摘要:

Synthesis, phase stability and electrical characterisation of BINAVOX solid solutions Craig J. Watson, Alison Coats and Derek C. Sinclair Chemistry Department, University of Aberdeen, Meston Walk, Aberdeen, UK AB24 3UE The compositional range of ‘Bi4V2O11’ solid solutions containing Na has been determined by means of a phase diagram study using X-ray diVraction and electron probe microanalysis. The locus of the solid solution indicates that Na replaces V in the crystal structure and the BINAVOX family has the overall formula Bi4+yV2-y-xNaxO11-y-2x. The solid solution limits are temperature dependent and for samples prepared and air-cooled from 830 °C the limits are defined as 0.00<x<0.18 and 0.05<y<0.19.BINAVOX compositions with x>0.10 are thermodynamically unstable below ca. 720 °C but can be stabilised kinetically by rapid cooling from temperatures above ca. 750 °C. ac Impedance measurements demonstrate that BINAVOX materials are electrically inhomogeneous, exhibit temperature- and time-dependent conductivities and do not oVer any significant advantages over the parent solid solution. The bismuth vanadate, Bi4+yV2-yO11-y has attracted a lot of to use.Reaction mixtures totalling 3–4 g were weighed and attention as the parent phase for a family of oxide ion mixed into a paste with acetone using an agate mortar and conductors known as BIMEVOX.1–4 Bi4+yV2-yO11-y is a pestle. In order to limit volatilisation of the reagents, approxisolid- solution phase whose compositional limits are highly mately two thirds of the powder for each sample was coldtemperature dependent, covering a compositional range from pressed into pellets, placed in Au foil boats and covered with ca. 0<y0.22 at 880 °C.5 Three polymorphs are known to the remaining powder. A heating sequence of 500 °C for 2 h, exist and the high-temperature c polymorph exhibits high 650 °C for 2 h, 800 °C for 12 h, regrinding and repelletising, conductivity mainly due to oxide ions.The explanation for the then 750 °C for 2 h and 830 °C for 12 h was found to be high conductivity is that the c polymorph has a layered adequate to obtain equilibrium. Pellets for EPMA and ac structure of alternating Bi2O22+ and VO3.52- layers and the impedance measurements were cold-pressed, covered with perovskite-like layers VO3.5%0.52- are oxygen deficient. powder of similar composition and sintered at 830 °C for a Transformation to the b and a polymorphs at lower tempera- further 12 h.tures results in ordering of the oxygen vacancies and the Phase purity was determined by X-ray diVraction using a conductivity decreases markedly. Partial substitution of Cu or Hagg Guinier camera or a Stoe-Stadi diVractometer using Cu- Ni for V suppresses the c�b/a transitions thus stabilising the Ka1 radiation.The final composition of the BiNaVOX matehigh- temperature c polymorph of Bi4+yV2-yO11-y, giving con- rials was determined using a Cameca SX51 electron microprobe ductivities as high as 1×10-3 V-1 cm-1 at 240 °C.1 analyser with an incident beam energy of 20 kV and a current The crystal chemistry of Bi4+yV2-yO11-y is still not clearly of 50 nA.All samples were polished to <1 mm and carbon resolved owing to problems associated with stoichiometry, coated. The standards used were Bi2CuO4 for Bi-La, V2O5 for polymorphism, twinning6 and incommensurate supercells.3 V-Ka and NaAlSi3O8 for Na-Ka. Oxygen was calculated by Recently, Huve et al.7 reported that the symmetry of the a stoichiometry.ac Impedance measurements from ca. 25 to polymorph depends on the purity of the V2O5 employed in 800 °C were conducted using a Hewlett Packard 4192A the synthesis. High-purity V2O5, free from Na/K contamiimpedance analyser over the frequency range 5 Hz–5 MHz nation gave rise to a monoclinic cell, space group A2/m, with an applied voltage of 100 mV. Electrodes were fabricated whereas reagents with trace amounts of Na/K result in the from Pt organopaste; electroded pellets were fired at 800 °C commonly reported orthorhombic cell, space group Amam.overnight to decompose the paste and harden the Pt residue. Despite the complexities associated with the crystal chemistry Pellets were attached to the Pt measuring leads of a conduc- of Bi4+yV2-yO11-y the number of BINAVOX materials has tivity jig and placed in a horizontal tube furnace whose grown considerably with over seventeen diVerent families being temperature was controlled and measured within ±3 °C.reported. In this paper we describe, for the first time, the BINAVOX family. The solid-solution limits at 830 °C are determined via a combination of X-ray diVraction (XRD) results and electron probe microanalysis (EPMA) and the thermal Results and Discussion stability and electrical properties of the BINAVOX materials The compositions studied have general formula are also discussed. Low levels of Na-doping, ca. 3 mol%, are Bi4+yV2-y-xNaxO11-y-2x. Although XRD was adequate to eVective for kinetic stabilisation of the c polymorph on rapid determine whether any given composition was single phase, cooling from 830 °C, but these materials are metastable and EPMA was employed to determine the final composition(s) of transform slowly to the a polymorph on post-annealing at, e.g.the phases present in selected samples. Combining both sets 350 °C. On slow cooling or annealing at ca. 650–720 °C for ca. of results enabled the solid-solution area at 830 °C in air to be 48 h the c-BINAVOX solid solutions decompose into a threedetermined, as shown on the composition triangle phase mixture containing a-BINAVOX.ac Impedance spec- Bi2O3–Na2O–V2O5 in Fig. 1. Closed circles represent the final troscopy measurements show the bulk conductivity behaviour compositions of single-phase BiNaVOX samples, half-filled of BINAVOX materials to be electrically inhomogeneous and circles represent the starting composition of samples which very dependent on thermal history.were determined to be phase mixtures and filled triangles represent the final composition of the BINAVOX phase within Experimental phase mixtures ( half-filled circles). EPMA established that there were no significant problems associated with volatility Bi2O3 (99.99%), V2O5 (99.6%) and Na2CO3 (99.99%) reagents were dried at 300 °C overnight and stored in a desiccator prior of the reagents.For example, a sample of nominal starting J. Mater. Chem., 1997, 7(10), 2091–2095 2091Fig. 1 Compositional extent of BINAVOX by EPMA for samples prepared and air-cooled from 830 °C. Closed circles represent phase- Fig. 2 X-Ray diVractograms of Bi4+yV2-x-yNayO11-x-2y solid solupure compositions, half-filled circles represent phase mixtures and tions prepared and air-cooled from 830 °C closed triangles represent the composition of BINAVOX in phase mixtures. The inset shows the loci of solid solutions for diVerent Bi4.10(2)V1.90(2)O10.90 (x=0, y=0.1). This is consistent with the substitution mechanisms. suggested substitution mechanism of replacing small V ions by larger Na ions.Compositions with 0.05<x<0.15 could not composition, Bi4.08V1.77Na0.15O10.62 was analysed to have a be indexed satisfactorily on any single polymorph, owing to final composition of Bi4.10(2)V1.76(2)Na0.15(2)Od. The results variations in peak intensities and problems associated with clearly indicate that Na has been incorporated into the crystal peak convergence/asymmetry, such as shown for the doublet lattice and that the solid-solution limits can be defined as at ca. 32° for x=0.06 and 0.10, Fig. 2(b) and (c), respectively. 0<x<0.18 and 0.05<y<0.19 for samples prepared and air- Compositions with x0.05 indexed as a polymorphs and cooled from 830 °C. diVerential thermal analysis (DTA) showed the presence of an Several interesting features emerge from the results. First, exotherm at ca. 450 °C on heating and an endotherm at ca. phase-pure, stoichiometric Bi4V2O11 (y=0) cannot be prepared 350 °C on cooling, consistent with a reversible a=c transition. in air at 830 °C. The parent phase of the samplwith the No thermal event was detected by DTA between 25 and 700 °C stoichiometric composition, y=0, was found by EPMA to be for samples with x>0.05. These results indicate a gradual Bi-rich with y=0.05 and trace amounts of a secondary phase change in polymorphism from a to c on increasing x.BiVO4 were detected. This confirms the work of Lee et al.4 XRD patterns obtained for BINAVOX solid solutions with who established that excess Bi is always required to prepare intermediate values of x, ca. 0.05<x<0.15 depend on thermal single-phase Bi4+yV2-yO11-y below ca. 850 °C. history, as shown for Bi4.13(2)V1.78(2)Na0.09(2)O10.69 (x=0.13, Secondly, on the basis of the solid solution locus, Fig. 1, the y=0.09) in Fig. 3. Rapid cooling from elevated temperatures phase diagram results suggest that Na is predominantly incorenables the c polymorph to be ‘preserved’ at ca. 25°C, as porated onto V rather than Bi sites within the crystal lattice. shown by the single peak at ca. 32°, Fig. 3(a). Annealing the There is considerable disorder in the coordination of oxygen same sample for 5 days at 350 °C leads to a clear doublet at ions around the V-sites within the lattice and the possibility ca. 31–33°, Fig. 3( b), consistent with the existence of the a exists that the large Na ions are located in distorted tetrahedral polymorph. No change in phase assemblage was detected by sites within the perovskite-like layers. Nevertheless, this result EPMA for any BINAVOX material prepared at 830 °C and is rather surprising given the large diVerence in ionic radii of subsequently annealed below ca. 600 °C.These results clearly V5+ and Na, 0.54 and 1.13 A ° ,8 respectively and needs confirdemonstrate that the c polymorph of BINAVOX materials mation by crystallographic studies. The low solubility limit of with x>0.05 in samples quenched from high temperature Na in Bi4+yV2-yO11-y makes it is diYcult to use the phase transform to the more stable a polymorph on low-temperature diagram to establish the precise substitution mechanism, howannealing, typically at ca. 350 °C. Without undertaking extens- ever, it is clear that the BINAVOX phase diagram diVers ive crystallographic studies on samples subjected to a variety significantly from those obtained with other large dopant of heat treatments, it is diYcult to determine whether the cations such as Sr2+ and Pb2+.Such diagrams9 show extensive change from a to c on increasing x is continuous or if a two- solid solutions extending in the -y direction, inferring that phase region exists at intermediate x which separates a and c large divalent cations substitute onto the Bi sites. regions at low and high x values, respectively. The a and c polymorphs can usually be distinguished by The above comments on polymorphism refer to materials XRD in the range 2h 31–33°.The presence of a doublet at ca. 32° is ascribed to (020) and (200) reflections of the lower symmetry, orthorhombic a polymorph whereas a singlet at ca. 32.5° is assigned to the (110) reflection of the higher symmetry, tetragonal c polymorph. For samples with y=0.1, it is clear from XRD that increasing the Na content (x value) is eVective in stabilising the c polymorph on air-cooling from 830 °C, as shown by the convergence of the (020)–(200) doublet at ca. 31–33° for x=0, Fig. 2(a), into a singlet for x=0.15, Fig. 2(d). In general, samples with high x values could be satisfactorily indexed as c polymorphs and the unit cell expands in the c direction compared with undoped materials. For example, the lattice parameters of c-Bi4.10(2)V1.75(2)Na0.15(2)O10.60 (x=0.15 and y=0.1) are a=3.941(9) and c=15.349(5) A ° , compared Fig. 3 X-Ray diVractograms for Bi4.13(2)V1.78(2)Na0.09(2)O10.69 quenched from 850 °C (a) and post-annealed for five days at 350 °C (b) with a=5.518(2), b=5.593(2) and c=15.239(5) A ° for a- 2092 J. Mater. Chem., 1997, 7(10), 2091–2095quenched from high temperature and subsequently annealed c polymorph and for the decomposition products including a- BiNAVOX.It is clear that Na is exsolved from c- at relatively low temperature, 350 °C. At intermediate temperatures, diVerent behaviour occurs. Bi4.10(2)V1.75(2)Na0.15(2)O10.60 on annealing at 650–720 °C to form a phase mixture including a-Bi4.02(5)V1.88(4)Na0.10(4)O10.78. Unlike many other BIMEVOX families, c-BINAVOX solid solutions with x>0.10 decompose on annealing between 650 Good quality analysis of the impurity phases was restricted because of their small grain size, Fig. 5, however, both are Bi- and 720 °C into a mixture of phases containing a-BINAVOX of low x values. XRD results for Bi4.10(2)V1.75(2)Na0.15(2)O10.60 rich with respect to the BINAVOX materials.Further details of these phases are discussed elsewhere.10 (x=0.15) quenched onto a brass block from 830 °C, Fig. 4(a), and annealed at 650 °C for 72 h, Fig. 4(b), clearly demonstrate The decomposition of BINAVOX materials with x>0.10 takes place only over a narrow temperature range, ca. a mixture of a-BINAVOX and peaks associated with secondary phases. 650–720 °C and is fully reversible: reheating above ca. 720 °C for 1 h results in reformation of single-phase BINAVOX. This EPMA revealed the presence of three phases in the decomposed sample as shown by the back-scattered image in Fig. 5. suggests that BINAVOX materials with x>0.10 are thermodynamically stable only above ca. 720 °C but may be preserved Table 1 lists the atom% of Bi, V and Na for the single-phase to room temperature, where they are kinetically stable, by rapid cooling.BINAVOX materials with x<0.10 did not decompose on prolonged annealing between ca. 650 and 720 °C, suggesting that a-BINAVOX materials are thermodynamically stable under the conditions tested. Given the obvious complexity associated with stability and polymorphism in the BINAVOX system, we chose not to indicate the polymorphs on the phase diagram, Fig. 1 and stress that this diagram is relevant only for materials which have been aircooled from 830 °C. Like many other BIMEVOX systems, the BINAVOX solid solution limits are highly temperature dependent with the maximum solubility in the c polymorph occurring close to melting temperatures. Three types of conductivity data set were collected for the BINAVOX solid solutions; (i) above 720 °C where all samples were thermodynamically stable c polymorphs, (ii) between ca. 150 and 650 °C for samples air-cooled from 800 °C, and (iii ) as a function of time at ca. 250 °C for samples quenched from Fig. 4 X-Ray diVractograms for Bi4.10(2)V1.75(2)Na0.15(2)O10.60 800 °C. quenched from 830 °C (a) and post-annealed at 650 °C for 72 h ( b).In general, impedance plane plots below 400 °C consist of a ×indicate unindexed peaks associated with secondary phases. single, semicircular arc and a low-frequency ‘spike’. There was no evidence of a grain boundary arc and the low-frequency eVects are attributable to ionic polarisation and diVusionlimited phenomena at the electrode and support the idea that conduction is mainly by means of ions.The inclined spike is similar to that expected for a Warburg impedance with an ideal slope of 45°. (i) Conductivity data above 720 °C Total conductivities for single-phase BINAVOX solid solutions were extracted from the inverse of the low-frequency intercept of the electrode spike with the real axis of impedance plane plots at 725, 750, 775 and 800 °C. The results demonstrated that all single-phase compositions had similar but slightly lower conductivity values and similar activation energies to the c polymorph in undoped materials, as shown for a variety Fig. 5 A back-scattered electron image of the composition of compositions in Fig. 6. Thus, Na-doping does not enhance Bi4.10(2)V1.75(2)Na0.15(2)O10.60 after being annealed at 650 °C for 72 h.Dark regions are a-Bi4.02(2)V1.88(2)Na0.10(2)O10.78; grey and white the high-temperature conductivity of c-Bi4+yV2-yO11-y regions are the minor phases 1 and 2, respectively. materials. Table 1 Composition analysis (atom%) before and after the decompo- (ii) Conductivity data between 150 and 650 °C sition of a c-BINAVOX solid solution at 650 °Ca The conductivity changes irreversibly on thermal cycling before between 150 and 600 °C, Fig. 7. The high conductivity values decomposition after decomposition obtained below ca. 400 °C on the initial heating cycle of c- BINAVOX materials are not reproduced on subsequent coolc- BINAVOX a-BINAVOX minor 1 minor 2 ing and reheating. Given the metastable nature of BINAVOX element av. s av. s av. av. materials with x>0.10 and the tendency to transform from the c to the a polymorph at 350 °C, as shown by the XRD Bi 24.68 0.07 23.98 0.25 25.63 26.87 V 10.56 0.04 11.20 0.18 9.75 9.35 results in Fig. 3, it is not surprising that conductivities depend Na 0.89 0.05 0.56 0.04 1.20 0.06 on thermal treatment. O 63.87 0.03 64.26 0.08 63.42 63.71 Conductivities of Bi4.13(2)V1.78(2)Na0.09(2)O10.69 and c- Bi4.10(2)V1.75(2)Na0.15(2)O10.60 are shown in Fig. 7 (conductivities aatom% shown here is normalised to give total %=100. Actual total for undoped a-Bi4.06(2)V1.94(2)O10.94 are indicated as a dotted mass% was 99.64 and 100.83 for the c- and a-BINAVOX phases, line for comparison). It is clear from the Arrhenius plots in respectively. Data of ten points were averaged in determining the composition of the BINAVOX phases; s is standard deviation.Fig. 7 that the samples have similar conductivities and acti- J. Mater. Chem., 1997, 7(10), 2091–2095 2093(iii) Conductivity data at ca. 250 °C for samples quenched from 800 °C The bulk conductivity at ca. 250 °C for a pellet of c- Bi4.10(2)V1.81(2)Na0.15(2)O10.60 which had been rapidly quenched from 800 °C was extracted from the inverse of the low-frequency intercept of the asymmetric semicircular arc on the real axis of impedance plane plots, Fig. 8. The associated capacitance of the arc was calculated to be 33 pF cm-1 using the relationship vRC=1 (where v=2pf and is the angular frequency) at the arc maximum and is consistent with a bulk or intragranular response. The bulk arc resistivity increased as a function of time and Fig. 9 shows the smooth decrease in bulk conductivity from 20 to 4.5 mS cm-1 over a period of ca. 150 h. Conductivity measurements are clearly very sensitive to oxygen ordering as the highly conducting metastable c polymorph transforms, albeit rather slowly, to the a polymorph. This decrease in conductivity is consistent with the observed line broadening/ splitting in XRD patterns when annealing rapidly quenched c-BINAVOX materials at moderate temperatures, as shown in Fig. 3 for Bi4.13(2)V1.78(2)Na0.09(2)O10.69. Fig. 6 Arrhenius plots for various BINAVOX materials above 725 °C In order to further investigate the bulk conductivity characteristics of BINAVOX materials, ac impedance data were replotted in the form of combined spectroscopic plots of the imaginary components of the complex impedance and electric modulus formalisms, Z and M, respectively. Such plots are well established for probing the electrical homogeneity of many electroceramics11 and reveal that the bulk characteristics of BINAVOX materials are inhomogeneous.Fig. 10 shows a Fig. 7 Arrhenius plots for various BINAVOX materials over the Fig. 8 Complex impedance plane plot for Bi4.10(2)V1.75(2)Na0.15(2)O10.60 temperature range 150 to 650 °C (–, cooling cycle: x=0, y=0.06; quenched from 800 °C and annealed at 253 °C for 143 h.Closed circles $, cooling cycle: x=0.09, y=0.13; #, initial heating cycle: x=0.15, identify selected frequencies on a logarithmic scale, e.g. 5=105 Hz. y=0.10) vation energies over the temperature range ca. 320–650 °C.Similar results were obtained for all single-phase BINAVOX compositions. Initial heating-cycle data for c-Bi4.10(2)V1.75(2)Na0.15(2)O10.60 below 300 °C demonstrate a characteristic feature of BINAVOX materials with x>0.10, namely enhanced, lowtemperature conductivity values which are time dependent and irreproducible on thermal cycling. For ac impedance measurements, samples are allowed to equilibrate for ca. 1 h between successive temperatures and are therefore maintained at elevated temperatures for long periods during thermal cycling.The conductivity of c-Bi4.10(2)V1.75(2)Na0.15(2)O10.60 below ca. 300 °C was over an order of magnitude higher on the initial heating cycle compared with that on any subsequent heating or cooling cycle and at 305 °C on the initial heating cycle decreased from 0.11 to 0.018 mS cm-1 overnight, ca. 12 h, Fig. 7. This time dependent decrease is clearly associated with the metastable nature of these materials and is probably associated with their tendency to undergo oxygen re-ordering at low temperatures. In order to further investigate this phenomenon, samples were Fig. 9 Variation in conductivity as a function of time for rapidly quenched from 800 °C and variations in the bulk Bi4.10(2)V1.75(2)Na0.15(2)O10.60 quenched from 800 °C and annealed at ca. 250 °C conductivity monitored as a function of time at ca. 250 °C. 2094 J. Mater. Chem., 1997, 7(10), 2091–2095extremely broad M spectra but such plots clearly show the bulk conductivity characteristics in these materials to be complex and inhomogeneous.In conclusion, although Na can be incorporated into the Bi4+yV2-yO11-y lattice the BINAVOX family exhibit complex thermal stability behaviour and are unsuitable for producing stable, high oxide-ion conducting materials for practical applications. The authors would like to thank the University of Aberdeen for an M.Sc. studentship for C.J.W., EPSRC for financial support for the EPMA facility and Professor Tony West for helpful discussions.References 1 F. Abraham, J. C. Boivin, G. Mairesse and G. Nowogrocki, Solid Fig. 10 A combined Z and M spectroscopic plot for State Ionics, 1990, 40/41, 934. Bi4.10(2)V1.75(2)Na0.15(2)O10.60 quenched from 800 °C and annealed at 2 T. Iharada, A. Hammouche, J. Fouletier, M. Kleitz, J. C. Boivin 253 °C for 143 h and G.Mairesse, Solid State Ionics, 1991, 48, 257. 3 K. B. R. Varma, G. N. Subbanna, T. N. Guru Row and C. N. R. Rao, J.Mater. Res., 1990, 5, 2718. combined Z and M plot for c-Bi4.10(2)V1.81Na0.15O10.60 after 4 C. K. Lee, B. H. Bay and A. R. West, J.Mater. Chem., 1996, 6, 331. ca. 143 h at 250 °C on quenching from 800 °C. For an ideal 5 C. K. Lee, D. C. Sinclair and A.R. West, Solid State Ionics, 1993, Debye response, the frequency maxima of Z and M peaks 62, 193. 6 F. Abraham, M. F. Debreuille-Gresse, G. Mairesse and should be coincident and the half-height peak widths 1.14 G. Nowogrocki, Solid State Ionics, 1988, 28/30, 529. decades on a log( f ) scale. Although it is not uncommon for 7 M. Huve, R. N. Vannier, G. Nowogrocki, G. Mairesse and the fmax values of Z and M to be separated by up to one G. V. Tendeloo, J.Mater. Chem., 1996, 6, 1339. order of magnitude in good solid electrolytes such as Na- 8 R. D. Shannon, Acta Crystallogr., Sect. A, 1976, 32, 751. b-Al2O3 12 it can clearly be seen that the M peak is exception- 9 C. K. Lee, G. S. Lim and A. R.West, J.Mater. Chem., 1994, 4, 1441. ally broad, with a half-height peak width of ca. 2.56 decades 10 C. J.Watson, M.Sc. Thesis, Aberdeen University, 1997. 11 D. C. Sinclair and A. R. West, J. Appl. Phys., 1988, 66, 3850. on a log( f ) scale. Such a response could be associated with 12 I. M. Hodge, M. D. Ingram and A. R. West, J. Electroanal. Chem., the metastability of the BINAVOX materials as they transform 1976, 74, 125. from the disordered c to the ordered a polymorph on low- 13 E. Pernot, M. Anne, M. Bacmann, P. Strobel, J. Fouletier, temperature annealing or it may be associated with the intrin- R. N. Vannier, G. Mairesse, F. Abraham and G. Nowogrocki, Solid sic, anisotropic nature of the conduction properties of State Ionics, 1994, 70/71, 259. BIMEVOX materials, as shown in single-crystal studies.13 Further studies are in progress to clarify the origin of the Paper 7/03629I; Received 27thMay, 1997 J. Mater. Chem., 1997, 7(10), 2091–2095 2095

ISSN:0959-9428    

DOI:10.1039/a703629i

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Preparation, crystal structure, and reducibility of K2NiF4 type oxides Sm2-xSrxNiO4+d Hui Lou,*a Yuping Ge,a Ping Chen,a Minghua Mei,a Futai Maa and Guanglie Lu� b aDepartment of Chemistry, Hangzhou University, Hangzhou 310028, China bCentral L aboratory, Hangzhou University, Hangzhou 310028, China K2NiF4 type compounds Sm2-xSrxNiO4+d (0.4x1.2) have been synthesized by a citric acid complex decomposition method.Rietveld refinement of the powder X-ray diVraction data showed that there was a crystal system transformation from orthorhombic Fmmm to tetragonal I4/mmm at around x=0.6. The Sr substitution caused a drastic shift of the OII ions along the c axis from Sm(Sr) towards Ni whilst scarcely aVecting the NiO4 network in the basal plane. The distortion of the NiO6 octahedron decreased, while the average bond valence of Ni increased with increasing x.K2NiF4 type rare-earth complex oxides, A2BO4, consist of Crystal structure characterization alternating layers of ABO3 perovskite and AO rock-salt struc- XRD data for Rietveld refinement were collected on a Rigaku tures along the c-axis.1–3 They were widely studied owing to D/max-3B powder diVractometer with Bragg–Brentano their high temperature stability, two-dimensional conductivity, geometry using graphite-monochromated Cu-Ka radiation antiferromagnetic properties, semiconductor–metal transition (40 kV×30 mA) and a scintillation detector.The intensity data properties and superconductivity.1–5 These properties are were collected at 25 °C over a 2h range 10–130° with a step closely related to the types of the ions A and B, interatomic interval of 0.02° and a counting time of 8 s per step.The d distances in the crystal, and the valence state of the transitionspacings were corrected for systematic errors by calibration metal ions B which can be altered by substitution of other with standard silicon powder (a=0.543 09 nm). The Rietveld ions such as Sr2+ for A.6,7 It is necessary to determine precise structure refinement was performed using the program crystal structures so as to further understand the relationships WYRLET (M.Schneider, a modified version of the program between properties and the crystal structure. The Rietveld by Wiles and Young, 1981).13 The starting model used in the method, a profile refinement method of analyzing X-ray powder refinement was the space group Fmmm or I4/mmm with diVraction data by fitting the entire diVraction pattern to a structural parameters of a K2NiF4 type compound. Peaks were calculated one, has been widely used to refine crystal structure.8 modeled using the pseudo-Voigt profile function in which a Recently, an application of the Rietveld method to refine the peak asymmetry parameter was included for peaks up to 2h= crystal structure of La2-xSrxNiO4 and Nd2-xSrxNiO4 has 50°.The background parameters were modeled using a refinbeen reported.9–11 able fourth-order polynomial. The occupation number of atoms A tolerance factor, t, is usually used to predict the possibility for all sites were fixed and were not refined. Table 1 lists the of K2NiF4 structure formation.Ganguli12 proposed a criterion variables and the number of these in the refinement. according to which the tetragonal structure is stable when the tolerance factor, (t=rA/rB), lies between 1.7 and 2.4. Because Determination of Ni3+ content of the limitation of the tolerance factor, neodymium is the smallest rare earth metal ion forming a K2NiF4 type lanthanide The percentage Ni3+ content, relative to the total nickel in the nickelate.11 Samarium, which is smaller than Nd, can not form sample, was determined iodometrically according to the nickelate compounds with the K2NiF4 type structure.By method described in ref. 6 and 14. In a three-necked spherical partial substitution of Sr2+ for Sm3+, however, K2NiF4 type flask kept under an N2 flow and containing 30 cm3 of 0.5 mol compounds Sm2-xSrxNiO4+d can also be prepared.In this dm-3 HCl and 600 mg of KI (both in excess), ca. 150 mg of study, we prepared Sm2-xSrxNiO4+d (0.4x1.2), refined accurately weighed sample was added and the flask was kept their crystal structure by X-ray Rietveld analysis, and measured in the dark for ca. 5 min. After dissolution of the solid, the their IR spectra and reducibility.The eVects of Sr substitution solution was quickly titrated with 0.01 mol dm-3 Na2S2O3. on the structural parameters are discussed. From the values of the mean oxidation number, the percentage Experimental Table 1 Parameters and number of parameters in the refinement Preparation parameters number Sm2-xSrxNiO4+d (x=0, 0.2, 0.4, 0.5, 0.6, 0.8, 1.0, 1.2, 1.4, 1.6) zero-point parameter 1 samples were prepared by evaporating aqueous solutions of background coeYcients 5 the mixed metal nitrates containing an equivalent amount of scale factor parameter 1 citric acid to obtain a gel, followed by decomposion at 920 K fractional atomic coordinates 2 for 4 h.Samples were then ground, pelletized and calcined in isotropic thermal parameters 4 air at 1473 K for 10 h.FWHM parameters 3 cell constants 2 for I4/mmm, 3 for Fmmm asymmetry parameter 1 Composition determination mixing parameter 1 The metal-ion composition was determined by inductively preferred orientation parameter 1 coupled plasma spectrometry (ICP) (USA Leeman, Plasma total parameters 21 for I4/mmm, 22 for Fmmm Spec I).J. Mater. Chem., 1997, 7(10), 2097–2101 2097Table 2 The mass% of Sm, Sr and Ni ions and compositions in Sm2-xSrxNiO4+d determined by ICP (data in parentheses are calculated values) x Sm Sr Ni compositiona 0.4 60.39(60.40) 9.05(8.80) 15.02(14.73) Sm1.59Sr0.41Ni2+0.55Ni3+0.45O4.02 0.5 57.06(57.53) 11.60(11.17) 15.09(14.97) Sm1.48Sr0.52Ni2+0.47Ni3+0.53O4.01 0.6 54.43(54.57) 13.84(13.63) 15.49(15.21) Sm1.39Sr0.61Ni2+0.37Ni3+0.63O4.01 0.8 48.42(48.35) 18.68(18.78) 15.58(15.73) Sm1.20Sr0.80Ni2+0.17Ni3+0.83O4.02 1.0 41.93(41.69) 24.55(24.29) 16.45(16.27) Sm1.00Sr1.00Ni2+0Ni3+1.00O4.01 1.2 34.80(34.56) 30.45(30.20) 16.59(16.86) Sm0.80Sr1.20Ni2+0Ni3+1.00O3.97 aThe contents of Ni2+,Ni3+ and oxygen were calculated from the mean oxidation number of Ni, as measured by iodometry.of Ni3+ in the sample was readily calculated and values are be obtained owing to geometric factors and the charge balance considerations: the excess of Sr and Ni remain as SrCO3 and given in Table 2. NiO, respectively. IR measurements Structure of Sm2-xSrxNiO4+d (0.4x1.2) The IR spectra were taken with a Pekin-Elmer 683 spectrophotometer. Samples were produced in the form of KBr pellets.The powder X-ray diVraction patterns of the samples can be indexed based on orthorhombic symmetry with space group Fmmm (x=0.4–0.6) and tetragonal symmetry with space group Temperature-programmed reduction (TPR) experiments I4/mmm (x=0.8–1.2). The agreement between the observed Samples of Sm2-xSrxNiO4+d (ca. 5 mg) were placed in a quartz and calculated powder X-ray diVraction profiles of one sample reactor which was connected to a conventional TPR apparin the final refinement is shown in Fig. 1. atus.15 In order to eliminate any possible contamination, samples were first pretreated with oxygen (flow rate 30 ml Table 4 Values of the tolerance factor (t=rA/rB) of Nd2+, Sm3+ and min-1) to 1073 K at 20 K min-1 for 15 min, and then cooled Sr2+ with diVerent valence values of Ni ions based on Shannon’s to room temperature in flowing oxygen.The reduction gas ionic radii employed was 10% (v/v) hydrogen in nitrogen (flow rate 20 ml min-1). After flushing, the reactor was heated to the final Nd2+ Sm3+ Sr2+ temperature (1173 K) at 20 ml min-1. Ni2+ 1.69 1.64 1.90 Ni3+ (low spin) 2.08 2.02 2.30 Results and Discussion Ni3+ (high spin) 1.94 1.89 2.20 Compositions and phases of Sm2-xSrxNiO4+d Table 2 lists the compositions of the compounds as measured by ICP.It can be seen that the results do not deviate significantly from the starting compositions. Table 3 lists the phase characterization results of Sm2-xSrxNiO4+d obtained from XRD. The single-phase K2NiF4 type compounds were obtained within the Sr2+ substitution range from x=0.4 to 1.2.Because of the limitation of the tolerae factor (t=rA/rB), the K2NiF4 type structure compounds can only be produced for t within the range 1.7–2.4. According to the values of the tolerance factor based on Shannon’s ionic radii16 ( listed in Table 4), neodymium is the smallest rare earth metal which can form a K2NiF4 type lanthanide nickelate Nd2NiO4.Since Sm is smaller than Nd, it can not form a K2NiF4 type compound. It is possible, however, for Sm to form K2NiF4 type compounds Sm2-xSrxNiO4+d by partial substitution of alkaline-earth metals (A) with larger ionic radii which increases rA, or by oxidizing B ions to a higher valence state to decrease Fig. 1 Rietveld refinement patterns for Sm1.4Sr0.6NiO4+d . The rB.For x=0.2, Sr2+ substitution is not suYcient to incorporate observed data are indicated by dots and the calculated by the solid all the Sm3+ and Ni2+ into the K2NiF4 type structure and line overlaying them. The short vertical lines mark the positions of some Sm2O3 and NiO remain as separate phases (Table 3). possible Bragg reflections and the lower curve shows the diVerence between the observed and calculated powder diVraction patterns.When Sr2+ substitution is >1.4, monophase samples can not Table 3 Crystallographic characterization of Sm2-xSrxNiO4+d space x phase composition a/A ° b/A ° c/A ° group 0.0 Sm2O3, NiO 0.2 Sm2O3, NiO, Sm2-xSrxNiO4+d(0.2x0.4) 0.4 Sm1.6Sr0.4NiO4+d 5.3356(1) 5.3588(1) 12.3014(2) Fmmm 0.5 Sm1.5Sr0.5NiO4+d 5.3103(1) 5.3534(1) 12.3415(2) Fmmm 0.6 Sm1.4Sr0.6NiO4+d 5.3257(1) 5.3172(1) 12.3471(1) Fmmm 0.8 Sm1.2Sr0.8NiO4+d 3.7727(1) 12.2795(2) I4/mmm 1.0 Sm1.0Sr1.0NiO4+d 3.7844(1) 12.2230(2) I4/mmm 1.2 Sm0.8Sr1.2NiO4+d 3.7922(1) 12.2017(1) I4/mmm 1.4 Sm2-xSrxNiO4+d(1.2<x<1.4), SrCO3 1.6 Sm2-xSrxNiO4+d(1.2<x<1.6), SrCO3, NiO 2098 J.Mater. Chem., 1997, 7(10), 2097–2101The lattice parameters a, b and c obtained are listed in region D values for NiMOI (d a-axis), Sm(Sr)MOI, Sm(Sr)MSm(Sr)(d c-axis) and Sm(Sr)MNi(d c-axis) are all Table 3.With increasing x, c reached a maximum in the range 0.5x0.6, similar to the systems La2-xSrxNiO4 10 and around 1%, whereas in the second region D for NiMOII (d caxis) and Sm(Sr)MOII (d c-axis), are above 5%. This indicates Nd2-xSrxNiO4.11 As indicated in Table 3 there is a crystal system transformation at x ca. 0.6. that there is a rigid NiO4 network parallel to the a,b plane which is aVected very little by Sr substitution whereas Sr Tables 5 and 6 list structural parameters and interatomic distances for the Sm2-xSrxNiO4+d system. From Table 6, substitution causes a drastic shift of the OII ions from Sm(Sr) sites to Ni while the unit-cell dimensions remain essentially it can be seen that the bond length of NiMOI (d a-axis) varies similarly to a and b parameters with increasing x, unaltered. The shift of the OII ions from Sm(Sr) towards Ni may arise and interatomic distance Sm(Sr)MSm(Sr) (d c-axis) and Sm(Sr)MNi (d c-axis) vary similarly to the parameter c.It is for two reasons. On the one hand, the bond strength of MMO can be presented in terms of the sum of the MMO bond of substantial interest that the NiMO bond length (d c-axis) monotonously decreases with increasing x without exhibiting formation enthalpies S(DHM-O) in the solids17 and Table 7 lists calculated results.It can be seen that S(DHA-O) decreases any anomaly at x=0.5–0.6 as expected from the variation of the parameter c, which indicates that the distortion of NiO6 with increasing x, which indicates that the acidity of the A ions decreases and the strength of AMO bond reduces with octahedra along the c-axis monotonously decreases with increasing x.This can also be confirmed by the bond length substitution by Sr. On the other hand, after Sr2+ partially substitutes for Sm3+, some Ni2+ ions are oxidized to Ni3+ in ratio of Ni–OII (d c-axis) to NiOI (d a-axis) as shown in Table 6.The % bond length change, D, defined as (dmax-dmin)/dmin, order to meet charge balance considerations. The acidity of the B ion then decreases, and its bond to oxygen becomes within the range 0.4x1.2, was used to clarify the extent of variation of interatomic distances, as shown in Table 6.The stronger. A combination of these eVects causes OII to shift from Sm(Sr) towards Ni ions. data can roughly be divided into two regions. In the first According to Pauling’s principles of bond valence,18 the average bond valence of Ni, S=(R/R1)-N, which has a close Table 5 Positional and thermal parameters for Sm2-xSrxNiO4+d relationship to the bond lengths in inorganic crystals, can be atom site x=0.4 x=0.5 x=0.6 calculated and results are listed in Table 7, where R is the bond length of NiMOI or NiMOII, R1 the length expected for Sm, Sr 8i z 0.3604(1) 0.3604(1) 0.3603(1) a bond of unit valence and N is an empirical constant.For B/A ° 2 0.74(2) 0.64(2) 0.52(2) Ni2+ R1=1.680 and N=5.4. Fig. 2 shows that the average Ni 4a B/A ° 2 0.19(7) 0.35(6) 0.28(5) OI 8e B/A°2 1.63(24) 1.34(22) 0.54(16) bond valence of Ni increases linearly with increasing x in the OII 8i z 0.1756(5) 0.1738(5) 0.1716(4) Sm2-xSrxNiO4+d system.The straight line has a break at x= B/A ° 2 2.76(26) 1.92(21) 1.96(17) 0.6, which is in accord with the fact that there is a crystal Rwp 11.62 11.49 10.11 system transformation at x ca. 0.6. residuals(%) Rp 8.36 8.32 7.33 IR spectra of Sm2-xSrxNiO4+d are shown in Fig. 3. There Re 0.72 0.72 0.65 are three main absorption bands at ca. 700, 520 and 400 cm-1, Ri 3.18 3.26 2.45 which are assigned to the symmetric stretching of site x=0.8 x=1.0 x=1.2 Sm(Sr)MOIIMNi (d c-axis), asymmetric stretching, and bending modes of the NiMOI linkages in the basal planes, respect- Sm, Sr 4e z 0.3607(1) 0.3608(1) 0.3604(1) ively.19 All the absorption bands become weaker with increased B/A ° 2 0.45(1) 0.40(1) 0.49(1) x, consistent with the decrease of the distortion of NiO6 Ni 2a B/A °2 0.32(4) 0.26(4) 0.47(3) octahedra and the increase of symmetry of the crystal structure, OI 4c B/A °2 0.38(11) 0.70(11) 0.76(8) OII 4e z 0.1689(4) 0.1666(4) 0.1639(3) as discussed above.B/A ° 2 1.01(11) 0.75(10) 0.95(8) Rwp 11.57 11.62 10.52 The valence state of nickel in Sm2-xSrxNiO4+d residuals(%) Rp 8.36 8.28 7.43 Re 0.69 0.72 0.67 The mean oxidation number of nickel in the samples is shown Ri 2.95 3.28 3.79 as a function of Sr content in Fig. 2. The Ni2+ and Ni3+ Table 6 Interatomic distances in Sm2-xSrxNiO4+d distance/A ° x=0.4 0.5 0.6 0.8 1.0 1.2 D (%) NiMOI (d a-axis) 1.891(1) 1.885(1) 1.881(1) 1.886(1) 1.892(1) 1.896(1) 0.80 NiMOII (d c-axis) 2.161(6) 2.145(6) 2.119(5) 2.074(5) 2.036(5) 2.000(4) 8.05 SmMOII (d c-axis) 2.272(6) 2.303(6) 2.330(5) 2.355(5) 2.374(5) 2.398(4) 5.55 SmMOI 2.554(1) 2.554(1) 2.553(1) 2.546(1) 2.545(1) 2.549(1) 0.35 SmMSm (d c-axis) 3.435(2) 3.446(2) 3.450(2) 3.421(2) 3.403(2) 3.407(2) 1.38 SmMNi (d c-axis) 4.433(6) 4.448(6) 4.449(5) 4.429(5) 4.410(5) 4.398(4) 1.16 NiMOII/NiMOI 1.143 1.138 1.127 1.100 1.076 1.055 Table 7 The average bond valence of Ni, reducing temperature and the sum of the MMO bond formation enthalpies S(DHMMO) in Sm2-xSrxNiO4+d average bond x valence of Ni Tmax, l/K Tmax, h/K S(DHAMO)/kJ mol-1 S(DHNiMO)/kJ mol-1 0.4 2.62 738 1033 521.2 70.13 0.5 2.68 798 1013 506.3 71.38 0.6 2.74 793 978 491.3 72.62 0.8 2.78 808 1038 461.4 73.45 1.0 2.83 813 1068 431.5 74.48 1.2 2.86 833 1103 401.6 75.10 J.Mater. Chem., 1997, 7(10), 2097–2101 2099Fig. 2 The average bond valence ($), measured mean oxidation number (#), calculated mean oxidation number (6) of Ni and d (() as functions of x in Sm2-xSrxNiO4+d Fig. 4 The TPR profiles of Sm2-xSrxNiO4+d higher temperature reduction peak leads to total reduction of the sample and the K2NiF4 structure is destroyed.20 The reducibility of the sample is probably determined by the stability of the NiO6 octahedra in the system.Table 7 lists the reduction peak temperatures (Tmax) of the peaks and the ratio of the two reduction peaks. The ratio of the peaks is not consistent with the relative amounts of Ni2+ and Ni3+ in the system, which implies that they do not relate to the reduction of the diVerent valence Ni ions, but rather relate to the reduction of oxygen with diVerent binding energies on diVerent sites.It can then be reasonably presumed that the lower temperature reduction peak is related to the reduction of active OII sites (with longer NiMO distances and higher thermal factor as indicated in Tables 5 and 6), while the higher temperature peak corresponds to reduction of less active OI (shorter NiMO distance and lower thermal factors).In the range 0.4x1.2, the peak temperature at lower temperature, Tmax, l increases with x, opposite to the changes of the NiMOII (d c-axis) distance, and consistent with the average bond valence of Ni. As x increases, Tmax, h increases in the range 0.4x0.6, and decreases in the range 0.6x1.2, which is opposite to the cell parameter c.This means that the structural stability of Sm2-xSrxNiO4+d is related to the distance between the layers of perovskite and rocksalt. Fig. 3 IR adsorption spectra of Sm2-xSrxNiO4+d Conclusion content and the oxygen non-stoichiometry (d) calculated from K2NiF4 type samarium nickelate can be synthesized by partial the mean oxidation state of nickel, and occupancy (equal to substitution of Sr2+ for Sm3+ to increase the radius of the unity), and the formula of Sm2-xSrxNiO4+d are listed in A ions and to decrease the radius of B ions by oxidizing the Table 2.It can be seen that the valence state of nickel in the latter into a higher valence state.The Sr2+ substitution causes samples does not deviate much from theoretical data, therefore, a drastic shift of the OII ions along the c-axis from Sm(Sr) the valence state of B ions in A2BO4 compounds can be easily towards Ni whilst scarcely aVecting the NiO4 in the basal controlled by substitution of A ions by A¾. In this study, the plane. The distortion of the NiO6 octahedron decreases with mean oxidation number of nickel can even be above three increasing x from x=0.4 to 1.2.The variation with x of the (3.02 and 3.14 for x=1.0 and 1.2, respectively). These properties TPR peak temperature is opposite to the variation of the should be important in catalysis studies. NiMOII bond length and the cell parameter c, which indicates that the structural stability of Sm2-xSrxNiO4+d is related to Reducibility of Sm2-xSrxNiO4+d the distance between the layers of perovskite and rocksalt layers.TPR profiles of Sm2-xSrxNiO4+d are shown in Fig. 4. There are two reduction peaks in the TPR profile for each sample. After the lower temperature reduction, the K2NiF4 structure This study was financially supported by the Natural Science Foundation of Zhejiang Province, China.still remains stable although some Ni ions are reduced. The 2100 J. Mater. Chem., 1997, 7(10), 2097–210111 Y. Takeda, M. Nishijima, N. Imanishi, R. Kanno, O. Yamamoto References and M. Takano, J. Solid State Chem., 1992, 96, 72. 1 P. Ganguly and C. N. R. Rao,Mater. Res. Bull., 1973, 8, 405. 12 D. Ganguli, J. Solid State Chem., 1979, 30, 353. 2 B.W. Arbuckle, K. V. Ramanujachary, Z. Zhang and 13 A. Sakthivel and R. A. Young, User’s Guide to Programs DBWSM. Greenblatt, J. Solid State Chem., 1990, 88, 278. 9006, 1981. 3 B. W. Arbuckle, K. V. Ramanujachary, A. M. Buckley and 14 B. E. Gushee, L. Katz and R. Ward, J. Am. Chem. Soc., 1957, M. Greenblatt, J. Solid State Chem., 1992, 97, 274. 79, 5061. 4 M. Sayer and P. Odier, J. Solid State Chem., 1987, 67, 26. 15 A. Jones and B. D. McNicol, T emperature-Programmed Reduction 5 C. N. R. Rao, P. Ganguly, K. K. Singh and R. A. Mohan Ram, for Solid Materials Characterization, Dekker, New York, 1986, J. Solid State Chem., 1988, 72, 14. p. 105. 6 A. K. Ladavos and P. Pomonis, J. Chem. Soc., Faraday T rans., 16 R. D. Shannon, Acta Crystallogr., Sect. A, 1976, 32, 751. 1991, 87, 3291. 17 A. K. Ladavos and P. J. Pomonis, Appl. Catal. B, 1992, 1, 101. 7 T. Nitadori, M. Muramatsu and M. Misono, Bull. Chem. Soc. Jpn., 18 I. D. Brown and K. K. Wu, Acta Crystallogr., Sect. B, 1976, 32, 1988, 61, 3831. 1957. 8 A. Albinati and B. T. M. Willis, J. Appl. Crystallogr., 1982, 15, 361. 19 K. K. Singh, P. Ganguly and J. B. Goodenough, J. Solid State 9 Y. Takeda, R. Kanno, M. Sakano, O. Yamamoto, M. Takano, Chem., 1984, 52, 254. Y. Banndo, H. Akinaga, K. Takita and J. B. Goodenough, Mater. 20 H. Lou, F. Ma and Y. Chen, React. Kinet. Catal. L ett., 1990, Res. Bull., 1990, 25, 293. 42, 151. 10 H. Zheng, S. Du, G. Lu, H. Lou and F. Ma, Acta Chim. Sinica, 1993, 51, 373. Paper 7/03271D; Received 12th May, 1997 J. Mater. Chem., 1997, 7(10), 2097–2101 2101

ISSN:0959-9428    

DOI:10.1039/a703271d

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Crystal structure of Ba2Li2/3Ti16/3O13 Christian Dussarrat,a R. Alan Howie,b Glenn C. Mather,b Leticia M. Torres-Martinezc and Anthony R.Westb aInstitut de Chimie de laMatie`re Condense�e de Bordeaux, Cha�teau Brivazac, Avenue du Dr. Albert Schweitzer, 33608 Pessac CEDEX, France bDepartment of Chemistry, University of Aberdeen, Meston Walk, Old Aberdeen, Aberdeen, Scotland, UK AB24 3UE cUniversidad Autonoma de Nuevo L eon, Facultad de Quimica, San Nicolas de los Garza,Mexico Single crystals of Ba2Li2/3Ti16/3O13 have been isolated and its crystal structure solved from X-ray diVraction data.The structure is similar to that of Ba2Ti6O13 and Na2Ti6O13 [monoclinic, space group C2/m (no.12), a=15.171(13), b=3.8992(18), c=9.106(4) A ° , b=98.64(6)°, Z=2], but one of the three independent octahedral sites contains partial substitution of Ti by Li, whereas the others are uniquely occupied by Ti.This phase contains only Ti4+, in contrast to Ba2Ti6O13, which has a mixture of Ti3+ and Ti4+. The structure is described in terms of both an octahedral framework and a cubic close-packing model of [Ba2%O13] layers (%: vacancy). The BaO–Li2O–TiO2 system is currently of interest due given in Table 1.Atomic coordinates and bond distances are summarised in Tables 2 and 3, respectively.† to the discovery of two phases which show relatively high Li+-ion conductivity. One is a hollandite-like phase, Ba3xLi2x+4yTi8-2x-yO16;1 the other is of uncertain compo- Results and Discussion sition but shows good conduction properties.2 Recently, Torres- Martinez et al. have carried out a phase diagram study of the The crystal structure of Ba2Li2/3Ti16/3O13 is essentially ternary system BaTiO3–Li2TiO3–TiO2, in which a number of analogous to that of Na2Ti6O13 and Ba2Ti6O13.The Ti(1)O6, ternary phase regions were established.3 A significant region Ti/Li(2)O6 and Ti(3)O6 octahedra form edge-sharing trimers, of Ba4Ti13O30 ternary solid solution and a smaller area of as shown in Fig. 1. These trimers corner-share with other BaTi5O11 were found. Other single-phase areas included a trimers to form layers parallel to the ac plane. Between region centred on the Li+-ion conducting hollandite-like phase, successive layers, the trimers edge-share forming infinite riba phase centred on BaLi2Ti6O14, a phase based on a previously bons with a zigzag arrangement which run parallel to b.A reported solid solution Ba2Ti10-xLi4xO224 and a new phase projection of the structure along the b-axis is shown in Fig. 2. which formed over a range of compositions close to the The Ba atoms are found within the tunnels which are formed composition Ba3Li2Ti8O20. As a consequence of the crystal between these ribbons of edge-sharing octahedra.structure studies reported here, the actual composition of this The structure may also be described by a close-packing latter phase is found to be Ba2Li2/3Ti16/3O13. Its crystal struc- model in which [Ba2%O13] layers form a cubic close-packed ture is discussed with reference to a series of titanate phases array. The Li and Ti cations lie in interstices between the of general formula AmM2nO4n+1 (A=alkali-metal, alkaline- close-packed layers.Fig. 3 shows the [Ba2%O13] layers which earth metal, M=octahedrally coordinated cation), and is also lie parallel to the (5 -1 -1) planes of the unit cell. Barium is compared with a number of similar structures in which the eleven-coordinate as a result of the oxygen vacancy in the octahedral cation composition is diVerent to that of the [Ba2%O13] layer which, if occupied, would confer a twelvetitle phase.coordinate cuboctahedral coordination on Ba, equivalent to the A-cation environment in the cubic perovskite structure. Fig. 4 shows the coordination environment of Ba; it is com- Experimental posed of an antiprism of oxygens together with five oxygens which lie in a [Ba2%O13] layer.Needle-shaped single crystals of the title phase were prepared by heating an intimate mixture of TiO2 and BaCO3 in 853 All three octahedra are distorted; the greatest distortion occurs for Ti(3)O6, in which the bond distances are spread molar ratio with an excess of Li2CO3 for one week at 1250 °C; the product was then slow-cooled to room temperature at a over the range 1.805 to 2.207 A ° .A large variation of TiMO distances is not uncommon in barium titanates, and is the rate of 5 °C h-1. A crystal suitable for single-crystal analysis was selected from the reaction mixture under a petrographic result of the surrounding O atoms experiencing diVering bond strengths.9 The mean TiMO distance, 1.982 A ° , is similar to microscope. Details of the data collection are summarised in Table 1.The atomic coordinates for Ba2Ti6O13 were adopted that found in most barium titanates. Similarly, the range of BaMO distances, 2.696–3.224 A ° , falls within the limits of as the starting point for the refinement,8 which was carried out by full-matrix least squares. On refinement, it became clear BaMO distances found in barium titanates. As mentioned previously, there are sizeable residual maxima that Ti(2) had a high (isotropic) thermal vibration parameter compared with Ti(1) and Ti(3) and it was concluded that the in the final diVerence map.One of these, 2.6 e A ° -3 in size, is situated at 0.84 A ° from Ba and is attributable to ripple. The Ti(2) site had undergone significant substitution of Li for Ti, compatible with the preparation conditions of the material.other, 3.7 e A ° -3 and 1.95 A ° from Ba, and almost directly above The refinement was then completed with variable occupancies of the Ti(2) site. All atoms were refined anisotropically, except † Atomic coordinates, thermal parameters, and bond lengths and for O(6) and O(7) which became non-positive definite when angles have been deposited at the Cambridge Crystallographic Data this was attempted.A further anomaly, discussed below, is the Centre (CCDC). See Information for Authors, J. Mater. Chem., 1997, presence in the final diVerence map of rather large residual Issue 1. Any request to the CCDC for this material should quote the full literature citation and the reference number 1145/46. maxima. Refinement conditions and final reliability factors are J.Mater. Chem., 1997, 7(10), 2103–2106 2103Table 1 Crystallographic data crystal data Dc=4.631 Mg m-3 Ba2Li2/3Ti16/3O13 Mo-Ka radiation Mr=742.76 l=0.710 73 A ° monoclinic cell parameters from 14 reflections space group C2/m h=7.7–14.5° a=15.171(13) A ° m=10.6 mm-1 b=3.8992(18) A ° T=298 K c=9.106(4) A ° platey needle b=98.64(6)° 0.30×0.12×0.04 mm V=532.5(6) A ° 3 colourless Z=2 data collection Nicolet P3 diVractometer h–2h scans absorption correction: y-scans5 Tmin=0.535, Tmax=0.648 h=-18�18 908 reflections with F>4s(F), n=4 k=0�5 Rint=0.027 l=0�12 hmax=30° 908 measured reflections 855 independent reflections frequency:c 1 in 50 refinement program used to refine structure: SHELX-76.7 refinement on F w=1/(s2F+0.0003F2) R=0.052 (D/s)max=0.001 wR=0.048 Drmax=3.74a, 2.62b GOF=2.092 Drmin=-4.64 776 reflections isotopic extinction coeYcient (SHELX76)7a 61 parameters scattering factors from ref. 7(b) data collection and cell refinement: Nicolet P3 software.6 a1.95 A ° from Ba, i.e. at x,D+y,z relative to Ba at x,y,z. b0.84 A ° from Ba. cStandard check frequency. Table 2 Atomic coordinates (×104) and thermal parameters (×103) for Ba2Li2/3Ti16/3O13 with e.s.d.s in parenthesesa atom x/a y/b z/c Ueq b occupancy Bac 514.8(6) 0 7676.7(8) 8.9(2) Ti(1) 3802(2) 0 9056(2) 7.9(6) Ti(2) 2583(2) 0 2284(3) 3.6(9) 0.640(13) Li(2) 2583(2) 0 2284(3) 3.6(9) 0.360(13) Ti(3) 3301(2) 0 5613(2) 4.9(6) O(1) 1306(8) 0 1091(10) 13(2) O(2) 2633(6) 0 7606(10) 6(2) O(3) 2010(6) 0 4313(9) 6(2) O(4) 3334(7) 0 845(10) 10(2) Fig. 1 Cation arrangement in trimer of edge-sharing octahedra O(5) 3724(8) 0 3863(11) 15(2) O(6) 4292(7) 0 7022(10) 10(2) sites are fully occupied with occupancies of 1/3 and 2/3, O(7) 5000 0 10000 16(3) respectively. The discrepancy is, however, only about 2×e.s.d. aAll sites are fully occupied unless otated. bUeq=1/3SiSj and may not be significant. It was not feasible to carry out UiUja*i a*j ai aj.cAll atoms are in position 4i, apart from O(7) which direct chemical analysis of the crystal, especially of its Li is in 2b of space group C2/m. content and so some slight ambiguity remains over the precise composition and structural model. Table 3 Bond distances (A ° ) for Ba2Li2/3Ti16/3O13 The stoichiometry of the title phase, Ba2Li2/3Ti16/3O13, diVers in its Li (and O) content from that proposed previously, Ti(1)MO(1) 1.9594(11) Li/Ti(2)MO(1) 2.074(12) Ba3Li2Ti8O20, but the Ba5Ti ratio is the same.This indicates Ti(1)MO(2) 2.046(9) Li/Ti(2)MO(2) 1.8920(18)×2 that loss of Li2O by volatilisation can be a significant problem Ti(1)MO(4) 1.873(9)×2 Li/Ti(2)MO(3) 2.157(9) during high temperature syntheses. This phase may exist over Ti(1)MO(6) 2.096(10) Li/Ti(2)MO(4) 1.861(10) Ti(1)MO(7) 1.890(2) Li/Ti(2)MO(5) 2.077(11) a limited stoichiometry range as indicated previously;3 one Ti(3)MO(2) 2.207(9) BaMO(1) 3.160(9) possibility is to have partial reduction of Ti4+ according to Ti(3)MO(3) 2.130(9) BaMO(1) 3.133(11) Li++2Ti4+u3Ti3+.This mechanism would retain full occu- Ti(3)MO(3) 2.009(3)×2 BaMO(2) 3.224(10) pancy of the octahedral sites, as is found in isostructural Ti(3)MO(5) 1.805(10) BaMO(4) 2.820(7)×2 Ba2Ti6O13.8 We think this mechanism to be unlikely in the Ti(3)MO(6) 1.824(10) BaMO(5) 2.755(8)×2 non-reducing conditions used here, however.BaMO(6) 2.696(7)×2 BaMO(7) 3.0632(6)×2 Two other means of creating non-stoichiometry can be imagined, assuming full occupancy of oxygen positions and the cation oxidation states to be +2(Ba), +1(Li), +4(Ti).(1) A replacement mechanism 4Li+>Ti4+, with creation of Ba in the direction of b, is not so easy to explain. One plausible explanation is that it results as a consequence of stacking vacancies on the octahedral sites, consistent with the structure of isostructural Ba2Ti5.5O13.10 (2) A slight substoichiometry in faults in the creation of the layers perpendicular to b.The Li and Ti occupancy factors for the shared site [0.360(13) and barium sites (4Ba2+u4Li++Ti4+) constrained by the limit of full occupancy of the octahedral sites. Then, similarly to the 0.640(13), respectively] suggest that the composition of the material is slightly Li rich compared with the ideal where all hollandite-like phase, Ba3xLi2x+4yTi8-2x-2yO16,1 a general 2104 J.Mater. Chem., 1997, 7(10), 2103–2106Fig. 4 Coordination environment of Ba of Li for Ti apparently exclusively in the Ti(2) site. In this way, charge balance is attained without the need for any reduction of Ti4+ to Ti3+. A number of other structurally related phases show variations in the occupancies of the octahedral sites.In isostructural Ba2Ti5.5O13, for example, the departure from the ideal composition is accommodated by Fig. 2 Projection of the Ba2Li2/3Ti16/3O13 structure along b. Ba atoms the presence of vacancies in the central octahedron of the are represented as circles. trimeric unit.10 A small number of phases with more than one type of octahedrally coordinated cation have been reported.In the case of Ba2Fe2Ti4O13,11 Fe and Ti are distributed over all three octahedral sites, although Fe shows a strong preference for one of the end sites in the trimer of octahedra. The octahedral cation site distribution in Ba2Ti4Cr2O13 could not be determined unambiguously by XRD due to the similar scattering factors of Cr3+ and Ti4+.12 Lattice energy calculations, nevertheless, predict that Cr3+ would prefer an end site in order to minimize Coulombic cation–cation repulsion eVects.In Ba2Ti5ZnO13, there is also a preference for Zn to occupy an end site in the trimer.13 As far as we are aware, Ba2Li2/3Ti16/3O13 is the only analogue in which the nontitanium cation occupies a single crystallographic site; in the others, with the possible exception of the Cr analogue, the two octahedral cations are distributed, but non-statistically, over more than one site.The Li+ ions are, thus, isolated and in fully occupied sites. This would appear to explain why Ba2Li2/3Ti16/3O13 shows no significant levels of Li+-ion conduction. The crystal structure of Ba2Li2/3Ti16/3O13 belongs to a series of titanates characterised by the formula AmM2nO4n+1, where A is an alkali-metal or alkaline-earth metal cation which fills a large (ten- or eleven-coordinate) site andMis an octahedrally coordinated cation. These titanates may be thought of as a series of tunnel structures built up from a network of Ticontaining octahedra.In general, m is equal to n/2 or (n+1)/2 for n even or odd, respectively, in order to minimize cation– cation interactions in the tunnels.For all members of this series, octahedra share edges within a layer, creating a unit which is n octahedra wide. Each octahedral unit shares corners with two other units in the same layer and edges with units in the layers above and below, to form a network of infinite chains of edge-sharing octahedra; these chains have a zigzag arrangement and are parallel to the b direction of the unit cell.Fig. 3 A close-packing layer of composition [Ba2%O13] where % is The large A cations are found in the tunnels which are formed a vacancy. Ba atoms are shown as shaded circles. by the network of infinite chains. The b parameter corresponds approximately to the diagonal of an octahedron (i.e. the width of a layer) and is about 3.7–3.8 A° in every structure of the formula covering the solid solution area can be written Ba2-xLi2/3+x-4yTi16/3+x/4+yO13. series (Table 4).The size of the a parameter is approximately 15 A ° in every structure, with the exception of Cs0.61Ti1.844O4. The results of this structure determination point to a comparatively simple explanation of the ease of preparation of this In this instance, the large size of caesium results in an a parameter that is greater than expected.Only the c parameter variant of the Ba2Ti6O13 structure type, namely the substitution J. Mater. Chem., 1997, 7(10), 2103–2106 2105Table 4 Unit-cell parameters and space group for members of the series AmM2nO4n+1 n phase space group a/A ° b/A ° c/A ° b/° ref. 2a BaTi4O9 C2/m 14.77 3.79 6.29 100.3 9 2+3b K2SrTi10O22 C2/m 15.317 3.787 15.439 102.68 15 3 Ba2Li2/3Ti16/3O13 C2/m 15.17 3.90 9.11 98.64 this work 3+4c Na2Ti7O15 C2/m 14.90 3.74 20.9 96.5 16 4 K2Ti8O17 C2/m 15.678 3.775 11.991 95.67 14 2 Cs0.61Ti1.844O4 Immm 17.012 3.829 2.962 90 17 aBaTi4O9 is a hypothetical structure.bIntergrowth between the members n=2 and n=3. cIntergrowth between the members n=3 and n=4. 8 W. H. Baur, E. Tillmanns and W. Hofmeister, Crystal Struct. depends on the number of octahedra in the units and corre- Commun., 1982, 11, 2021. sponds to the distance between opposite corners in a unit. 9 E. Tillmanns, W. Hofmeister and W. H. Baur, J. Solid State Chem., Assuming the octahedra are regular, the length of c can be 1985, 58, 14. estimated from the relation: c=Ó2(n2+1)dMMO. 10 V. W. Hofmeister and E. Tillmanns, Acta Crystallogr., Sect. B, 1979, 35, 1590. 11 T. A. Vanderah, Q. Huang, W. Wong-Ng, B. C. Chakoumakos, References R. B. Goldfarb, R. G. Geyer, J. Baker-Jarvis, R. S. Roth and A. Santoro, J. Solid State Chem., 1995, 120, 121. 1 C. Suckut, R. A. Howie, A. R. West and L. M. Torres-Martinez, 12 S. Mo�hr and Hk. Mu�ller-Buschbaum, Z.Naturforsch. T eil B, 1994, J.Mater. Chem., 1992, 2, 993. 49, 911. 2 W. J Zheng, R. Okuyuma, T. Esaka and H. Iwahara, Solid State 13 R. S. Roth, C. J. Rawn, C. J. Lindsay and W. Wong-Ng, J. Solid Ionics, 1989, 35, 235. State Chem., 1993, 104, 99. 3 L. M. Torres-Martinez, C. Suckut, R. Jimenez and A. R. West, 14 M. Le Granvalet and L. Brohan, J. Solid State Chem., 1993, 107, J.Mater. Chem., 1994, 4, 5. 127. 4 E. Tillmanns and I.Wendt, Z. Kristallogr., 1976, 144, 16. 15 A. D.Wadsley and W. G. Mumme, Acta Crystallogr., Sect. B, 1968, 5 A. C. T. North, D. C. Phillips & F. S. Mathews, Acta Crystallogr., 24, 392. Sect. A, 1968, 24, 251. 16 T. Sasaki, M. Watanabe, Y. Fujiki, Y. Kitami and M. Yokoyama, 6 Nicolet P3/R3 Data Collection Operator’s Manual, Net XRD J. Solid State Chem., 1991, 92, 537. Corporation, 1980. 17 I. E. Grey, C. Li, I. C. Madsen and J. A.Watts, J. Solid State Chem., 7 (a) G. M. Sheldrick, SHELX-76 Program for Crystal Structure 1987, 66, 7. Determination, University of Cambridge, 1976; (b) International T ables for X-Ray Crystallography, Kynock Press, Birmingham, 1974, vol. 4. Paper 7/03252H; Received 12th May, 1997 2106 J. Mater. Chem., 1997, 7(10), 2103–2106

ISSN:0959-9428    

DOI:10.1039/a703252h

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

New preparation method of Lan+1NinO3n+1-d (n=2, 3) Maria Deus Carvalho,*a Fernanda Madalena A. Costa,a Isabel da Silva Pereira,a AlainWattiaux,*b Jean Marc Bassat,b Jean Claude Grenierb and Michel Pouchardb aDepartamento de Quý�mica da Faculdade de Cie�ncias da Universidade de L isboa, Rua Ernesto de Vasconcelos, C1–5°, 1700 L isboa, Portugal bInstitut de Chimie de laMatie`re Condense�e de Bordeaux (ICMCB), CNRS, Avenue du Docteur Albert Schweitzer, 33 608 Pessac Cedex, France Samples Lan+1NinO3n+1-d (n=2,3) have been prepared by two diVerent methods, which lead to diVerent oxygen stoichiometry values.Materials obtained by the citrate route always show higher content of Ni3+, which can be rationalized by the high reactivity and the original morphology of the precursors obtained by this method.The ternary nickel oxides with the general formula 2 Nitrate route Lan+1NinO3n+1, similar to the Ruddlesden–Popper series, show Stoichiometric amounts of NiO (99.99%) and La2O3 (99.999%, dried at 1273 K prior to use) were dissolved in nitric acid a multilayered crystal structure which can be described by the solution (ca. 1/3 HNO3, 2/3 H2O). stacking along the c axis of n finite LaNiO3 perovskite layers This solution was fired to dryness and the resulting nitrates separated by LaO rocksalt-like layers.1,2 slowly decomposed on a hot plate.The dark gray products Generally, it has been diYcult to obtain pure phases of the were ground and heated in air at 1173 K for 2 h. The black Lan+1NinO3n+1 (n=2,3) series and often the referred phases powder was annealed in air at 1423 and 1353 K in order to exhibit intergrowth problems influencing their physical obtain La3Ni2O7-d and La4Ni3O10-d, respectively. Several properties.3–5 intermediate grindings were necessary and single-phase mate- Recently, Zhang et al.prepared the n=2 member6,7 starting rials were obtained after heating for ca. 3 days for La3Ni2O7-d from an organic precursor and obtained a single-phase oxygenand 14 days for La4Ni3O10-d.deficient La3Ni2O6.92. The stoichiometric phase with respect to the oxygen content was made by heating the as-prepared 3 Citrate route sample at high oxygen pressure. These authors had also prepared the stoichiometric n=3 member.8 Stoichiometric amounts of NiO (99.99%) and La2O3 (99.999%, dried at 1273 K prior to use) were dissolved in a slight excess Sreedhar et al.4 have observed that intergrowths occur only of nitric acid solution followed by the addition of an equivalent as isolated line defects for n=2 and n=3 members.These molar proportion of citric acid with respect to NiO and La2O3. results have definitively shown the strong influence of the The solution was heated to dryness and the formation of a synthesis method on the stacking ordering as well as on the green gel was observed during this thermal treatment.Then, oxygen stoichiometry, which results in diVerent Ni3+ contents the product was gently heated on a sand bath which induces and consequently in various physical and chemical properties auto-combustion. Final thermal treatment was the same as for (for the stoichiometric phases the mean oxidation state of the nitrate route to allow a comparative study. A drastic nickel is 2.50 (n=2) and 2.67 (n=3), respectively).decrease in reaction times could be observed by this route, It is well established that decomposition of precursors with especially for La4Ni3O10-d for which only 2 days were necessmall particle size and high surface area often allows one to sary to obtain a single phase.prepare homogeneous pure phases inaccessible by conventional solid-state reactions. The samples were characterized by powder X-ray diVraction The aim of this work was the preparation of the n=2 and using a Philips diVractometer with Cu-Ka radiation. Highn= 3 members of this series, without intergrowth phenomena, resolution transmission electron microscopy (HRTEM) was using diVerent methods.performed with a JEOL 2000FX electron microscope. SEM A careful chemical analysis (lanthanum, nickel and oxygen images were obtained with a JEOL JSM 35C and EDS analysis contents) was also carried out in order to determine their exact carried out with a Noran-Voyager apparatus. Nickel and formulation.lanthanum contents were measured by atomic absorption with a Unicam 929 AA spectrometer and using a Radiometer ionselective electrode (ISE25F), respectively. The oxygen non-stoichiometry content (d) of all the samples Experimental was deduced from chemical analyses of trivalent nickel (t) by The samples were prepared as follows. a well known titration method [d=(n-1-t)/2] according to the formulation Lan+1Ni2+(n-t)Ni3+ tO3n+1-d.Electrical resistivity measurements were carried out in 1 Solid state route the range 4.2–300 K on sintered pellets using a standard four-probe technique.9 Stoichiometric quantities of La2O3 (99.999%, dried at 1273 K prior to use) and NiO (99.99%) were mixed and heated in air at 1423 and 1353 K for the preparation of La3Ni2O7-d and Results and Discussion La4Ni3O10-d , respectively. 1 Characterization of the precursors However, this conventional solid-state reaction always resulted in the formation of a phase mixture, even after SEM images of the precursors revealed significant diVerences in their morphology, depending on the preparation. SEM prolonged heating as previously reported.6 J. Mater. Chem., 1997, 7(10), 2107–2111 2107images of the precursors obtained after auto-combustion (citrate route) and after decomposition on a hot plate (nitrate route) for n=2 and n=3 samples are shown in Fig. 1 and 2. It is clear that, for both phases, the morphologies of the precursors obtained by the citrate and nitrate routes are completely diVerent. The materials obtained by the citrate route are extremely thin, ‘fly-wing’ like particles without grains and agglomerates [Fig. 1, 2], in contrast to those prepared by the nitrate route.After heating for 2 h at 1173 K the SEM characteristics show similar diVerences for the two samples. However, the grain size of the powders obtained by the nitrate route is obviously bigger and agglomerates can be seen, which was not observed when using the citrate method.XRD studies of these precursors after annealing 2 h at 1173 K revealed that the citrate route leads to a mixture of diVerent members of the series Lan+1NinO3n+1 for n=2 (nominal composition), but a single phase for n=3 [Fig. 3, 4]. Conversely, by the nitrate route, both precursors seem to be a mixture of diVerent members of the series, La2O3 and NiO [Fig. 3, 4]. These results show the higher reactivity observed when preparing the samples by the citrate route; this is due to the original morphology of the precursors obtained by this rapid thermal procedure.10 2 Lan+1NinO3n+1 (n=2,3) compounds The change in the X-ray pattern of each sample with time of reaction revealed that the kinetics in the citrate route is faster than in the nitrate route.After a heat treatment of 12 h at the annealing temperature (T=1423 K for n=2, T=1353 K for n=3), the X-ray data revealed single phases for the samples obtained with the citrate precursors, while by the nitrate route, Fig. 2 SEM images of the precursor of La4Ni3O10-d they showed a mixture of diVerent phases. This feature was observed for both prepared phases (n=2 and n=3).The XRD patterns for the final materials prepared by both routes are identical as can be seen in Fig. 5 and 6, with only Fig. 3 X-Ray powder diVraction patterns of the precursor of La3Ni2O7-d after heating for 2 h at 1173 K Fig. 4 X-Ray powder diVraction patterns of the precursor of La4Ni3O10-d after heating for 2 h at 1173 K Fig. 1 SEM images of the precursor of La3Ni2O7-d 2108 J.Mater. Chem., 1997, 7(10), 2107–2111Fig. 5 X-Ray powder diVraction patterns of the final material La3Ni2O7-d Fig. 7 High resolution electron microscopy image of La3Ni2O7-d (citrate route) ([110] zone axis) Fig. 6 X-Ray powder diVraction patterns of the final material La4Ni3O10-d the reaction times diVering (Table 1). The compounds are pure and well crystallized. The observed XRD data were used for least-squares refinement of the unit-cell parameterse Fmmm space group.The refined unit-cell parameters are given in Table 1; they are in good agreement with those previously reported for the n=2 phase3,5,6,11 and also for the n=3 member.3,8,12,13 However, a careful examination using high-resolution electron microscopy revealed a more complex situation.For n=2, one can still observe some disordered intergrowths even in the samples obtained by the citrate route. Fig. 7 shows stacking defects, which is confirmed by streaking observed in the electron diVraction pattern. For n=3, the citrate route leads to a well ordered structure (Fig. 8); only a few defects were observed, which was not the case for the sample prepared by the nitrate route, which showed a disordered intergrowth. Electron microscopy seems to reveal that these phases do not crystallize in the Fmmm space group. 3 Chemical analysis Quantitative chemical analyses of lanthanum and of nickel Fig. 8 High resolution electron microscopy image of La4Ni3O10-d (citrate route) ([110] zone axis) indicate a La/Ni ratio very close to 1.5 and 1.3 for La3Ni2O7-d Table 1 Cell parameters and non-stoichiometry content of the prepared compounds reaction dd t route time/h aa/nm ba/nm cb/nm V c/nm3 [(n-1-t)/2] (% Ni3+) formulation La3Ni2O7-d nitrate 72 0.5393 0.5451 2.054 0.604 +0.07 43 La3Ni2O6.93 citrate 48 0.5400 0.5452 2.052 0.604 -0.03 53 La3Ni2O7.03 La4Ni3O10-d nitrate 336 0.5413 0.5468 2.795 0.827 +0.25 50 La4Ni3O9.75 citrate 28 0.5415 0.5467 2.797 0.828 -0.02 68 La4Ni3O10.02 a±0.0002 nm. b±0.001 nm.c±0.001 nm3. d±0.02 J. Mater. Chem., 1997, 7(10), 2107–2111 2109Fig. 10 Temperature dependence of the electrical resistivity for Fig. 9 Temperature dependence of the electrical resistivity for La4Ni3O10-d La3Ni2O7-d preparation. From our XRD study, no structural change was and La4Ni3O10-d , respectively; these results are confirmed by evidenced at low temperature (77<T298 K).We suggest SEM/EDS analysis. that the nature of this transition is electronic. Thus, according Values of d are reported in Table 1. For the citrate route to Taniguchi et al.,11 the properties of these compounds can stoichiometric compounds (d=0) were obtained, while the be interpreted by a model of charge ordering in the NiO2 nitrate route samples are oxygen deficient (d>0). This result planes induced by oxygen vacancy ordering, although it does is still not very well understood since the nitrate route should not explain the observed anomaly.be more oxidative than the citrate one. This confirms the For La4Ni3O10-d, metallic behavior is always observed with strong influence of the morphology of the precursors on the a small variation in the resistivity values. These resistivity preparation of these phases.However, it seems that the occurvalues are one order of magnitude smaller for n=3 compared rence of some oxygen non-stoichiometry, d, does not aVect the to those for n=2. Moreover, the inflexion point seen at ca. cell parameters, which is in accordance with the results of 140 K in the n=3 sample (Fig. 10) is quite diVerent from the Zhang et al.6 metal–insulator transition seen at ca. 120–140 K in the n=2 Occasionally, it has been shown that some samples prepared sample (Fig. 9). One should point out that all the previous via the citrate route showed a small excess of oxygen, which data for this compound present some controversy. The most corresponds to an oxygen over-stoichiometry likely resulting recent work of Zhang et al.8 do not reveal any resistivity from the occurrence of interstitial oxygen atoms in the La2O2 anomaly but Shreedhar et al.4 report an inflexion point in the layers, as was previously described for La2NiO4+d .14,15 curve at 140 K, but with thermal behavior diVering from ours, while Tkalich et al.13 report a very marked anomaly in 4 Electrical resistivity their work.Additional investigations are necessary in order to give an Fig. 9 and 10 show the thermal dependence of the electrical resistivity, r(T ), for the La3Ni2O7-d and La4Ni3O10-d samples, interpretation of the observed anomalies. respectively. For both phases, resistivity values are always higher for the Conclusion nitrate route samples than for the citrate ones.Thus, it seems that resistivity decreases with increased Ni3+ content, in agree- The n=2 and n=3 members of the series Lan+1NinO3n+1-d, have been successfully synthesized using, for the first time, the ment with previous work.3,6,11 La3Ni2O7-d (Fig. 9) clearly shows a diVerent electrical citrate method. This process leads to powders which exhibit small particle sizes, which is responsible for the high reactivity behavior depending on the preparation method, which results from the diVerent oxygen stoichiometry of the compounds and of the precursors.This work clearly shows that the preparation method has a the value of the carrier concentration is directly correlated to the Ni3+/Ni2+ ratio. Above 120 K, the metallic behavior of strong influence on the oxygen stoichiometry of these compounds.It is very important to carefully measure the Ni3+ the citrate material agrees well with the electrical resistivity data published earlier by Moham-Ram et al.3 and by Zhang content, and consequently, the total oxygen content in each sample. et al.6 for samples with similar composition. However, there is some discrepancy between these results and those of Sreedhar The samples obtained by this new method are always more oxidized than those prepared by other methods and have a et al.4 who found a non-metallic behavior for this phase, but no data were given about the oxygen content of their sample.higher Ni3+ content. These observations are very important since they lead to These results clearly show that the electrical resistivity is strongly dependent on d as previously reported by Taniguchi variations in the physico-chemical properties of the diVerent samples as has been shown by electrical resistivity behavior.et al.11 However, our results are somewhat diVerent with respect to those of these authors. In this work, a marked These studies will be discussed in future work.inflexion is visible for both samples (d=0.0 and 0.1). These inflexions occur at diVerent temperatures: 120 K for the sample The authors are grateful to ‘Embaixada de Franc�a em Portugal’ and ‘JNICT’ (Portugal) for financial support. obtained by the citrate preparation and 140 K for the nitrate 2110 J. Mater. Chem., 1997, 7(10), 2107–211110 P.Barboux, P. Griesmar, F. Ribot and L. Mazerolles, J. Solid State References Chem., 1995, 117, 343. 1 P. Lacorre, J. Solid State Chem., 1992, 97, 495. 11 S. Taniguchi, T. Nishikawa, Y. Yasui, Y. Kobayashi, J. Takeda, 2 P. Lacorre, Actual. Chim., 1995, 3, 16. S. I. Shamoto and M. Sato, Phys. Soc. Jpn., 1995, 64, 1644. 3 R. A. Mohan-Ram, L. Ganapathi, P. Ganguly and C. R. N. Rao, 12 C. Brisi, M. Vallino and F. Abbattista, J. L ess Common Met., 1981, Solid State Commun., 1986, 63, 139. 79, 215. 4 K. Sreedhar, M. Mc Elfresh, D. Perry, D. Kim, P. Metcalf and 13 K. Tkalich, V. P. Glazkov, V. A. Somenkov, S. Shil’Shteoin, M. Honig, J. Solid State Chem., 1994, 110, 208. A. E. Kar’Kin and A. V. Mirmel’Shtein, Superconductivity, 1991, 5 J. Drennan, C. P. Tavares and B. C. H. Steele, Mater. Res. Bull., 4, 2280. 1982, 17, 621. 14 A. Demourgues, F. Weill, J. C. Grenier, A. Wattiaux and 6 Z. Zhang, M. Greenblatt and J. B. Goodenough, J. Solid State M. Pouchard, Physica C, 1992, 192, 425. Chem., 1994, 108, 402. 15 A. Demourgues, F. Weill, B. Darriet, A. Wattiaux, J. C. Grenier, 7 Z. Zhang and M. Greenblatt, J. Solid State Chem., 1994, 111, 141. P. Gravereau and M. Pouchard, J. Solid State Chem., 1993, 106, 8 Z. Zhang and M. Greenblatt, J. Solid State Chem., 1995, 117, 236. 317. 9 P. Dordor, E. Marquestaut and G. Villeneuve, Rev. Phys. Appl., 1985, 20, 795. Paper 7/02424J; Received 9th April, 1997 J. Mater. Chem., 1997, 7(10), 2107–2111 21

ISSN:0959-9428    

DOI:10.1039/a702424j

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Preparation and characterization of uniform, spherical particles of Y2O2S and Y2O2S:Eu Ligia Delgado da Vila, Elizabeth Berwerth Stucchi* and Marian Rosaly Davolos Departamento de Quý�mica Geral e Inorga�nica, Instituto de Quý�mica, Universidade Estadual Paulista, C.P. 355, CEP 14801-970, Araraquara, SP, Brasil The preparation of spherical Y2O2S and Y2O2S:Eu particles using a solid–gas reaction of monodispersed precursors with elemental sulfur vapor under an argon atmosphere has been investigated.The precursors, undoped and doped yttrium basic carbonates, are synthesized by aging a stock solution containing the respective cation chloride and urea at 82–84 °C. Y2O2S and Y2O2S:Eu were characterized in terms of their composition, crystallinity and morphology by chemical analysis, X-ray powder diVraction (XRD), IR spectroscopy, and scanning electron microscopy (SEM).The Eu-doped oxysulfide was also characterized by atomic absorption spectrophotometry and luminescence spectroscopy. The spherical morphology of oxysulfide products and of basic carbonate precursors suggests a topotatic inter-relationship between both compounds. Rare-earth oxysulfides have been known for a long time as and Hsu8 and Sordelet and Akinc9 was adapted by Santos.10 Sulfidization of the basic carbonate precursor compounds was excellent phosphor host materials.When activated with trivalent europium, Y2O2S becomes an important red phosphor carried out in a horizontal tube furnace following the procedure outlined by Luiz.7 Precursors were loaded into the principal in color TV picture tubes1 because it has high brightness, short decay time, and exhibits long-term stability in poly(vinyl furnace in an alumina crucible.The furnace was completely sealed to ensure an oxygen-free atmosphere. Argon gas was alcohol). Recently, yttrium oxysulfide has attracted a great deal of attention because of its electroluminescent properties.2 flushed through the system during both the reaction and cooling time.Sulfur was heated at 220 °C in an auxiliary In modern technology, the necessity of dispersed powders consisting of uniform particles in size and shape is widely furnace. Then the main furnace temperature was gradually increased to 770 °C (heating rate 2 °C min-1) and sulfur vapor recognized. For example, in ceramic manufacture, there is a significant reduction in the sintering time and temperature if was carried by the argon stream across the system (flow rate 116 cm3 min-1).The end of the reaction was accompanied by monodispersed powders are used as starting materials.3 Phosphors used in both CRT and X-ray screens with uniform bubbling of acid vapor products in alkaline solution until it size distribution of the particles results in the best screen reached constant pH.After the desired reaction time the surfaces.4 In materials science, the particle shape is equally auxiliary furnace was shut oV and the product was heated at important. Spherical morphology is interesting because most 800 °C for at least 2 h. The choice of synthesis temperature theoretical models dealing with fine particle properties and was based on DTA curve data.interactions are based on spherical particles. Furthermore, the size distribution of relatively uniform spherical particles can be determined by optical techniques without altering or Characterization destroying the system.5 Lanthanide content was assayed by EDTA titration using Several methods are known for the preparation of rare-earth xylene orange as indicator. Atomic absorption spectrophoto- oxysulfides.6 However, it seems that few investigations have metry was performed using an INTERLAB AA-1475 atomic been reported in the literature on spherical particle preparation absorption spectrophotometer. X-Ray powder diVractograms of these phases.In a previous paper Luiz et al.7 described the were recorded on an XRD HZG-4B diVractometer equipped preparation process of lanthanum and yttrium oxysulfide with Cu-Ka or Co-Ka radiation (36 kV, 20 mA).IR spectra in powders by using a solid–gas reaction of oxalates with elemen- KBr pellets and Nujol in CsI windows were measured with tal sulfur vapor under an argon atmosphere. This paper reports Nicolet FT-730 and Impact 400 FTIR spectrometers.The an analogous method, using monodispersed basic carbonates particle morphology was examined in a scanning electron as precursors. The present method is very attractive because microscope JEOL-JSM-T330 A. Powders were pre-coated with the particle morphology of oxysulfide reproduces that of the Au using a cathodic sputtering Edwards S 150 B instrument. basic carbonates.This fact indicates a topotatic inter-relation- Emission spectra of Eu-doped oxysulfide were obtained in a ship between both groups of compounds. Thus, desired mor- Fluorolog Spex 212 I fluorescence spectrometer equipped with phology for the oxysulfides can be achieved through the an R928 Hammamatsu photomultiplier. The samples were morphological control of precursor particles.excited by a 450 W xenon lamp. Experimental Synthesis of doped and undoped yttrium oxysulfides Results and Discussion Yttrium and europium oxides (99.99 and 99.999% pure, EDTA titration and atomic absorption spectrophotometry Aldrich) were used as starting materials. Other chemicals used EDTA titration and atomic absorption spectrophotometry were grade reagents. Monodispersed yttrium basic carbonate results of the oxysulfides are shown in Table 1.As can be seen, particles were produced by heating a stock solution containing the agreement between calculated and obtained results are a cationic chloride and urea at 82–84 °C under continuous stirring for 2 h. This procedure, based on work by Matijevic satisfactory. J. Mater. Chem., 1997, 7(10), 2113–2116 2113Table 1 EDTA titration and atomic absorption spectrophotometry results for yttrium oxysufide particles Ln (mass%) atomic absorption EDTA titration spectrophotometry compound found calc.found calc. Y2O2S 73.97 73.52 — — (Y0.964Eu0.036)2O2S 75.30 74.00 4.26 4.44 X-Ray diVraction The dhkl data of the oxysulfides are shown in Table 2. The precursors were non-crystalline and the products of their thermodecomposition under sulfur and Ar atmosphere yielded an X-ray pattern identified as crystalline Y2O2S.11 The results Fig. 2 FTIR spectra for yttrium oxysulfides in Nujol–CsI windows did not exhibit patterns due to oxysulfate and/or oxide phases (1500 to 220 cm-1 range): (a) yttrium oxysulfide; (b) Eu-doped in the oxysulfide samples. yttrium oxysulfide FTIR spectroscopy precursor.In the low-frequency region, absorptions in the IR spectra of Y2O2S and Y2O2S:Eu, obtained between 4000 range 800–220 cm-1 can be attributed to LnMO or LnMS and 400 cm-1 (high-frequency region) and 1500 to 220 cm-1 vibrations. This assignment is very diYcult because of the (low-frequency region) are shown in Fig. 1 and 2, respectively. proximity of the bands.However, if this sample spectrum is In the high-frequency region, the absence of the SMO stretchcompared with that of yttrium oxide (Fig. 3), oxysulfide phase ing band at 1200 cm-1 attributed to sulfate,12 is noted. This formation is evidenced. The oxysulfide spectrum exhibits less suggests that the oxygen was entirely removed from the broad and strong bands than those observed for the oxide (see reaction system.Moreover, the bands corresponding to the Table 3). This behavior could be understood considering that carbonate group were also missing, suggesting the absence of Table 2 X-Ray diVraction powder data for yttrium oxysulfide this work lit.11 Y2O2S Y2O2S5Eu d/A ° I/Io(%) d/A ° I/Io(%) d/A ° I/Io(%) 3.280 35 3.255 31 3.283 32 2.930 100 2.942 100 2.920 100 2.320 60 2.320 40 2.312 34 2.190 16 — — — — 1.892 65 1.885 38 1.889 34 1.824 45 1.817 23 1.819 32 1.642 30 1.639 16 1.640 16 1.591 25 1.587 18 1.589 15 Fig. 3 FTIR obtained in Nujol–CsI windows (1500 to 220 cm-1 range): (a) yttrium oxysulfide; (b) Eu-doped yttrium oxysulfide; (c) yttrium oxide Table 3 Comparative frequency ranges of yttrium oxysulfide, Eudoped yttrium oxysulfidettrium oxide bands wavenumber/cm-1 Y2O2S Y2O2S5Eu Y2O3 — 235m (sp) 238w (sp) 269m (sh)a 270w (sh) 270w (sh) 290m (sh) — — — — 306s (sp) 310m (sh) 312m (sh) — — — 339s (sp) 391m (sh) 394m (sh) 380s (br) 467s ( br) 462s ( br) 466s (sp) — — 563s ( br) Fig. 1 FTIR spectra for yttrium oxysulfides in KBr pellets (4000 to 400 cm-1 range): (a) yttrium oxysulfide; (b) Eu-doped yttrium aBand features: br=broad, sp=sharp; sh=shoulder. Band intensity; s=strong, m=medium, w=weak.oxysulfide 2114 J. Mater. Chem., 1997, 7(10), 2113–2116there are only contributions from C3v symmetry for Y2O2S,13 while in Y2O3 there are contributions from C2v and S6 symmetry. 14 Therefore, the IR spectra suggest that the LnMS vibrations are present. This result corroborates the XRD data, showing the formation of the oxysulfide phase.The absorptions at 1460, 1376 and 723 cm-1 are characteristic of Nujol vibrational modes. Scanning electron microscopy Scanning electron micrographs of the basic carbonate precursors and oxysulfides are shown in Fig. 4 and 5, respectively. The precursor and oxysulfide particles were relatively uniform in size and spherical in shape.The average particle size of both undoped basic carbonates and oxysulfides was 0.4 mm, while Eu-doped precursors and oxysulfides showed a medium particle size of 0.2 mm. This suggests that the dopant can influence the size of the particles, but further study is required to confirm this eVect. Fig. 4(b) and 5( b) show that no change in morphology took place during the sulfidization procedure.This suggests the occurrence of a topotatic reaction from basic carbonates to oxysulfides. It can be observed that slight sintering occurs during thermal processing of the precursors. Luminescence spectroscopy Emission spectra of Eu-doped oxysulfide obtained at room temperature (lexc=311 nm) and 77 K (lexc=303 nm) are shown in Fig. 6(a) and (b), respectively. The main signals are found in the region between 580 and 630 nm.The assignments agree Fig. 4 Electron micrographs of: (a) yttrium basic carbonate; (b) yttrium oxysulfide Fig. 6 Emission spectra of Eu-doped yttrium oxysulfide (400–750 nm range): (a) room temperature (lexc=311 nm); ( b) 77 K (lexc=303 nm) Fig. 5 Electron micrographs of: (a) Eu-doped yttrium basic carbonate; Fig. 7 Emission spectra of Eu-doped yttrium oxysulfide (400–750 nm range, room temperature, lexc=265.5 nm) (b) Eu-doped yttrium oxysulfide J.Mater. Chem., 1997, 7(10), 2113–2116 2115p. 315; (b) K. Sowa, M. Tanabe, S. Furukawa, Y. Nakanishi and with that expected for C3v symmetry.13 The dominant emission Y. Hatanaka, Jpn. Appl. Phys., 1993, 32, 12A, 5601. arises from the 5D0�7F2 transition and is observed at 626 nm. 3 Y. Her, E. Matijevic and W. R. Wilcox, J. Mater. Sci. L ett., 1992, According to crystal-field theory, the J=0 state does not split 11, 1629. (A1), and the J=1 state splits into two Stark levels (A2 and 4 L. H. Brixner,Mater. Chem. Phys., 1987, 16, 253. E) resulting in one and two emission bands, respectively. This 5 (a) E. Matijevic, CHEMTECH, 1991, 21, 176; (b) E.Matijevic, Annu. Rev. Mater. Sci., 1985, 15, 483. observed behavior indicates only one emitting Eu3+ symmetry 6 (a) M. Leskela�, Res. Pap., 1980, 64, 39; (b) O. Kanehisa, T. Kano site. The emission spectrum obtained at 265.5 nm excitation, and H. Yamamoto, J. Electrochem. Soc., 1985, 132, 2023. characteristic of Eu3+ in yttrium oxide, is shown in Fig. 7. No 7 J.M. Luiz, E. B. Stucchi and N. Barelli, Eur. J. Solid State Inorg. peak arising from a 5D0�7F2 transition at 611 nm is observed, Chem., 1996, 33, 321. suggesting the absence of oxide impurity. 8 E. Matijevic and W. P. Hsu, J. Colloid Interface Sci., 1987, 118, 506. 9 D. Sordelet and M. Akinc, J. Colloid Interface Sci., 1987, 122, 47. 10 M. F. Santos, Estudo de Precursores paraMateriais L uminescentes: The authors would like to thank to CNPq for a scholarship Hidroxicarbonatos de ý� trio e de Gadolý�nio Dopados, M. Sc. (L.D.V.). Dissertation, Instituto de Quý�mica, UNESP, Araraquara, 1993. 11 JCPDS 24-1424. 12 J. R. Ferraro, L ow-Frequency V ibrations of Inorganic and Coordination Compounds, Plenum Press, New York, 1971. References 13 O. J. Sovers, T. Yoshioka, J. Chem. Phys., 1968, 49, 4945. 14 H. Forest, G. Ban, J. Electrochem. Soc., 1969, 116, 474. 1 P. N. Yocom, US Pat., 3 418 247, 1968. 2 (a) K. Sowa, M. Tanabe, S. Furukawa, Y. Nakanishi and Y. Hatanaka, Electroluminescence, Proc. Int. Workshop 6th, 1992, Paper 7/01540B; Received 4thMarch, 1997 2116 J. Mater. Chem., 1997, 7(10), 211

ISSN:0959-9428    

DOI:10.1039/a701540b

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Fabrication of nanometer-sized anatase particles by a pulsed laser ablation method K. Kawasaki,a J. F. Despres,a S. Kamei,b M. Ishikawab and O. Odawara*a aDepartment of Electronic Chemistry, T okyo Institute of T echnology, 4259 Nagatsuta,Midori-ku, Yokohama 226, Japan bMitsubishi Research Institute, Inc., 2-3-6 Otemachi, Chiyoda-ku, T okyo 100, Japan TiO2 particles of<10 nm have been prepared by ablating a rutile target rod with a pulsed laser beam, which was assisted by the synchronized injection of an Ar–O2 mixture.By controlling the wavelength of laser beam, injecting pressure ( p) and O2 concentration in the gas mixture, ultrasmall anatase particles could be obtained. In the present work, l=532 nm, p=5×105 Pa and O2 concentration<5% resulted in anatase particles of <6 nm.TiO2 has been used in various optical materials including Experimental semiconductor photocatalysts since the pioneering work by Laser ablation experiment Honda and Fujishima.1 Nanometer-sized TiO2 particles show size dependence on optical properties (‘quantum size eVect’), Laser ablation was carried out on a TiO2 rod target prepared which corresponds to a shift of optical absorption spectrum.2–4 by pressing TiO2 powder and sintering at 1400 °C for 2 h.The The bandgap shift (DEg) relates to the radius (R) of nanometer- apparatus was a combination of a molecular beam generator17 sized particles as follows;3,5 and an ablation system as shown in Fig. 1. The vacuum system employed in the ablation experiments consisted of a stainlesssteel chamber pumped by three pumps; a 900 l s-1 18 in.DEg= h2 8R2C1 me + 1 mhD- 1.8e2 4peR (1) diVusion pump, a 3000 l s-1 10 in diVusion pump, and a 1500 l s-1 turbomolecular pump. The pressure in the chamber was <1.2×10-8 Pa. where h is Planck’s constant, me and mh are the eVective mass The TiO2 target was irradiated by a Q-switched Nd5YAG of the electron and hole, e is the electron charge, and e is the laser.The wavelength of irradiating light was varied from the relative permittivity of the semiconductor. Ultrasmall TiO2 fundamental to the second and fourth harmonic of the particles are useful materials for highly activated photocata- Nd5YAG laser. In this study, 266 nm (5 ns pulse width, 70 lysts. Anpo et al.2 reported that the photocatalytic activity of mJ pulse-1), 532 nm (5 ns, 260 mJ pulse-1) and 1064 nm (6 ns, TiO2 particles become stronger with decrease in particle size, 455 mJ pulse-1) were used.The repetition time was 10 Hz. and rutile particles below 12 nm and anatase particles below The target TiO2 rod was rotated (1 rpm) so that the laser 8.5 nm result in a large blue shift and large quantum yield. could irradiate the clear surface of the target constantly during There have been many studies concerning other ultrasmall the rotation. The target irradiated by the laser was excited and semiconductor particles. For example, CdS and ZnS crystalproduced a plasma plume on its surface.A high pressure line5–8 and colloidal9 nanoparticles of <10 nm have attracted Ar–O2 gas mixture (purity >99.9% for both gases) was particular interest owing to the large blue shift of their absorpinjected into this plume for ca. 200 ms to cool it, injection being tion spectra and photoluminescence. synchronized with the laser pulse. The synchronization was Nanometer-sized particles have been mainly prepared by very sensitive because of a time delay for the injected gas to sol–gel techniques, and TiO2 particles have been prepared by reach the plume on the surface of the target and the gas was hydrolysis of TiCl4.2–4 Other methods such as laser induced injected 1 ms earlier than the pulse.The injection pressure of reactions, pyrolysis of titanium isopropoxide10 and titanium the mixture gas was varied up to 5×105 Pa and six conditions alkoxide,11 and ignition of TiCl4–H2–O2,12 have also been were tested.The concentration of O2 was varied from 1 to 5 reported. In these studies, a low photon energy CO2 laser (l= to 10%. 10.6 mm) has been used to provide heat for reactant gases to produce a very high reaction temperature over a short time which prevents dissociation of gas molecules. To discuss the quantum size eVect more clearly, nanometersized particles require high purity since impurities lead to a serious influence on the bandgap of the semiconductor.Laser ablation is a novel technique for fabrication of homogeneous nanometer-sized particles incorporated with ultrahigh vacuum systems. The technique is to evaporate the ceramic target by irradiation of a high-energy pulsed laser assisted with injection of high-pressure noble or reactive gas synchronized with the laser pulse.For laser ablation, a high photon energy Nd5YAG laser (fundamental wavelength=1064 nm) and ArF (193 nm) or KrF (248 nm) excimer lasers have been used. There have been few studies about TiO2 fabrication except for thin films.13–16 In this study, ultrasmall TiO2 particles of <10 nm have been fabricated by pulsed Nd5YAG laser ablation and conditions of fabrication and parameters con- Fig. 1 Schematic diagram of the apparatus trolling particle size are discussed. J. Mater. Chem., 1997, 7(10), 2117–2120 2117Analysis of the products and Ti ion intensity is shown in Fig. 3. The TiO2 ion intensity increased with injection gas pressure, whereas the Ti ion The cluster ions of the plume were analyzed by a Q-mass intensity decreased.This result indicates that TiO2 species are analyzer. Results from the analyzer may be somewhat diVerent generated by collisions with high-injection pressure gas. For from the actual distributions because of ionization ratio eVects. the limitation of vacuum capability, the injection pressure was However, in this study, the ionization source was a highlimited to 5×105 Pa, and examination of laser wavelength and energy laser and laser excitation was not influenced by coexist- O2 concentration dependence was carried out on products at ent elements.Therefore, the ionization ratio of each ion was 5×105 Pa. estimated to be equal and the ion intensity analysis is thus believed to be accurate. Laser wavelength dependence Particles were captured on carbon coated grids (Fig. 2) and investigated by transmission electron microscopy (TEM). Fig. 4 shows the laser wavelength dependence of ion intensity of TiO2 and Ti in the plasma plume. There was a large amount Particles sizes were determined using the image mode and crystal structures have been identified using selected area of Ti ions and the TiO2 ion intensity using a wavelength of 266 nm was much lower than of Ti.On the other hand, TiO2 electron diVraction (SAED). Dependence of injection pressure, laser wavelength, and O2 ion intensity with 532 and 1064 nm wavelength light was four orders of magnitude higher than that at 266 nm and was concentration in the gas mixture were examined to determine the optimum fabrication conditions. Dependence of injection almost equal to the Ti ion intensity.The influence of laser power of each wavelength was not considered because the pressure was examined by ion intensity analysis using a Qmass analyzer. Laser wavelength dependence was related to behavior of Ti and TiO2 was obviously diVerent. The laser wavelength dependence might be related to the the ion intensity and the captured particle size and structure. The O2 concentration dependence was related to the size diVerence of the excitation mechanism of the target material.The diVerence between 266, 532 and 1064 nm light is due to of particles. the photon energy (4.66, 2.33 and 1.17 eV respectively). The crystal structure of the target was rutile-type, and its maximum Results and Discussion absorption wavelength is at 413 nm, so the target absorbs only 266 nm laser light.The target which absorbs laser light is Injection pressure dependence electronically highly excited and entirely decomposed into Ti TiO2 nanocrystals were obtained by injection gas cooling of and O ions. In this process, the probability of Ti and O ion the plasma plume. Thus pressure is one of the most important recombination is very low.Therefore, 266 nm wavelength laser parameters for the fabrication and it was directly related to the cooling rate. The injection pressure dependence of TiO2 Fig. 4 Laser wavelength dependence on TiO2 and Ti ion intensity (gas Fig. 2 Schematic diagram of the capture of particles mixture: Ar 90%–O2 10%, injection pressure: 5×105 Pa) Fig. 3 Injection gas pressure dependence on TiO2 and Ti ion intensity (gas mixture: Ar 90%–O2 10%) 2118 J.Mater. Chem., 1997, 7(10), 2117–2120Fig. 5 TEM and SAED images of spherical particles prepared by Ar 90%–O2 10% gas mixture injection. Wavelength (a) 532 nm, (b) 1064 nm. Fig. 6 TEM image of TiO2 particles prepared by Ar 95%–O2 5% gas mixture injection using a 532 nm wavelength laser was not appropriate for generating TiO2 species. 532 and 1064 nm wavelength laser breaks TiMO bonds by thermal excitation and in this case, many titanium oxide ions exist in the plasma (mainly TiO ions14,15). In all cases, ions in the plasma can react by injection of an Ar and O2 gas mixture. Oxygen atoms of oxygen gas attach to titanium ions or titanium oxide ions to form TiO2 species [the TiMO bond (6.92 eV) is stronger than OMO (5.12 eV)18,19].From these results, 532 or 1064 nm wavelength laser was required for formation of fine TiO2 particles. The most suitable excitation wavelength (532 or 1064 nm) could not be diVerentiated by ion intensity analysis. The particles prepared using 532 and 1064 nm light were captured and investigated by TEM and SAED as shown in Fig. 5. At 1064 nm, the average particle size was large with a wide distribution (particles>100 nm were found).The crystal structure of these particles corresponded to the rutile-type according to SAED. 1064 nm light breaks only a few TiMO bonds of the target because its photon energy is very low compared with the bond energy of Ti–O; the big particles are simply fragments of the target, and do not arise from the process of recombination of oxygen and Ti ions or titanium oxide ions.By contrast, in the image of the product obtained with 532 nm wavelength, few large particles were observed, and almost all particles were small with comparatively uniform size. SAED reveals a ‘powder pattern’ whose peaks correspond to an anatase-type structure, which is strongly suggested to be stoichiometric.From these results, 532 nm wavelength laser light is appropriate for formation of homogeneous TiO2 particles. O2 concentration dependence Fig. 7 Size distributions of TiO2 particles prepared by 532 nm wave- The oxygen content in the gas mixture aVects the reaction in length laser. Gas mixture: (a) Ar 99%–O2 1%, (b) Ar 95%–O2 5%, (c) Ar 90%–O2 10%. the plasma plume.No reaction could occur in the case of Ar gas injection. Therefore, O2 content plays an important role in the product particle size and particles were prepared at low O2 concentration gas injection. The sizes of almost all particles all experiments. The oxygen concentration becomes lower when the collision frequency of oxygen decreases, which leads prepared using 5% O2 were <10 nm as shown in Fig. 6. The size distributions of the particles prepared by injection of 1, 5 to a slower reaction rate with titanium and titanium oxide ions, therefore, in the short reaction time, injection of a low and 10% O2 concentration in the gas mixture are shown in Fig. 7. The data at 1 and 5% diVered little but the distributions O2 concentration (<5% in the gas mixture) corresponds to a small amount of additional oxygen atoms leading to the were narrower and the average particle size smaller than that at 10% O2.These results indicate that small and uniform formation of ultrasmall TiO2 particles of <10 nm. particles are formed by injection of a low O2 concentration gas mixture. Addition of oxygen to titanium and titanium Conclusions oxide ions occurs only in the plasma over a very short time and limited area.In the plasma, not only titanium and titanium We have succeeded in the fabrication of ultrasmall anatase particles of <10 nm by a pulsed Nd5YAG laser ablation oxide ions but also a large amount of oxygen ions could be identified by Q-mass analysis. Oxygen ions are lighter than method. 532 nm wavelength laser light and high-pressure gas mixture injection were required for formation of small and titanium ions, and they primarily fly out from the reaction zone.The collision frequency does not depend on O2 concen- uniform TiO2 particles. The O2 concentration of the injected gas mixture was an important parameter in controlling particle tration in this work since the injection pressure was equal in J.Mater. Chem., 1997, 7(10), 2117–2120 21197 R. Rossetti, R. Hull, J. M. Gibson and L. E. Brus, J. Chem. Phys., size. The average particle size was ca. 6 nm at <5% O2 1985, 82, 552. concentration in the gas mixture. 8 A. P. Alivisatos, J. Phys. Chem., 1996, 100, 13 226. Laser ablation is one of the most useful techniques to 9 S. Baral, A. Fojtik, H. Weller and A.Henglain, J. Am. Chem. Soc., fabricate ultrasmall semiconductor particles of <10 nm. We 1986, 108, 375. are now studying the spectroscopic properties of plasma to 10 G. W. Rice, J. Am. Ceram. Soc., 1987, 70, C-117. 11 L. E. Depero, P. Bonzi, M. Zocchi, C. Cesale and D. D. Michele, understand the fabrication mechanism and are also investiga- J.Mater. Res., 1993, 8, 2709.ting the optical properties of TiO2 nanoparticles synthesized 12 T. Oyama, Y. Iimura, K. Takeuchi and T. Ishii, J.Mater. Sci. L ett., in this study for application as a highly activated photocatalyst. 1996, 15, 594. 13 H. Funakoshi, K. Fumoto, M. Okuyama and Y. Hamakawa, Jpn. J. Appl. Phys., 1994, 33, 5262. This research is conducted as a part of the industrial technology 14 H.A. Durand, J. H. Brimaud, O. Hellman, H. Shibata, S. Sakuragi, development promotion program of the Research Institute of Y. Makita, D. Gesbert and P. Meyrueis, Appl. Surf. Sci., 1995, Innovative Technology for the Earth. 86, 122. 15 N. Lobstein, E. Millon, A. Hachimi, J. F. Muller, M. Alnot and J. J. Ehrhardt, Appl. Surf. Sci., 1995, 89, 307. 16 C. Garapon, C. Champeaux, J. Mugnier, G. Panczer, P. Marchet, References A. Catheriot and B. Jacquier, Appl. Surf. Sci., 1996, 96–98, 836. 1 A. Fujishima and K. Honda, Nature (L ondon), 1972, 238, 37. 17 S. Kamei, M. Ishikawa, T. Hashimoto, S. Ishizuka and 2 M. Anpo, T. Shima, S. Kodama and Y. Kubokawa, J. Phys. Chem., I. Kusunoki, Bull. Res. Inst. Sci. Meas., T ohoku Univ., 1995, 44, 41. 1987, 91, 4305. 18 B. C. Guo, K. P. Kerns and A. W. Castleman, Jr., Int. J. Mass 3 E. Joselevich and I. Willner, J. Phys. Chem., 1994, 98, 7628. Spectrom. Ion Processes, 1992, 117, 129. 19 D. E. Clemmer, J. L. Elkind, N. Aristov and P. B. Armentrout, 4 N. Serpone, D. Lawless and R. Khairutdinov, J. Phys. Chem., 1995, J. Chem. Phys., 1991, 95, 3387. 99, 16 646. 5 L. E. Brus, J. Phys. Chem., 1986, 90, 2555. 6 L. E. Brus, J. Chem. Phys., 1984, 80, 4403. Paper 7/03816J; Received 2nd June, 1997 2120 J. Mater. Chem., 1997, 7(10), 2117–2120

ISSN:0959-9428    

DOI:10.1039/a703816j

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Synthesis and structure of a new oxynitride Ba3W2O6N2 † P. Subramanya Herle,a M. S. Hegde*a and G. N. Subbannab aSolid State and Structural Chemistry Unit and bMaterials Research Centre, Indian Institute of Science, Bangalore, 560012, India A new oxynitride Ba3W2O6N2 has been synthesised from the ammonolysis of Ba3W2O9. This compound crystallises in a hexagonal structure with a=5.993(2) and c=21.40(4) A ° .Transmission electron microscopy (TEM) studies were carried out to elucidate the structure of this new compound. IR and Raman data are consistent with the C3v site symmetry of the (WO3N)3- unit. This compound is isostructural with Ba3V2O8 reported in the literature. In recent years there has been a spurt of interest in exploring Raman spectra of the samples were recorded using a Spectra Physics SPEX 1403 double spectrometer (Ar-ion laser, l= new oxynitrides because of interesting structural and catalytic properties.1,2 It has been known that oxygen atoms can 514.5 nm) series 2000.Electron diVraction and microscopy (TEM) were carried out using a JEOL 200CX transmission substitute nitrogen atoms in monometallic nitrides due to similarity in their ionic radius as well as in polarizability.electron microscope to elucidate the microstructural features. Of late a number of new ternary nitrides and oxynitrides have been discovered, e.g., LiMN2 (M=Mo,W),3,4 Results and Discussion Mn2(MnTa3)N6-d (0d1)5 and Ln2Ta2O5N2 (Ln=lanthanide). 6 Bimetallic oxynitrides containing strongly electroposi- BaWO4 was heated in ammonia at diVerent temperatures. The product obtained when heated in ammonia at 900 °C for 12 h, tive elements such as LaTaO2N, Na3WO3N7 have substantial ionic character.Their limiting compositions have been was black and did not contain any starting material. The powder X-ray diVraction pattern of this sample (Fig. 1) was described by the normal rule of valency.8 There are reports on the ternary oxynitrides of rare-earths such as Ln2WO6-xNx9 diVerent from that of BaWO4.When heated in O2 atmosphere in the TPR system, this black material yielded only N2 as the and alkali-metal ions with tungsten. However, there are no reports on alkaline-earth metal and tungsten oxynitride phases. gaseous product. Since the sample was not highly crystalline and also because there was an indication of a small amount We wondered if partial substitution of N3- for O2- in the Ba–W–O system would yield any interesting material.Here of a W2N impurity phase, we were not sure whether the N2 came from the W2N phase alone or from any other unknown we report our work on the synthesis and structure of the new oxynitride Ba3W2O6N2. nitride product. To understand this reaction further, Ba2WO5 was heated in ammonia at 900 °C for 12 h.Interestingly, the powder pattern strongly resembles that of the black product Experimental obtained from BaWO4 heated in ammonia. A small amount of W2N impurity was also observed. From this it is clear that The previously reported oxides BaWO4, Ba2WO5, Ba3WO6 and Ba3W2O910 were synthesised by taking stoichiometric the starting oxide is decomposing to give rise to a new phase.When Ba3WO6 was heated in ammonia, a white compound amounts of BaO2 (99.5%, Fluka) and WO3 (99%, Koch Light) and heating in a muZe furnace and checking for product was obtained. The powder diVraction pattern of this compound contained the same unknown phase and BaO as impurity but formation by powder X-ray diVraction.About 1.5 g of these oxides was loaded in a quartz tube in an alumina boat. The there was no W2N impurity. The TPR of oxidation of this white product gave N2 as the gaseous product emanating samples were heated in flowing ammonia gas (flow rate ca. 120 ml min-1) at diVerent temperatures. The products were above 375 °C. The amount of N2 liberated was not as high as would be expected for a pure nitride phase indicating that analysed employing a JEOL-8P powder X-ray diVractometer (Cu-Ka radiation).The samples were heated in O2 atmosphere the new phase is an oxynitride. From this it is clear that the unknown phase has a Ba/W composition between that of in a temperature programmed reaction (TPR) system attached to a VG QXK300 quadrupole mass spectrometer11 to check Ba2WO5 and Ba3WO6.When Ba3W2O9 was heated in ammonia, a colourless single phase compound was obtained. for nitride phase formation. In a typical experiment about 200 mg of the sample was loaded and the reactor was evacuated This also liberated N2 on heating in O2. Experimental details to 10-6 Torr. Oxygen gas was admitted at ca. 20–25 mmol s-1 and the reactor was heated from 30 to 700 °C at a rate of 15 °C min-1 and the gaseous products were analysed.Nitrogen estimation was carried out using a home-built thermogravimetric analyser. EDX analyses of these samples were conducted using a Cambridge scanning electron microscope (SEM) S- 360, equipped with a LINK systems AN10000 X-ray analyser. FTIR spectra of the samples were recorded in polyethylene pellets in the range 100–700 cm-1 employing a Bruker IFS- 113V FTIR spectrometer and a Nicolet Impact 400D FTIR instrument in the range 700–1400 cm-1 using KBr pellets.Fig. 1 Powder X-ray diVraction pattern of BaWO4 heated in ammonia. † Contribution no. 1263 from Solid State and Structural Chemistry Unit. The asterisks indicate W2N impurity. J. Mater. Chem., 1997, 7(10), 2121–2125 2121Table 1 Summary of ammonolysis of diVerent ternary oxides phase in oxygen atmosphere is shown in Fig. 4. The oxidised product contained a mixture of Ba2WO5 and BaWO4 as the ammonolysis major phases and small amounts of Ba3W2O9. There was a starting compound conditionsa products 2.2% mass gain in the TG experiment, which could be attributed to the loss of N2 along with the uptake of oxygen. The BaWO4 900 °C Ba3W2O6N2+W2N Ba2WO5 900 °C Ba3W2O6N2+W2N molecular formula of the oxynitride from the TG studies is Ba3WO6 900 °C Ba3W2O6N2+BaO Ba3W2O9 800 °C Ba3W2O6N2 Ba3V2O8 900 °C no reaction aDuration 12 h.of all these studies are summarised in Table 1. When Ba3W2O9 was heated at diVerent temperatures, we found that the colour of the sample changed slowly from white to grey above 850 °C for 12 h.This colour change may be due to a partial reduction of WVI. From this point, we focused our attention on the white product obtained from Ba3W2O9. The powder X-ray diVraction pattern of this new phase is shown in Fig. 2 and the TPR of the new product in an oxygen atmosphere is shown in Fig. 3. We can see that the N2 was liberated at 350 °C with simultaneous uptake of oxygen.The TG of this oxynitride Fig. 2 Powder X-ray diVraction pattern of Ba3W2O6N2 Fig. 3 TPR of oxidation of Ba3W2O6N2 in O2 Fig. 5 (a) Electron diVraction pattern along the [0001] zone axis for Ba3W2O6N2; (b) electron diVraction pattern along the [01190] zone Fig. 4 TG of oxidation of Ba3W2O6N2 axis; (c) bright-field image of Ba3W2O6N2 crystallites 2122 J.Mater. Chem., 1997, 7(10), 2121–2125Ba3W2O6N2.00(±0.01). It is interesting that after the 2.2% mass Ba3W2O9 belongs to the B cation vacancy-ordered perovgain, the sample started losing mass above 500 °C. Although skite system with 2/3 of the octahedral sites in every layer the mass gain was expected, the mass loss could not be filled with tungsten ions.10 It can be recalled that, for Ba3W2O9, accounted for.To clarify this, Ba3W2O9 was heated in flowing the (W2O9)6- unit is in D3 symmetry and the tungsten trigonal O2 and it was found that this oxide does indeed lose mass prisms share faces. However, in the case of Ba3V2O8, vanadium above 500 °C, and gains its original mass on cooling. ions are in tetrahedral coordination and there is no bridging Scanning electron microscopy (SEM) of these samples was oxygen between the two tetrahedra.In Ba3W2O6N2, the conducted to confirm their metal composition. Spot mode (WO3N)3- unit should be in tetrahedral coordination and analysis showed that the Ba/W ratio is 352. The above therefore the local site group of tungsten tetrahedra is assumed observations lead to the following chemical equations for the to be C3v.To elucidate the local site geometry of oxide and ammonolysis of Ba3W2O9 and subsequent heating of the nitride ligands around the tungsten atom, FTIR and Raman product in oxygen atmosphere: spectroscopic investigations were carried out for these samples. Fig. 6 shows the FTIR spectra of the Ba3W2O9 and Ba3W2O9+4NH3 �Ba3W2O6N2+3H2O+N2+3H2 (1) Ba3W2O6N2 phases.The IR spectrum of Ba3W2O9 matches 2Ba3W2O6N2+3O2�2Ba2WO5+2BaWO4+2N2 (2) very well with the reported spectrum.10 The bands in the region 1000–650 cm-1 have been assigned to the terminal All the peaks in the powder X-ray diVraction pattern (Fig. 2) stretching modes and the region 650–450 cm-1 contains bridg- could be indexed to a hexagonal cell with a=5.993(2) and c= ing stretching modes.The far-IR bands can be attributed to 21.40(4) A ° . The intensity pattern of this sample was generated the bending modes. using the Lazy-Pulverix program with Ba3V2O812 as the struc- Fig. 7(a), (b) and (c) show Raman spectra of Ba3W2O9, tural model. The observed and calculated intensity patterns of Ba3V2O8 and Ba3W2O6N2 respectively. The FTIR and Raman Ba3W2O6N2 are given in Table 2.The calculated intensities spectra of the oxynitride closely resemble those of the Ba3V2O8 match the observed intensities quite well. For comparison phase. There are no bands in the region 650–450 cm-1, which powder diVraction data of Ba3V2O8 is also given in Table 2. means that there are no bridging modes in the oxynitride. This Electron diVraction and microscopy was carried out in order observation clearly confirms that the bridging oxygen atoms to confirm the crystal structure of this phase.Fig. 5(a) and (b) were lost during the reaction and that the nitridation process show electron diVraction along [0001] and [01190] zone axes, is not topotactic in nature because the local coordination confirming the hexagonal symmetry of the phase. The d-values geometry around tungsten is completely changed.On the basis obtained from electron diVraction agree with the observed dof intensities and on comparison with the spectra of (MO3N)n- parameters from the powder X-ray diVraction. Fig. 5(c) shows molecules (M=Re and Os),13,14 we assign these bands to the bright field image of the powder particles. The particles fundamentals of the (WO3N)3- unit.According to the irreduc- are in the range of submicrometre regime with irregular morphology. ible representations of the C3v point group, C=3A1+3E. All Table 2 X-Ray powder diVraction data for Ba3W2O6N2 [cell parameters a=5.993(2), c=21.40(4) A ° ] and Ba3V2O8 [cell parameters a=5.7845(1), c=21.317(1) A ° ]a Ba3W2O6N2 Ba3V2O8 h k l dobs/A ° dcalc/A ° I/I0(obs.) I/I0 (calc.) dobs/A ° I/I0 (obs.) 0 0 3 7.131 7.136 1 <1 7.09 4 1 0 1 5.039 5.057 5 6 4.878 11 0 1 2 4.670 4.680 2 2 4.537 3 0 1 4 3.738 3.731 <1 <1 3.651 12 0 0 6 3.560 3.568 3 3 — — 0 1 5 3.302 3.306 100 100 3.247 100 1 1 0 2.998 3.005 89 90 2.893 75 0 2 1 2.585 2.583 2 1 2.487 4 2 0 2 — 2.528 — <1 2.439 7 0 1 8 — 2.380 — 1 2.353 3 0 0 9 2.370 2.378 2 2 2.369 9 0 2 4 — 2.340 — <1 2.267 11 1 1 6 — 2.298 — <1 2.243 6 2 0 5 2.222 2.224 37 41 2.16 40 0 2 7 — 1.982 — <1 1.934 4 1 0 10 1.980 1.980 21 24 1.962 25 2 1 1 1.951 1.959 6 9 1.886 3 1 2 2 1.923 1.934 2 5 — — 1 1 9 — 1.865 — <1 1.833 10 2 1 4 — 1.846 — <1 1.785 3 1 2 5 1.787 1.787 19 17 1.731 25 3 0 0 1.733 1.735 14 18 1.670 13 0 2 10 1.653 1.653 11 12 1.624 12 1 2 8 — 1.585 — <1 1.543 1 3 0 6 — 1.560 — <1 1.511 2 2 0 11 1.559 1.558 1 3 — — 2 2 0 1.503 1.502 11 15 1.446 10 2 1 10 1.448 1.448 12 9 1.416 13 1 3 1 — 1.440 — <1 1.386 1 0 0 15 — 1.427 — 2 1.421 4 3 0 9 — 1.401 — <1 1.3647 3 1 3 4 — 1.393 — <1 1.3495 1 3 1 5 1.366 1.368 7 7 1.3208 9 aThe values for Ba3V2O8 are from JCPDS file 29–211.J. Mater. Chem., 1997, 7(10), 2121–2125 2123these six modes are IR- and Raman-active.There are six IR and Raman bands in the region 100–1400 cm-1 clearly con- firming the C3v site symmetry of the (WO3N)3- unit. We can assign the strong absorption band at 857 cm-1 [in Fig. 7(c)] to the symmetric stretching mode (n2). The less intense bands at 976, 707, 433, 370 and 296 cm-1 can be attributed to stretching (n1), antisymmetric stretching (n4) and three bending modes n6, n3 and n5.In the IR spectrum of Ba3W2O6N2, the broad unsymmetrical band centred around 350 cm-1 is composed of bands at 290, 340 and 391 cm-1 and three bands at 760, 815 and 983 cm-1 are assigned to n5, n3, n6, n4, n2 and n1 modes, respectively. The diVerence between FTIR and Raman spectra of the corresponding modes may be due to lack of coincidence arising from a coupling between the tetrahedral ions in the unit cell.15 Our assignments also agree with the general rule that symmetrical stretch vibration gives the most intense Raman line.16 Table 3 presents a comparison of IR spectrum of K2ReO3N and the Raman spectrum of [OsO3N]1- ion in solution with the FTIR and Raman spectrum of Ba3W2O6N2.It is clear from the table that the transition-metal and the nitrogen ligand bond is stronger than the corresponding metal–oxygen bond. The general usefulness of vibration spectroscopy for confirmation and identification Fig. 6 FTIR spectra of (a) Ba3W2O9 and (b) Ba3W2O6N2 of structural features in polytypes of mixed metal oxides such as Ba3(B,B¾)2O9-y (B,B¾=Mo, W, V and Ti) has been systematically studied.17 On the basis of above studies, we propose a structural model for this new oxynitride as shown in Fig. 8. From the literature we observe that in most of the oxynitrides which contain alkali/alkaline-earth metal the tetrahedrally coordinated transition metal in the highest possible oxidation state (i.e. NbV, TaV, MoVI, WVI, ReVII and OsVIII). The stability Fig. 8 Projection of Ba3W2O6N2 structure on the (010) plane Fig. 7 Raman spectra of (a) Ba3W2O9, (b) Ba3V2O8 and (c) Ba3W2O6N2 emphasising the (WO3N)3- tetrahedral polyanions. The unit cell is indicated by the dashed lines. Table 3 Fundamental frequencies of vibration (cm-1) of (MO3N)n- (M=W, Re and Os) groupsa assignment and description n1(A1) n2(A1) n3(A1) n4(E) n5(E) n6(E) compound n(MMN) ns(MO3) d(MO3) nas(MO3) d(OMO) d(OMN) K2(ReO3N) (IR solid) 1022 878 315 830 273 380 Na(OsO3N) (R soln.) 1021 897 309 871 309 372 Ba3W2O6N2 (IR solid) 983 815 340 760 290 391 Ba3W2O6N2 (R solid) 976 857 370 707 296 433 R=Raman; soln.=solution. 2124 J. Mater. Chem., 1997, 7(10), 2121–2125of Mn+ (M=transition metal) ion in the tetrahedral environ- W6+ atom changes from six to four and correspondingly, the local site group symmetry changes from D3 to C3v.This ment could be due to eVective p bonding formation by the nitrogen atom in contrast with this type of bonding by the oxynitride is isostructural with the Ba3V2O8 phase. The corresponding molybdenum analogue Ba3Mo2O6N2 was obtained oxygen atoms. This is in agreement with predictions based on the simple consideration that, owing to the formal negative but could not be synthesised in a pure form.charge of N3- and lower electronegativity of nitrogen than The authors thank Dr H. N. Vasan for collection of FTIR and oxygen, N3- will have greater tendency to donate electrons to Raman data of the samples. One of the authors (P. S. H.) the central metal atom through back bonding.13 thanks the Council of Scientific and Industrial Research It is interesting that when Ba3V2O8 is heated in ammonia, (CSIR), New Delhi for a fellowship.Financial support from there is no change in the powder diVraction pattern and also the Department of Science and Technology, Govt. of India is that there is no change in colour. This indicates a high gratefully acknowledged. thermodynamic stability of this oxide.The TPR of the product also did not show N2 released from the sample. It is worth noting that, in the case of ammonolysis of Ba3W2O9 and other References barium tungsten oxide systems, Ba3W2O6N2 was the only 1 Rainer Niewa and Herbert Jacobs, Chem. Rev., 1996, 96, . stable phase obtained. 2 S. Sellem, C. Potvin, J. M. Manoli, R. Contant and G. Dje�ga- Attempts were made to obtain isostructural oxynitride Mariadassou, J.Chem. Soc., Chem. Commun., 1995, 359. phases containing molybdenum, Sr3W2O9, Sr3Mo2O9 and 3 S. H. Elder, L. H. Doerrer, F. J. DiSalvo, J. B. Parise, D. Guyomard Ba3Mo2O9 oxide phases not having been reported in the and J. M. Tarascon, Chem. Mater., 1992, 4, 928. literature, in order to explore new possible oxynitrides belong- 4 P. Subramanya Herle, N.Y. Vasanthacharya, M. S. Hegde, ing to Ba3W2O6N2 family. We prepared a series of oxide J. Gopalakrishnan and G. N. Subbanna, J. Solid State Chem., 1994, 112, 208. mixtures containing BaO2, Sr(NO3)2 (Fluka, 99.5%), MoO3 5 J. Grins, P.-O. Ka�ll and G. Svensson, J. Solid State Chem., 1995, (Aldrich 99.9%) and WO3 with (Ba,Sr):(Mo,W) in the ratio 117, 48. 352, and heated them to 600 °C for 24 h with intermittent 6 F. Pors, R. Marchand and Y. Laurent, J. Solid State Chem., 1993, grindings. Ammonolysis of these oxides was subsequently 107, 39. carried out at diVerent temperatures and the products were 7 S. H. Elder, F. J. DiSalvo, J. B. Parise, J. A. Hriljac and analysed. Only for Ba–Mo oxide did X-ray patterns show the J. W. Richardson, Jr., J.Solid State Chem., 1994, 108, 73. 8 D. S. Bem and H.-C. Zur Loye, J. Solid State Chem., 1993, 104, 467. presence of an oxynitride phase related to Ba3W2O6N2 and 9 R. Marchand, P. Antoine and Y. Laurent, J. Solid State Chem., BaMoO3. In the other cases a perovskite-like phase was the 1993, 104, 34. major phase. It is known in the literature that the oxynitride 10 K.R. Poeppelmeier, A. J. Jacobson and J. M. Longo, Mater. Res. SrMoO2.6N0.418 can be synthesised by heating SrMoO4 in Bull., 1980, 15, 339. ammonia. For the BaMoO3 related phase which we obtained 11 M. S. Hegde, S. Ramesh and G. S. Ramesh, Proc. Indian Acad. Sci. as one of the product phases, it is quite possible that partial (Chem. Sci.), 1992, 104, 591. 12 P. Su� sse and M. J. Buerger, Z. Kristallogr., 1970, 131, 161. substitution of oxygen with nitrogen can take place. 13 B. Krebs and A. Mu� ller, J. Inorg. Nucl. Chem., 1968, 30, 463. 14 L. A.Woodward, J. A. Creighton and K. A. Taylor, T rans. Faraday Soc., 1960, 56, 1267. Conclusions 15 R. H. Busey and O. L. Keller, Jr., J. Chem. Phys., 1964, 41, 215. We have synthesised a new oxynitride Ba3W2O6N2 from the 16 K. Nakamoto, Infrared and Raman Spectra of Inorganic and Co-ordination Compounds, John Wiley & Sons Inc., New York, 3rd ammonolysis of Ba3W2O9 at 800 °C for 12 h. This confirms edn., 1978. the generally observed trend that the ternary oxides containing 17 B. Mo� ssner and S. Kemmler-Sack, J. L ess-Common Met., 1985, the most electropositive ions (alkaline, alkaline earth, rare 114, 333. earth) will not form ternary nitrides by ammonolysis. They 18 Guo Liu, Xinhua Zhao and H. A. Eick, J. Alloys Compd., 1992, will perhaps form oxynitrides, or decompose into the electro- 187, 145. positive metal oxide and binary transition-metal nitride. From Ba3W2O9 to Ba3W2O6N2 the coordination number around the Paper 7/02969A; Received 30th April, 1997 J. Mater. Chem., 1997, 7(10), 2121&ndash

ISSN:0959-9428    

DOI:10.1039/a702969a

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Ordering of nitrogen and oxygen in nitrogen-containing melilites Y2Si3O3N4 and Nd2Si2.5Al0.5O3.5N3.5 Pei-LingWang,*a Per-Eric Werner,b Lian Gao,a Robin K. Harrisc and Derek P. Thompsond aState Key L ab of High Performance Ceramics and SuperfineMicrostructure, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China bDepartment of Structural Chemistry, Arrhenius L aboratory, University of Stockholm, S-106 91 Stockholm, Sweden cDepartment of Chemistry, University of Durham, South Road, Durham, UK DH1 3L E dMaterials Division, Department of Mechanical, Materials andManufacturing Engineering, University of Newcastle, Newcastle upon T yne, UK NE1 7RU Ordering of N and O atoms in nitrogen-containing melilite Y2Si3O3N4 (Y-M) was investigated by Rietveld refinement in the space group P4 : 21m using neutron powder diVraction data. The results show that, contrary to previous work using X-ray data, the occupancies by N and O atoms can be refined to show that about 1.7 N and 0.3 O atoms occupy 2c (inWyckoV notation, space group no. 113) sites in the Y-M unit cell, while the remaining N atoms in the unit cell are disordered at 8f (inWyckoV notation) sites.The similar ordering of N and O atoms in Nd2Si2.5Al0.5O3.5N3.5 (Nd-M¾) was also confirmed, by refinement from neutron data, which means that similar distributions of N and O atoms occurred at 2c sites in spite of the fact that the total numbers of N atoms per unit cell were eight and seven in Y-M and Nd-M¾ respectively. The present results give a more detailed picture of O,N ordering in these structures than was provided by the previous work ofWang andWerner based on X-ray data alone.Nitrogen-containing melilite and solid-solution phases, of gen- eight N atoms were located at 8f sites (eight-fold) of the unit cell4 (Fig. 1). The O,N ordering in Y-M was studied5 with eral formula R2Si3-xAlxO3+xN4-x (x1, R=Y, Nd, Sm, Gd, Dy, etc.), frequently occur at grain boundaries when yttria and MAS NMR by two of the authors and colleagues recently.Readers are referred to that paper for further details of the rare-earth oxides are used as sintering aids in the preparation of a-sialon and mixed a,b-sialon ceramics.1,2 The suggested3 structure. It was found that N atoms probably occupy the 2c (bridging, two-fold) sites and O atoms occupy 4e (terminal, structure of Y-M has space group P4 : 21m and is derived from akermanite (Ca2MgSi2O7) by substituting Y for Ca, Si for Mg four-fold) sites in the Si2(O,N)7 units, while the remaining six N and two O atoms are distributed at 8f sites.This is in and N for eight of the fourteen O in the unit cell (Z=2) for N. Further structure refinements based on the Rietveld whole- accordance with Pauling’s rules for charge distribution that O atoms are located at the 4e sites as these are the only ones in pattern fitting technique for Y2Si3O3N4 (Y-M) and Nd2Si2.5Al0.5O3.5N3.5(Nd-M¾) confirmed the model, in which the structure not coordinated to two silicon atoms and they have the lowest total positive valencies from surrounding cations (ca. 2). In order to determine the N,O ordering at the 2c and 8f sites of the Y-M unit, it is clearly more powerful if two techniques (Rietveld refinement and NMR) are used in combination.Since the scattering factors of N and O atoms for neutrons show a greater contrast than those for X-rays, neutron diVraction data were used to investigate the N,O ordering in Y-M to supplement the X-ray results.This paper describes the conclusions regarding N,O ordering from the refinement of the occupancies by N and O atoms together with other structural and profile parameters based on the distributions of N,O atoms of Y-M obtained by NMR.5 The N,O ordering in the Nd-M¾ unit cell was also derived and compared in this work. Experimental Samples of the Y-M and Nd-M¾ phases were prepared by the hot-pressing technique.The starting powders used were silicon nitride (LC12, H.C. Starck, Berlin), aluminium oxide (Wu Song Chemical Works, China, 99.9%) and aluminium nitride (prepared at Shanghai Institute of Ceramics), while Y2O3 and Nd2O3 were the products of Yaolung Chemical Works, China. Fig. 1 N-Melilite crystal structure, (001) projection. Large circles Powder mixtures were prepared with the overall nominal represent Y/Nd cations and small filled circles, at tetrahedron centres, composition R2Si3-xAlxO3+xN4-x (x=0 for R=Y and x=0.5 correspond to Si (and Al) in 4e and 2a sites.The remaining circles for R=Nd), taking into account surface oxygen on the particles correspond to 8f (filled) and 4e (open) non-metal sites, while open squares indicate 2c sites.of both Si3N4 (1.8 mass%) and AlN (2.0 mass%). The powders J. Mater. Chem., 1997, 7(10), 2127–2130 2127were mixed in absolute alcohol and milled in an agate mortar cell dimensions (a,c), one peak shape parameter and five background parameters (so-called profile parameters). No for 1.5 h. The compacted samples were fired by the hot-pressing technique (20 MPa) in a flowing nitrogen atmosphere for 1 h absorption correction was applied, and a common isotropic temperature factor was used for all the atoms. However, at the at 1750 and 1675 °C for Y-M and Nd-M¾ phases, respectively.6 The synthesised samples were characterised by powder X-ray beginning of the refinement procedure, an attempt was made to refine thermal parameters together with occupancies (30 diVraction (XRD) and it was found that the major phase was melilite, with a very small amount of the J-phase R4Si2O7N2 variables in total ).Although the error parameters of the fitting were lowered by this procedure, some of the thermal param- (R=Y, Nd) also being present in both samples. The unit cells of the Y-M and Nd-M¾ phases were refined, using Si powder eters were found to be negative, which is unacceptable.For this reason, we prefer to quote results obtained by constraining as an internal standard, from X-ray Guinier-Ha�gg camera diVraction patterns (Cu-Ka1 radiation, l=1.540 5981 A ° ) evalu- all thermal parameters to be equal (see Table 3). The refinement was terminated when all shifts in the parameters were <10% ated with a computer-controlled film scanner and associated programs.7,8 The neutron data collections (l=1.470 A ° ) of Y- of the corresponding standard deviations.At the final stage, the occupancies of N and O atoms for Y-M at 2c and 8f sites M and and Nd-M¾ were performed at 275 K at the Swedish Studsvik R2 reactor. Results were obtained in steps of 0.08° in were refined together with 24 parameters, keeping the total number of N and O atoms constant.The final R values for the ranges for 2h of 15–130 and 10–109.4° for Y-M and Nd- M¾ respectively, with a measuring time of 3 min per step. 195 reflections, together with some essential data, are shown in Table 1. For comparison, the refinement results of Y-M from neutron and X-ray data,4 in which two O and eight N Results atoms were located at 2c and 8f sites respectively, are also listed in Table 1.Fig. 2 shows the observed and computer-fitted neutron diVraction pattern of Y-M. The unit-cell dimensions obtained after Because of the higher scattering amplitudes of N,O atoms in the neutron case, the distributions of N,O atoms were much least-squares refinements are: a=b=7.6137(2), c=4.9147(2) A ° for Y-M and a=b=7.7462(5), c=5.0390(4) A ° for Nd-M¾.The more sensitive indicators than in the X-ray case. It is shown in Table 1 that all reliability index R values were obviously Rietveld refinements were performed with a version of the refinement program written by Wiles et al.9 The background reduced by fixing two N atoms at 2c sites, together with six N and two O atoms at 8f sites.The results were even better when intensity Ybi at the ith step was described by the polynomial the occupancies of N,O atoms at 2c and 8f sites were refined Ybi=SBm[2hi/BKPOS-1]m together. An unreasonable occupancy of N atoms at 2c sites was obtained for a similar refinement from X-ray data. where Bm are parameters to be refined and BKPOS is the origin of the polynomial for the background.e peak shape For refinements of Nd-M¾ the models were more complicated, because there are three kinds of possible substitutions used was a Pearson VII function. The extent of a peak was taken to be 3.0 times the FWHM (full-width at half-maximum), of Al for Si, i.e. one Al atom distributed in either 4e or 2a or in both of these sites, in addition to the diVerent distributions Hk, on either side of the peak centre.Hk was given by of anions. Similar refinement procedures as used for Y-M were Hk2=Utan2hk+Vtanhk+W performed on Nd-M¾ neutron data under the diVerent distributions of Si, Al atoms. Table 2 gives the final R values for where U, V, W are the width parameters and k is the reflection index. The atomic coordinates of Y-M and Nd-M¾ obtained the diVerent distributions of N,O,Si,Al in Nd-M¾.The atomic coordinates, isotropic thermal and occupancy parameters for in our previous work1 were used as the starting parameters. In the first trial for Y-M, two N atoms were located at 2c sites both Y-M and Nd-M¾ structures are shown in Table 3. Some selected interatomic distances for Y-M and Nd-M¾ are summar- of the space group P4 : 21m, and the remaining six N atoms were taken to occupy 8f sites of the unit cell together with two ised in Table 4.Similar refinement levels to those for Y-M were obtained O atoms. The convergent refinement involved twelve structural parameters and twelve profile parameters, i.e. one scale, ten for the Nd-M¾ phase. The R values were much lower when N,O atoms were refined at (2c,8f ) sites or were fixed at those atomic coordinates, and one isotropic temperature factor (socalled structural parameters), together with the zero-point sites (see Table 2).The results shown in Table 3 confirm that the N atoms occupy the two sites 2c and 8f in both Y-M and position, three peak half-width parameters (U, V, W ), the unit Nd-M¾ cases. The occupancies of N, O atoms further indicate Table 1 Final R values and refinement details for the diVerent distributions of O,N in Y-M N,O O(2c), N(8f ) N(2c) refined at N,O(8f ) (2c,8f ) X-ray4 neutron neutron neutron RF (%)a 5.25 6.37 4.00 3.94 RB (%) 6.55 11.91 6.81 6.67 RP (%) 6.76 6.88 4.97 4.89 RWP (%) 8.45 9.06 6.77 6.69 U 0.047(0) 1.54(8) 1.41(6) 1.41(6) V -0.014(1) -0.83(7) -0.74(5) -0.74(5) W 0.024(1) 0.20(1) 0.18(1) 0.18(1) no.of structural parameters 17 12 12 13 no. of profile parameters 11 12 12 12 Fig. 2 The neutron diVraction pattern for Y-M. Top: profile calculated by least-squares Rietveld refinement. Middle: experimental pattern. aReliability index R can be defined as RF, RB, RP, RWP. RF: R value for structure amplitudes; RB: R value for Bragg intensities; RP: the Bottom: diVerence pattern.The asterisks indicate peaks assigned to Y4Si2O7N2, which is the principal impurity in Y-M. pattern R factor; RWP: the weighted pattern R factor. 2128 J. Mater. Chem., 1997, 7(10), 2127–2130Table 2 Final R values and refinement details for the diVerent distributions of N,O,Si,Al in Nd-M¾ (x=0.5) (l=1.470 A° ) 1 Al and 3 Si at (4e) sites 4 Si at (4e) sites 1 A1, 5 Si refined at 2 Si at (2a) sites 1 A1 and 1 Si at (2a) sites (4e, 2a) sites 0(2c) N(2c) N,O O(2c) N(2c) N,O O(2c) N(2c) N,O N(8f ) N,O(8f ) refined N(8f ) N,O(8f ) refined N(8f ) N,O(8f ) refined RF (%) 6.52 3.46 3.52 6.59 3.50 3.52 6.54 3.44 3.49 RB (%) 11.26 5.65 5.53 11.24 5.69 5.55 11.24 5.60 5.49 RP (%) 7.24 5.16 5.06 7.25 5.20 5.11 7.24 5.16 5.06 RWP (%) 9.27 6.81 6.71 9.28 6.84 6.74 9.26 6.81 6.70 Table 3 Final refinement results of atomic coordinates, isotropic thermal (B) and occupancy (N) parameters in Y-M and Nd-M¾ (x=0.5) (l=1.470 A ° ) Nd-M¾ (N, O refined at (2c, 8f ) sites) Y-M WyckoV N, O refined 1 Al, 3 Si at 1 Al, 1 Si at 1 Al, 5 Si refined atom notation x,y,z, B, N at (2c, 8f ) sites (4e) sites (2a) sites at (4e, 2a) sites Y/Nd 4e x 0.3365(4) 0.3361(4) 0.3359(4) 0.3360(4) y 0.1635(4) 0.1639(4) 0.1641(4) 0.1640(4) z 0.5038(8) 0.5019(9) 0.5022(9) 0.5020(9) N 0.5 0.5 0.5 0.5 Si(1)/Al(1) 4e x 0.1456(8) 0.1429(9) 0.1423(8) 0.1427(9) y 0.3544(8) 0.3571(9) 0.3577(8) 0.3573(9) z 0.9464(13) 0.9473(14) 0.9474(14) 0.9473(14) Si(1) N 0.5 0.375 0.5 0.407(26) Al(1) N 0.125 0.093(26) Si(2)/Al(2) 2a x 0.0 0.0 0.0 0.0 y 0.0 0.0 0.0 0.0 z 0.0 0.0 0.0 0.0 Si(2) N 0.25 0.25 0.125 0.218(26) Al(2) N 0.125 0.032(26) N(1)O(1) 2c x 0.5 0.5 0.5 0.5 y 0.0 0.0 0.0 0.0 z 0.1882(12) 0.1785(11) 0.1783(11) 0.1784(11) N(1) N 0.213(6) 0.211(6) 0.212(6) 0.212(6) O(1) N 0.037(6) 0.039(6) 0.038(6) 0.038(6) N(2)/O(2) 8f x 0.0852(3) 0.0847(3) 0.0849(4) 0.0848(3) y 0.1600(4) 0.1641(4) 0.1641(4) 0.1641(4) z 0.7976(6) 0.8047(6) 0.8049(6) 0.8047(6) N(2) N 0.787(6) 0.664(6) 0.663(6) 0.663(6) O(2) N 0.213(6) 0.336(6) 0.337(6) 0.337(6) O(3) 4e x 0.1416(4) 0.1397(5) 0.1396(5) 0.1397(5) y 0.3584(4) 0.3603(5) 0.3604(5) 0.3603(5) z 0.2783(11) 0.2842(11) 0.2786(11) 0.2784(11) N 0.5 0.5 0.5 0.5 Ba 0.83(4) 0.24(4) 0.24(4) 0.24(4) a Isotropic thermal parameter, see text that ca. 1.7 N and 0.3 O atoms occupy 2c sites in both Y-M and Nd-M¾ units in spite of the fact that the total numbers of Table 4 Some selected interatomic distances in Y-M and Nd-M¾ N atoms are eight and seven in Y-M and Nd-M¾ respectively; the remaining N atoms in the unit cells are disordered to Y-M (N, O refinedNd-M¾ (N, O, Al, occupy 8f sites.at 2c, 8f sites) Si refined) It was diYcult to distinguish between Al and Si atoms in M(Y/Nd)M[N( l ) O(1)] 2.346(5) 2.426(5) Nd-M¾ because of the small diVerence in scattering factors, as MO(3) 2.360(5) 2.428(6) stated in our previous work.4 However, the atomic ratio of M[N(2), O(2)] (×2) 2.397(4) 2.472(5) Al5Si at 4e sites (0.1950.81) appears to be marginally higher MO(3) (×2) 2.551(5) 2.607(5) than the one (0.1350.87) at 2a sites.M[N(2), O(2)](×2) 2.754(4) 2.806(5) average 2.514(5) 2.578(5) This work was partly supported by a Royal Society Joint (Si/Si,Al)MO(3) 1.661(8) 1.669(9) M[N(1), O(1)] 1.702(7) 1.687(7) project between Shanghai Institute of Ceramics and the M[N(2), O(2)] (×2) 1.714(7) 1.720(8) Nitrogen Ceramics Research Group in the University of average 1.698(7) 1.699(7) Newcastle, UK. We also acknowledge financial support from the National Natural Science Foundation of (Si/Si,Al)M[N(2), O(2)] (×4) 1.701(3) 1.737(3) China.J. Mater. Chem., 1997, 7(10), 2127–2130 21296 P. L. Wang, H. Y. Tu, W. Y. Sun, D. S. Yan, M. Nygren and References T. Ekstro�m, J. Eur. Ceram. Soc., 1995, 15, 689. 1 S. Slasor, K. Liddell and D. P. Thompson, Br. Ceram. Proc., 1986, 7 K. E. Johansson, T. Palm and P.-E. Werner, J. Phys. E.: Sci. 37, 51. Instrum., 1980, 13, 1289. 2 P. L. Wang, W. Y. Sun and T. S. Yen, (D. S. Yan,), Mater. Res. Soc. 8 P.-E.Werner, Ark. Kemi, 1969, 31, 513. Symp. Proc., ed. I.-W. Chen, P. F. Becher, M. Mitomo, G. Petzow 9 D. B. Wiles, A. Sakthivel and R. A. Young, User’s Guide to Program and T. S. Yen, MRS Pittsburgh, PA, 1993, vol. 287, p. 387. DBW32s for Rietveld Analysis of X-Ray and Neutron Powder 3 A. W. J. M. Rae, D. P. Thompson, N. J. Pipkin and K. H. Jack, DiVraction Pattern, Version 8804, School of Physics, Institute of Special Ceram., 1975 6, 347. Technology, Atlanta, USA, 1988. 4 P. L. Wang and P.-E. Werner, J.Mater. Sci., 1997, 32, 1925. 5 A. Koroglu, D. C. Apperley, R. K. Harris and D. P Thompson, Paper 7/02842C; Received 25th April, 1997 J.Mater. Chem., 1996, 6, 1031. 2130 J. Mater. Chem., 1997, 7(10), 2127–21

ISSN:0959-9428    

DOI:10.1039/a702842c

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Preparation of the layered double hydroxide (LDH) LiAl2(OH)7·2H2O, by gel to crystallite conversion and a hydrothermal method, and its conversion to lithium aluminates M. Nayak,a T. R. N. Kutty,*a V. Jayaramanb and G. Periaswamy,b aMaterials Research Centre, Indian Institute of Science, Bangalore 560 012, India bMaterials Chemistry Division, Indira Gandhi Centre for Atomic Research, Kalpakkam 603 102, India A layered double hydroxide (LDH) with chemical composition LiAl2(OH)7·2H2O was prepared via a wet chemical route of gel to crystallite (G–C) conversion at 80 °C involving the reaction of hydrated alumina gel, Al2O3·yH2O (80<y<120) with LiOH (Li2O/Al2O30.5) in presence of hydrophilic solvents such as ethanol under refluxing conditions.The hydrothermal synthesis was carried out using the same reactants by heating to140 °C in a Teflon-lined autoclave under autogenerated pressure (20 MPa).Transmission electron microscopy showed needle-shaped aggregates of size 0.04–0.1 mm for the gel to crystallite conversion product, whereas the hydrothermal products consisted of individual lamellar crystallites of size 0.2–0.5 mm with hexagonal morphology. The LDH prepared through the gel to crystallite conversion could be converted into LiAl(OH)4·H2O or LiAl(OH)3NO3·H2O by imbibition of LiOH or LiNO3, respectively, under hydrothermal conditions.Thermal decomposition of LDH above 1400 °C gave rise to LiAl5O8 accompanied by the evaporation of Li2O. LiAl(OH)4·H2O and LiAl(OH)3NO3·H2O decomposed in the temperature range 400–1000 °C to a- or b-LiAlO2.The compositional dependence of the product, the intermediate phases formed during the heat treatment and the possible reactions involved are described in detail. The compounds LiAl5O8 and LiAlO2 are luminescent hosts The gel can be converted directly to crystallites in the presence of an organic solvent owing to instability of the gel caused by when doped with Fe3+, emitting in the red spectral region.1,2 These red-emitting phosphors are useful for artificial illumi- influx of aliovalent ions.The merits of this process over the conventional ceramic processing are the increased homogeneity nation in plant growth applications.2–4 One of the polymorphic forms of LiAlO2, viz. c-LiAlO2, has received much attention of the products and the reduction in the processing temperature.G–C conversion can take place even with coarser gels so because of the possibility of its use as a tritium breeding material in fusion reactors5,6 and as an electrolyte matrix for that the raw materials need not include expensive organometallics or alkoxides. The general reaction involved in this molten carbonate fuel cells.7,8 High surface area a-LiAlO2 is used as a catalyst support9 and for the preparation of the technique is the breakdown of the gel network owing to the ionic pressure caused by the influx of aliovalent ions. delithiated transitional alumina compounds.10 The present work aims to decipher routes to obtain phase-pure LiAlO2 Hydrothermal synthesis is based on supersaturated solvents under elevated pressure–temperature (P–T) conditions; accord- (diVerent polymorphic forms) and LiAl5O8 wherein LDH [LiAl2(OH)7·2H2O] is used as a precursor and also has many ingly the end product may diVer.applications of its own. Layered lithium dialuminium hydroxide, LiAl2(OH)7·2H2O, which is analogous to the mineral Experimental hydrotalcite, [Mg6Al2(OH)16]CO3·4H2O, has received a lot of attention because of its potential application in the field of The principle involved and the experimental details of G–C sensors, and also as antacid, by way of selective sorption of conversion technique have been presented in our previous weak acids (H2S, CO2, etc.,).It also finds uses in ion exchange publications.18–20 Hydrated alumina gel was prepared through for poisonous anions such as [Fe(CN)6]4-, solid-state anion precipitating Al3+ (aq) with 30 mass% ammonia solution, conductors and in catalysis.11–17 Moreover, the compound is washed free of anion contaminants using hot water (tested for a precursor for the preparation of LiAl5O8 and LiAlO2.the absence of sulfate by adding Ba(OH2) solution and for A number of publications exist in the literature on the chloride ions by adding AgNO3 solution to the filtrate) and synthesis and physicochemical properties of these com- suspended in a conical flask containing ethanolic lithium pounds.1–17 Poppelmeir et al.17 prepared LiAl2(OH)7·2H2O by hydroxide.The presence of anionic contaminants such as the insertion of LiOH into Al(OH)3 and studied the phase SO42- and Cl- impedes the reaction.The reaction vessel was relations and stability of the compounds at diVerent tempera- fitted with a water-cooled condenser and an alkali guard tube ture regimes. Their papers do not deal with the phases stabilised to prevent the ingress of CO2 and refluxed for 5–6 h at 80 °C at temperatures >1200 °C and the possible conversion of while continuously stirring using a magnetic stirrer.The solid LiAlO2 to LiAl5O8. The diYculty in the imbibition technique product obtained was washed free of unreacted LiOH and airto maintain the Li2O/Al2O3 molar ratio, often leads to the dried in a desiccator. The lithium content in the washings was presence of unreacted reactants and phase purity is diYcult to monitored to measure the extent of reaction and to compare attain. Most of the literature is on the ion exchange, intercal- the composition of the solid product obtained via the wetation products and the possible applications thereof.We chemical route (AAS). prepared LiAl2(OH)7·2H2O through the novel route of gel to Hydrothermal preparation was carried out in a Teflon crystallite (G–C) conversion18–20 as well as via a hydrothermal (PTFE)-lined Morey-type autoclave SS318.The autoclave was method,21–23 which acted as a precursor for the preparation of charged with the reactants, viz. hydrated alumina gel mixed LiAlO2 and LiAl5O8. The advantages of the G–C conversion with LiOH in the desired molar ratio; deionised water was include procedural simplicity and economy of the method as added to the required percentage so as to autogenerate pressure in the range 5–40 MPa and heated in the temperature range the starting materials are cheap water-soluble inorganic salts.J. Mater. Chem., 1997, 7(10), 2131–2137 2131100–240 °C for 12 h. The temperature was varied to check the phase stability region. Hydrothermal imbibition was performed by charging the autoclave with previously prepared LDH and LiOH or LiNO3 in the desired molar ratio.Imbibition was carried out at 140 °C for 12 h. The chemical compositions of the products were determined by wet chemical analyses using atomic absorption spectroscopy (AAS). Thermal analyses were performed on a simultaneous thermogravimetry–diVerential thermal analysis (TG–DTA) instrument from Polymer Laboratory STA-1500 at a heating rate of 5 °C min-1.Phase identification of the powders was carried out by X-ray powder diVraction using a Scintag/USA diVractometer. IR absorption spectra were recorded on a BIORAD FTIR spectrometer in the range 4000–400 cm-1. Solid-state 27Al MAS NMR spectra were obtained at 78.2 MHz using a high-resolution NMR spectrometer (BRUKER 300 MHz) at room temperature fitted with a magic angle spinning probe (MAS) for rotating the sample at a frequency of 7 kHz. 27Al MAS NMR chemical shifts (d) were referenced to 1 M Al(H2O)6Cl3 (d=0). The morphology and the particle size were determined by the intercept method on the micrographs obtained from a JEOL 200 CX, 200 kV transmission Fig. 1 Variation of Li2O/Al2O3 molar ratio in the product as a electron microscope (TEM) having 2 A ° resolution.function of initial Li2O/Al2O3 molar ratio in the reaction mixture Results and Discussion Gel to crystallite conversion EVect of initial composition on phase formation. The composition of the solid phases prepared through G–C conversion were dependent on the initial Li2O/Al2O3 mol ratio in the reaction mixture (Fig. 1). It is evident that as the Li2O/Al2O3 ratio is increased in the reaction mixture, the lithium retained in the product is also increased and attained a limiting value of ca. 0.5. Above this concentration, the excess lithium added was washed out. Table 1 gives the chemical composition of the as-prepared samples for various starting ratios and the resultant products on subsequent heat treatment at 900 or above Fig. 2 TG–DTA traces of LiAl2(OH)7·2H2O prepared through G–C 1400 °C. At Li2O/Al2O30.05, the phase formed is pseudo- conversion boehmite which decomposed on heat treatment to give a-Al2O3 and LiAl5O8.For Li2O/Al2O3 ratios between 0.05 and 0.5, a mixture of nordstrandite [Al(OH)3] and A mass loss (16%) between 100 and 200 °C accompanied by a very strong endothermic peak in the DTA is due to the removal LiAl2(OH)7·2H2O was formed.Above a Li2O/Al2O3 ratio of 0.5, a voluminous mass having the chemical composition of structural water. Between 200 and 500 °C, the mass loss is ca. 32% which can be attributed to dehydroxylation. DTA LiAl2(OH)7·2H2O was obtained. The formation of LDH is confirmed by X-ray diVraction (XRD) (Fig. 4, later). Preparing shows a corresponding broad and shallow endotherm centred around 270 °C.A subsequent endotherm at 640 °C a phase of composition LiAl(OH)4·H2O was not possible via this route owing to leaching of LiOH. accompanied by a mass loss of 6% is due to the dehydroxylation of the residual hydroxy groups. Continuous mass loss between 1000 and 1300 °C, and the endotherm which shows a Thermal analyses. Fig. 2 shows thermal analyses traces of LiAl2(OH)7·2H2O prepared by G–C conversion.The sample peak at 1143 °C, is due to the evaporation of Li2O. Samples prepared at higher initial Li2O/Al2O3 ratios (>0.5) show the showed a total mass loss of 53% up to 540 °C in agreement with the literature,17 however the sequence of mass losses same trends as above during thermal analyses. Isothermal mass loss measurements were also carried out at diVer.The TG curve shows that the major mass loss occurred below 540 °C in three steps. The initial mass loss (5%) below various temperatures until a constant mass was obtained and the results are in good agreement with the dynamic mass loss 100 °C is probably due to the loss of physically absorbed water. Table 1 Chemical compositions of the as-prepared LDH and the phases obtained upon calcination above 1400 °C Li2O/Al2O3 molar ratio after calcination at reaction mixture product as-prepared 900 °C 1400 °C 0.05 0.04 pseudoboehmite a-Al2O3+LiAl5O8 a-Al2O3+LiAl5O8 0.19 0.18 Al(OH)3+LiAl2(OH)7·2H2O c-LiAlO2+LiAl5O8 c-LiAlO2+LiAl5O8 0.22 0.2 Al(OH)3+LiAl2(OH)7·2H2O LiAl5O8 LiAl5O8 0.29 0.29 Al(OH)3+LiAl2(OH)7·2H2O (major) c-LiAlO2+LiAl5O8 LiAl5O8 0.5 0.49 LiAl2(OH)7·2H2O c-LiAlO2+LiAl5O8 LiAl5O8 0.72 0.49 LiAl2(OH)7·2H2O c-LiAlO2+LiAl5O8 LiAl5O8 1 0.5 LiAl2(OH)7·2H2O c-LiAlO2+LiAl5O8 LiAl5O8 2132 J.Mater. Chem., 1997, 7(10), 2131–2137measurements. Static mass loss measurements shows a total were observed between 1000 and 400 cm-1. These sharp peaks are attributed to AlO6 groups.24 Dm of ca. 58% between room temperature and 1300 °C.Based on the thermal analyses results, the following reaction scheme can be proposed for the formation of LiAl5O8 from X-Ray diVraction. XRD patterns as a function of calcination temperature are shown in Fig. 4. An oven dried (105 °C) sample LiAl2(OH)7·2H2O: does not show any diVerence in its XRD pattern compared LiAl2(OH)7·2H2O CA 150 °C LiAl2(OH)7+2H2O Dm=17%(1) with as-prepared specimens.The sample annealed at 150 °C shows a similar pattern but with broadened diVraction peaks and diminished intensity. The reason for this is the removal of LiAl2(OH)7CA 500 °C LiAl2O3(OH)+3 H2O Dm=30%(2) interlayer water, a fact confirmed by TG–DTA studies and IR absorption spectra. Further heating of the sample to 250 °C, resulted in a complete change of the diVraction pattern, which 4 LiAl2O3(OH) CA 600 °C 2 ‘Li2Al4O7’+2H2O Dm=7% consists of very broad and weak diVraction peaks corresponding to LiAlO2 with small crystallite size.The change in the (�3 LiAlO2+LiAl5O8) (3) XRD pattern is due to the complete destruction of brucite type layers. As the temperature of calcination is increased to 5 LiAlO2 CA 1400 °C 24 h LiAl5O8+2Li2OF Dm=18% (4) 450 °C, b-LiAlO2 started nucleating, and become a major phase at 600 °C.At 1000 °C the phases stablised are c-LiAlO2 (major) and LiAl5O8 (minor). Upon further increase in tempera- IR absorption spectra. Samples isolated from the isothermal ture, c-LiAlO2 decomposed and the formation of LiAl5O8 was mass loss studies were analysed by IR absorption spectra for favoured.This change is associated with the evaporation of the presence of hydroxy groups in the intermediates. Fig. 3 lithia, Li2O, as detected from the eZuent gas when purged shows the IR absorption spectra of the sample as a function with dry argon while the sample is heated in a tubular furnace. of temperature. The sample heated to 105 °C shows a broad Prolonged heat treatment above 1400 °C yielded monophasic absorption band centred around 3450 cm-1 owing to the LiAl5O8.Further phase separation occurred on raising the OMH stretching frequencies from hydrogen bonded as well as temperature above 1600 °C. The resultant phases are abridged hydroxy groups. Sharp peaks at 1025, 825 and Al2O3(minor) and LiAl5O8. 550 cm-1 and shoulders at 675 and 650 cm-1 are characteristics of AlO6 octahedra.On heating the sample to 250 °C, the Hydrothermal preparation of LiAl2(OH)7·2H2O peak at 3450 cm-1 is considerably reduced in intensity, indicat- The products from the hydrothermal preparative runs are ing that the hydroxy groups are removed from the sample. shown in Table 2. The temperature of the hydrothermal treat- Also, the peak at 1375 cm-1 vanishes, and a doublet appears ment was varied in order to study the conditions of preparation at ca. 1550 cm-1. A new peak was observed at 1200 cm-1. at which the phase is stable. Essentially, our results indicate Also, the number of peaks between 1000 and 500 cm-1 reduces that, along with the composition, P–T conditions have great to one (centred around 550 cm-1). When the sample was influence on the phase stability.LDH is stabilised only if heated at 800 °C, the absorption band at 3450 cm-1 was completely lost indicating the disappearance of all hydroxy groups. On further heating at 1400 °C, sharp multiple bands Fig. 4 XRD traces of the LiAl2(OH)7·2H2O as a function of Fig. 3 IR spectra of LiAl2(OH)7·2H2O as a function of temperature: temperature: (a) 105 °C, (b) 150 °C, (c) 250 °C, (d) 450 °C, (e) 1000 °C and ( f ) 1400 °C (6, c-LiAlO2; *, LiAl5O8) (a) 105 °C, (b) 250 °C, (c) 1000 °C and (d) 1400 °C J.Mater. Chem., 1997, 7(10), 2131–2137 2133Table 2 Results of the hydrothermal preparative runs at three diVerent temperatures as-prepared Li2O/Al2O3 in the reaction mixture 140 °C 180 °C 240 °C 1 LiAl2(OH)7·2H2O b-LiAlO2+c-AlOOH b-LiAlO2 0.5 LiAl2(OH)7·2H2O b-LiAlO2+c-AlOOH b-LiAlO2+c-AlOOH 0.33 LiAl2(OH)7·2H2O+c-AlOOH c-AlOOH c-AlOOH 0.25 c-AlOOH+LiAl2(OH)7·2H2O c-AlOOH c-AlOOH 0.2 c-AlOOH c-AlOOH c-AlOOH Li2O/Al2O30.5, at temperatures below 140 °C.Larger crystallite size (0.5 mm) and better crystallinity were obtained using higher concentrations of LiOH. As the temperature of the hydrothermal preparation was increased above 150 °C, irrespective of the composition, the phases stabilised were b- LiAlO2+c-AlOOH (boehmite) at Li2O/Al2O30.5 and only boehmite at Li2O/Al2O3<0.5.Pure phase b-LiAlO2 was obtained for Li2O/Al2O3=1 at a temperature 240 °C. Fig. 5 shows the X-ray diVraction patterns of LDH prepared through the hydrothermal route at 140 °C. All the diVraction peaks due to the basal planes are split, as a relt of anion insertion in the intermediate layer.DiVraction peaks arising Fig. 6 TG–DTA traces of LiAl2(OH)7·2H2O prepared via the hydrofrom the (330) and (600) reflections are totally absent. Fig. 5 thermal route (inset) shows enlarged portions of (002) and (004) reflections. On increasing the lithium concentration, the extent of splitting planes and the peaks are broadened accompanied by a shift was not very much increased. to higher angles, yet retaining the LDH pattern.Shallow, DTA traces of the hydrothermal products (Fig. 6) show broad endotherms at 250 and 275 °C arise from the removal multiple thermal events in the temperature range 30–400 °C of hydroxy groups attached to the Al and Li in the brucite unlike the G–C prepared sample.Sharp endotherms at 46 and type layers. X-Ray analysis of the heat-treated sample showed 77 °C are due to the desorption of water and endotherms at a completely amorphous pattern. A very broad and shallow 132 and 192 °C arise from the removal of interlayer water. endotherm centred around 950 °C is due to the loss of residual XRD analysis shows diminished intensity of all the basal hydroxy groups; the phases formed thereupon are mixtures of c-LiAlO2 and LiAl5O8. The intermediate phases, on heat treatment, are the same as that of the sample prepared through G–C conversion; above 1400 °C the phase stabilised is LiAl5O8.Table 3 shows TG data of the intermediate products formed. TG curves indicate that the major mass loss occurred below 500 °C.Between 500 and 1050 °C, the mass loss is only 3% owing to removal of residual hydroxy groups. The continuous mass loss above 1050 °C without any thermal events is due to the evaporation of Li2O which is responsible for the formation of LiAl5O8 above 1400 °C, according to reaction (4). Hydrothermal imbibition to prepare LiAl(OH)4·H2O To achieve higher Li2O/Al2O3 molar ratios in the LDH prepared through G–C conversion or hydrothermally, samples were further hydrothermally imbibed with LiOH and LiNO3 at 140 °C.Hydrothermal imbibition of LiOH into the LDH prepared through the hydrothermal route resulted in the formation of a-LiAlO2 and boehmite (c-AlOOH) [Fig. 5(c)]. Further insertion of LiOH created instability in the compound resulting in phase separation into a-LiAlO2 and boehmite.In Table 3 TG analyses data of the LDH prepared through hydrothermal route at 140 °C and the intermediate phases stabilised temperature/°C mass loss (%) phase formed RT–100 °C 6 LiAl2(OH)7·2H2O 100–200 °C 17 LiAl2(OH)7·2H2O (intensity of all the basal diVraction peaks diminished) 200–500 °C 30 very broad peaks corresponding to a-LiAlO2 , amorphous Fig. 5 XRD patterns of LiAl2(OH)7·2H2O prepared via the hydrother- background indication of mal route at 140 °C with initial Li2O/Al2O3 ratios of (a) 1, (b) 0.5 and poorly crystallised second phase (c) sample on further imbibition with LiOH. Inset shows the enlarged 500–1050 °C 3 b-LiAlO2+LiAl5O8 portions of the basal reflections of hydrothermally prepared samples 1050–1300 °C 7 LiAl5O8+c-LiAlO2 (a) (002) and (b) (004).(#, Boehmite; 6, a-LiAlO2). 2134 J. Mater. Chem., 1997, 7(10), 2131–2137contrast, G–C prepared samples allow imbibition. The imbibed IR absorption spectra (Fig. 8) of the imbibed samples show clear diVerences from that of the starting composition in terms sample has a Li2O/Al2O3 ratio of ca. 1, corresponding to the composition LiAl(OH)4·H2O.Fig. 7 shows the XRD patterns of the intensity and sharpness of the absorption bands for the LiOH imbibed sample and splitting of the bands and broaden- of the imbibed samples in relation to the starting composition (G–C prepared). The major diVerence is that some of the basal ing are observed for the LiNO3 imbibed sample. However, the essential pattern is the same, indicating that the basic structure reflections, (004) and (006), are split, whereas splitting is not observed for (002) reflection.The splitting of the XRD peaks remains undisturbed. Both samples show absorption due to hydrogen-bonded, bridged hydroxy groups similar to that of is due to the insertion of LiNO3, LiOH or [Li(OH)2]- ions into the intermediate layer. Splitting is not very pronounced the starting composition.The LiOH imbibed sample shows a very strong and sharp peak at 1380 cm-1 band and the for LiNO3 imbibed samples, whereas LiOH inserted samples show clear splittings. On further increase of LiOH, the phases shoulders at 1618 and 1491 cm-1 almost the same as in the starting material. A gain in intensity of the 1380 cm-1 band is stabilised are LiAl2(OH)7·2H2O and b-LiAlO2.Table 4 summarises the decomposition products of these compositions and due to the intercalation of Li(OH)2- in the intermediate layer. An increase in sharpness as well as the absorption intensity the polymorphic forms of LiAlO2 stabilised at diVerent temperatures. The decomposition product is c-LiAlO2 which is a beyond 1000 cm-1 (characteristics of the AlO6 group) indicate cation ordering in the octahedral sites.In contrast, insertion luminescent host for Fe3+ ions having an emission maximum at ca. 708 nm. of LiNO3 leads to intensification of bands at 1618, 1491 (clearly split) and 1380 cm-1 along with a new band at 1436 cm-1 corresponding to the asymmetric stretching of the NO3- group (n3) while the absorption band due to symmetric stretching (n1) is observed at 1112 cm-1; absorption bands beyond 1000 cm-1 are much broadened. 27Al MAS NMR spectra of diVerent polymorphs of LiAlO2 are shown in Fig. 9.The polymorphs show clear diVerences in chemical shift. The low temperature form (<400 °C) of a- LiAlO2 shows a chemical shift of d 9.5 corresponding to 27Al Fig. 7 XRD tracings of the hydrothermally imbibed samples: (a) asprepared LiAl2(OH)7·2H2O prepared through G–C conversion; (b) imbibed with LiNO3; (c) imbibed with LiOH, LDH/LiOH Fig. 8 IR spectra of LiAl2(OH)7·2H2O: (a) before imbibition, (b) imbibi- =1; (d) imbibed with LiOH, LDH/LiOH=0.5 [LDH= LiAl2(OH)7·2H2O]; #, b-LiAlO2; 6, LiAl2(OH)7·2H2O tion with LiOH and (c) imbibition with LiNO3 Table 4 Phases prepared through hydrothermal imbibition of LiAl2(OH)7 prepared through G–C conversion and the polymorphic forms of LiAlO2 composition of the charge as-prepared 400 °C 500 °C 1000 °C 1300 °C LDH51LiNO3 LiAl(OH)3NO3·H2O a-LiAlO2 b-LiALO2 c-LiAlO2 c-LiAlO2 LDH52LiNO3 LDHa+b-LiAlO2 b-LiAlO2+amorphous b-LiAlO2+LiAl5O8 c-LiAlO2+LiAl5O8 c-LiAlO2+LiAl5O8 background LDH51LiOH LiAl(OH)4·H2O a-LiAlO2 b-LiAlO2 c-LiAlO2 c-LiAlO2 LDH52LiOH LDH+b-LiAlO2 b-LiAlO2+amorphous b-LiAlO2+LiAl5O8 c-LiAlO2+LiAl5O8 c-LiAlO2+LiAl5O8 background aLDH=LiAl2(OH)7·2H2O.J. Mater. Chem., 1997, 7(10), 2131–2137 2135SAED photographs [Fig. 10(c)] show spotty rings indicating the polycrystalline nature of the sample. Hydrothermally prepared samples show lamellar crystals [Fig. 10(b)] with sizes ranging from 0.2 to 0.5 mm and SAED [Fig. 10(d)] shows a spotty single crystalline pattern. The results indicate that G–C produces aggregates of fine particles, whereas hydrothermal treatment yields lamellar individual single crystallites. LiAl2(OH)7·2H2O prepared either through G–C conversion or hydrothermally acts as a precursor for the preparation of LiAl5O8. Conversion is made possible owing to Li2O evaporation at higher temperatures, a fact not previously clarified in the literature.LiAl5O8 was also prepared from a reaction mixture having Li2O/Al2O3=0.22. LiAl5O8 exists in two forms viz. an ordered low-temperature form and a disordered hightemperature form. The ordered form is primitive cubic (P4332) while the disordered form has a true spinel structure (F41/d32m).The order–disorder transformation occurs at 1295±5 °C according to the literature.26 However, samples prepared through the present route did not undergo this type of transformation even at 1600 °C and the phase remained ordered. This indicates that the order–disorder transformation is dependent on the preparative route through which the phase is formed. When fast cooled (600 °C min-1), the samples were more disordered as shown by the decreased intensity of (100) and (110) reflections in comparison with (111) and (220) reflections.Monophasic LiAlO2 was prepared hydrothermally and also by the imbibition of LiAl2(OH)7·2H2O prepared through G–C conversion by LiNO3 or LiOH. The structure of the LDH consists of the positively charged brucite-like layers [LiAl2(OH)6]+ bridged by interlayer anions as well as water molecules.Depending upon the radius of the anion in the intermediate layer, the width of the interlayer varies. The existence of this intermediate layer is responsible for the unique properties of the material viz. anion exchange and intercalation, Fig. 9 Solid-state 27Al MAS NMR of (a) a-LiAlO2, (b) b-LiAlO2 and (c) c-LiAlO2 which makes preparation of LiAlO2 possible.LiAlO2 exists in diVerent polymorphs a-, b-, and c-LiAlO2, which diVer in the structure and cation coordination as established by solid-state NMR. The low-temperature form, a-LiAlO2, is stable below in an octahedral site. By contrast, b-LiAlO2 shows a chemical shift of d 78.6 and the c-form shows a minor splitting at d 70 400 °C and is prepared via hydrothermal imbibition.The bform is stable in the temperature range 500–650 °C prepared and 66. 27Al chemical shifts in the range d 50–85 are characteristic of 27Al in tetrahedral sites.25 The diVerence between the by heating the a-form at 500 °C or by hydrothermal synthesis at 240 °C. The compound c-LiAlO2 is prepared by heating b and c polymorphs is that the latter shows a clear splitting indicating site occupancy of 27Al in two diVerent types of the b-form at 1000 °C and is very stable in contrast to the material prepared via the sol–gel method11 which decomposes tetrahedra, one of which may diVer from the other in the degree of distortion.around 600 °C to LiAl5O8 and c-LiAlO2. The above results indicate that the two preparative techniques, G–C conversion and hydrothermal treatment, are Conclusions potential routes to synthesise LiAl5O8 and LiAlO2.The basic reactions in both the methods are the same; however the Gel to crystallite conversion and the hydrothermal method preparation conditions diVer. Continuous influx of aliovalent can be easily adapted for the preparation of important mateions into the gel network in the presence of hydrophilic solvents rials such as LiAl2(OH)7·2H2O, LiAl(OH)4·H2O, LiAl5O8 and such as ethanol will bring about rapid disintegration of the LiAlO2.The products formed are superior, in terms of homogel network. Hydrophilic solvents prevent H2O molecules re- geneity and phase content, to those prepared by conventional entering into the gel network and, thereby impede the reaction.solid-state methods. The samples prepared through G–C conversion always have very low crystallite size (0.04–1 mm), whereas hydrothermally prepared samples yield higher crystallite sizes (0.2–0.5 mm). References LDH [LiAl2(OH)7·2H2O] prepared through the two routes 1 N. T. Melamed, F. de S. Barros, P. J. Viccaro and J. O. Artman, diVer in their characteristics although the decomposition prod- Phys. Rev.B, 1972, 5, 3377. uct is the same both cases. The discrepancy may be due to the 2 J. G. Rabatin, J. Electrochem. Soc., 1978, 25, 920. particle size and cation distribution. The diVerences in the 3 J. Van Broekhoven, Illuminating Engineers Society Meeting, New particle morphology and crystallite size are further established York, July 1, 1973. 4 M. W. Parker and H. A. Borthwick, Plant Phys., 1949, 24, 345. by TEM studies. 5 B. Rasneuer, Fusion T echnol., 1985, 8, 1909. Fig. 10 shows TEM micrographs and corresponding selected 6 J. Jimnez-Beceril, P. Bosch and S. Bulbulian, J. Nucl.Mater., 1991, area diVraction patterns (SAED) of LiAl2(OH)7·2H2O pre- 185, 304. pared via G–C and hydrothermal routes. The morphology 7 A. J.Appleby and F. R. Foulkes in Fuel Cell Handbook, Van and crystallite size are diVerent in each case. The sample Nostrand Reinhold, New York, 1989, p. 297. prepared by G–C conversion shows [Fig. 10(a)] fibrous needle- 8 K. Kinoshita, J. W. Sim and J. P. Ackerman, Mater. Res. Bull., 1978, 13, 445. shaped aggregates, with particle sizes in the range 0.04–0.1 mm. 2136 J. Mater. Chem., 1997, 7(10), 2131–2137Fig. 10 TEM micrographs of LiAl2(OH)7·2H2O prepared through (a) G–C conversion, (b) hydrothermal route, the white patches are due to the defect in the carbon grid used to hold the sample; (c) and (d) are SAED patterns of the samples prepared through G–C conversion and the hydrothermal method, respectively 9 K. R. Poeppelmeier, C. K. Chiang and D. O. Kipp, Inorg. Chem., 20 T. R. N. Kutty, V. Jayaraman and G. Periaswami, Mater. Res. 1988, 27, 4523. Bull., 1996, 31, 1159. 10 K. R. Poeppelmeier and D. O. Kipp, Inorg. Chem., 1988, 27, 766. 21 T. R. N. Kutty, R. Jagannathan and R. P. Rao, Mater. Res. Bull., 11 J.-M. Jung and S.-B. Park, J.Mater. Sci. L ett., 1996, 15, 2012. 1990, 25, 1355. 12 S. Miyata, Clays Clay Miner., 1980, 28, 50. 22 T. R. N. Kutty, Mater. Res. Bull., 1990, 25, 343. 13 S. Miyata, Clays Clay Miner., 1983, 31, 305. 23 C. J. Serna, J. L. Rendon and J. Iglesias, Clays Clay Miner., 1982, 14 I. Sissoko, E. T. Eyagba, R. Sahai and P. Biloen, J. Solid State 30, 180. Chem., 1985, 60, 283. 24 M. P. Tarte, Acad. Des. Sci., 1962, 254, 2008. 15 J. E. Moneyron, A. De Roy and J. P. Besse, Sens. Actuators B, 25 D. Mu� ller, G. Gessner, H. J. Behrens and G. Scheler, Chem. Phys. 1991, 4, 189. L ett., 1981, 79, 59. 16 S. Miyata and A. Okaida, Clays Clay Miner., 1977, 4, 189. 26 K. Datta and R. Roy, J. Am. Ceram. Soc., 1963, 46, 388. 17 K. R. Poppelmeir and S. J. Hwu, Inorg. Chem., 1987, 26, 3297. 18 T. R. N. Kutty and P. Padmini, Mater. Chem. Phys., 1995, 39, 200. 19 T. R. N. Kutty and M. Nayak,Mater. Res. Bull., 1995, 30, 325. Paper 7/02065A; Received 25thMarch, 1997 J. Mater. Chem., 1997, 7(10), 2131–2137 21

ISSN:0959-9428    

DOI:10.1039/a702065a

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Non-stoichiometry, structural defects and properties of LaMnO3+d with high d values (0.11d0.29) J. A. Alonso,*a M. J. Martý�nez-Lope,a M. T. Casais,a J. L. MacManus-Driscoll,b P. S. I. P. N. de Silva,c L. F. Cohenc and M. T. Ferna�ndez-Dý�azd aInstituto de Ciencia deMateriales de Madrid, C.S.I.C. Cantoblanco, E-28049 Madrid, Spain bDepartment of Materials, Imperial College, L ondon, UK SW7 2BZ cBlackett L aboratory, Imperial College, L ondon, UK SW7 2BZ cInstitut L aue-L angevin, B.P. 156, F-38042 Grenoble Cedex 9, France Strongly oxygenated LaMnO3+d perovskites, with nominal MnIV contents up to 58%, have been prepared by thermal decomposition of metal citrates followed by annealings either in air or under high oxygen pressure (200 bar). A high-resolution neutron powder diVraction study of four representative samples with 0.11d0.29 reveals the presence of both La and Mn vacancies.Contrary to previous studies, it is found that there are a substantially higher proportion of Mn vacancies, depending rather sensitively on the oxidation conditions. The oxidation state for Mn calculated for the refined stoichiometry La1-xMn1-yO3 is in good agreement with the d values previously determined by thermal analysis.Further to this, it is also found that as d increases the MnMO bond lengths shorten, the MnMOMMn angles progressively increase and the perovskite structure becomes more regular, which is consistent with the incorporation of MnIV cations. The presence of Mn vacancies (as much as 13% in samples prepared under high oxygen pressure) perturbs the conduction paths for the transport of holes across MnMOMMn, weakening the double-exchange interaction.This structural disorder explains the observed decrease of the ferromagnetic Curie temperature (TC) as d increases. The structure and properties of the perovskite LaMnO3 and temperature) from orthorhombic to rhombohedral.10,11 In fact, the transition temperature from the orthorhombic (low-tem- related materials have been extensively studied in the past.1,2 More recently, the observation of large negative magnetoresist- perature form) to the rhombohedral (high-temperature form) was found to change10 from 600 °C for LaMnO3.00 to ca.ance eVects in mixed-valence manganites3 based on LaMnO3 has renewed the interest in the study of these systems.In -90 °C or LaMnO3.15. The crystal structure of the oxygen excess rhombohedral particular, the strongly correlated magnetic and electrical properties are being widely studied. The magnetoresistance is phases has been investigated by neutron diVraction.12,13 Tofield and Scott12 and, later on, Van Roosmalen et al.13,14 concluded associated with the insulator-to-metal transition which occurs in the vicinity of the Curie point, for ferromagnetic composi- that instead of incorporating oxygen interstitials, as directly suggested by the formula LaMnO3+d, this perovskite is defec- tions such as La0.7Ca0.3MnO3 which, thus, simultaneously exhibit metallic conductivity and ferromagnetism at low tem- tive in both La and Mn positions: for instance, a sample with a formal composition of LaMnO3.158 was found to exhibit an peratures.The presence of divalent alkaline earth cations on the La sites of the perovskite induces a MnIII–MnIV mixed- actual crystallographic formula of La0.95Mn0.95O3. Recently15 we have prepared LaMnO3+d phases with a wide valence state with the formation of holes which undergo fast hopping between the two oxidation states.The temperature of range of d values, 0.11d0.31, by annealing finely divided precursors at moderate temperatures, either in air or under both transitions depends on the MnIV content (for instance, in the La1-xCaxMnO3 system TC varies between 182 and 278 K high oxygen pressure. The samples were characterized by Xray diVraction and thermal analysis, allowing us to determine for MnIV contents between 14 and 38%, respectively.4) The parent perovskite LaMnO3.00 (where only MnIII is their oxygen contents (d) very reliably.This paper complements those previous studies giving details on the actual way of present) is antiferromagnetic below TN=140 K and shows a semiconducting behaviour over the whole temperature range.5 incorporation of the non-stoichiometry into the crystal structures, refined from high-resolution neutron powder diVraction.LaMnO3 crystallizes in the orthorhombic GdFeO3 structural type,6,7 with a=5.5392(6), b=5.6991(7), c=7.7175(9) A ° , in the Our main result is that the probability of creating equal numbers of La and Mn vacancies depends very strongly on Pbnm setting. A strong Jahn–Teller distortion of the oxide octahedra around the d4 MnIII cations has been identified from the preparative oxidation conditions.The influence of the observed defective structure on the magnetic properties is also a neutron diVraction study.7 A MnIII–MnIV mixed-valence state can also be induced in discussed. phases of nominal stoichiometry LaMnO3+d (with MnIV content of 2d per formula unit). These phases also show ferromag- Experimental netic behaviour for a suYciently high content of MnIV.For Four selected LaMnO3+d samples were prepared in polycrys- instance, LaMnO3.08 (with 16% of MnIV) was reported to be talline form by a citrate technique. Stoichiometric amounts of ferromagnetic with a Curie temperature TC=125 K).5 In fact, analytical grade La(NO3)3 6H2O and Mn(NO3)2 4H2O were GMR behaviour has been reported8,9 in samples with moderate dissolved in citric acid.The metal solution was slowly evapor- d values, showing overall compositions LaMnO3.17 and ated and the resultant resin was subsequently decomposed at LaMnO3.12. 600 °C for 12 h. A second treatment at 800 °C for 6 h enabled For LaMnO3+d the oxidation of some MnIII to MnIV ions the total elimination of the nitrates and organic materials.The reduces the driving force for Jahn–Teller distortion. Above d= 0.105 (i.e. 21% MnIV) there is a change in symmetry (at room finely divided precursor powders were annealed either in air J. Mater. Chem., 1997, 7(10), 2139–2144 2139(samples 3 and 4) or under 200 bar of oxygen pressure (samples 1 and 2), at temperatures ranging from 800 to 1100 °C.Then, the samples were slowly cooled to room temperature. The final materials were characterized by X-ray powder diVraction (XRD) with Cu-Ka radiation. The determination of d was performed by thermogravimetric (TG) analysis in a reducing H2–N2 flow, as indicated elsewhere.15 Neutron powder diVraction (NPD) data for LaMnO3+d were collected at room temperature in the high-resolution D2B diVractometer at the ILL, Grenoble, with a wavelength of 1.594 A ° , selected from a Ge monochromator.About 5 g of each sample were contained in a cylindrical vanadium can. The time consumed in each data collection was about 3 h. The thermal evolution of samples 2, 3 and 4 was studied in the multidetector DN5 diVractometer at the Siloe� reactor of the Centre d’Etudes Nucle�aires, Grenoble, with a wavelength of 2.488 A° in the temperature range 2–250 K.The Rietveld method16 was used to refine the crystal structures, using the FULLPROF program.17 The line shape of the diVraction peaks was generated by a pseudo-Voigt function, and the background refined to a fifth-degree polynomial. The coherent scattering lengths for La, Mn andO were, respectively, 8.24, -3.73 and 5.803 fm.In the final run the following parameters were refined: six background coeYcients, zeropoint, half-width, pseudo-Voigt and asymmetry parameters for Fig. 1 XRD patterns of LaMnO3+d , for samples 1 (above) to 4 (below) the peak shape; scale factors, positional, thermal isotropic factors and unit-cell parameters. The occupancy factors of La Table 2 Atomic parameters for rhombohedral LaMnO3+d (samples and Mn were also allowed to vary in the last stages of the 1–3) after the Rietveld refinement of NPD data at 295 K refinement.In the final cycle the shifts in the atomic parameters were zero up to the fourth decimal place. sample no. 1 2 3 Magnetization data were collected in an Oxford Instrument 3001 vibrating-sample magnetometer, in a remnant field of ca.atom La B/A 35(6) 1.38(6) 0.97(6) 30 Oe, between 10 and 300 K. focc 0.95(1) 0.97(1) 0.969(9) Mn B/A ° 2 0.3(1) 0.2(1) 0.26(10) Results focc 0.89(2) 0.87(2) 0.93(2) O x 0.4561(2) 0.4545(2) 0.4489(2) Structural features of LaMnO3+d B/A ° 2 1.48(3) 1.41(3) 1.22(4) focc 3.0 3.0 3.0 The XRD patterns of samples 1–4 were characteristic of aa/A ° 5.4905(1) 5.4951(1) 5.5222(1) monophase perovskites with rhombohedral (samples 1–3) or ba/A° 13.3077(1) 13.3030(1) 13.3317(1) orthorhombic (sample 4) symmetry, as shown in Fig. 1. V /A ° 3 347.42(1) 347.88(1) 352.08(1) discrepancy factors Observe that there is an increase of the rhombohedral splitting Rp (%) 4.13 4.12 4.12 of the peaks under less oxidizing conditions (i.e. high tempera- Rwp (%) 5.20 5.22 5.26 ture, air).Table 1 summarizes the preparation conditions, d Rexp (%) 3.96 4.19 4.26 values obtained by TG and unit-cell parameters determined x2 1.73 1.55 1.52 by XRD. For samples 1 to 3 both hexagonal (a, c) and RI (%) 7.34 7.73 6.89 rhombohedral (ar, ar) descriptions of the unit cell are given. aThe unit-cell parameters diVer slightly from those determined from The room-temperature NPD patterns of samples 1–3 were XRD (Table 1) owing to a small inaccuracy of the neutron wavelength.refined in the space group R39c, hexagonal description (Z=6), Note: space group R39c, Z=6. La atoms are at 6a positions, (0,0,1/4); taking as starting model that of the perovskite LaNiO3, which Mn at 6b, (0,0,0); O at 18e, (x,0,1/4). exhibits the same symmetry.18 La atoms are at 6a, (0,0,1/4) positions; Mn at 6b, (0,0,0); and O at 18e, (x,0,1/4).The final atomic parameters after the refinements are shown in Table 2. For sample 4, the profile refinement was performed in the orthorhombic space group Pbnm (Z=4), according to the The good matching of the fits is illustrated in Fig. 2(a) for sample 3. No extra lines or additional splitting of the peaks GdFeO3 structural model.La is placed at 4c, (x,y,1/4); Mn at 4b, (1/2,0,0); O(1) at 4c and O(2) at 8d, (x,y,z). Subsequently, was observed in any case. The refinement of the occupancy factor of La and Mn led to values significantly lower than 1. a detailed analysis of the profile made it necessary to consider Table 1 Preparation conditions, unit-cell parameters and volume per formula of LaMnO3+d , samples 1–4, determined from XRD data.Samples 1–3 are rhombohedral, space group R39c, Z=6; the main phase in sample 4 is orthorhombic, space group Pbnm, Z=4. sample no. 1 2 3 4 prep. conditions 800 (200 bar O2) 1000 °C (200 bar O2) 1000 °C (air) 1100 °C (air) d(TG) 0.29 0.26 0.15 0.11 a/A ° 5.4898(4) 5.4951(3) 5.5239(2) 5.54002(2) b/A° 5.4898(4) 5.4951(3) 5.5239(2) 5.4963(2) c/A ° 13.311(1) 13.3061(8) 13.3349(7) 7.7876(4) ar/A ° 5.453(1) 5.453(1) 5.470(1) — ar/° 60.450(5) 60.509(5) 60.64(5) — Vf/A° 3 57.90(1) 57.99(1) 58.73(1) 59.28(1) 2140 J.Mater. Chem., 1997, 7(10), 2139–2144Table 3 Atomic parameters for sample 4 [annealed at 1100 °C(air)] after the refinement of NPD data at 295 K. The sample consists of a mixture of a main orthorhombic phase (Pbnm, 64.4%) and a minor rhombohedral phase (R39c, 35.6%)a atom Pbnm R39c La x 0.9970(5) y 0.0171(4) B/A ° 2 0.83(5) 0.84(8) focc 0.978(2) 0.978(2) Mn B/A ° 2 0.41(7) 0.31(12) focc 0.946(4) 0.946(4) O(1) x 0.0645(5) 0.4466(4) y 0.4955(7) B/A° 2 1.09(6) 1.06(6) focc 1.0 3.0 O(2) x 0.7398(4) y 0.2705(4) z 0.0337(2) B/A ° 2 1.23(5) focc 2.0 a/A ° 5.5388(2) 5.5348(1) b/A° 5.4958(2) 5.5348(1) c/A ° 7.7843(2) 13.3438(3) RI 6.39 4.81 aDiscrepancy factors: Rp=6.32%, Rwp=4.18%, Rexp=2.38%, x2= 2.18.Note: for the orthorhomic phase: space group Pbnm, Z=4; La at 4c, (x,y,1/4); Mn at 4b, (1/2,0,0); O(1) at 4c and O(2) at 8d, (x,y,z). For the rhombohedral phase: space group R39c, Z=6; La at 6a (0,0,1/4); Mn at 6b (0,0,0); O at 18e, (x,0,1/4).additionally a minor rhombohedral LaMnO3+d phase (with space group R3 : c) to correctly fit the spectra. The strong overlapping between the reflections of both polymorphs had prevented the detection of the minor rhombohedral phase Fig. 2 Room temperature observed (crosses), calculated (solid line) prior to the Rietveld analysis of the neutron pattern. From the and diVerence (at the bottom) NPD profiles for (a) rhombohedral refined scale factors a composition of 63.6% (orthorhombic)+ LaMnO3.26, sample 2; and (b) orthorhombic LaMnO3.11, sample 4: 36.4% (rhombohedral) was determined for the mixture.The the second series of tick marks indicate the reflections of a minor La and Mn occupancy factors were constrained for both rhombohedral phase.For the sake of clarity only half of the experimenphases. Table 3 includes the results of the refinement. Fig. 2(b) tal points are represented. shows the observed and calculated NPD profiles for sample 4. Bond distances and angles are listed in Table 4. Observe that in the orthorhombic structure the MnO6 octahedra do not show any appreciable Jahn–Teller distortion, as expected for lengths progressively increase from 1.9494(6) A ° , for the perovskite prepared in the most oxidizing conditions, sample 1, the relatively high proportion of MnIV in the crystal (for d= 0.11, [Mn4+]=22%), which prevents the cooperative distor- to an average value of 1.976 A ° for the orthorhombic phase present in sample 4.tion of the octahedra observed7 in stoichiometric LaMnO3.00.As shown in Table 1, there is a net increase of the cell Some relevant parameters determined from the refinements are listed in Table 5. The crystallographic formulae of the volume per formula (Vf) from 57.90 to 59.28 A ° 3 as d decreases from 0.29 (sample 1) to 0.11 (sample 4), which is consistent LaMnO3+d phases, according to the refined occupancy factors for La and Mn, can be written as La1-xMn1-yO3, as indicated with the decreasing amount of MnIV cations in the structure.This is directly related to the regular variation of the MnMO in Table 5. Under more oxidizing conditions (from sample 4 to sample 1) the amount of Mn vacancies increases more distances, quoted in Table 4. Observe that MnMO bond Table 4 Selected bond distances (A ° ) and angles (degrees) for LaMnO3+d (samples 1–4) sample no 1 2 3 4 space group R39c R39c R39c R39c Pbnm MnMO(1) 1.9494(6) ×6 1.9514(7) ×6 1.9634(6) ×6 1.969(1) ×6 1.9788(6) ×2 MO(2) — — — — 1.975(2) ×2 MO(2) — — — — 1.975(2) ×2 LaMO(1) 2.504(1) ×3 2.497(1) ×3 2.479(1) ×3 2.472(1) ×3 2.891(5) MO(1) 2.7367(5) ×6 2.7377(5) ×6 2.7492(4) ×6 2.7543(5) ×6 2.656(5) MO(1) — — — — 2.432(4) MO(2) — — — — 2.639(3) ×2 MO(2) — — — — 2.500(3) ×2 MO(2) — — — — 2.811(3) ×2 LaMO 2.6591(2) 2.6576(2) 2.6591(2) 2.6602(3) 2.653(1) O(1)MMnMO(1) 90.84(6) 90.91(6) 91.13(5) 91.24(1) 180.0 O(2)MMnMO(2) — — — — 91.45(17) MnMO(1)MMn 165.782(7) 165.264(7) 163.477(7) 162.72(1) 159.15(5) MnMO(2)MMn — — — — 162.12(9) J.Mater. Chem., 1997, 7(10), 2139–2144 2141Table 5 Relevant parameters obtained from the structural data: crystallographic formulae, valence of Mn determined from the occupancy factors ( focc), thermogravimetric data (TG) and the bond valence theory ( b.v.); and tolerance factors, t Mn valence sample nominal actual no.composition stoichiometry focc TG b.v.a tb 1 LaMnO3.29 La0.95(1)Mn0.89(1)O3 3.54 3.58 3.65 0.965 2 LaMnO3.26 La0.97(1)Mn0.87(2)O3 3.55 3.52 3.61 0.963 3 LaMnO3.15 La0.97(1)Mn0.93(1)O3 3.33 3.30 3.50 0.958 4 LaMnO3.11 La0.978(2)Mn0.946(4)O3 3.24 3.22 3.26 0.949 aThe valence is the sum of the individual bond valences (si) for MnMO bonds within the MnO6 octahedra.Bond valences are calculated as si=exp[(r0-ri)/B]; B=0.37, r0=1.760 for the MnIIIMO2- pair, from ref. 22. Individual MnMO distances (ri) are taken from Table 3.bTolerance factors are calculated as t=LaMO /Ó2MnMO . Fig. 4 Magnetization vs. temperature plot for LaMnO3+d (H#30 Oe). Annealing temperature and atmosphere: (a) 1000 K, 200 bar O2; (b) 800 K, 200 bar O2; (c) 1000 K, air; (d) 1100 K air. Fig. 3 Variation of the La and Mn contents with d, determined from the neutron diVraction refinements. The full lines are guides for the eye.quickly than the La vacancies, as shown in Fig. 3. The very oxidizing conditions (under 200 bar O2) are more easily able to create Mn vacancies than La vacancies. The final valences for Mn calculated from the metal vacancy concentration agree quite well with those determined by thermal analysis, as shown in Table 5. The oxidative non-stoichiometry determined for these products is thus able to explain the macroscopic behav- Fig. 5 Thermal evolution of the NPD patterns for LaMnO3.11, sample iour, represented by the d values. 4. Below TC=165 K the magnetic contribution to the peaks of nuclear Comparing the observed MnMO distances to the ionic radii origin corresponds to the three-dimensional ferromagnetic ordering. sums19 for a random occupancy of MnIII/MnIV in the manganese positions, the observed values are systematically lower.In Fig. 6 the Curie temperatures are plotted vs. d, determined For instance, for sample 1, [MnMO]calc=1.983 A ° (observed from both magnetization data and low temperature NPD 1.949 A ° , Table 4); for sample 4, orthorhombic phase, experiments. Values of TC obtained by NPD are systematically [MnMO]calc=2.017 A ° (observed 1.976 A ° on average).This fact lower than those determined from the magnetization curves, can be understood as an eVect of the additional contraction as a consequence of a gradual transition from short- to longof the lattice due to the presence of metal vacancies. For the range magnetic order, the latter taking place at temperatures same reason, the calculation of the valence of Mn cations about 5–15 K lower.within the MnO6 coordination octahedra by means of Brown’s bond valence model20,21 systematically leads to values higher than observed (Table 5). Discussion In their pioneering work, Tofield and Scott12 showed that Comparison of low-temperature NPD and magnetization the defect chemistry of LaMnO3+d is better described with measurements randomly distributed La and Mn vacancies.Later on, Van Roosmalen et al.13,14 concluded from a neutron diVraction The magnetization and transport measurements have been studied for a more complete set of samples and will be experiment combined with density measurements that La and Mn vacancies are present in equal amounts in the solid, in described in detail elsewhere.22 Fig. 4 shows the magnetization vs.temperature curves for the samples discussed in the present such a way that the crystallographic formula should be written as La1-xMn1-xO3, with x=d/(3+d). Our neutron diVraction study. They show a ferromagnetic ordering of the Mn spins below a critical temperature. The ferromagnetic contribution study on samples with high d values, prepared at relatively low temperatures from finely divided precursors, confirms the to the neutron scattering is presented in Fig. 5, including the thermal evolution of the NPD patterns for sample 4. In all presence of significant amounts of La and Mn vacancies. Final Fourier synthesis did not yield any identifiable peaks which cases the magnetic structures can be described in terms of a single propagation vector, k=(0,0,0).23 could suggest the presence of additional oxygen atoms in 2142 J.Mater. Chem., 1997, 7(10), 2139–2144observation by To� pfer et al.,11 who described an orthorhombic perovskite-type structure for LaMnO3+d at d0.10, while for d>0.10 they identified a rhombohedral cell. Sample 4, with d=0.11, is at the boundary of the phase transformation, consistent with the observed phase separation.It is worth mentioning that, in spite of the change in crystal symmetry, the determined La and Mn vacancy concentration follows the trend shown in Fig. 3 for the three pure rhombohedral compositions. Some perovskites with oxidative non-stoichiometry have been reported to show ferromagnetic behaviour with Curie points in a comparable temperature range to those observed in the present work for the samples with lower d values.For instance, Ranno et al.24 described a TC of 125 K for LaMnO3.15, or (LaMn)0.95MnO3, whereas To� pfer et al.11 obtained TC= 171 K for LaMnO3.14; in both cases the composition is close to that of sample 3 (also d=0.15), with a refined stoichiometry La0.97Mn0.93O3, and TC (from magnetization data) of 155 K. Fig. 6 Variation of the Curie temperature with d. Full symbols from Fig. 6 shows that TC decreases as d increases, i.e. as the MnIV magnetization data, for samples prepared, from left to right, at 1100 K in air, 1000 K in air and 800 K in 200 bar O2. Open symbols from concentration increases. A priori, an opposite trend would be NPD data, for samples prepared, from left to right, at 1100 K in air, expected, since the introduction of MnIV cations (t2g3) into the 1000 K in air, 1000 K in 200 bar O2 (from ref. 23). MnIIIO6 array (high-spin MnIII: t2g3e2g1) of LaMnO3 by hole doping (or as a result of deviations from the stoichiometric composition) creates empty eg orbitals which favour the ferro- interstitial positions, even in the more oxygenated samples. Our previous suggestion15 that the structure of the samples magnetic double exchange.This would lead to increasing Curie temperatures with increasing d, as demonstrated in Ca-doped prepared under oxygen pressure would probably involve oxygen interstitials, is not confirmed. manganites La1-xCaxMnO3 up to the ‘optimum doping level’, of 33% MnIV. In addition, for a constant doping level, TC has The present NPD study shows definitively that the concentrations of both kinds of metal vacancies are not equal. been shown to increase as the tolerance factor of the perovskites is closer to 1, i.e.when the structure becomes more regular, Moreover, the concentration of Mn vacancies increases at a higher rate when the annealing conditions of the samples with MnMOMMn angles closer to 180° favouring the double-exchange interactions, as shown25 for the series become more oxidizing (i.e.lower temperatures and higher oxygen pressure), as shown in Fig. 3. The La and Mn vacancy (La0.7-xYx)Ca0.33MnO3. In the present case the observed trend in the variation of TC concentration only become closer for low d values, i.e. for samples annealed in air at higher temperatures.Note that the vs. d can be interpreted as a consequence of the presence of significant amounts of Mn vacancies in the structure. The final samples described by Tofield and Scott12 and Van Roosmalen et al.13,14 had undergone thermal treatments at even higher TC value and the strength of the ferromagnetic behaviour result from an interplay between defect density, MnIV content and temperatures of 1200 or 1300 °C, respectively. The soft-chemistry synthesis procedure described in the MnMOMMn bond angles. Even though the tolerance factors (Table 5) of the perovskite structures of samples 1–4 increase present work seems to favour the formation of highly defective perovskites, especially in the Mn sublattice.We show that the with d (which is related to the progressive shortening of the MnMO distances), the relatively high proportion of Mn vac- non-stoichiometry of LaMnO3 cannot be simply denoted with a single parameter, d or x, but requires the specification of two ancies (up to 13% in sample 2) perturbs the connecting paths for the holes to transport across MnMOMMn, and leads to parameters, x and y in La1-xMn1-yO3.The relative values of x and y depends dramatically on the preparation procedures weaker ferromagnetic coupling, via the double-exchange interaction.The presence of random vacancies may also have a of the samples, in such a way that ceramic synthesis at higher temperatures favours equal values of x and y. strong localising eVect on the available charge carriers, which reduces their mobility.24 Our results suggest that at moderate temperatures, in the range 800–1100 °C, the higher mobility of Mn cations allows them to predominantly migrate across the perovskite structure, Conclusions giving rise to a dtive Mn sublattice and consequently a higher oxidation state of the remaining Mn cations. The higher A room-temperature, high-resolution neutron diVraction study of four selected LaMnO3+d samples with high d values mobility of Mn is probably related to the smaller size of Mn vs.La cations. At higher temperatures the mobility of La (0.11d0.29), prepared by low-temperature treatments of citrate precursors, shows highly defective perovskite structures cations also becomes significant, giving rise to final products in which the amounts of both metal vacancies are closer.containing both La and Mn vacancies. The Mn vacancies are present in a substantially higher proportion with respect to In all cases the samples were prepared from stoichiometric mixtures of nitrates, with La5Mn=151 ratio. The observed the La vacancies in all the samples, although the diVerence in the number of both kinds of defects decreases under less deviation of the 151 stoichiometry in the final perovskite phases implies the presence of minor Mn-rich phases in the oxidizing conditions.(i.e. as d decreases). There is a good agreement between the Mn valences determined from thermal products. No diVraction peaks other than those corresponding to the perovskite oxides could be detected in either the XRD analysis and those estimated from the defect stoichiometry La1-xMn1-yO3.The cell volume, MnMO distances, or ND patterns. The segregated Mn-rich minor impurities must be present in a poorly crystallized or amorphous form, MnMOMMn angles and tolerance factors, t, of the perovskites vary regularly with d, i.e. with the oxidizing power of the probably as a thin layer covering the surface of the perovskite crystallites.materials. The high concentration of defects, mostly in the Mn sublattice, explains the anomalies observed in the magnetic The perovskites prepared at temperatures up to 1000 °C show rhombohedral symmetry, whereas that prepared at properties: the samples are ferromagnetic below TC which decreases when d increases, even though t increases. We 1100 °C (sample 4) could be identified as a mixture with a main orthorhombic phase.This fact is consistent with the conclude that the preparation by soft-chemistry methods of J. Mater. Chem., 1997, 7(10), 2139–2144 2143Ramakrisnan, R. Mahesh, N. Raganvittal and C. N. R. Rao, Phys. LaMnO3+d materials with high d contents starting from Rev. B, 1996, 53, 3348. La5Mn=151 mixtures leads to highly Mn defective phases 9 C.N. R. Rao and A. K. Cheetham, Science, 1996, 272, 369. with anomalous magnetic properties. The synthesis of undoped 10 A. Wold and R. J. Arnott, J. Phys. Chem. Solids, 1959, 9, 176. La–Mn–O materials with a perfect Mn sublattice would 11 J. To� pfer, J. P. Doumerc and J. C. Grenier, J. Mater. Chem., 1996, require starting mixtures with La5Mn<1 ratios, in such a way 6, 1511. 12 B. C. Tofield and W. R. Scott, J. Solid State Chem., 1974, 10, 183. that the selective creation of La vacancies allows Mn cations 13 J. A. M. Van Roosmalen, E. H. P. Cordfunke, R. B. Helmholdt and to reach the valence corresponding to the oxidation potential H. W. Zandbergen, J. Solid State Chem., 1994, 110, 100. given by the preparative conditions. 14 J. A. M. Van Roosmalen and E.H. P. Cordfunke, J. Solid State Chem., 1994, 110, 106. 15 J. A. Alonso, M. J. Martý�nez-Lope and M. T. Casais, Eur. J. Solid The authors acknowledge the financial support of the Spanish State Inorg. Chem., 1996, 33, 331. DGICyT to the project PB94–0046, and the Engineering and 16 H. M. Rietveld, J. Appl. Crystallogr., 1969, 2, 65. 17 J. Rodrý�guez-Carvajal, Physica B, 1993, 192, 55.Physical Science Research Council of the United Kingdom. 18 J. L. Garcý�a-Mun� oz, J. Rodrý�guez-Carvajal, P. Lacorre and J. B. Torrance, Phys. Rev. B, 1992, 46, 4414. 19 R. D. Shannon, Acta Crystallogr. Sect. A, 1976, 32, 751. References 20 I. D. Brown, in Structure and Bonding in Crystals, ed. M. O’Keefe and A. Navrotsky, New York, 1981, vol. 2, pp. 1–30. 1 G. H. Jonker and J. H. 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ISSN:0959-9428    

DOI:10.1039/a704088a

出版商:RSC    

年代:1997

数据来源: RSC