ge1.5] 0.45 2.1 (Co2+) Li2CuTi3O8 cubice 8.39e 0.3e 15.6 (Cu2+) Li2NiGe3O8 cubic (P4332) 8.1799(2) Li[(Ni0.5)Ge1.5] 0 22.8 (Ni2+) aIn A[(B0.5)B¾1.5]O4, A denotes tetrahedral 8c sites, B octahedral 4b sites, and B¾ octahedral 12d sites.bY changes in the formula Li1-YMY[(LiYM0.5-Y)M¾1.5]O4, 0Y0.5. cAfter McClure, ref. 2. dAfter Hernandez et al., ref. 13. eAfter Blasse, ref. 6. Y=0.3 for Li2CuTi3O8 has been recently confirmed by the Rietveld method.19 1274 J.Mater. Chem., 1998, 8(5), 1273–1280convergence, the atomic positions were refined in the order: obtained using powder XRD data, thus supporting the validity of refinement results using our powder XRD data. Structure heavy atoms (M and M¾), oxygen, and lithium; isotropic thermal vibration parameters were then refined in the same refinement parameters from neutron data are given for Li2ZnGe3O8 in Table 7.It is seen clearly from the comparison order. In some cases, isotropic thermal parameters for oxygen and sites containing high concentrations of lithium would not of the estimated error values for atomic positions and thermal parameters that neutron data oVer more precise refinement converge or became negative during refinement. These thermal parameters were then fixed at 0.05.When the model allowed results than XRD data. more than one kind of cation to occupy the same sites, their positions were refined with the same atomic coordinates and Structure descriptions and cation distributions thermal parameters. Of the various models tested, the model with minimum R values18 was used for the final stage of The refined atomic positions are in good agreement with those reported previously.16,17 The cation distributions may refinement.At this stage, thermal parameters were fixed at the refined values, and atomic positions and occupancy of cation be expressed by the formula, Li1-YMY[(LiYM0.5-Y)M¾1.5]O4 (0Y0.5). Li and M are distributed over tetrahedral 8c sites were allowed to refine, but with the total occupancy fixed at unity and the composition fixed at Li2MM¾3O8.In situations and octahedral 4b sites according to Y. Y therefore denotes the degree of cation mixing in tetrahedral sites. The refine- where the occupancy of cation sites was modified, atomic positions and temperature parameters were refined again, with ments revealed that M¾, i.e., Ti and Ge are situated only in octahedral 12d sites.For Li2NiGe3O8, Y=0; tetrahedral sites the occupancy fixed at the modified values. Therefore, the final parameter listings do not have estimated error data for the are occupied only by Li; Ni and Ge are located in octahedral sites with 153 order, Fig. 2. For Li2MgTi3O8, Li2 CoTi3O8 occupancies, but list estimated error values for the thermal parameters. The final structure refinement parameters are given and Li2CoGe3O8, the previous refinements16,17 concluded that Y=0.5, but our refinements showed minimum R values in Tables 3–7, with R values and bond lengths.The refined cation distributions are listed in Table 2. The final profile fit is when Y=0.45; Li and M are randomly distributed in tetrahedral sites, whereas Li/M and M¾ show 153 order in shown in Fig. 1 for Li2MgTi3O8 as an example. The final RP and RWP values for Li2NiGe3O8 were considerably greater octahedral sites. For Li2ZnGe3O8, Y=0.5; there is complete 153 order of Li and Ge in octahedral sites; Li and Zn are than those of the other phases. This is attributed to the inadequate profile fit probably caused by diYculty in describing distributed randomly in tetrahedral sites.Bond lengths calculated were appropriate for spinel oxides, the peak shape, owing to a possible interaction between Ni and Cu-Ka1 radiation. supporting the structure model refined. In situations where low-scattering Li fully occupies tetrahedral 8c sites, i.e., For Li2ZnGe3O8, Rietveld refinement was additionally carried out using powder neutron diVraction data, following the Li2NiGe3O8, Y=0, the 8c atomic coordinates and the corresponding LiMO bond lengths should be viewed with same refinement procedure; the structure refined to the same cation distribution and similar atomic coordinates to the one care, however, Table 6.Table 3 Structure refinement parameters and bond lengths for Li2MgTi3O8 from X-ray data. Space group P4332; a=8.3774(9) A° ; RP=2.41%; RWP=3.26% atom site x/a y/b z/c Uiso occupancy Li(1) 8c 0.998(2) 0.998(2) 0.998(2) 0.027(3) 0.55 Mg(1) 8c 0.998(2) 0.998(2) 0.998(2) 0.027(3) 0.45 Li(2) 4b 0.625 0.625 0.625 0.025(15) 0.9 Mg(2) 4b 0.625 0.625 0.625 0.025(15) 0.1 Ti 12d 0.125 0.3687(5) 0.8812(5) 0.023(1) 1 O(1) 8c 0.391(2) 0.391(2) 0.391(2) 0.013(7) 1 O(2) 24e 0.107(2) 0.129(1) 0.390(1) 0.021(3) 1 distance to O(1)/A ° distance to O(2)/A ° average distance to oxygen/A ° Li(1)/Mg(1) 2.02(2) 1.97(2) (×3) 1.98 Li(2)/Mg(2) 2.14(1) (×6) 2.14 Ti 2.02(2) (×2) 1.90(1) (×2) 1.97 1.98(2) (×2) Table 4 Structure refinement parameters and bond lengths for Li2CoTi3O8 from X-ray data.Space group P4332; a=8.3766(12) A ° ; RP=2.30%; RWP=3.05% atom site x/a y/b z/c Uiso occupancy Li(1) 8c 0.001(4) 0.001(4) 0.001(4) 0.013(6) 0.55 Co(1) 8c 0.001(4) 0.001(4) 0.001(4) 0.013(6) 0.45 Li(2) 4b 0.625 0.625 0.625 0.06(6) 0.9 Co(2) 4b 0.625 0.625 0.625 0.06(6) 0.1 Ti 12d 0.125 0.368(2) 0.882(2) 0.020(4) 1 O(1) 8c 0.375(7) 0.375(7) 0.375(7) 0.05 1 O(2) 24e 0.100(8) 0.127(4) 0.387(5) 0.05 1 distance to O(1)/A ° distance to O(2)/A ° average distance to oxygen/A ° Li(1)/Co(1) 1.83(7) 1.98(6) (×3) 1.94 Li(2)/Co(2) 2.12(5) (×6) 2.12 Ti 2.15(6) (×2) 1.83(4) (×2) 2.00 2.01(7) (×2) J.Mater. Chem., 1998, 8(5), 1273–1280 1275Table 5 Structure refinement parameters and bond lengths for Li2CoGe3O8 from X-ray data. Space group P4332; a=8.2098(4) A ° ; RP=1.88%; RWP=2.46% atom site x/a y/b z/c Uiso occupancy Li(1) 8c 0.005(1) 0.005(1) 0.005(1) 0.027(3) 0.55 Co(1) 8c 0.005(1) 0.005(1) 0.005(1) 0.027(3) 0.45 Li(2) 4b 0.625 0.625 0.625 0.05 0.9 Co(2) 4b 0.625 0.625 0.625 0.05 0.1 Ge 12d 0.125 0.3759(6) 0.8742(6) 0.016(1) 1 O(1) 8c 0.382(6) 0.382(6) 0.382(6) 0.03(1) 1 O(2) 24e 0.096(7) 0.123(1) 0.395(2) 0.015(8) 1 distance to O(1)/A ° distance to O(2)/A ° average distance to oxygen/A ° Li(1)/Co(1) 1.95(5) 2.00(4) (×3) 1.99 Li(2)/Co(2) 2.06(3) (×6) 2.06 Ge 1.99(5) (×2) 1.83(1) (×2) 1.91 1.90(6) (×2) Table 6 Structure refinement parameters and bond lengths for Li2NiGe3O8 from X-ray data.Space group P4332; a=8.1799(2) A ° ; RP=8.03%; RWP=12.14% atom site x/a y/b z/c Uiso occupancy Li 8c 0.018(8) 0.018(8) 0.018(8) 0.05 1 Ni 4b 0.625 0.625 0.625 0.012(4) 1 Ge 12d 0.125 0.3777(7) 0.8723(7) 0.013(1) 1 O(1) 8c 0.385(3) 0.385(3) 0.385(3) 0.01(1) 1 O(2) 24e 0.094(3) 0.127(2) 0.394(2) 0.012(5) 1 distance to O(1)/A ° distance to O(2)/A ° average distance to oxygen/A ° Li 2.17(7) 1.94(7) (×3) 2.00 Ni 2.08(2) (×6) 2.08 Ge 1.95(3) (×2) 1.82(2) (×2) 1.89 1.90(2) (×2) Table 7 Structure refinement parameters and bond lengths for Li2ZnGe3O8 from neutron data.Space group P4332; a=8.1961(5) A ° ; RP=2.49%; RWP=2.10% atom site x/a y/b z/c Biso occupancy Li(1) 8c 0.00488(14) 0.00488(14) 0.00488(14) 0.203(12) 0.5 Zn 8c 0.00488(14) 0.00488(14) 0.00488(14) 0.203(12) 0.5 Li(2) 4b 0.625 0.625 0.625 0.96(3) 1 Ge 12d 0.125 0.37464(4) 0.87536(4) 0.226(3) 1 O(1) 8c 0.38653(4) 0.38653(4) 0.38653(4) 0.302(8) 1 O(2) 24e 0.09906(4) 0.13104(5) 0.39490(4) 0.321(4) 1 distance to O(1)/A ° distance to O(2)/A ° average distance to oxygen/A ° Li(1)/Zn 2.008(1) 1.953(1) (×3) 1.967 Li(2) 2.1156(3) (×6) 2.1156 Ge 1.9619(4) (×2) 1.8416(5) (×2) 1.9008 1.8989(3) (×2) As shown in the following sections, magnetic susceptibility and conductivity data demonstrated that M is divalent for all the spinel phases obtained, i.e., the formulae are Li2M2+M¾4+3O8.Li and M have similar ionic radii15 and diVer in oxidation state only by unity.The electrostatic energies, i.e., the sum of the Coulomb and Born repulsion energies, are thus expected to be similar throughout the range 0Y0.5 in any one phase. The diVerence in Y between the spinel phases is then attributable mainly to either the diVerence in individual site preference of M or oxygen polarisation. McClure2 and simultaneously, Dunitz and Orgel3 calculated site preference energies of transition-metal ions using crystal field theory; the diVerence between the energy values for octahedral and tetrahedral stabilisation was used as a measure of the octahedral- site preference.It is seen from Table 2 that Y, i.e., the concentration of M in tetrahedral 8c sites decreases roughly with increasing octahedral-site preference energy of M, consistent with the expectations from crystal field theory.Zn2+ Fig. 1 Observed and diVerence powder XRD profiles for Li2MgTi3O8 is usually regarded as forming covalent bonds in tetrahedral 1276 J. Mater. Chem., 1998, 8(5), 1273–1280GeMO distances (12d sites) are all around 1.90 A ° , significantly smaller than both Li/MMO distances (4b sites), average 2.10 A ° and TiMO distances (12d sites), average 1.99 A ° .Thus, these size diVerences make less likely the disordering of Ge and Li/Mover the octahedral sites. By contrast, the small diVerence in size of the Ti and Li/M sites makes disordering easier. The disorder can be introduced, either by raising the temperature or by varying the composition by means of the solid solution mechanism, 3M=2Li+M¾.As listed in Table 1, all the titanates have melting temperatures higher than the germanates. The melting temperatures for the germanates are similar to the order–disorder phase transition temperatures for the titanates, in the temperature range 900–1000 °C. Magnetic properties Fig. 2 Projection of the spinel structure of Li2NiGe3O8 showing the Magnetic properties were examined for the samples containing 4b (shaded NiO6) and 12d (unshaded GeO6) octahedra. Each NiO6 Co and Ni.For all the samples, inverse molar susceptibility, shares edges with six GeO6: for the Ni octahedron indicated, four GeO6 octahedra: 1, 2, 3 and 5, are shown; the other two, 4 and 6, are xm-1, changes linearly with temperature in the range 5–293 K, above the plane of the projection. according to the Curie–Weiss law, xm-1=(T-h)/C, where h is the Weiss constant, and C is a constant, Fig. 3. Even at 5 K, magnetizations, M, are small, and increase linearly with increasing applied field, H, without showing any evidence of coordination,20 consistent with Y=0.5 for Li2ZnTi3O8 and spontaneous magnetization, Fig. 4. Li2CoTi3O8, Li2 CoGe3O8 Li2ZnGe3O8.and Li2NiGe3O8 are therefore paramagnetic down to 5 K. EVective magnetic moments per Co or Ni, meff, estimated Order–disorder phenomena and solid solution formation from the susceptibility data, Fig. 3, are listed in Table 8. meff for Li2CoTi3O8 and Li2CoGe3O8 are comparable to that of After reaction to give single phase products, all samples the spinel Co[Rh2]O4 in which the only paramagnetic ion, were annealed at 800 °C, then cooled stepwise to 300 °C in Co2+, is located in tetrahedral sites: meff/mB=4.55 for intervals of 100 °C, with at least 12 h at each temperature. Co[Rh2]O4.21 Li2NiGe3O8 possesses meff similar to that of the There were no phase changes such as decomposition or spinel Ge[Ni2]O4 which has the only paramagnetic ion, Ni2+, additional ordering induced by the annealing.The spinel in octahedral sites: meff/mB=3.24 for Ge[Ni2]O4.22 Co and Ni phases with 153 cation ordering are thus thermodynamically are thus likely to be divalent in Li2CoTi3O8, Li2CoGe3O8 and stable at low temperatures. Li2NiGe3O8. For all the titanates, DTA traces showed a set of reversible In a spinel A[B2]X4, three types of magnetic interactions peaks on heating and cooling, Table 1.Exotherms on cooling between paramagnetic ions should be taken into consideration, were 15–35 °C lower than the corresponding endotherms on i.e., A–A, A–B and B–B interactions. For Li2CoTi3O8, heating, attributable to thermal hysteresis. Sharp peaks were Li2CoGe3O8 and Li2NiGe3O8, every 4b octahedron occupied observed which imply a first-order phase transition.There by paramagnetic Co2+ or Ni2+ is surrounded by six 12d were no mass changes through the peaks, at least up to ca. octahedra containing diamagnetic Ti4+ or Ge4+, as a conse- 1200 °C, Table 1. The thermal anomalies are probably due quence of 153 cation ordering, Fig. 2. No B–B interactions to an order–disorder phase transition involving cations in between Co2+ or Ni2+ therefore occur in these ordered spinels.the octahedral sites, as shown previously for Li2ZnTi3O8.13 Tetrahedral A sites are occupied only by diamagnetic Li+ in The high-temperature disordered phases were, however, not Li2NiGe3O8. Neither A–A nor A–B interactions between Ni2+ retained by air quenching from above the peak temperatures; instead, the low-temperature ordered phases formed.No DTA peaks or mass changes were detectable for any of the germanates, Table 1. The ordered phases were thus maintained until melting temperatures, without undergoing any phase transitions. For the two titanates, a solid solution Li2-2XM1+3XM¾3-XO8 formed on both sides of the Li2MM¾3O8 stoichiometry (X=0) as shown for two compositions, X=0.11 and -0.067, Table 1.By contrast, with one exception there was no evidence of solid solution formation on either side of X=0 for the germanates. For Li2ZnGe3O8, the composition X=0.11 gave a mixture of Li2ZnGe3O8 and Zn2GeO4 (phenacite), and the composition X=-0.067 a mixture of Li2ZnGe3O8 and an unidentified phase(s). For Li2NiGe3O8, the composition X=-0.067 was a mixture of Li2NiGe3O8 and Li4Ge5O12; composition X=0.11 gave two spinels, Li2NiGe3O8 and Ni2GeO4.Li2 CoGe3O8 was the only exception, having a range of solid solution towards Co2GeO4; composition X=-0.067 gave a mixture of Li2CoGe3O8 and Li4Ge5O12. The general absence of both order–disorder transitions and solid solution formation in the germanates, but their presence in the titanates, probably has a common origin in the relative sizes of cations and in particular, in the small size of Ge4+, in Fig. 3 Magnetic susceptibility data for spinel phases Li2MM¾3O8 in the temperature range 5–293 K octahedral sites, compared with Ti4+, Li+ and M2+. The J. Mater. Chem., 1998, 8(5), 1273–1280 1277Fig. 4 Field dependence of magnetization for spinel phases Li2MM¾3O8 Fig. 5 Impedance data for Li2NiGe3O8 at 243 °C at 5 K partial or complete ion blocking and either a rough interface Table 8 Magnetic data for spinel phases Li2MM¾3O8 or Warburg diVusion.Similar, inclined spikes (not shown), at angles of 30–40° to the horizontal were also observed for composition meff/mB a h/K remarks Li2MgTi3O8, Li2 CoTi3O8 and Li2CoGe3O8: their associated capacitances were ca. 10-6, 10-2 and 10-2–10-3 F, respect- Li2CoTi3O8 4.70(3) -35(4) paramagnetic down to 5 K Li2CoGe3O8 4.40(1) -20(1) paramagnetic down to 5 K ively.Such high capacitance values for the latter two must be Li2NiGe3O8 3.09(1) 7(1) paramagnetic down to 5 K caused by ionic conduction together with significant amounts of electrochemical discharge at the metal electrodes.23 For amB denotes the Bohr magneton.Li2ZnGe3O8, the impedance response (not shown) was scattered at low frequencies in the measured temperature range 562–838 °C, Table 9; a poorly developed spike probably associ- thus take place in it. Ni2+ are then regarded as free ions ated with ion blocking at the sample–electrode interface was, without any magnetic interactions. This accounts for the very however, recognisable at high temperatures, e.g. 838 °C.In small h value for Li2NiGe3O8, Table 8. Li2CoTi3O8 and conclusion, the conductivities in the measured (high) tempera- Li2CoGe3O8 have significantly negative h values, Table 8. ture ranges appear to be dominated by Li+ conduction for the A–B interactions between Co2+ are expected to be negative. spinels studied here; the conductivity of Li2ZnGe3O8 is, how- However, the concentration of Co2+ in the octahedral B sites ever, extremely low, and the nature of the conducting species is very small and appears to be inadequate to cause even is not certain.The absence of measurable electronic conduction short-range A–B interactions between Co2+. Negative A–A is further evidence for the absence of mixed-valence states for interactions between Co2+ have been found in some spinels M and M¾, as proposed in the magnetic results section.such as CoRh2O4.21 The negative h values for Li2CoTi3O8 and Bulk conductivity data were extracted from the complex Li2CoGe3O8 are thus ascribable to weak, negative A–A impedance plots, and are shown in Fig. 6 as a function of interactions between Co2+. temperature.The conductivities, s, fit the Arrhenius equation, sT=s0 exp(-DH/kT ), where s0 is the pre-exponential factor, Electrical properties and k Boltzmann’s constant. The activation energies, DH, are In the ideal normal spinel structure A[B2]X4, space group characteristic of moderate to poor ionic conductors, Table 9. Fd3m, each 8a tetrahedron occupied by A-cations shares Because the structures are based on cubic close packing of common faces with four neighbouring empty 16c octahedra.oxygens, the bottleneck size for Li+ hopping through the A-cations can, therefore, move three-dimensionally through tetrahedral site-empty octahedral site conduction pathway the possible diVusion pathway, 8a�16c�8a�16c�. For struc- proposed above is small, leading to the high observed actitures with Y=0, i.e., Li2 NiGe3O8, only Li+ ions occupy the vation energies.Li2ZnGe3O8 possesses by far the highest DH 8a/16c sites. For Y>0, however, M2+ ions partially occupy (2.14 eV) of all the spinels studied here, and its conductivity is the 8a sites and these may act to block the pathways for Li+ extremely low. Li+ conduction through the pathway would be ion conduction. For all Y, the cubic close-packed nature of blocked significantly by Zn2+ located in tetrahedral sites for anions in the spinel structure with full occupancy makes Li2ZnGe3O8, Y=0.5.It may be seen that DH decreases with anionic conduction through spinel lattices unlikely. Impedance data for Li2NiGe3O8 at 243 °C are shown in Table 9 Conductivity Arrhenius parameters for spinel phases Fig. 5. Two partly resolved arcs and a well formed spike are Li2MM¾3O8 seen. The high frequency arc has an associated capacitance, calculated from the relation vRC=1 at the arc maximum of log10 measured temperature 7.9×10-12 F. This value is typical of the bulk (intragranular) composition DH/eV (s0/S cm-1) range/°C response of the sample. The low frequency arc, with an Li2ZnGe3O8 2.14 8.08 562–838 associated capacitance of 1.2×10-10 F corresponds to a grain Li2MgTi3O8 0.71 3.99 163–440 boundary (intergranular) resistance.The spike at low frequen- Li2CoTi3O8 1.33 6.94 349–640 cies is inclined at ca. 30° to the horizontal axis, with an Li2CoGe3O8 1.49 6.86 440–638 associated capacitance of #6×10-6 F, and is characteristic of Li2NiGe3O8 0.55 3.46 63–268 electrode–electrolyte interfacial phenomena associated with 1278 J.Mater. Chem., 1998, 8(5), 1273–1280neutrality. For all the spinels prepared in this work, step (i), i.e., Li+ migration is possible through the tetrahedral siteempty octahedral site pathway proposed above, although for Y>0, partial occupancy of tetrahedral sites by M may act as channel blocking agents.Cyclic voltammograms show a set of reversible peaks at ca. 1.5 V vs. Li/Li+ for the two titanates. The rest potential was ca. 3.5 V, and therefore the peaks in the cyclic voltammograms correspond to insertion/deinsertion of Li, Fig. 7. Redox reactions M2++e=M+ are unlikely to take place, especially for M=Mg, Fig. 7, since the monovalent state is thermodynamically unstable for the M species.Lithium insertion/ deinsertion appears, therefore, to be possible, but to be linked to the reaction, Ti4++e=Ti3+. There were no detectable peaks on cyclic voltammograms for the germanates, consistent with the expectation that, unlike Ti4+, Ge 4 + is not readily reduced. Li insertion is thus prevented by the absence of step (ii ) for the germanates. Attempts to extract Li electrochemically from either titanates or germanates up to 5 V vs.Li/Li+ were unsuccessful: no peaks were observed between ca. 3.5 and 5.0 V in the cyclic voltammo- Fig. 6 Conductivity Arrhenius plots for spinel phases Li2MM¾3O8 grams. Among the spinels studied here, oxidation of Co2+ and Ni2+ is certainly feasible and Li has been extracted electrochemically from the spinel Li2NiMn3O8, involving the oxi- decreasing Y, reaching 0.55 eV foe3O8, Y=0, which dation reaction Ni2+�Ni4++2e.24 In that case, electrons has perfectly connected three-dimensional conduction pathproduced by the oxidation probably migrate through octa- ways without any channel blocking agents in tetrahedral sites; hedral interstices occupied by Ni and Mn.For the titanates further studies on possible correlations between DH and Y are and the germanates studied here, octahedral 12d sites contain- in progress. ing Ti4+(d0) or Ge4+(d10) completely surround each octahedral 4b site containing M2+.Even though electrons were produced Electrochemical properties by oxidation of M2+, the succeeding electronic migration Insertion/extraction of Li into/from Li2M2+M¾4+3O8 requires through octahedra would be blocked by the 12d octahedra three-steps: (i) Li+ migration into/out of Li2M2+M¾4+3O8; (ii), occupied by Ti4+ or Ge4+.Long-range electronic migration redox reactions for either M2+ or M4+ and (iii) electron/hole from octahedral M sites to tetrahedral M sites or vice versa migration from/to the counter electrode to maintain charge would be diYcult, because of very small (or no) concentration of M in either tetrahedral or octahedral sites.Moreover, the relatively long distances between tetrahedra would render electronic conduction through tetrahedral M sites unlikely. It is concluded that, although step (i) may be diYcult to attain in the above samples, owing to the insulating behaviour at room temperature, Fig. 6, Li extraction could be prevented by the lack of long-range electronic conduction and the diYculty of step (iii) for the titanates and the germanates, including those with M=Co or Ni. H. K. thanks CVCP for an ORS Award. We acknowledge the assistance of J. M. S. Skakle (Rietveld refinement of powder neutron diVraction data) and R. A. Howie (bond length calculations). References 1 E. J. W. Verwey and E. L. Heilmann, J. Chem. Phys., 1947, 15, 174. 2 D. S. McClure, J. Phys. Chem. Solids, 1957, 3, 311. 3 J. D. Dunitz and L. E. Orgel, J. Phys. Chem. Solids, 1957, 3, 318. 4 A. Navrotsky and O. J. Kleppa, J. Inorg. Nucl. Chem., 1967, 29, 2701. 5 A. Durif and J. C. Joubert, Compt. Rend., 1962, 255, 2471. 6 G. Blasse, J. Inorg. Nucl. Chem., 1963, 25, 743. 7 J. C. Joubert and A. Durif, Bull. Soc. Fr.Mineral Crist., 1963, 86, 92. 8 J. C. Joubert and A. Durif, Compt. Rend., 1963, 256, 4403. 9 G. Blasse, Philips Res. Rep. Suppl., 1964, 3, 1. 10 G. Blasse, J. Inorg. Nucl. Chem., 1964, 26, 1473. 11 J. C. Joubert and A. Durif, Compt. Rend., 1964, 258, 4482. 12 G. Blasse, Philips Res. Rep., 1965, 20, 528. 13 V. S. Hernandez, L. M. T. Martinez, G. C. Mather and A. R. West, J.Mater. Chem., 1996, 6, 1533. 14 I. M. Hodge, M. D. Ingram and A. R. West, J. Electroanal. Chem., 1976, 74, 125. 15 R. D. Shannon, Acta Crystallogr., Sect. A, 1976, 32, 751. 16 K. Hirota, M. Ohtani, N. Mochida and A. Ohtsuka, J. Ceram. Soc. Fig. 7 Cyclic voltammograms for (a) Li2MgTi3O8 and (b) Li2CoTi3O8 Jpn., 1988, 96, 92. J. Mater. Chem., 1998, 8(5), 1273–1280 127917 T. Saito, N. Mochida and A. Ohtsuka, Yogyo-Kyokai-Shi, 1987, 22 G. Blasse and J. F. Fast, Philips Res. Rep., 1963, 18, 393. 23 H. H. Sumathipala, M. A. K. L. Dissanayake and A. R. West, 95, 604. 18 H. M. Rietveld, J. Appl. Crystallogr., 1969, 2, 65. J. Electrochem. Soc., 1995, 142, 2138. 24 Q. Zhong, A. Bonakdarpour, M. Zhang, Y. Gao and J. R. Dahn, 19 H. Kawai and A. R. West, unpublished data. 20 R. C. Evans, An Introduction to Crystal Chemistry, Cambridge J. Electrochem. Soc., 1997, 144, 205. University Press, London, New York, 1964. 21 G. Blasse and D. J. Schipper, Phys. L ett., 1963, 5, 300. Paper 8/00234G; Received 7th January, 1998 1280 J. Mater. Chem., 1998, 8(5), 1273–12

ISSN:0959-9428    

DOI:10.1039/a800234g

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Tin sulfide clusters in zeolite Y, Sn4S6-Y Carol L. Bowes and GeoVrey A. Ozin* Materials Chemistry Research Group, L ash Miller Chemical L aboratories, University of T oronto, 80 St. George St., T oronto, Canada, M5S 3H6 The synthesis of a tin sulfide cluster array, denoted Sn4S6-Y, using the quantitative sequential surface anchoring and reaction of tetramethyltin with hydrogen sulfide (metal organic chemical vapor deposition, MOCVD, reagents) within the supercages of acid zeolite-Y is detailed. The tethered methyltin species and their transformation to the encapsulated tin sulfide clusters are elucidated through gravimetry, coupled with mid-IR and 119Sn Mo�ssbauer spectroscopy.Transmission electron microscopy (TEM) and Rietveld refinement of synchrotron powder X-ray diVraction (PXRD) data show that the clusters are internally confined and homogeneously dispersed in the supercages of the host zeolite, while optical spectroscopy show a cluster size dependent blue-shift of the absorption edge with respect to the bulk tin sulfide phase. A geometry is proposed for the encapusulated Sn4S6-Y cluster.Introduction Experimental Materials Various approaches to the fabrication of semiconductor quantum dots have been explored, both ‘top-down’ and ‘bottom- Acid zeolite Y was prepared by repeated ion-exchange of up’.1 The former starts with two-dimensional quantum well calcined and defect-removed Na56Y (UOP Y-52 lot no. 13076- structures formed usually by molecular beam epitaxy or chemi- 81) with ammonium nitrate.Crystallinity was confirmed by cal vapour deposition, followed by lithography, etching and powder XRD; 27Al MAS NMR indicated no occluded alumimilling techniques to reduce the lateral dimensions and form nous material and 29Si MAS NMR confirmed the Si/Al ratio.16 cylindrical dots.2 Great gains have been made in recent years Dehydrating at 430 °C resulted in deammination, leaving proand the minimum size of such structures has decreased.Dots tons as charge-balancing cations. At this stage, the resulting on the order of tens of nanometers can be created using material was very moisture sensitive, and was maintained modified STM, but not on a practical production scale, and under high vacuum (10-5 Torr) for subsequent steps. Elemental by sub-optical lithography, although the thicknesses of dots analysis (Galbraith, ICP) indicated a unit cell formula of made using this technique tend to be many times this size.3 H44Na12[(AlO2)56(SiO2)136] corresponding on average to 5.5 Variations of this form of fabrication include the construction protons and 1.5 Na+ ions per a-cage and b-cage.Tetraof nanoelectrodes within the quantum well structure so that methyltin (Aldrich, 99+%) was stored over a molecular sieve electric fields can be produced to ‘squeeze’ the electrons and and degassed by a freeze–evacuate–thaw cycle.H2S (Matheson, holes electrostatically.4 In this way coupling between the dots 99.5%) was passed through freshly dehydrated sodium Y can be controlled and adjusted via an applied gate-voltage.zeolite before use. Structures have been made in which layers are created ‘epitaxially’ such that they are not well lattice-matched, resulting in a Synthetic methods strained interface. Selective etching of the straining layer relieves the compressive stress locally and results in confine- Most reactions involving the zeolite were carried out on a selfment of the charge carriers in that area.5 supporting pressed disk of acid zeolite Y (approximately 40 mg) The ‘bottom-up’ route involves chemical synthesis in which in an in situ cell which allowed dehydration, incremental attempts are made to tailor the size, shape and distribution of Me4Sn adsorption, and had NaCl and quartz windows to semiconductor nanoclusters using self-limiting techniques.6 allow mid-IR and UV–VIS spectroscopy.17 Some experiments This was initially approached by solution-phase synthesis were performed in a small reaction cell on larger amounts of with organic capping7 or through sol–gel-type synthesis.8 acid zeolite Y to allow gravimetric measurements. In general, Alternatively, synthesis within structured media was employed the cluster was produced by dehydrating and deaminating the to constrain size.This was originally attempted in micelles, zeolite wafer at 430 °C, followed by cooling under dynamic vesicles and lipid bilayers, polymers, glasses, clays, and zeo- vacuum. Degassed tetramethyltin was introduced incremenlites. 9–12 Organic-passivated clusters are now prepared in tally by allowing the vapour above the liquid to expand into macroscopic quantities with tunable sizes of 15–100 A ° with a a small ‘titration’ volume and then allowing this volume of standard deviation of less than 4%, and can be manipulated gas to enter the main body of the closed in situ cell.After to form superlattices of close-packed nanoclusters as faceted allowing the gas to adsorb into and react with the zeolite for crystals or thin films.13 Colloidal chemistry is used to produce a brief time (ca. 15 min), the wafer was heated in the quartz layered nanocrystals, for example CdS/HgS/CdS, with epitaxial end of the sample cell to 150 °C for 2 h, after which it remained matching between the core and layers.14 It has also been found white. The sample cell was then returned to the vacuum line, that nanoclusters will self-assemble at high-strain regions when evacuated and filled to 150 Torr with dry H2S.The wafer was direct epitaxy is performed on the slopes of the valleys of a again heated for 2 h at 150 °C, whereupon it turned a pale corrugated layer. Such nanoclusters even self-align to form golden yellow colour, and the cell was evacuated. Any further quasi-ordered arrays.15 handling of the sample was performed under inert atmosphere.In this paper a self-limiting synthetic approach is described which takes advantage of the host–guest chemistry of acid Characterization zeolite Y, in which the quantitative sequential anchoring and FT-mid-IR spectra were collected on a Nicolet 20SXB (reso- reaction of tetramethyltin with hydrogen sulfide forms tin sulfide clusters within the zeolite supercages.lution 2 cm-1) by co-adding 100 interferograms. UV–VIS J. Mater. Chem., 1998, 8(5), 1281–1289 1281diVuse reflectance spectra were recorded with respect to BaSO4 standards on a Perkin-Elmer 330 spectrophotometer using an integrating sphere attachment. Data were digitized and converted to absorbance using Kubelka–Munk theory.18 Powder X-ray diVraction data collected for Rietveld structure refinement were obtained at Brookhaven National Laboratories on beam-line X7A of the synchrotron.A germanium double crystal monochromator was used to select wavelengths of about 0.7 A ° , which was detected by a Kevex solid-state detector. Samples were sealed in Lindemann capillary tubes, cooled to 15 K, and rocked by about 2° while data were collected in 0.005° intervals between 2 and 50° 2h, with collection times increasing with increasing 2h for a total of 12 h.Data were analyzed using GSAS software.19 119Sn Mo�ssbauer data were collected in constant acceleration mode using a Ranger Scientific MS-1200 instrument. The radiation source was 5 mCi 119Sn in a CaSnO3 matrix. Data were collected at room temperature or at 77 K relative to SnO2.Semi-quantitative Fig. 1 Acid zeolite Y showing 13 A ° a-cage formed by 6 A ° b-cages (2.5 atom%) scanning electron microscopy with energy disper- linked by double six-rings. Oxygen atoms bridge between Si and Al sive X-ray microanalysis, SEM–EDX, was performed by tetrahedral centres located at each vertex. Cation sites within the a- Imagetek, using a Hitachi 800 analytical transmission electron cage are shown, as well as Brønsted acid sites, Ha, Hb.microscope operating at 100 kV. The probe size was 50 nm in area and penetrated greater than 1 mm in depth. A LINK X- specific crystallographic positions of the cations, considered ray microanalyser was employed for data anas, and SnS oxide coordination sites in the ‘zeolate’ coordination chemistry and SnS2 were used as standards.Independent elemental model,23,24 are indicated in Fig. 1. analysis was performed by Galbraith Laboratories, Inc., using The adsorption of tetramethyltin is observed in the mid-IR ICP for Si, Al, and Sn determinations and for C, the IR spectrum between 4000 and 1200 cm-1, using an in situ intensity of CO resulting from sample combustion.TEM reaction cell as described in the experimental section. Fig. 2(a) images of the lattice at more than 1.2×106 times magnification shows the spectrum of HY dehydrated at 430 °C. The bands were obtained also by Imagetek using a Hitachi 7000 trans- at 3640 and 3540 cm-1 have been assigned to the a and b mission electron microscope, operating at an accelerating bridging nOH modes.25,26 The small peak at 3740 cm-1 repvoltage of 100 kV.The resolution was estimated conservatively resents terminal nOH stretches, i.e. those OH existing on the at 4.5 A ° . Exposure times were limited to 2 s as electron beam external surface, or in hydroxyl nests and defect sites. Fig. 2(b) induced sample alteration was observed at times of 45 s or shows the final eVect of titration of the a-proton with numerous more.Extended Hu�ckel molecular orbital, EHMO, calculations incremental additions of thoroughly degassed tetramethyltin. were undertaken for isolated tin sulfide clusters, Sn4S6, using The IR intensity of the a-proton was used as the indicator for the ICONCL software.20 The program used the standard the titration. The a-protons reacted, evolving methane and EHMO theory wherein atomic ionization potentials were anchoring a trimethyltin moiety in the a-cage as follows: used for Hii and Hjj, and Hijs were calculated using the Z-O-H+Me4SnAZ-O-SnMe3+CH4 Wolfsberg–Helmholtz approximation, Hij=1/2K(Hii+Hjj)Sij.Overlap integrals, Sij, were calculated for all atomic orbitals Heating to 150 °C for 2 h resulted in the consumption of all rather than just nearest neighbours, using Slater-type orbit- reactive a- and b-protons, as is shown in Fig. 2(c). The resulting als. A weighted Wolfsberg–Helmholtz approximation, K= methane IR spectrum (n3 and n4 modes) was observed in the k¾+D2+D4(1-k¾) where D=(Hii-Hjj)/(Hii+Hjj) and k¾ is a atmosphere of the sealed in situ cell and its intensity was constant,21 was used in this program to account for the measured after each step.An intensity–pressure calibration diVering ‘diVuseness’ characteristic of complexes containing was made, in order to ensure that Beer’s law holds over the unoccupied high-energy basis functions. In addition, a distance pressure range studied. Acid zeolite Y is known not to adsorb dependence for K was included, according to the work of methane, unlike other zeolites or other alkanes.27,28 The further Calzaferri et al.22 Results and Discussion The internal surface of acid zeolite Y is ideal for MOCVD because the protons, which are charge balancing in the structure, can participate in surface anchoring thereby providing control over reaction stoichiometry.The reaction and regeneration of protons as well as their solvation and migration are all observable by mid-IR spectroscopy. Thus the progress of the host–guest chemistry can be monitored in situ providing control over the cluster assembly process.23 Synthesis details Anchoring of tetramethyltin.Zeolite Y has an open-framework topology comprised of a- and b-cages, roughly 13 A ° and 6 A° diameter, respectively, interconnected through 6 T-atom windows.In the illustration of the structure of zeolite Y shown Fig. 2 The titration of protons in zeolite Y with tetramethyltin. Midin Fig. 1, each vertex represents a tetrahedral (T-atom) site IR spectra of (a) dehydrated H44Na12Y, (b) spectrum following awhich is either SiO4 or AlO4-. Extra framework cations proton titration and tetramethyltin anchoring, (c) spectrum following reaction at 150 °C.balance the framework charge due to the AlIII centres, and the 1282 J. Mater. Chem., 1998, 8(5), 1281–1289reaction of the anchored methyltin species with H2S (150 Torr, The combination of restricted space in the a-cage and unfavourable b-proton reactivity limits the extent of Me4Sn 150 °C, 2 h) released essentially all of the remainder of the methyl groups bound to tin as methane, the IR intensity of loading at room temperature. It is critical to this ‘deductive’ reaction characterization that which was measured.By equating the final methane intensity measured with the the tin centres remain in the same oxidation state of +IV throughout. Confirmation that the anchored tin centre maximum of four methyl groups per tin, it was possible to conclude that, on average, 1.0±0.1 methyl groups per tin were remained in the original oxidation state of tetramethytin was ascertained by 119Sn Mo�ssbauer spectroscopy, Fig. 3(a). The released from the titration of the four a-cage protons at room temperature, and therefore that four trimethyltin species were anchored precursor and the cluster product, which will be discussed later, have isomer shifts of 1.27 and 1.25 anchored per a-cage.By the same calculations, after reaction at 150 °C, 1.5±0.1 methane molecules per tin were present in (±0.01) mm s-1, respectively. Isomer shifts of less than 1.5 define tin in oxidation state IV.31 the atmosphere of the in situ cell. This corresponds to a total of 6 protons per a-cage/b-cage unit (48 per unit cell ), in reasonable agreement with the 5.5 (44 per u.c.) determined by Hydrogen sulfide treatment. After the Me4Sn reaction and anchoring were complete, the sample was subjected to approxi- elemental analysis.After reaction at 150 °C, there is some mately 150 Torr H2S at room temperature followed by 150 °C small, remaining b-proton intensity which was thought to be for 2 h.Fig. 4( b) displays the eVect of room temperature H2S either inaccessible or weakly acidic defect proton sites, lying treatment, while Fig. 4(c), (d) show the results of heating at within the b-proton bandwidth. The reaction of essentially all 150 °C and evacuating at room temperature, respectively. protons was taken to indicate the ‘homogeneous dispersion’ Features T, A and B are familiar; the terminal and b-proton of the reactants throughout the zeolite.Although the protons are essentially unaVected while all remnants of the a-proton are to some extent mobile, they cannot delocalize from the have disappeared. Vestiges of the methyl groups remain also, unit cell entirely because it would involve creating chargefeature E. In experiments in which the hydrogen sulfide was separated regions in the zeolite lattice which is expected to be introduced in small, incremental volumes, features C and D energetically unfavourable.Therefore, the disappearance of all grew in together with increasing amounts of H2S. Features C protons implies the homogeneity of the anchored precursor. and D, centred at about 3000 and 2370 cm-1, have been The reaction pathway for the tetramethyltin adsorption may assigned to the solvated a-protons and its solvators, respect- be summarized as follows ively: ZMOMHa,(SH2)n.Earlier work32 documented similar broad and intense peaks resulting from solvation (hydrogen Ha32Hb12Na12Y+32Me4Sn bonding) of zeolite Y a-cage protons by anhydrous hydrogen CA RT (Me3Sn)32Hb12Na12Y+32CH4 halides, of which the vibration was similarly bathochromically shifted.In the earlier work, the degree of shifting of the vibration was found to depend upon the electronegativity of (Me3Sn)32Hb12Na12Y the halide and the number of solvating molecules. In all cases a second broad band was observed, corresponding to band D CA 150 °C (Me3Sn)20(Me2Sn)12Na12Y+12CH4 in this experiment, which also shifted with electronegativity, but to a lesser extent.It was bathochromically shifted from It is interesting to consider the factors contributing to the the expected frequency of the free molecule nHX and in this reaction path described above. One factor is the relative reactivity of the methroups on the tin centre. It is clear that the first methyl group reacts with a proton with ease at room temperature, while the second is less readily removed.It is expected that the nucleophilicity of the second methyl carbon will be less than that of the first because an electron releasing methyl group of Me4Sn has been replaced by an electron withdrawing oxide-type ligand of ZO-SnMe3. Therefore the reactivity towards an electrophile (proton) will incrementally decrease as first, second or third methyl groups react.Also, electrostatics may contribute to the diminished reactivity. For a second methyl group to react, the proton must eVectively interact with a Me3Sn+ centre bearing a formal 1+ charge. For a third, it must approach a Me2Sn2+ centre bearing a formal 2+ charge. Electrostatics are expected to be an important factor, based on decationization and dehydrohalogenation kinetic experiments in zeolite Y.29 A second factor is the steric limitation imposed by the size of the a-cage.It was observed, through IR spectroscopy of the sample pellet and sample-cell atmosphere, that Me4Sn added after completion of the titration of the a-protons remained in the atmosphere of the cell (outside the zeolite), despite the presence of further unreacted b-protons.This information should be coupled with the idea that there are expected to be thermodynamic and kinetic barriers to the reaction of the b-protons. It must be so, or the initial, room temperature uptake of tetramethyltin would aVect a- and bprotons equally which is not observed (RT decrease of the bproton intensity corresponds only to 0.15 protons).Because Me4Sn is too large to enter the b-cages, the b-protons must leave the b-cage in order to react; plus, b-protons are less Fig. 3 Mo�ssbauer spectra of (a) the precursor material (CH3)2Sn-Y acidic and less mobile than their a-counterparts, as shown and (b) the product material Sn4S6-Y. Isomer shifts are reported with respect to SnO2. through adsorption and desorption energies of given bases.30 J. Mater.Chem., 1998, 8(5), 1281–1289 1283hydrosulfurization–deanchoring (proton regeneration): Me2Sn-Y+H2SAMe2(SH)Sn-Y+H-Y Me2(SH)Sn-Y+H2SAMe2(SH)2Sn+H-Y dehydrosulfurization–condensation (cluster assembly): (SH)4Sn+(SH)3Sn-YA(SH)3Sn-S-Sn-(SH)2-Y+H2S The alternative reaction for the last step is an intramolecular self-dehydrosulfurization, in which a SnNS species is formed following H2S elimination.While this step could not be ruled out a priori, it is not thought to contribute in light of the remainder of the data, as will be discussed. Thus the proposed synthesis in which four anchored methyltin species per a-cage are demethylated, hydrosulfurized, and self-assembled via dehydrosulfurization–condensation reactions accounts adequately for the formation of a tin sulfide cluster product, except for the question of cluster stoichiometry.Several methods were employed for the determination of the Fig. 4 Adsorption and reaction of hydrogen sulfide in methyltintin to sulfur ratio in the nanocluster product. The first was a loaded zeolite Y. (a) Mid-IR spectrum of anchored dimethyl- and quantitative gravimetric experiment in which the entire syn- trimethyltin in zeolite Y, (b) spectrum following room temperature thesis was carried out on a large sample (approximately introduction of 150 Torr H2S, (c) the product of 150 °C reaction, (d) spectrum following room temperature evacuation of the reaction 300 mg) in a small cell (approximately 60 g) which was accucell. rately weighed after each step in the reaction process.The second was semiquantitative SEM–EDX using SnS and SnS2 standards. A conservative estimate of the error is ±2.5 atom% for ratio measurements with standards and within the thin crystal limit for micro-X-ray fluorescence.34 The third was case the vibration may be assigned also to the nSH stretch for independent elemental analysis (Galbraith) using ICP and with solvating hydrogen sulfide.standards to establish errors. The results were as follows: Finally, the peaks F and G at about 2562 cm-1 have been assigned also to nSH vibrations. In the literature,33 the n3 gravimetric: S/Sn=1.63±0.05 STEM–EDX: S/Sn=1.42±0.04 Galbraith: S/Sn=1.50±0.15 mode of molecularly adsorbed hydrogen sulfide as well as the nSH stretching mode of a hydrosulfide group resulting from dissociatively adsorbed H2S in zeolites, have been reported at giving an average result of 1.5 for the S/Sn ratio and therefore about 2560 cm-1.While the former may be present (some H2S a likely cluster stoichiometry of Sn4S6, with an overall charge molecules not involved in hydrogen bonding to a-cage protons of 4+ based on SnIV and S-II oxidation states, Fig. 3(b), see but rather physisorbed to the oxygen framework), the latter is Mo�ssbauer discussion above. almost certainly present, as will be considered in the discussion The absence of any significant contribution from tin sulfide of the reaction path below. material aggregating on the external surface of the zeolite Heating to 150 °C has little eVect on the IR spectrum, except crystals was determined by TEM lattice imaging, Fig. 5. for a decrease in the D-band and a shift to higher frequency Aggregations of the heavier guest species generally appear as of the C-band, both associated with a decrease in the number darker spots on the crystal surface which was not observed in of solvating H2S molecules. Also, the intensity of and ratio any TEM images obtained from many of these materials.In between bands F and G have changed, possibly indicating the addition, a general impression of the homogeneity of the reaction of, for example, H2S to form HS- species. The sample interior structure can be obtained which suggests that the tin appearance, at this stage, changed from the previously white sulfide clusters are encapsulated and homogeneously distribcolour to a pale golden colour.uted. In fact, because not all protons are regenerated in the After 150 °C reaction, the excess hydrogen sulfide and methtin sulfide cluster formation process, it is expected that the ane reaction product were pumped out of the in situ cell at clusters are anchored to the lattice, performing some frameroom temperature, leaving a material whose IR spectrum is work charge-balancing function, in accord with the expected shown in Fig. 4(d). The most noticeable aspect of this spectrum charge on Sn4S64+. is that protons, both a- and b-cage, were regenerated. It is not Finally, the regeneration of unsolvated protons has importclear at what point in the reaction scheme the protons are ant ramifications with respect to cluster growth within the regenerated, because in the presence of H2S the a-protons are zeolite. Preliminary molecular graphics structural models indisolvated and their number cannot be easily quantified from cated that a Sn4S6 cluster would not fill the a-cage void space their IR intensity.Nevertheless, in repeated experiments, on and this idea is supported in that there exist a-protons average 40% and up to 50% of the original intensity of these unsolvated by the cluster in the final product.In principle, it bands has returned after evacuation. Band D disappeared should be possible to re-titrate the regenerated protons and in completely, while some portion of the protons remain solvated, a second, controlled loading step add tin and sulfur to the indicating that all of the H2S was removed but that the tin existing clusters in order to increase their size.Doping might sulfide product species solvates the neighbouring protons. also be achieved in this way. Finally, the bands F and G disappeared, suggesting that no Growing the clusters was accomplished as shown in Fig. 6, hydrosulfide species remained. Reactions of the following type, where each IR spectrum represents a complete Me4Sn titration in which Y represents the zeolite framework, are thought to and anchoring/H2S adsorption and reaction/evacuation cycle.contribute to the process observed in the IR: Important aspects of these spectra are the decrease of both hydrosulfurization–demethylation (methane evolution): solvated and lvated protons and the lack of increase of methyl or hydrosulfide group intensity.The former indicated Me3Sn-Y+H2SAMe2(SH)Sn-Y+CH4 that more of the cation sites are occupied ‘capping’ the clusters, the latter suggests that methyl groups are decreasingly required Me2Sn-Y+H2SAMe(SH)2Sn-Y+CH4 1284 J. Mater. Chem., 1998, 8(5), 1281–1289most significant. Having control over the process, the ability to stop, and change the reaction conditions at any point in the synthesis is what is most unusual, and most useful about this method.Therefore it is interesting to consider what species are present at various points; in short, to consider the reaction pathway. The reaction of methyltin moieties with protons is fairly straightforward, as was described above, the first loss of methane occurring easily at room temperature, and the second with a higher activation energy due to the change in nucleophilicity and electrostatics. Thus the anchoring of Me4Sn in HY finished with a mixture of anchored trimethyl- and dimethyltin, there being 5.5–6.0 reactive protons and 4 tin centres per a-cage.At room temperature, upon H2S addition, methane is released corresponding to almost one equivalent per tin (0.85±0.10) so that 2.4±0.1 of the four methyls per tin have been released.This is curious, since the reacting species were mixed di- and trimethyltin species. No argument based on relative SnMC bond strengths will easily account for this observation. The outcome of the reaction must depend upon which other factor was dominant.First, the degree to which the HMS bond of coordinating H2S is polarized by the tin centre will aVect the reactivity of the proton. The replacement of methyl groups by hydrosulfide groups would increase the ability of tin to polarize that bond. Second, the susceptibility of the methyl carbon to electrophilic attack will be decreased by the introduction of hydrosulfide groups about the tin centre that are more electron withdrawing, having the eVect of attracting electron density away from carbon and decreasing its nucleophilicity.Fig. 5 Typical high resolution TEM lattice image at 1.2×106 magnifi- If the reactivity of the H2S proton were the limiting step, cation of zeolite Y with encapsulated tin sulfide clusters, showing lack tin nuclei with fewer electron donating methyl groups and of external tin sulfide and homogeneity of sample more electron withdrawing hydrosulfide groups would be more reactive.After a dimethyltin species reacted, it would become even more reactive, and one equivalent of methane could result from those centres alone. Alternatively, if the reactivity of the carbon were limiting, then the result might be that a centre, having gained one hydrosulfide group, would be invulnerable to further attack at room temperature, so that methane produced represents one methane from each centre.Neither scheme accounts for interference by de-anchoring proton regeneration reactions which may compete for hydrogen sulfide protons. Described in reaction (2) above, the dissociative deanchoring must occur in order to form a tin sulfide cluster. The IR band corresponding to solvated protons in Fig. 4(a) seems too large to be entirely due to a few ‘leftover’ protons, yet it is not possible to quantify the amount, as the intensity is dependant both on the number of solvated protons and on their extent of solvation, as described above. However, it is likely that there is a suYciently large excess of hydrogen sulfide so that the methane production and de-anchoring can proceed without mutual interference. Fig. 6 Mid-IR spectra of (a) Sn4S6 clusters in zeolite Y, (b) product The second process is that of cluster formation self-assembly, following a second loading of tetramethyltin and hydrogen sulfide, described in reaction (3). After thermal treatment, Fig. 4(c), a and (c) following third loading strong nSH vibration remained (feature G), along with some significant number of solvated protons (C).This is taken to indicate that many, if not all, species are in the hydrosulfide to satisfy tin coordination sites and that hydrosulfide condensation is occurring. Possibly, the new material is incorporated form, and only condense and assemble when the excess hydrogen sulfide is pumped from the cell. to form larger clusters rather that forming new, smaller clusters which would leave many tin atoms coordinatively unsaturated.Using the methods of methane quantification described earlier, Proposed cluster geometry of Sn4S6-Y. There exist in the literature a number of tin sulfide clusters of various stoichio- the reactions of second and third Me4Sn loadings were shown to produce as much as 118% as much methane as the first, metries and cluster charges having specific geometries.35 However, the clean chemistry leading to the cluster and various suggesting eight Sn atoms or more per a-cage.The UV–VIS spectra of the materials illustrate the cluster generation and pieces of evidence have led us to propose a single Sn4S6 cluster moiety.Firstly, the only example of a cluster of this stoichi- growth nicely, and are discussed in a later section. ometry in the literature is the recurring adamantanoid geometry with a 4+ charge. Because the clusters in the zeolite Reaction pathway. Although the reactions presented above are thought to represent those participating in cluster forma- are partially charge balancing with respect to the anionic framework, creating larger clusters would involve separation tion, it is the step-wise nature of this process which is perhaps J.Mater. Chem., 1998, 8(5), 1281–1289 1285of charge between the anionic lattice and cationic clusters Spectroscopic characterization of precursors and products which would be energetically unfavourable. Further evidence IR Spectroscopy of anchored methyltin reagents in zeolite Y.was obtained from a Rietveld refinement of the structural A key tool in determining the product of the first step anchoring model. Rietveld PXRD refinement is the most widely used reaction has been in situ mid-IR spectroscopy. Yet, it was not method for obtaining crystal structures of materials when it is possible to identify the products directly from their spectra, not possible to obtain crystals of suYcient size for single crystal both because of the occurrence of mixed species, and because structure determination, provided a good starting model is the spectra of the various candidates are quite similar.Fig. 7 available. In the case of tin sulfide in zeolite Y the quality of shows the spectra in question.Trace (a) is the nCH stretching the refinement was insuYcient to completely define the cluster, region of Me4Sn chemisorbed in Na56Y, any excess physisorbed due to significant symmetry-related disorder intrinsic in the Me4Sn having been pumped away. This represents a saturation problem, and possibly dynamic disorder, therefore the inforloading of four Me4Sn per a-cage, in which the Me4Sn moieties mation gleaned must be taken cautiously due to the known are most likely adsorbed to the Na+ cations with interactions pitfalls of locating extraframework species in zeolites.36 between the methyl hydrogens and the oxygen framework Nevertheless some further evidence for its structure and sites.Trace (b) of Fig. 7 is the nCH IR spectrum of the location was suggested.Using the low temperature (15 K) corresponding situation in HY whereby trimethyltin acts as synchrotron PXRD data, a preliminary Rietveld structure an extraframework cation, probably situated above the four refinement of the zeolite framework only was made in the tetrahedrally arranged, puckered oxygen six-ring sites, with up space group Fd39m, in order to locate electron density on a to three of the adjacent oxygens coordinating the tin centre.Fourier diVerence map. All electron density was located in the Trace (c) of Fig. 7 is the corresponding nCH IR spectrum of a-cage, the b-cage being completely empty, as was expected the product of thermal treatment of the sample in (b), expected due to the size of the precursor species.This confirms the ‘ato be roughly a 50–50 mixture of trimethyltin and dimethyltin cage specific’ nature of this kind of intrazeolite MOCVD anchored species (due to 5.5–6 total protons reacting with 4 synthesis. The most significant concentration of electron dentetramethyltin species). sity was located over the site II positions (see Fig. 1), centred Tetramethyltin is a regular tetrahedral molecule belonging above the three prominent oxygens of the six-ring in the ato the point group Td.The set of twelve displacement vectors cavity, at a distance of about 2.4 A ° . Introducing a tin atom at representing the CMH stretches generate an irreducible repthis location significantly reduced the refinement residuals and resentation Cvib=A1+E+T1+2T2. Of these, only the T2 are x2.The occupancy of this site, determined to be slightly less IR active. The predicted two major T2 nCH bands of Me4Sn than one quarter, is indicative of a single species symmetry in NaY were observed at 2975 and 2902 cm-1. These values disordered over the four crystallographically inequivalent sites, correspond to the reported nasym and nsym modes of liquid rather than multiple species each occupying one of the sites Me4Sn observed at 2982 and 2914 cm-1 respectively.37 In that which would lead to occupancies of one half or higher.No work it was assumed that the vibrational coupling between significant electron density was observed at any of the other the nCH methyl group and SnMC skeletal modes was small. common cation sites of the a-cages of zeolite Y.It was possible On the basis of this argument, it might be expected that the to introduce a second distinct tin at a location of high electron diVerence in nCH vibrational frequencies between Me4Sn, density which, due to the threefold symmetry of the site, formed Me3Sn and Me2Sn species anchored in the a-cage of zeolite Y a tetrahedron with the first with tin–tin distances of approxishould be small.This is the case in practice. mately 3.5 A ° , consistent with the sulfur-bridged tins of the The nine internal coordinates representing CMH stretches proposed adamantanoid geometry. This also resulted in sigfor a C3v pyramidal, anchored ZO-Sn(CH3)3 moiety generate nificant reduction in residuals. Beyond this point, at which an irreducible representation Cvib=2A1+A2+3E.Of these A1 x2=2.4, Rp=9.1 and weighted residual wRp=12.6 (Table 1), no further improvement in the refinement could be made, despite the observation of electron density on the Fourier diVerence map at locations suitable for sulfur positions. Consequently, no sulfur atoms are included in the model. Although the Rietveld refinement of the data was not able to unambiguously define the structure of the tin sulfide cluster, the information that was obtained was consistent with the proposed adamantane geometry and location exclusively inside the zeolite a-cages.Table 1 Atom positions and parameters of the best model for a single Sn4S64+ cluster in zeolite Y x y z Uiso/A ° 2 Si/Al(1/2) -0.0549(3) 0.1328(2) 0.0356(3) 0.0039 O(3) 0 0.1111(6) -0.1111(6) 0.0366 O(4) -0.0023(5) -0.0023(5) 0.1459(6) 0.0050 O(5) 0.0752(4) 0.0752(4) -0.0317(7) 0.0176 O(6) 0.0713(6) 0.0713(6) 0.3224(8) 0.0068 Sn(7) 0.2421(7) 0.2421(7) 0.2421(7) 0.1377 Sn(8) 0.411(13) 0.411(13) 0.205(15) 0.1292 statistics of refinement x2=2.4 wRp=12.6 Rp=9.1 lattice parameter, a/A ° 24.6537(5) r[Sn(7)MO(4)]/A ° 2.398(20) r[Sn(7)MSn(8)]/A° 3.5(3) r[Sn(8)MSn(8)]/A ° 4.7(8) Fig. 7 Mid-IR spectra of nCH region for (a) tetramethyltin impreg- [Sn(8)MSn(7)MSn(8)]/° 83(6) nated NaY, (b) anchored trimethyltin in acid zeolite Y, (c) anchored [O(4)MSn(7)MO(4)]/° 101.3(9) dimethyltin and trimethyltin in acid zeolite Y 1286 J.Mater. Chem., 1998, 8(5), 1281–1289and E are IR active and so five modes are expected. The two to values such as 3.31 mm s-1 (ref. 40) {Me2Sn[N,N¾-(2- hydroxytrimethylene)bis(salicylaldamine)]}, agreeing best with major bands observed are blue-shifted from those of chemisorbed (CH3)4Sn,NaY by roughly 15 cm-1 for the symmetric the splitting determined here. On going to more symmetric, trigonal bipyramidal structures with methyl groups both equa- stretch and 4 cm-1 for the asymmetric stretch. These nCH bands, however, remain within the linewidths of the tetra- torial {Me2Sn[N-(2-hydroxyphenyl)salicylaldimine]}, the quadrupole splitting drops to 3.04.41 This suggests that the methyltin vibrations.It is possible, through consideration of the characteristic motions which compose the nCH normal dimethyltin species, anchored over the three prominant oxygens of site II, probably has a structure intermediate modes of ZO-Sn(CH3)3, to separate the five normal nCH modes into two sets having their origins in the symmetric and between the latter two examples, a distorted ‘octahedron’ with one missing ligand, or a distorted trigonal bipyramid, with asymmetric nCH modes of uncoupled CH3 groups.The result is that the lower frequency band may be assigned as A1+E one methyl group equatorial and one axial. It is also interesting to note in this data the possiblity of a (A1, CH3), while the higher frequency major band includes the components A1+2E (E, CH3), the A2 contribution being Goldanskii–Karyagin eVect, reflected in the diVerence in areas of the two peaks of the quadrupole doublet while similar peak IR inactive. Upon heating the ZO-Sn(CH3)3 samples containing b- widths (FWHH) are maintained.42 This eVect arises when there is an anisotropy in the Mo�ssbauer recoil-free fraction, i.e., the protons to 150 °C, for 2 h, some ZO-Sn(CH3)2 species are formed.The spectrum shows a further shift of the band maxima mean-square vibrational amplitude is not cubically symmetric. This is quite reasonable in a situation in which the recoiling to higher frequency, and a change in the relative intensities.The lower frequency band shifts by 3 cm-1 and the higher by species is eVectively anchored to a surface site (i.e. the zeolite framework) such that the vibrational amplitude perpendicular 5 cm-1 with respect to the anchored ZO-Sn(CH3)3 species. As above, under C2v, the irreducible representation Cvib= to the ‘wall’ will likely be very diVerent from that parallel to it.Further work would be required to prove such a point, but 2A1+A2+2B1+B2 can be split into two groups, A1+B1 (A1, CH3) for the lower frequency band and A1+B1+B2 (E, CH3) it remains an interesting possibility. Fig. 3(b) shows the Mo�ssbauer spectrum of Sn4S6-Y. The for the higher frequency band. In these considerations, a rigid model for the methyl groups isomer shift of 1.25 mm s-1 indicates that the oxidation state of tin is IV, as mentioned earlier, and the quadrupole splitting has been assumed.While it is also possible to predict the active modes using a non-rigid Bunker-type model in which the is 2.28 mm s-1, considerably smaller than that of the precursor material. However, octahedral or tetrahedral symmetries have methyl groups are free to rotate, in the sterically hindered environment of the a-cage containing a total of four ZO- no intrinsic quadrupole splitting, and no alkyl groups are expected to contribute, as only sulfide and oxide ligands are Sn(CH3)3 and ZO-Sn(CH3)2 species, it is unlikely the methyl groups would truly be free rotors, and therefore the rigid expected.Therefore, an unsymmetric distribution of the diVerent ligands about a tetrahedral or octahedral tin(IV) centre will oscillator approach was preferred. There are very weak satellite bands or shoulders in the nCH spectra which are, in every probably describe the structure, consistent with the proposed adamantane structure. The isomer shift of the encapsulated case, split oV from the major bands by about 35 to 130 cm-1, precluding the simple assignment of these bands as components cluster, 1.25 mm s-1, agrees very closely with that measured for the methyl-capped cluster Me4Sn4S6, 1.28 mm s-1.of the nCH ZO-Sn(CH3)3/ZO-Sn(CH3)2 modes.38 Such splittings are not readily accounted for through arguments of However, the quadrupole splitting of the encapsulated cluster is considerably bigger than that of Me4Sn4S6 (2.28 vs.multiple anchoring site eVects, splitting of degeneracy through symmetry eVects, or even correlation coupling eVects, 1.34 mm s-1), probably because the latter has a more rigid symmetry, and alkyl groups more closely match the electrone- which are all typically 20 cm-1 or less. No Raman modes of Me4Sn occur at appropriate frequencies such that a change of gativity of sulfide ligands than zeolite oxygens do.The observation of a single kind of tin(IV) Mo� ssbauer site is consistent selection rules due to site symmetry eVects could result in IRactive modes at these positions. The best interpretation is that with the Rietveld PXRD structural conclusion, that the Sn4S64+ cluster is four-fold positionally disordered over the these shoulders and weak bands be assigned as overtones and combinations with low frequency fundamental modes, consist- four site II oxygen six-ring positions in the a-cage of zeolite Y.Variable temperature Mo�ssbauer studies (10–300 K) will be ent with the assignment of similar shoulders in the spectrum of Me4Sn.37 required to distinguish static from dynamic Sn4S64+ disorder in this system.Mo�ssbauer spectroscopy. The isomer shift, 1.27 mm s-1, measured in the Mo�ssbauer spectrum of dimethyltin anchored in Optical spectroscopy. In Fig. 8 the optical reflectance data for four forms of tin sulfide are presented. Qualitatively, the zeolite Y, Fig. 3(a), indicates that the oxidation state of tin is IV. The large quadrupole splitting, 3.44 mm s-1, speaks of the spectrum of the encapsulated clusters, Sn4S6-Y, is considerably blue-shifted as compared to the bulk semiconductor berndtite, low symmetry surrounding the tin(IV) centres.The sample of pure dimethyltin in zeolite Y was prepared by titrating the SnS2. This is as expected for confinement of electrons within a quantum dot. However, it is quite similar to the spectrum of protons while the sample was held at 150 °C, so that the reactivity was similar for the a- and b-protons.Thus every tin the molecular species, (CH3)4Sn4S6, indicating that the excitonic confinement regime was overshot. This species is within centre reacted with two protons immediately, and therefore about three tin centres, uniformly dimethyl species, were the very strong confinement size regime and is essentially molecular.Without good information about the eVective anchored in each a-cage. Quadrupole splittings of this size (3.44 mm s-1) are usually indicative of low symmetry, such masses of the electrons and holes in SnS2 it is not possible to calculate the optimum cluster size for excitonic confinement, as trigonal bipyramidal stereochemistry, for compounds (R2SnX2)n or (R3SnX)n where X is F-, Cl-, Br-, I-, or O2- but it is clear that the size obtained was too small.An EHMO calculation for the isolated cluster was consistent with the donor.31 Among alkyltin species with oxygen containing ligands, the size of the quadrupole splitting is dominated by qualitative molecular orbital diagram shown in Fig. 9, and indicated that the HOMO and LUMO were composed mainly the number and position of the alkyl ligands. For example, for dimethylated amino acid and SchiV base derivatives with of sulfur 3p and tin 5s and 5p orbitals, respectively, and give rise to a ligand-to-metal charge-transfer electronic transition octahedral structure and methyl groups trans, quadrupole splittings of up to 4.02 mm s-1 have been reported39 as seen in the optical spectrum, Fig. 8.43 In considering the possibility of intercluster coupling, the EHMO method was [Me2Sn(acac)]. Distorting the octahedra such that the CMSnMC bond angle is 160° reduces the quadrupole splitting used to determine that clusters would need to have centre-to- J. Mater. Chem., 1998, 8(5), 1281–1289 1287Fig. 10 UV–VIS diVuse reflectance spectra of the products of (a) the Fig. 8 UV–VIS absorbance spectra of four forms of tin sulfide, first loading of Sn4S6 clusters in zeolite Y, (b) a second loading, (a) molecular (CH3)4Sn4S6, (b) zeolite encapsulated Sn4S6-Y, titrating regenerated protons in order to add more tin sulfide material (c) [(CH3)4N]2Sn3S7 denoted TMA-SnS-1, (d) bulk SnS2, berndtite to each cluster, and (c) a third loading, increasing the cluster again, (d) bulk SnS2 tion edge on going from Sn4S6 clusters, (a), to clusters resulting from a second loading, (b), and from a third loading, (c).The red-shift is probably indicative of less confined charge carriers, due to an increase in the cluster size.47 This supports the idea that the titration of regenerated protons and formation of more tin sulfide material in each a-cage does increase the nuclearity of the initial tin sulfide clusters.On a final note, one would expect that adamantanoid Sn4S6-Y might be grown from intrazeolite Me4Sn and H2S reagents, as it is well known that the solution phase reaction of MeSnCl3 with H2S yields adamantane clusters, Me4Sn4S6, with terminal methyl-capping groups. The four nucleophilic methyl groups of Me4Sn4S6 are simply substituted in the zeolite encapsulated analogue by charge-balancing framework oxygens to give (ZO)4Sn4S6 according to the zeolate model of bonding.23 Conclusions Fig. 9 Extended Hu� ckel molecular orbital, EHMO, energy diagram representing a zeolite-encapsulated Sn4S6 cluster The application of size-limiting intrazeolitic MOCVD techniques has been demonstrated to be a method of choice for the synthesis of arrays of zeolite encapsulated IV–VI semi- centre distances of less than 9 A° before any modification of their electronic structure was calculated.While the clusters conductor clusters. Step-wise, molecule-by-molecule control over the intrazeolite reaction processes is possible by employing were not expected to occupy the centres of the a-cages, which are 13 A ° apart, there would nevertheless be an average inter- non-intrusive, quantitative, in situ methods of observing them, as shown in this work.Charge-balancing tin sulfide clusters of cluster distance of 13 A ° over the whole material, which puts it outside this range for potential intercluster coupling. This of average Sn4S64+ stoichiometry have been synthesised in the acages of zeolite Y, possibly with an adamantanoid geometry.course does not account for any through-bond coupling, possible via the zeolite lattice. In comparison, in a microporous It is expected that the methodologies presented in this paper will be generally applicable to other semiconductor clusters layered tin(IV) sulfide having the formula R2Sn3S7 with Sn4S3, broken cube clusters linked by double sulfur bridges,44 the within other media in syntheses that are equally controllable.It should be noted that, at the time this work was done, clusters are less than 8 A° apart and are anticipated to be significantly coupled through the double sulfur bridges. clusters in zeolites had an upper limit of 13 A ° imposed by the restricted choice of large-cage zeolites.Since then the world The intermediate value of the absorption edge of R2Sn3S7, Fig. 8, with respect to those of the bulk and molecular cluster of size-controlled large-aperture mesoporous silicas, beginning with MCM-41,48 has opened up, and ordered channels of forms of tin sulfide is noteworthy both in terms of connectivity arguments that have been put forward regarding the dimen- 20–100 A ° are available as hosts for this type of CVD topotaxy.49 sionality of such materials,45 and in the context of semiconductor quantum confinement and cluster-framework topological complements.46 The authors are indebted to Dr Robert Broach and Dr Robert L.Bedard for valuable discussion regarding the Rietveld As was discussed above, further growth of the tin sulfide clusters was possible and resulted in red-shifting of the absorp- analysis.The PXRD measurements were carried out at the National Synchrotron Light Source, Brookhaven National tion maximum, but the material was insuYciently homogeneous for good characterization. Nevertheless, Fig. 10 shows Laboratories which is supported by the US Department of Energy, Division of Materials Sciences and the Division of the results of cluster growth: a definite red-shift of the absorp- 1288 J.Mater. Chem., 1998, 8(5), 1281–128921 J. , H.-B. Burgi, J. C. Thibeault and R. HoVman, J. Am. Chemical Sciences. We acknowledge the Natural Sciences and Chem. Soc., 1978, 100, 3686. Engineering Research Council of Canada (NSERC) as well as 22 G.Clazaferri, L. Forss and I. Kamber, J. Phys. Chem., 1989, 93, the Canadian Space Agency (CSA) and UOP for financial 5366. support of this endeavour. In addition, C.B. thanks NSERC 23 G. A. Ozin, C. L. Bowes and M. R. Steele, Mater. Res. Soc. Symp. and the University of Toronto for financial support during her Proc., 1992, 277, 105. 24 G. A.Ozin and S. Ozkar, Chem. Mater., 1992, 4, 51; G. A. Ozin, graduate work. A. Kuperman and A. Stein, Angew. Chem., Int. Ed. Engl., 1989, 101, 373; G. A. Ozin, S. Ozkar and R. A. Prokopowicz, Acc. Chem. Res., 1992, 25, 553; G. A. Ozin, Adv. Mater., 1992, 4, 612; G. A. Ozin, Adv.Mater, 1992, 4, 11. References 25 D. H. Olsen and E. Dempsey, J. Catal., 1969, 13, 221. 26 J. W. Ward, J. Catal., 1967, 9, 396. 1 P. Day, Chem. Br., 1996, July, 29. 27 G. A. Ozin, S. Ozkar and L. McMurray, J. Phys. Chem., 1990, 2 M. A. Reed, J. N. Randall, R. J. Aggarwal, R. J. Metyi, T. M. Moore 94, 8297. and A. E.Wetsel, Phys. Rev. L ett., 1988, 60, 535. 28 T. Yamazaki, I. Watanuki, S. Ozawa and Y. Ogino, L angmuir, 3 R. F. Pease, in Nanostructures and Mesoscopic Systems, ed. 1988, 4, 433. W.P. Kirk and M. A. Reed, Academic Press, Toronto, 1992. 29 G. A. Ozin, S. Ozkar and L. McMurray, J. Phys. Chem., 1990, 4 F. R. Waugh, M. J. Berry, D. J. Mar, R. M. Westervelt, 94, 8289. K. L. Campman and A. C. Gossard, Phys. Rev. L ett., 1995, 75, 705. 30 L. Quanzi, A. Ruiming and X. Zhiyuan, in New Developments in 5 K. Kash, D. D. Mahoney, B. P. Van der Gaag, A. S. Gozdz, Zeolite Science and T echnology, ed.Y. 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Beck, Nature, 1992, 359, 710. Alamos, NM 87545, USA. 49 C. L. Bowes, A. Malek, G. A. Ozin, Chem. Vap. Deposit. (Adv. Mater.), 1996, 2, 97. 20 ICONCL software: OS/2 Version, MS-Fortran 5.0, 30 November 1989, Modified ICON8 from QCPE 344, Calzaferri Group, Institute for Inorganic Chemistry, University of Bern, Switzerland. Paper 7/08093J; Received 11th November, 1997 J. Mater. Chem., 1998, 8(5), 1281–1289 12

ISSN:0959-9428    

DOI:10.1039/a708093j

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials A variable temperature neutron diVraction study of the layered perovskite YBaMn2O5 Judith A. McAllister and J. Paul Attfield*† Department of Chemistry, University of Cambridge, L ensfield Road, Cambridge, UK, CB2 1EW, and the Interdisciplinary Research Centre in Superconductivity,Madingley Road, Cambridge, UK, CB3 0HE A variable temperature neutron diVraction study has been carried out on the layered perovskite YBaMn2O5 between 100 and 300 K.A broad peak indexed as (1/2, 1/2, 1) on the nuclear cell is consistent with short range MnII/MnIII valence ordering. The magnetic structure below the ferrimagnetic ordering temperature of 167 K has been determined and is in agreement with a previously proposed model. Changes in lattice parameters, bond lengths and angles show evidence of an exchange striction at TC.Introduction Perovskite type Ln1-xAxMnO3 phases (Ln=trivalent lanthanide, A=Ca, Sr, Ba) exhibit giant magnetoresistances (GMR).1 The features important to this phenomenon are the threedimensional MnMOMMn network and Mn of mixed valency with the mean Mn oxidation state between +3 and +4, i.e.hole-doped MnIII. This gives rise to ‘double-exchange’,2 a transfer of spin-polarised electrons from Mn3+ to Mn4+, below the Curie temperature. A new perovskite-related Mn oxide, YBaMn2O5, was recently reported.3 The structure has a layered arrangement with oxygen vacancies in the yttrium plane and the manganese –oxygen network consists of double layers of identical MnO5 square pyramids linked through their apical oxygens.This structure is similar to those of YBaCuFeO5+d4 and Fig. 1 Proposed3 valence and magnetic ordering model for YBaMn2O5 YBaCo2O5+d.5 Magnetic measurements on YBaMn2O5 showing alternating MnII with S=5/2 (up) spins and MnIII with S=2 showed a ferrimagnetic ordering temperature of 167 K with (down) spins the measured ferromagnetic component of about 0.5 mB, this is consistent with a magnetic and valence (charge) order of MnII and MnIII, as shown in Fig. 1. Each MnII is linked via were assigned to Ba2SiO4 resulting from reaction with the oxygen to five MnIII ions and vice versa which if extended over silica tube, due to a slight rupture in the Au foil. long distances would give a Ó2a×Ó2a×c superstructure. In Constant wavelength neutron powder diVraction data were this study we have investigated the magnetic and crystal collected at 100, 140, 190 and 300 K on instrument D2B at structures and the extent of superstructure formation in the Institut Laue-Langevin, Grenoble, France. The wavelength YBaMn2O5 using variable temperature powder neutron selected for this experiment was from the (335) Bragg reflection diVraction.of the Ge crystals, with l=1.594 A ° . Rietveld refinement6 was carried out using the General Structure Analysis System (GSAS).7 Experimental A stoichiometric mixture of YMnO3, BaO and MnO was used Results in the preparation of a 10 g sample of YBaMn2O5 for neutron diVraction. YMnO3 was prepared by heating a mixture of The starting model for YBaMn2O5 in this refinement was Y2O3 (Aldrich 99.99%) and MnO2 (Aldrich 99.99%) at 1350 °C taken from the previous X-ray study,3 the structural parameters for 8 h.BaO and MnO were prepared by heating BaO2 are given in Table 1. The impurity phases were fitted using (Aldrich 99.99%) and MnO2 under flowing H2–N2 at 900 °C models were taken from Yakel et al.,8 Grosse et al.,9 and Sasaki for 8 h. These precursors were then ground together, pressed et al.10 for YMnO3, MnO and Ba2SiO4 respectively.A good into 13 mm pellets, wrapped in gold foil and sealed in an fit to the data was obtained using a pseudo-Voigt peak shape evacuated silica tube which was heated at 1000 °C for 4 days. function and a cosine Fourier series background function, This procedure was repeated twice. After each heating the tube despite secondary phases being present. The refined phase was quenched into air to avoid possible decomposition of proportions by mass were YBaMn2O5 32.1%, YMnO3 36.3%, YBaMn2O5 on cooling. Despite many attempts with varying Ba2SiO4 24.9% and MnO 6.7%.The refined YBaMn2O5 experimental conditions it has not been possible to prepare a structural model is in good agreement with that of the previous phase pure sample of this phase.YMnO3 and MnO were X-ray diVraction study.3 Fig. 2 shows the calculated, observed observed in the X-ray powder diVraction pattern and and diVerence plots for the 300 K data. The oxygen content additional peaks subsequently observed in the neutron profile was confirmed as being stoichiometric. Results of the refinement are given in Table 1 and Fig. 4 and 5 ( later). Both the 300 and 190 K data showed a broad peak at 2h= †E-mail: jpa14@cam.ac.uk J. Mater. Chem., 8(5), 1291–1294 1291Table 1 Refinement results for YBaMn2O5 a with e.s.d.s in parentheses T/K 100 140 190 300 x2 3.2 3.9 2.9 2.5 RWP(%) 6.9 7.6 6.5 6.0 cell parameters a/A ° 3.9114(2) 3.9121(2) 3.9146(2) 3.9186(2) c/A ° 7.6241(1) 7.6287(6) 7.6351(6) 7.6540(5) atomic parameters Uiso/A° 2 0.0009(6) 0.0067(5) 0.0017(5) 0.0041(6) Mn z 0.234(1) 0.233(1) 0.236(1) 0.236(1) m/mB 1.78(2) 1.17(6) — — O(1) z 0.1868(5) 0.1870(5) 0.1858(4) 0.1870(4) distances BaMO(1)×8 3.087(3) 3.087(3) 3.096(3) 3.095(3) BaMO(2)×4 2.7658(1) 2.7662(1) 2.7681(1) 2.7709(1) YMO(1)×8 2.419(2) 2.421(2) 2.418(2) 2.426(2) MnMO(1)×4 1.989(7) 1.988(2) 1.995(2) 1.995(2) MnMO(2)×1 2.026(9) 2.034(10) 2.012(8) 2.020(8) Mn,Mn×1 3.57(2) 3.56(2) 3.61(2) 3.61(2) angles O(1)MMnMO(1)×4 88.10(9) 88.19(10) 87.85(9) 87.97(8) O(1)MMnMO(1)×2 159.0(5) 159.5(6) 157.6(5) 158.3(5) O(1)MMnMO(2)×4 100.5(3) 100.2(3) 101.2(2) 100.9(2) MnMO(1)MMn 159.0(2) 159.5(6) 157.6(5) 158.3(5) MnMO(2)MMn 180 180 180 180 aSpace group P4/mmm, Y 1(c) 1/2,1/2,0, Ba 1(d) 1/2,1/2,1/2, Mn 2(g) 0,0,z, O(1) 4(i) 1/2,0,z, O(2) 1(b) 0,0,1/2.Fig. 2 Calculated, observed and diVerence neutron diVraction plots for 300 K. The reflection markers from bottom to top represent YBaMn2O5, YMnO3, Ba2SiO4 and MnO respectively. 20.6°, which is not attributable to any impurity phase and can be indexed as (1/21/2 1) on the nuclear cell. The extent of the ordering was calculated using the Scherrer equation (1): t= 0.9l Ó(C2 pos-C2 std)cosh (1) where t is the domain size, l is the neutron wavelength, Cpos is the full width at half maximum (FWHM) of the peak, Cstd is the FWHM of a standard peak from the diVraction pattern (which is assumed to be instrumentally resolution-limited) and h is the peak position.This gave t#50 A ° . Fig. 3 DiVracted neutron intensity at the 1/21/2 1 position (2h=20.5°) Below 167 K a magnetic peak appears at the (1/21/21) for YBaMn2O5 as a function of temperature.The Rietveld fits to the position showing that the magnetic cell parameters are background and magnetic peaks are shown. The intensity scale for Ó2a×Ó2a×c, where a and c are of those of the nuclear model. the 140, 190 and 300 K data is twice that for the 100 K data.In the 140 and 100 K data the sharp peak due to long range magnetic order is superimposed on the broad valence order peak as shown in Fig. 3. The magnetic model in the refinement Discussion is as shown in Fig. 1, but with the Mn moments constrained to have equal magnitude, as attempts to refine the Mn2+ and The basic structure of YBaMn2O5 and the proposed valence and magnetic order models are confirmed by this variable Mn3+ moments independently were unsuccessful. 1292 J. Mater. Chem., 8(5), 1291–1294temperature neutron diVraction study. The oxygen content of YBaMn2O5 was confirmed by the refinement to be five oxygen atoms per unit cell. The presence of a broad peak above TC and a sharp peak below, indexing on a Ó2a×Ó2a×c supercell, give evidence for long range magnetic order below TC but a valence ordering that extends only over ca. 50 A ° domains. Fig. 4 shows the variation of YBaMn2O5 lattice parameters a and c with temperature. Both parameters show a decrease with decreasing temperature which is significantly steeper about TC. Fig. 5(a) shows the variation of MnMO(1) and MnMO(2) bond lengths with temperature.MnMO(1) decreases below TC whereas MnMO(2) increases. These changes in bond lengths arise due to a change in the MnMO(1)MMn bond angle which shows an anomalous Fig. 6 Schematic one-dimensional representation of the valence and magnetic ordering of Mn2+ and Mn3+ ions in YBaMn2O5 through oxygen atoms (not shown). The broken lines represent domain walls due to disruption of the valence ordering; (a) shows the antiferromagnetic ordering of adjacent Mn2+ (5 mB) and Mn3+ (4 mB) moments, (b) shows the same situation with the moments separated into the average 4.5 mB and diVerence (±0.5 mB) components.The former order antiferromagnetically and are unaVected by the domain walls, whereas the latter order ferromagnetically and the direction of the magnetisation is switched at the domain boundary. Fig. 4 Variation with temperature of a and c parameters of YBaMn2O5 increase as the temperature decreases below TC [Fig. 5(b)]. These changes evidence an exchange striction at TC. The changes in lattice parameters, bond lengths and angles promote superexchange via the MnMO(1)MMn bridges mainly by increasing the MnMO(1)MMn bond angle towards the most favourable value of 180 °.Fig. 6 shows how the valence ordering can be disrupted without aVecting the long range antiferromagnetic ordering. The ideal average Mn moment is 4.5 mB and for the Mn2+ moment there is an additional 0.5 mB in the same direction as the average, but for the Mn3+ ion the extra 0.5 mB opposes the average. Antiferromagnetic MnMOMMn interactions lead to ferromagnetic ordering of the 0.5 mB components with unit cell a×a×c, but antiferromagnetic ordering of the average 4.5 mB with unit cell Ó2a×Ó2a×c.If the valence ordering is disrupted, leading to two adjacent Mn3+ or Mn2+ ions, Fig. 6 shows that the direction of the ferromagnetic component is reversed but the antiferromagnetic order is ideally unaVected. However, the average antiferromagnetic moment at 100 K is only 1.8 mB, showing that valence disorder also causes some disruption of this order.The ferromagnetic component should give rise to weak broad magnetic reflections at the hkl positions of the structural cell, and so are not seen beneath the structural Bragg intensities. In conclusion, YBaMn2O5 shows long range structural order (periodicity a×a×c) and short range valence order with a Ó2a×Ó2a×c superstructure up to at least 300 K.Below the ferrimagnetic ordering temperature of 167 K, these structural features respectively give rise to a long range antiferromagnetically ordered component (periodicity Ó2a×Ó2a×c) and a short range ferromagnetic component that is commensurate with the basic cell. Fig. 5 Variation with temperature of (a) MnMO(1) and MnMO(2) bond lengths and (b) MnMO(1)MMn bond angle in YBaMn2O5 We thank EPSRC for financial support for J.A.M.and pro- J. Mater. Chem., 8(5), 1291–1294 12935 W. Zhou, Adv.Mater., 1993, 5, 735. vision of neutron facilities at the ILL, and Dr Paolo Radaelli 6 H. M. Rietveld, J. Appl. Crystallogr., 1969, 35, 1. for technical assistance with data collection. 7 A. C. Larson and R. B. Von Dreele, Los Alamos National Laboratory Rep. No. LAUR-86-748, 1987. 8 H. L. Yakel, W. C. Koehler, E. F. Bertaut and E. F. Forrat, Acta References Crystallogr., 1963, 16, 957. 9 H. P. Grosse and E. Tillmanns, Cryst. Struct. Commun., 1974, 1 G. H. Jonker and J. H. Van Santen, Physica (Utrecht), 1950, 16, 337. 3, 599. 10 S. Sasaki, K. Fujino, Y. Takeuchi and R. Sadanaga, Acta 2 C. Zener, Phys. Rev., 1951, 82, 403. 3 J. P. Chapman, J. P. Attfield, M. Molgg, C. M. Friend and Crystallogr., 1980, 36, 904. T. P. Beales, Angew. Chem., 1996, 35, 2482. 4 L. Er-Rakho, C. Michel, P. Lacorre and B. Raveau, J. Solid State Paper 8/00605I; Received 22nd January, 1998 Chem., 1988, 73, 531. 1294 J. Mater. Chem., 8(5), 1291–1294

ISSN:0959-9428    

DOI:10.1039/a800605i

出版商:RSC    

年代:1998

数据来源: RSC