摘要:

J O U R N A L O F C H E M I S T R Y Materials Synthesis and structural characterization of a novel tin(II ) oxyphosphate, [NH4+]2[Sn3O(PO4)2]2-·H2O, containing one-dimensional chains constructed from tin phosphate cages Srinivasan Natarajan* Chemistry and Physics of Materials Unit, Jawaharlal Nehru Centre for Advanced Scientific Research, Jakkur P.O., Bangalore 560 064, India. E-mail: raj@jncasr.ac.in Received 13th August 1998, Accepted 30th September 1998 The hydrothermal synthesis and single crystal structure of a novel tin(II) oxyphosphate, [NH4+]2[Sn3O(PO4)2]2-·H2O, made from the networking of distorted square-pyramidal SnO4 and tetrahedral PO4 moieties, is presented.Crystal data: a=7.240(1), b=19.552(3), c=8.438(1) A° ; V=1194.5(3) A° 3; space group Cmc21 (no. 36) and Z=2. The structure, dominated by the presence of a large number of three-coordinated oxygen atoms and SnKOKSn linkages, creating an oxy-phosphate unit, consists of capped three-membered rings (cages) connected to each other via phosphate groups forming infinite one-dimensional chains. These chains are related to each other via multipoint hydrogen bonding involving the protonated ammonium and water molecules.atoms. The structure of this material consists of capped three- Introduction membered rings (cages) connected to each other via phosphate The design and synthesis of new open-framework materials groups forming infinite chains. These chains are held together having micro- and meso-porosity is a challenge to the synthetic via multipoint hydrogen bonding involving the protonated chemist, as such materials are being used in the areas of ammonium and water molecules.catalysis and separation processes.1 The synthesis under hydro/solvo-thermal conditions provides a facile route for Experimental making materials with new open architectures as well as complex organic–inorganic composites. This methodology Synthesis leads to the development of new classes of materials that can The title compound, synthesized hydrothermally employing exploit the ability of polar organic molecules to direct the guanidium carbonate as the structure directing agent, is crystallization of inorganic frameworks through multipoint described below.In a typical experiment, 2.0671 g of tin hydrogen bonding.2 The hydrogen bonded interactions oxalate (Aldrich) was dispersed in 10 ml of water.To this between the inorganic and organic moieties of a framework mixture 1.153 g of phosphoric acid (Aldrich) was added drop- structure are all the more important when dealing with low wise and the mixture stirred vigorously. Then 1.262 g of dimensional materials as is becoming more apparent by the guanidium carbonate was added to the above very slowly large number of available literature dealing with such materials.under continuous stirring. The final mixture was transferred The chemistry of bivalent tin and its related compounds, and sealed in a PTFE-lined stainless steel autoclave (Parr, especially the phosphates and oxalates, synthesized in the USA), and heated at 175 °C for 3 days under autogeneous presence of organic structure directing agents (amines) pressure.The final composition of the mixture was continues to yield unexpected results with the isolation of 1 SnC2O451H3PO450.7 guadinium carbonate555H2O. The materials having one- [Sn2(PO4) (C2O4)0.5]3 (chain), tworesulting product, which contained a few single crystals along {[Sn2(PO4)2]2-[C2N2H10]2+·H2O}4 ( layer) and threewith some white powder, was filtered oV and washed thor- {[Sn4P3O12]-0.5[NH3(CH2)2NH3]2+, [Sn4 P3O12]-0.5- oughly with de-ionized distilled water (yield ca. 50%). The [NH3(CH2)4NH3]2+}5,6 dimensionally extended networks. powder X-ray diVraction pattern of both the crushed single The basic building unit, four-membered rings constructed by crystals as well as the white powder, taken independently, two Sn atoms and two P atoms (Sn2P2O4 unit), present in all were found to be identical and indicated that the product was these materials, has also been isolated.7 All these compounds homogeneous and a new material; the pattern is entirely contain either trigonal pyramidal SnO3 or distorted square consistent with the structure determined by single crystal X-ray pyramidal SnO4 units, vertex linked with tetrahedral PO4 units diVraction.Thermogravimetric analysis (TGA) was carried to form the network. These solids, based on Sn(II ), in addition out in nitrogen atmosphere from room temperature up to to having novel architectures also provide a basis for evaluating 600 °C. the influence of the Sn(II) lone pair of electrons on the structural features as well as the importance of multipoint hydrogen Structure determination bonding in the stability of such phases.The synthesis and structure of a novel tin(II ) oxyphosphate A suitable colorless single crystal was carefully selected under material having a 10-membered star shaped channel8 has a polarizing microscope and glued to the tip of a glass fiber been reported recently.In the present work, the synthesis using Superglue (cyanoacrylate). Crystal structure determiand structure of another novel tin(II) oxyphosphate, nation by X-ray diVraction was performed at room-tempera- [NH4+]2[Sn3O(PO4)2]2-·H2O, made by the networking of ture on a Siemens Smart-CCD diVractometer equipped with distorted square-pyramidal SnO4 and tetrahedral PO4 units is a normal focus, 2.4 kW sealed tube X-ray source (Mo-Ka presented.This is the first report of a Sn(II ) phosphate material radiation, l=0.71073 A° ) operating at 50 kV and 40 mA. A formed entirely by four-coordinate Sn(II) atoms; all the pre- hemisphere of intensity data were collected in 1321 frames viously reported Sn(II ) phosphate3–8 and phosphonate9,10 with v scans (width of 0.30° and exposure time of 40 s per frame).The final unit-cell constants were determined by a materials contain both three- and four-coordinated Sn(II) J. Mater. Chem., 1998, 8, 2757–2760 2757Table 1 Summary of crystal data, intensity measurements and structure refinement parameters for [NH4+]2[Sn3O(PO4)2]2-·H2O Empirical formula Sn3P2O10N2H10 Crystal system Orthorhombic Space group Cmc21 (no. 36) Crystal size/mm 0.04×0.04×0.125 a/A° 7.240(1) b/A° 19.552(3) c/A° 8.438(1) V/A° 3 1194.5(3) Z 2 Formula mass 616.1(1) Dc/g cm-3 1.713 l(Mo-Ka)/A° 0.71073 m/mm-1 3.27 Temperature of 298 measurement/K h range/° 2.08–23.28° Total data collected 4995 Index ranges -7h7, -21k20, -6l9 Unique data 2299 Observed data [s>2s(I )] 749 Rint 4.68 Refinement method Full-matrix least-squares on |F 2| R indices [s>2s(I )] R=3.67; Rw=7.83 R (all data) R=4.79; Rw=8.47 Goodness of fit/S 1.13 No.of variables 103 Largest diVerence map peak 0.699, -0.975 and hole/e A° -3 least-squares fit of 1379 reflections in the range 42h46.5° and are presented in Table 1. A total of 4995 reflections were collected in the range -7h7, -21k20, -6l9 and these were merged to give 2299 unique reflections (Rint.= 0.0468) of which 749 were considered to be observed [I>2s(I )].The pertinent experimental conditions for the structure determination are listed in Table 1. The structure was solved by direct methods using SHELXS-8611 and diVerence Fourier syntheses. Hydrogen atoms of the water and ammonia molecules were found in the diVerence Fourier map and held in the riding mode.Refinement of the Flack parameter to ca. 1.0 indicated that the incorrect absolute structure had been established in the initial model. The polarity of the model was reversed by changing the sign of the polar-axis z coordinate for all the atoms and re-refining the model, which resulted in a final value for the Flack parameter of-0.05(8) and residuals of R=0.0367 and Rw=0.0783. Re-refinement of this transformed coordinate set with the Flack parameter set to 1.00 led to residuals of R=0.0382 and Rw=0.0902.The last cycles of refinement included atomic positions and anisotropic thermal parameters for all non-hydrogen atoms and isotropic thermal parameters for all the hydrogen atoms. Full-matrixleast- squares structure refinement against |F2| was carried out using the SHELXTL-PLUS12 package of programs.The final Fourier map had a minimum and maximum peaks of -0.975 and 0.699 e A° -3, respectively. Further details of the crystal structure investigation may be obtained from the Fig. 1 (a) Asymmetric unit of [NH4+]2[Sn3O(PO4)2]2-·H2O. Thermal Fachinformationszentrum, Karlsruhe, 76344 Eggensteinellipsoids are shown at 50% probability.(b) The basic building block Leopoldshafen (Germany), on quoting the depository number showing the cage-like structure with the terminal phosphate groups. CSD-408705. Full crystallographic details, excluding structure (c) View showing the connectivity between the cages, forming factors, have been deposited at the Cambridge six-membered rings along the b axis, through the phosphate groups.Crystallographic Data Centre (CCDC). See Information for Authors, J. Mater. Chem., 1998, Issue 1. Any request to the CCDC for this material should quote the full literature citation new anionic tin(II) oxyphosphate structure built up from the vertex linking between distorted square pyramidal SnO4 and and the reference number 1145/125.tetrahedral PO4 building blocks. The framework has the formula [Sn3O(PO4)2]2- and charge neutrality is achieved by Results and discussion the incorporation of protonated ammonium molecules; there are two [NH4]+ ions per framework formula unit. The entire The asymmetric unit consists of 13 non-hydrogen atoms [Fig. 1(a)] and final atomic coordinates for all the non- architecture is constructed by four-coordinate Sn(II) atoms and tetrahedral PO4 units.To our knowledge, this is the first hydrogen atoms are given in Table 2. This materials forms a 2758 J. Mater. Chem., 1998, 8, 2757–2760Table 2 Atomic coordinates (×104) and equivalent isotropic displacement parameters (A° 2×103) for [NH4+]2[Sn3O(PO4)2]2-·H2O Atom x y z Ueq a Sn(1) 10000 8675(1) 9852(2) 26(1) Sn(2) 7568(1) 7410(1) 7759(4) 22(1) P(1) 10000 8627(3) 5938(8) 17(2) P(2) 10000 6357(3) 5104(10) 18(2) O(1) 10000 7696(6) 8954(16) 16(3) O(2) 10000 9070(6) 7469(15) 22(4) O(3) 8237(13) 6574(4) 4197(12) 22(3) O(4) 8249(12) 8167(4) 5937(11) 18(2) O(5) 10000 6724(7) 6741(14) 20(4) O(6) 10000 9128(7) 4533(19) 27(5) O(7) 10000 5581(6) 5379(18) 26(4) N(100) 7697(28) 5155(11) 8001(43) 59(8) O(100) 10000 5932(10) 625(32) 73(7) aEquivalent isotropic U defined as one third of the trace of the orthogonalized Uij tensor.Fig. 3 Structure of [NH4+]2[Sn3O(PO4)2]2-·H2O showing the chains and the interactions between the ammonium and water molecules time an open-framework tin phosphate material is constructed along the c axis (hydrogens on the ammonium and water molecules by four-coordinated Sn(II) atoms alone.The exclusive presence are omitted). of four-coordinate Sn(II), in the current material, leads to the presence of a large number of SnKOKSn type linkages. These linkages create a situation where more than one type of oxygen Decomposition of amine molecules during hydrothermal atom has three-coordination. Oxygen atom O(1) bonds to reactions has been observed before.8 three tin atoms and two other oxygens, O(3) and O(5), are The SnKO bond lengths are in the range 2.058–2.457 A° (av.also three coordinate but bonded to two tin atoms and one 2.276 A ° ), and the OKSnKO bond angles lie between 72.7 and phosphorus atom. However, it should be noted that the three- 146.2° (av. 91.7°). These values are in excellent agreement with coordinate oxygen atoms do not create a situation where those for other tin phosphate materials where Sn atoms are PKOKP bondings occur, in accord with the principle that lies four coordinate3,4 with respect to oxygen.However, the at the heart of Lo� wenstein’s rule.13 longest bond distance [Sn(1)KO(3) 2.460 A ° ] and the largest The entire framework is constructed from a cage-like unit bond angles [O(3)KSn(1)KO(3) 145.1° and O(5)KSn(2)KO(3) as shown in Fig. 1(b). This unit, viz., the capped three-ring, 146.2°] are observed for linkages involving the three-coordiseen for the first time, is built up from three two-membered nate oxygen atoms. The PKO distances are in the range rings formed by the SnKOKSn linkages creating a bowl with 1.536–1.555 A ° (av. 1.546 A ° ) and the OKPKO bond angles are the oxygen [O(1)], bonding to three tin atoms, forming the in the range 106.7–111.6° (av. 109.5°). These values are comparbase of the bowl. The bowl is closed with a capping phosphate able to those observed in other phosphate materials. Important group forming a cage [Fig. 1(b)]. The individual cages, in bond distances and angles are presented in Table 3.turn, are connected to each other through another phosphate Thermogravimetric analysis (TGA) of the title compound group [Fig. 1(c)]. The linkage between the capped three- was carried out in the presence of nitrogen from room temperamembered rings and the phosphate groups occurs via a three- ture to 600 °C. Only one weight loss was observed in the coordinate oxygen atom [Fig. 1(b) and (c)]. This type of region of 300–400 °C, which corresponds to about 11.5% of bonding creates a heavily distorted six-membered ring along the total mass of the sample and can be directly correlated to the b axis [Fig. 1(c)]. the decomposition of very strongly hydrogen bonded water The linkages between the cage and the phosphate groups and ammonium molecules (calc. ca. 10%).The powder X-ray create zigzag infinite one dimensional chains along the a axis (Fig. 2). Along the c axis the structure presents continuous Table 3 Selected bond distances (A° ) and angles (°) for ribbons made of the caged three-rings and phosphate groups [NH4+]2[Sn3O(PO4)2]2-·H2Oa (Fig. 3). The ammonium and water molecules occupy spaces in between these ribbons/chains held together by strong multi- Sn(1)KO(1) 2.058(11) Sn(1)KO(2) 2.153(13) Sn(1)KO(3)a 2.457(10) Sn(1)KO(3)b 2.457(10) point hydrogen bonded interactions (Fig. 2 and 3). The Sn(2)KO(1) 2.104(8) Sn(2)KO(4) 2.190(10) ammonium molecules arise from the decomposition of the Sn(2)KO(5) 2.375(9) Sn(2)KO(3)b 2.400(9) starting amine used for the synthesis, viz., guanidium carbonate. P(1)KO(6) 1.54(2) P(1)KO(2) 1.555(14) P(1)KO(4) 1.554(9) P(1)KO(4)c 1.554(9) P(2)KO(7) 1.536(13) P(2)KO(3) 1.547(10) P(2)KO(3)c 1.547(10) P(2)KO(5) 1.56(2) O(1)KSn(1)KO(2) 89.4(5) O(1)KSn(1)KO(3)a 74.5(2) O(2)KSn(1)KO(3)a 82.0(2) O(1)KSn(1)KO(3)b 74.5(2) O(2)KSn(1)KO(3)b 82.0(2) O(3)aKSn(1)KO(3)b 145.1(4) O(1)KSn(2)KO(4) 88.2(4) O(1)KSn(2)KO(5) 72.7(3) O(4)KSn(2)KO(5) 87.8(4) O(1)KSn(2)KO(3)b 75.0(4) O(4)KSn(2)KO(3)b 81.4(3) O(5)KSn(2)KO(3)b 146.2(3) O(6)KP(1)KO(2) 106.7(8) O(6)KP(1)KO(4) 111.6(5) O(2)KP(1)KO(4) 108.8(5) O(6)KP(1)KO(4)c 111.6(5) O(2)KP(1)KO(4)c 108.8(5) O(4)KP(1)KO(4)c 109.3(7) O(7)KP(2)KO(3) 110.2(6) O(7)KP(2)KO(3)c 110.2(6) O(3)KP(2)KO(3)c 111.2(9) O(7)KP(2)KO(5) 108.8(10) O(3)KP(2)KO(5) 108.3(5) O(3)cKP(2)KO(5) 108.3(5) aSymmetry transformations used to generate equivalent atoms: Fig. 2 Structure of [NH4+]2[Sn3O(PO4)2]2-·H2O showing the ax-1/2, -y+1/2, z+1/2; b-x+1/2, -y+1/2, z+1/2; c-x, y, z. one-dimensional chains along the a axis. J. Mater. Chem., 1998, 8, 2757–2760 2759diVraction pattern of the decomposed sample indicates a poorly References crystalline phase, for which all the lines correspond to those 1 D. W. Breck, Zeolite Molecular Sieves: Structure, Chemistry and of the crystalline phase Sn2P2O7 (JCPDS: 35-28); it seems Use, Wiley and Sons, London, 1974; R. M.Barrer, Hydrothermal likely that an amorphous phase with a Sn5P ratio greater than Chemistry of Zeolites, Academic Press, London, 1982; R. Szostak, 352 is also present. Molecular Sieves: Principles of Synthesis and Identification, Van The synthesis of a novel tin oxyphosphate material Nostrand Reinhold, New York, 1989.containing one-dimensional chains constructed from capped 2 M. E. Davis and R. F. Lobo, Chem. Mater., 1992, 4, 759. 3 S. Natarajan, J. Solid State Chem., 1998, 139, 200. tin phosphate cages is accomplished. The exclusive presence 4 S. Natarajan and A. K. Cheetham, J. Solid State em., 1998, of four-coordinate SnII atoms along with a large number of in press.three-coordinate oxygens makes this material unique amoung 5 S. Natarajan, M. P. Attfield and A. K. Cheetham, Angew. Chem., framework tin phosphates. This new phase, together with Int. Ed. Engl., 1997, 36, 978. previously reported Sn(II) phosphate solids, illustrates pro- 6 S. Natarajan and A. K. Cheetham, Chem. Commun., 1997, 1089. found structural influences of relatively minor modifications 7 S.Ayyappan, A. K. Cheetham, S. Natarajan and C. N. R. Rao, in reaction conditions and/or changes in the starting source J. Solid State Chem., 1998, 139, 207. 8 S. Natarajan and A. K. Cheetham, J. Solid State Chem., 1997, for the tin. While the isolation of a one-dimensional solid with 134, 207. a [SnO4] distorted-square pyramidal core along with three- 9 G. H. Bonavia, R. C. Haushalter, S. Lu, C. J. O’Conner and coordinate oxygen provides information about the stereochem- J. Zubieta, J. Solid State Chem., 1997, 132, 144 . ical consequences of the Sn(II ) lone pair electrons, further 10 P. J. Zapf, D. J. Rose, R. C. Haushalter and J. Zubieta, J. Solid evaluation is required to exploit the structure-directing influ- State Chem., 1996, 125, 182; 1997, 132, 438 and references therein. ences of this unit in the presence of other organic amines in 11 G. M. Sheldrick, SHELXS-86 Program for Crystal Structure Determination, University of Go� ttingen, 1986; Acta. Crystallogr., the synthesis of potentially open-framework phosphate Sect. A, 1990, 35, 467. materials. 12 G. M. Sheldrick, SHELXS-93 Program for Crystal Structure solution and refinement, University of Go� ttingen, 1993. 13 W. Lo� wenstein, Am. Mineral., 1954, 39, 92. Acknowledgments The author thanks Professor C.N.R. Rao, FRS for his keen interest, help and encouragement. Paper 8/06412A 2760 J. Mater. Chem., 1998, 8, 2757&ndas

ISSN:0959-9428    

DOI:10.1039/a806412a

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Room-temperature synthesis of monodisperse mixed spinel (CoxMn1-x)3O4 powder by a coprecipitation method Young-Il Jang, HaifengWang and Yet-Ming Chiang* Department of Materials Science and Engineering, Massachusetts Institute of Technology, Cambridge, MA 02139, USA Received 24th August 1998, Accepted 25th September 1998 Spinel oxide Co3O4 has traditionally been synthesized by thermal decomposition of cobaltous salts at temperatures of 250–900 °C under oxidizing conditions.Here, we report a solution synthesis route that yields mixed spinel oxide (CoxMn1-x)3O4 at room temperature in air. (Co0.75Mn0.25)3O4 solid solutions were synthesized by coprecipitation and oxidation of Co(OH)2 and Mn(OH)2 from mixed aqueous solutions of Co- and Mn-salts.The spinel oxide phase was stabilized by crystallization of Mn3O4 resulting from the oxidation of Mn(OH)2 in the mother liquor. While undoped Mn3O4 is tetragonal, the (Co0.75Mn0.25)3O4 solid solution has a cubic symmetry due to the reduced Mn3+ concentration. The resulting powder is spherical with an average particle diameter of ca. 0.1 mm and a narrow size distribution. Transmission electron microscopy (TEM) showed that individual particles are nearly single crystalline but consist of a mosaic of multiple nanocrystallites. This observation supports an aggregation mechanism of formation for the nearly monodispersed particles.We therefore anticipated that under autoxidation conditions Introduction where Mn3O4 is favored, a predominantly Co(OH)2 material Co3O4 has extensively been investigated as an electrocatalyst doped with Mn(OH)2 would exhibit enhanced autoxidation for both O21–4 and Cl2 evolution.5 Although it has the draw- of Co(OH)2 to Co3O4 spinel.The Mn3O4 was expected to act back of lower long-term stability compared to recently devel- as a stabilizer for Co3O4 formation. The direct precipitation oped RuO2 catalysts,6 Co3O4 has the advantage of lower cost of mixed metal spinel oxides has previously been reported for while retaining adequate performance.This spinel compound ferrites,21 including those doped with Co,22 Mn,23 and Ni.24 is also receiving wide attention as a promising electrochromic material.7 There are numerous reported procedures for the preparation of Co3O4, all essentially based on the thermal Experimental decomposition of cobaltous salts at temperatures varying between 250 and 900 °C under oxidizing conditions as reviewed Co(NO3)2 6H2O (Alfa Aeasar, 99.5%), Mn(NO3)3 6H2O by Sugimoto and Matijevic.8 High temperature firing typically (Aldrich, 98%) and LiOH H2O (Alfa Aeasar, 98%) were used results in oxides with surface area of only a few square meters for preparing the starting aqueous solutions.A special aspect per gram.9 Because oxides are used in finely divided form, of our synthesis procedure is the use of LiOH as precipitating electrocatalytic properties often depend strongly on morpho- reactant rather than previously used reactants such as NH4OH. logical as well as electronic factors. Therefore, particles of The ammonium ion is known to retard the rate of autoxidation uniform shape with a large surface area and a narrow size of Mn(OH)2,20 mitigating against use of NH4OH.Preliminary distribution are generally required.10 examinations to determine appropriate autoxidation con- EVorts have previously been made to synthesize fine Co3O4 ditions for Mn3O4 were performed with a 0.1 M Mn(NO3)3 powder at low temperatures by solution methods. Sugimoto solution at room temperature in ambient air atmosphere.and Matijevic8 obtained Co3O4 particles of cubic shape with Results from these experiments were used to determine the uniform edge length (ca. 0.1 mm) by heating cobalt acetate reaction conditions for obtaining (Co0.75Mn0.25)3O4. solution at 100 °C in the presence of O2 gas.Recently, 0.1 M aqueous solutions of Co(NO3)2 and Mn(NO3)3 in a Furlanetto and Formaro10 reported that mixtures of Co3O4 Co5Mn=351 atomic ratio were added dropwise to a LiOH and CoOOH can be obtained by mixing Na3Co(NO2)6 solusolution that was vigorously stirred and monitored to keep tion with NH4NO3 solution while purging with N2 gas. Upon the pH at ca. 11. This is near the minimum solubility conditions heating at 70–100 °C, they observed that CoOOH transformed for both Co(OH)2 and Mn(OH)2.25 For comparison, the to Co3O4, resulting in spherical particles with diameter of precipitation of Co3O4 was also attempted from 0.1 M 0.2–0.3 mm.Co(NO3)2 solution under the same conditions. In each The present work describes a novel room-temperature instance, the precipitate was aged in the mother solution for coprecipitation synthesis method which yields mixed spinel 12 h at room temperature in air with continuous stirring, (CoxMn1-x)3O4.This composition is of interest for several followed by rinsing to remove Li+ and NO3- species with reasons. Like the endmembers Co3O41–5 and Mn3O4,11,12 distilled water as described in ref. 26. Finally the precipitate (CoxMn1-x)3O4 solid solutions can be applied as catalysts for was atomized into liquid nitrogen, and the frozen droplets oxygen reduction,13 and considered to be amongst the mixedwere freeze-dried (VirTis Consol 12LL, Gardiner, NY). spinel candidates for active catalysts.4,14–18 Like Co3O4, The present synthesis method has also been used to prepare Mn3O4 can easily be obtained by thermal decomposition of lithium intercalation compounds, in which case LiOH solution manganous salts.19 However, this compound can also be is added after the rinsing step and before freeze-drying.26–29 prepared by the so-called ‘autoxidation’ of Mn(OH)2 in The eVect of residual LiOH on the present materials was also alkaline aqueous solution.It is known that mixtures of MnO2, studied, as an extreme case where the precipitating reactant is Mn2O3 or Mn3O4 can be obtained by oxidation of Mn(OH)2, not completely removed.The powders were characterized by depending on many factors such as temperature, pH, nature of ions in the solution, and the rate of air or oxygen flow.20 X-ray diVraction (XRD) using a Rigaku diVractometer J. Mater. Chem., 1998, 8, 2761–2764 2761(RTP500RC) with Cu-Ka radiation and by transmission electron microscopy (TEM) using a JEOL-200CX instrument.Results and discussion The XRD pattern of the oxide precipitated from Mn(NO3)3 solution is shown in Fig. 1. Mn3O4 appears as the predominant phase after freeze-drying with no other Mn-containing minor phases being detectable. Miller indices (hkl) are indexed for the tetragonal phase (space group I41/amd) in Fig. 1. Lattice parameters calculated from the XRD data using Cohen’s least squares method30 are: a=5.764 A° , c=9.452 A° . These values agree well with the data in JCPDS (#24-734).31 The tetragonal symmetry of Mn3O4 spinel is known to result from the cooperative Jahn–Teller distortion of Mn3+ ions in octahedral sites.32 According to ref. 32, the c/a ratio decreases from 1.64 Fig. 2 Powder XRD pattern of (Co0.75Mn0.25)3O4 (with hkl in Fd39m) to 1.50 (equivalently from 1.157 to 1.054 in space group obtained by coprecipitation from Co(NO3)2 and Mn(NO3)3 solution F41/ddm) when Mn3O4 is lithiated to LiMn3O4, indicating mixed with LiOH solution (%: Li2CO3). that the degree of tetragonal symmetry decreases when lithiation reduces Mn3+ to Mn2+.Since the Mn3O4 obtained the same conditions. As shown in Fig. 3, Co(OH)2 appears as in this study has a c/a ratio of 1.64, it appears to be largely the major phase (space group P3m1), and CoOOH as a minor unlithiated after freeze-drying. This is consistent with TEM phase (space group R39m). It is thus clear that when Co(OH)2 results on LiCoO2 produced by the same method,26 where the is coprecipitated with Mn(OH)2, the spinel phase is stabilized LiOH remains as a separate amorphous phase after freeze- by Mn3O4 resulting from the autoxidation of Mn(OH)2.drying. One can notice from Fig. 1 that Li2CO3 exists as a Miller indices (hkl) are indexed in Fig. 2 for the cubic mixedminor phase. This is due to minor reaction of residual LiOH spinel phase.The lattice parameter calculated from the XRD with CO2 to form the carbonate phase during sample handling; data is a=8.168 A° , inbetween that of Co3O4 (8.084 A° )34 and LiOH is a well-known CO2 absorbent.33 However, the residual MnCo2O4 (8.226 A° ),35 as is expected from its intermediate lithium salts LiOH or Li2CO3 are easily removed from the Mn concentration.Fig. 4 shows that the variation in cubic present powder by washing as they are water soluble. spinel lattice parameter is approximately linear (Vegard’s law) Having confirmed that Mn3O4 is obtained under the present with Mn concentration, increasing upon substituting the larger precipitation and autoxidation conditions, we employed the Mn ions for Co ions.36 same experimental conditions to synthesize (Co0.75Mn0.25)3O4.Fig. 2 shows the XRD pattern of powder obtained by coprecipitation and autoxidation of Co(OH)2 and Mn(OH)2 from their mixed nitrate solution. The diVraction peaks match very well with the XRD patterns of either Co3O4 or MnCo2O4 as reported in JCPDS (#43-1003 and #23-1237).34,35 Considering the fact that MnCo2O4 in ref. 35 was synthesized by heating Co- and Mn-nitrate precursors at 760 °C, one can conclude that the present oxide obtained at room temperature is nearly as well crystallized as high temperature fired materials.The slight peak broadening in Fig. 1 and 2 is consistent with fine crystallite size, as discussed later. Note that (Co0.75Mn0.25)3O4 has a cubic symmetry (space group Fd39m), indicating that the concentration of Mn3+ is lower than the critical concentration for the cooperative Jahn–Teller distortion to a tetragonal symmetry.The formation of Co3O4 spinel was not observed in Fig. 3 Powder XRD pattern of Co(OH)2 (with hkl in P3m1) obtained precipitation experiments with Co(NO3)2 solution alone under by precipitation from Co(NO3)2 mixed with LiOH solution (%: Li2CO3; 1: CoOOH).Fig. 1 Powder XRD pattern of Mn3O4 (with hkl in I41/amd) obtained Fig. 4 Lattice parameter of (Co0.75Mn0.25)3O4 in comparison with by precipitation from Mn(NO3)3 solution mixed with LiOH solution (%: Li2CO3). Co3O434 and MnCo2O4.35 2762 J. Mater. Chem., 1998, 8, 2761–2764direct observations with higher resolution of a single particle using TEM. Fig. 6 shows a bright field TEM image of a single oxide particle and corresponding selected area diVraction pattern.We observe diVraction contrast variations within each particle showing clearly that each particle is composed of much smaller nanocrystallites. However, sharp crystalline diVraction spots are observed in the selected area diVraction pattern in Fig. 6, indicating that each particle is overall singlecrystalline. It appears that the nanocrystallites have aggregated to form a particle in which the crystallites are suYciently misoriented with respect to one another to provide diVraction contrast, yet are well aligned enough to yield sharp reflections in the selected area diVraction pattern (mosaic structure). These results strongly support the aggregation mechanism for the present system. Conclusion A cobalt-rich mixed spinel oxide (Co0.75Mn0.25)3O4 has been synthesized at room temperature by the coprecipitation of Mn Fig. 5 TEM micrograph of (Co0.75Mn0.25)3O4 particles obtained by and Co hydroxides from mixed nitrate solutions. Mn3O4 coprecipitation from Co(NO3)2 and Mn(NO3)3 solution mixed with resulting from the autoxidation of Mn(OH)2 stabilizes the LiOH solution. mixed spinel phase under conditions where the undoped Co precursor produces Co(OH)2.The resulting particles are Fig. 5 shows a TEM bright-field image of the nearly single-crystalline spheres (diameter ca. 0.1 mm), and (Co0.75Mn0.25)3O4 obtained in this study. Spherical particles have a narrow size distribution. Each particle is composed of of ca. 0.1 mm with a narrow particle size distribution are nanometer-scale crystallites, supporting an aggregation mechobserved. The uniform particle size obtained in this study is anism of particle formation.This new route is simple and noteworthy because it has generally been believed that it is cost-eVective, and is suitable for the synthesis of chemically diYcult, if not impossible, to obtain monodispersed metal homogeneous mixed cobalt spinel oxides of high surface area (hydrous) oxide colloids by the reaction of a strong base with and uniform particle size distribution for such applications as a metal salt solution.37 The much smaller particle size that we catalysis and electrochromic materials.obtained compared to that of the Co3O4 of ref. 10 (0.2–0.3 mm) is probably due to our lower synthesis temperature (25 °C vs.Acknowledgments 70–100 °C). The Li2CO3 and LiOH phases resulting from LiOH added after the rinsing step and before freeze-drying This work has been funded by the INEEL University Research are seen surrounding the spherical particles. However, without Consortium. The INEEL is managed by Lockheed Martin addition of LiOH after rinsing, those lithium-containing phases Idaho Technology Company for the U.S.Dept. of Energy, were rarely observed with TEM. Idaho Operations OYces, under contract no. DEIn order to explain the formation of precipitates with AC07–94ID13223. We used instrumentation in the Shared uniform size, the LaMer model,38,39 based on a short single Experimental Facilities at MIT, supported by NSF Grant burst of nucleation followed by uniform growth has been No. 9400334-DMR. Y.I.J. also acknowledges a fellowship previously used. However, more recently, an aggregation mech- from the Ministry of Education, Korea. anism in which each particle forms a large number of smaller subunits has been proposed by Matijevic, as reviewed in refs. References 37 and 40. In order to address this matter, we have conducted 1 C.Iwakura, A. Honji and H. Tamura, Electrochim. Acta, 1981, 26, 1319. 2 P. Rasiyah and A. C. C. Tseung, J. Electrochem. Soc., 1983, 130, 365. 3 B. E. Conway and T. C. Liu, Mater. Chem. Phys., 1989, 22, 163. 4 R.-N. Singh, M. Hamdani, J.-F. Koenig, G. Poillerat, J. L. Gautier and P. Chartier, J. Appl. Electrochem., 1990, 20, 442. 5 R. Boggio, A. Carugati, G. Lodi and S.Trasatti, J. Appl. Electrochem., 1985, 15, 335. 6 P. Siviglia, A. Daghetti and S. Trasatti, Colloids Surf., 1983, 7, 15. 7 F. Svegl, B. Orel, M. G. Hutchins and K. Kalcher, J. Electrochem. Soc., 1996, 143, 1532. 8 T. Sugimoto and E. Matijevic, J. Inorg. Nucl. Chem., 1979, 41, 165. 9 D. Pope, D. S. Walker and R. L. Moss, J. Colloid Interface Sci., 1977, 60, 216. 10 G. Furlanetto and L.Formado, J. Colloid Interface Sci., 1995, 170, 169. 11 T. Yamashita and A. Vannice, J. Catal., 1996, 163, 158. 12 A. Maltha, H. F. Kist, T. L. F. Favre, H. G. Karge, F. Asmussen, H. Onishi, Y. Iwasawa and V. Ponec, Appl. Catal. A: General, 1994, 115, 85. 13 M. Sugawara, M. Ohno and K. Matsuki, J. Mater. Chem., 1997, 7, 833. Fig. 6 Bright field TEM image and corresponding selected area 14 M.Hamdani, J. F. Koenig and P. Chartier, J. Appl. Electrochem., 1988, 18, 561. diVraction pattern of a (Co0.75Mn0.25)3O4 particle obtained by coprecipitation from Co(NO3)2 and Mn(NO3)3 solution mixed with 15 J. L. Gautier, A. Restovic and P. Chartier, J. Appl. Electrochem., 1989, 19, 28. LiOH solution. J. Mater. Chem., 1998, 8, 2761–2764 276316 D. Panayotov, M.Khristova and D. Mehandjiev, J. Catal., 1995, 29 B. Huang, Y.-I. Jang, Y.-M. Chiang and D. R. Sadoway, J. Appl. 156, 219. Electrochem., in press. 17 J. Ziolkowski, A. M. Maltha, H. Kist, E. J. Grootendorst, 30 B. D. Cullity, in Elements of X-Ray DiVraction, Addison and H. J. M. de Groot and V. Ponec, J. Catal., 1996, 160, 148. Wesley, Massachusetts, 2nd edn., 1978, p. 363. 18 J.Ghose and K. S. R. C. Murthy, J. Catal., 1996, 162, 359. 31 JCPDS-International Centre for DiVraction Data, #24-734. 19 O. Bricker, Am. Mineral., 1965, 50, 1296. 32 M. M. Thackeray, W. I. F. David, P. G. Bruce and 20 A. R. Nichols, Jr. and J. H. Walton, J. Am. Chem. Soc., 1942, J. B. Goodenough, Mater. Res. Bull., 1983, 18, 461. 64, 1866. 33 W. A. Hart and O. F. Beumel, Jr., in Comprehensive Inorganic 21 T.Sugimoto and E. Matijevic, J. Colloid Interface Sci., 1980, Chemistry, ed. A. F. Trotman-Dickenson, Pergamon, New York, 74, 227. 1973, p. 352. 22 H. Tamura and E. Matijevic, J. Colloid Interface Sci., 1982, 90, 34 JCPDS-International Centre for DiVraction Data, #43-1003. 100. 35 JCPDS-International Centre for DiVraction Data, #23-1237. 23 Z. X. Tang, C. M. Sorensen, K. J. Klabunde and 36 R. D. Shannon and C. T. Prewitt, Acta Crystallogr., Sect. B, 1969, G. C. Hadjipanayis, J. Colloid Interface Sci., 1991, 146, 38. 25, 925. 24 C.-L. Huang and E. Matijevic, Solid State Ionics, 1996, 84, 249. 37 E. Matijevic, Chem.Mater., 1993, 5, 412. 25 C. F. Baes, Jr. and R. E. Mesmer, in The Hydrolysis of Cations, 38 V. K. LaMer, Ind. Eng. Chem., 1952, 44, 1269. Krieger, Florida, 1986, pp. 219, 238. 39 V. K. LaMer and R. J. Dinegar, J. Am. Chem. Soc., 1950, 72, 26 Y.-M. Chiang, Y.-I. Jang, H. Wang, B. Huang, D. R. Sadoway 4847. and P. Ye, J. Electrochem. Soc., 1998, 145, 887. 40 E. Matijevic, Langmuir, 1994, 10, 8. 27 G. Ceder, Y.-M. Chiang, D. R. Sadoway, M. K. Aydinol, Y.-I. Jang and B. Huang, Nature, 1998, 392, 694. 28 Y.-I. Jang, B. Huang, Y.-M. Chiang and D. R. Sadoway, Electrochem. Solid-State Lett., 1998, 1, 13. Paper 8/06653A 2764 J. Mater. Chem., 1998, 8, 2761–2764

ISSN:0959-9428    

DOI:10.1039/a806653a

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Microwave synthesis of LaCrO3 Masato Iwasaki,a Hirotsugu Takizawa,*a Kyota Uheda,a Tadashi Endoa and Masahiko Shimadab aDepartment of Materials Chemistry, Graduate School of Engineering, Tohoku University, Aoba-yama 07, Sendai 980–8579, Japan. E-mail: takizawa@aim.che.tohoku.ac.jp bInstitute for Advanced Materials Processing, Tohoku University, Katahira, Sendai 980–8577, Japan Received 2nd June 1998, Accepted 14th September 1998 The synthesis of LaCrO3 from a stoichiometric mixture of Cr2O3 and La2O3 powder by microwave irradiation was examined using a multi-mode 28 GHz microwave heating system.In the La2O3–Cr2O3 system, Cr2O3 strongly absorbs microwaves while La2O3 is transparent to microwaves. The reaction proceeded rapidly and LaCrO3 could be synthesised within 15 min of irradiation.Surprisingly, the temperature required for the microwave synthesis (440 °C) is much lower than that for conventional synthesis (>1200 °C) using an electric furnace. The eVect of the partial pressure of oxygen in the microwave chamber on the formation of LnCrO3 was investigated. The formation of LnCrO3 is enhanced with increase in the partial pressure of oxygen for Ln=La and Nd, while opposite behavior is seen for Ln=Gd, Tb, Dy and Ho. 1 Introduction 2 Experimental The increase in temperature of Cr2O3 and La2O3, and the Lanthanum chromite, LaCrO3, has received much interest as an electrode or interconnector for solid oxide fuel cells (SOFC) possible synthesis of LaCrO3 by microwave irradiation were determined. Multi-mode microwave heating equipment (Fuji and a heating element for high-temperature electric furnaces.1 Generally, LaCrO3 is synthesised by solid state reaction of the Dempa Kogyo Co., LTD, Japan; FMS-10–28) was used.Appropriate amounts of Cr2O3 and La2O3 powder were mixed component oxides at high temperatures (T>1200 °C) in air. Since the diVusion of reactants is very slow in the solid state and pressed into a pellet (10 mm in diameter and 5 mm in thickness).A hole (1.6 mm in diameter and 3 mm in depth) reaction, the reaction requires a long heating period and intermediate grinding to achieve good homogeneity. Attempts was drilled at the centre of pellet to insert a thermocouple. In the microwave chamber (multi-mode cavity; 0.8 m in diameter to synthesise LaCrO3 at lower temperature have been made by hydrothermal reaction2 and thermal decomposition of a and 1.2 m in length), a platinum sheathed Pt–Pt/Rh10% thermocouple (1.5 mm in diameter) was inserted into the mixed organometallic complex of La and Cr.3 sample pellet and then the sample was surrounded by quartz Microwave heating has been used for the syntheses of wool as a heat insulating material which is transparent to various inorganic materials, e.g., YBa2Cu3O7-d,4,5 La2CuO4,5 microwaves.A platinum sheath was used to achieve an eVective b-SiC,6 BaWO4,7 Bi4V2O11, PbV2O6.8 If some constituent shielding from the microwaves. A schematic illustration of the materials of a chemical reaction system strongly absorb microsample set up is shown in Fig. 1. The temperature was moni- waves, the resulting heat generated can be used to drive a tored during microwave irradiation. We considered that moni- solid state reaction with another component. This is especially tored temperature was not aVected by microwave dielectric applicable for solid state reaction systems containing C, V2O5, fields as suggested by Rowly et al.13 After irradiation, the Cr2O3, CuO, MnO2, PbO2, WO3 or Fe3O4.9 Detailed aspects of microwave processing have been reviewed by Sutton.10 The above mentioned examples were investigated by using microwaves with a frequency of 2.45 GHz which is widely applied for heating and cooking foodstuVs.It is known by experiment that the required size of the microwave chamber to achieve a uniform electric field distribution is directly proportional to the wavelength of the microwaves.11 The temperature rise of a material due to coupling with the microwaves is theoretically proportional to the dielectric loss factor of the material.9 However, because the dielectric loss factor correlates with temperature and microwave frequency,9 diVerent temperature profiles may be obtained even for the same material if the frequencies of microwaves used are diVerent.Chromium sesquioxide (Cr2O3) absorbs 2.45 GHz microwave radiation very strongly.12 If Cr2O3 absorbs 28 GHz microwave radiation in a similar manner, such microwaves can be utilised for the syntheses of various chromium double oxides. This paper reports the rapid and simple procedure for the synthesis of LaCrO3 by 28 GHz microwave irradiation.The oxygen pressure dependence on the formation of LnCrO3 Fig. 1 Schematic illustration of the cross section of the microwave chamber. (Ln=La, Nd, Gd, Tb, Dy and Ho) is also investigated. J. Mater. Chem., 1998, 8, 2765–2768 2765sample was cooled to room temperature in the microwave chamber and the obtained sample was characterised by X-ray diVraction analysis and scanning electron microscopy. 3 Results and discussion Fig. 2 shows the temperature–time profiles of (a)Cr2O3 and (b)La2O3 during 28 GHz microwave irradiation (0.3 kW output) in air. It is apparent that La2O3 is transparent to 28 GHz microwave radiation. In contrast to La2O3, the temperature of Cr2O3 rose rapidly. It was reported that the temperature of Cr2O3 rose rapidly above 1200 °C on exposure to a few minutes (referred to ‘thermal runaway’) of 2.45 GHz microwave irradiation.12 However, the drastic temperature rise was completed within 2 min and the temperature stabilized at ca. 500 °C over longer periods of irradiation when 28 GHz microwaves was irradiated on Cr2O3. The diVerence is thought to be due to the diVerence in the dielectric properties of Cr2O3 at 28 GHz and 2.45 GHz.X-Ray diVraction analysis of the irradiated sample did not reveal any phase changes. Fig. 3 shows the temperature–time profiles of a Fig. 4 X-Ray diVraction patterns of Cr2O3+La2O3 stoichiometric mixture after microwave irradiation of (a) 0.3 kW, (b) 0.5 kW and (c) 1 kW in air. Cr2O3+La2O3 stoichiometric mixture during 28 GHz microwave irradiation of (a) 0.3 kW, (b) 0.5 kW and (c) 1 kW in air.The profile of the Cr2O3+La2O3 mixture is similar to that of Cr2O3 with the maximum temperature reached for the Cr2O3+La2O3 mixture being slightly lower than that of Cr2O3. The profiles at various microwave power are quite similar except for diVerences in saturation temperature. X-Ray diVraction patterns of Cr2O3+La2O3 mixtures after microwave irradiation of various power outputs (a) 0.3 kW, (b) 0.5 kW and (c) 1 kW in air for 15 min are shown in Fig. 4. As seen, LaCrO3 was readily formed from a mixture of Cr2O3+La2O3 within a short time (15 min) at each microwave power, but the amount of LaCrO3 formed at 0.3 and 0.5 kW was slightly larger than that at 1 kW. This result is due to the Fig. 2 Temperature–time profiles of (a) Cr2O3 and (b) La2O3 during 0.3 kW microwave irradiation in air.Fig. 5 X-Ray diVraction patterns of Cr2O3+La2O3 stoichiometric Fig. 3 Temperature–time profiles of Cr2O3+La2O3 stoichiometric mixture during microwave irradiation of (a) 0.3 kW, (b) 0.5 kW and mixture after 0.3 kW microwave irradiation (a) in O2(3.7 atm), (b) in air and (c) in vacuo. (c) 1 kW in air. 2766 J. Mater. Chem., 1998, 8, 2765–2768Scanning electron microscopy showed that the particle size of the raw material powders, Cr2O3 and La2O3, were 0.39 and 0.79 mm, respectively. Fig. 6 shows a scanning electron micrograph of LaCrO3 obtained by microwave heating. In this case, microwave heating was carried out using a 0.3 kW output in 3.7 atm oxygen pressure for 15 min. After microwave irradiation, the average particle size of the resultant LaCrO3 was about 1.6 mm, with a size distribution of 0.4–3.2 mm.The synthesis of other lanthanide chromites, LnCrO3(Ln= Nd, Gd, Tb, Dy, Ho), from stoichiometric mixtures of Cr2O3 and Nd2O3, Gd2O3, Tb4O7, Dy2O3, or Ho2O3 have also been carried out. Fig. 7 shows the X-ray diVraction patterns of Cr2O3+Gd2O3 mixtures after 0.3 kW microwave irradiation for 30 min in various atmospheres.It is clear that GdCrO3 was formed in each atmosphere. However, in contrast to LaCrO3, the amount of GdCrO3 increased with decreasing the partial pressure of oxygen and single phase GdCrO3 was obtained in vacuo. During our extensive study, it was seen that the eVect of the atmosphere on the formation of the LnCrO3 Fig. 6 Scanning electron micrograph of LaCrO3 obtained by 0.3 kW phase by microwave irradiation can be divided into two groups microwave irradiation in O2(3.7 atm) atmosphere. i.e., enhancement or inhibition of reactivity upon increasing the partial pressure of O2. The former behavior is seen for LaCrO3 and NdCrO3, and the latter for GdCrO3, TbCrO3, diVerence in the maximum temperature at each microwave DyCrO3 and HoCrO3.Considering the lanthanide contraction, power and high microwave power is thus not required for the we can suppose that it is easy to synthesise LnCrO3 under synthesis of LaCrO3. It is notable that it takes only 15 min to high partial pressure of O2 if the ionic radius of Ln is relatively synthesise LaCrO3 by microwave irradiation, which is a very large (Ln=La, Nd), while the use of a low partial pressure short time as compared with conventional heating methods.14 of O2 is favoured if the ionic radius of Ln is relatively small Fig. 5 shows the X-ray diVraction patterns of a (Ln=Gd, Tb, Dy, Ho). This diVerence must be due to Cr2O3+La2O3 mixture after 0.3 kW microwave irradiation diVerences in the diVusion mechanism for the formation of (a) in O2(3.7 atm), (b) in air and (c) in vacuo for 15 min.It LnCrO3 from a Ln2O3 and Cr2O3 mixture for diVerent Ln. can be seen that LaCrO3 was formed in each atmosphere and Although the detailed diVusion mechanism under microwave the higher the partial pressure of O2, the smaller was the irradiation has, as yet, not been clarified, it should be pointed amount of the unreacted La2O3 and Cr2O3.When the out that the diVerence in crystal structure of Ln2O3, i.e., irradiation was carried out at a pressure of 3.7 atm oxygen, hexagonal A-type for Ln=La, Nd and cubic C-type for Ln= single phase LaCrO3 was obtained. Generally, conventional Gd, Tb, Dy and Ho, would aVect the diVusion mechanism for synthesis of LaCrO3 does not require a high oxygen pressure.the formation of LnCrO3 via solid state reaction with Cr2O3. Synthesis under high oxygen pressure leads to the formation In conclusion, a new simple procedure for the synthesis of of volatile chromium(VI) oxide.15 However, in the microwave LaCrO3 has been developed in the present study. Microwave synthesis, chromium(VI) oxides was not formed even at 3.7 atm heating is eVective for the synthesis of complex oxides contain- oxygen.This is due to considerably lower reaction temperature ing a component oxide strongly coupled with microwaves. for the microwave synthesis as compared with the conventional There is still an unsolved problem as to why the reaction synthesis. proceeded rapidly and completely within a short time at low temperature (440 °C).This observation suggests an existence of a ‘non-thermal eVect’ of the microwave electric field on diVusion of chemical species. Acknowledgement This work was supported in part by the new Energy and Industrial Technology Development Organization (NEDO) for the R&D Program (No. A-061) and the National Industrial Research Institute of Nagoya (NIRIN), Japan. Thanks are due to Mr.T. Saji and Mr. T. Kuge, Research and Development Department, Fuji Dempa Kogyo, Co., LTD, Tsukuba, Japan, for helpful discussions on experimental set up. References 1 A. Furusaki, H. Konno and R. Furuichi, J. Mater. Sci., 1995, 30, 2829. 2 M. Yoshimura, S. T. Song and S. Somiya, J. Ceram. Soc. Jpn., 1982, 90, 91. 3 S. Nakayama and M. Sakamoto, J. Ceram. Soc. Jpn., 1992, 100, 342. 4 K. G. K. Warrier, H. K. Varma, T. V. Mani and A. D. Damodaran, J. Am. Ceram. Soc., 1992, 75, 1990. 5 D. R. Baghurst, A. M. Chippindale and D.M. P.Mingos, Nature, Fig. 7 X-Ray diVraction patterns of Cr2O3+Gd2O3 stoichiometric 1988, 332, 311. 6 P. D. Ramesh, B. Vaidhyanathan, M. Ganguli and K. J. Rao, mixture after 0.3 kW microwave irradiation (a) in O2(7.7 atm), (b) in air and (c) in vacuo. J. Mater. Res., 1994, 9, 3025. J. Mater. Chem., 1998, 8, 2765–2768 27677 D. R. Baghurst and D. M. P. Mingos, J. Chem. Soc., Chem. 12 L. M. Sheppard, Ceram. Bull., 1988, 67, 1656. 13 A. T. Rowley, R. Wroe, D. Vazquez-Navarro, W. Lo and Commun., 1988, 829. 8 B. Vaidhyanathan, M. Ganguli and K. J. Rao, Mater. Res. Bull., D. A. Cardwell, J. Mater. Sci., 1997, 32, 4541. 14 N. Sakai and S. Stølen, J. Chem. Thermodyn., 1995, 27, 493. 1995, 30, 1173. 9 D. M. P. Mingos and D. R. Baghurst, Chem. Soc. Rev., 1991, 15 D. B. Meadowcroft and J. M. Wimmer, Am. Ceram. Soc. Bull., 1979, 58, 610. 20, 1. 10 W. H. Sutton, Ceram. Bull., 1989, 68, 376. 11 T. Saji, New Ceramics, 1995, 5, 21 (in Japanese). Paper 8/04139C 2768 J. Mater. Chem., 1998, 8, 2765–2768

ISSN:0959-9428    

DOI:10.1039/a804139c

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Gas-phase formation of zinc/cadmium chalcogenide cluster complexes and their solid-state thermal decomposition to form II–VI nanoparticulate material Nigel L. Pickett,a† Steven Lawson,a W. Gregor Thomas,a Frank G. Riddell,a Douglas F. Foster,a David J. Cole-Hamilton* and John R. Fryerb aSchool of Chemistry, The University of St. Andrews, St. Andrews, Fife, UK KY16 9ST.E-mail: djc@st-and.ac.uk bDepartment of Chemistry, Glasgow University, Glasgow, UK G12 8QQ Received 14th August 1998, Accepted 30th September 1998 Gas-phase reactions between R2Zn (R=Me and Et) and tBuSH produce cluster complexes of the type [RZnStBu]n. These clusters, along with [MeZnStBu(py)]2 (py=pyridine), have been characterised by 13C{1H} solid-state NMR. On heating to 100 °C in the solid-state, the complexes [MeZnStBu]5 and [MeZnStBu(py)]2 release dimethylzinc (Me2Zn) to form the zinc bis(thiolate) compound, [Zn(StBu)2]n, with further heating (>200 °C) leading to the formation of ZnS.The ethyl analogue, [EtZnStBu]5, does not lose Et2Zn on heating and thermogravimetric analysis (TGA) suggests a diVerent decomposition pathway, one which mainly involves loss of the organic moieties without the concurrent loss of volatile Zn or S compounds, although ZnS is again the final thermal decomposition product.The decomposition of the involatile pentamers, [MeZnStBu]5 and [EtZnStBu]5, and the dimer, [MeZnStBu(py)]2, proceeds at higher temperature (200–350 °C) to give agglomerates of ME nanoparticulate material, with the individual particles having diameters of 2–20 nm in all cases.The mechanistic pathway by which these clusters decompose appears to be highly dependent upon the R group (Me or Et) present within the cluster. Preliminary results suggest that complexes of the type [RMEtBu]n are also produced from the gas-phase reactions of Me2Zn with tBuSeH and from Me2Cd with tBuSH.others,11 allowed us to propose a mechanism by which particle Introduction nucleation, growth, and suppression of growth by pyridine The use of metal organic vapour phase epitaxy (MOVPE) to can be explained.10c produce single-crystal layers of wide band gap 12–16 (II–VI ) In the absence of a Lewis base (pyridine), particle growth materials (ZnS, ZnSe, CdS and CdSe) is now an established occurs by initial association between R2M and H2E leading technique.The most common method still employs the ther- to the formation of [RMEH]2, which grows rapidly into mally controlled reaction between a group 12 dialkyl com- clusters of the type, HEn(ME)xMRm, consisting of a central pound, R2M (R=Me, Et etc.; M=Zn and Cd) and a group ME core with E–H and M–R fragments on the surface of the 16 hydride (H2S or H2 Se).1 However, a premature reaction particles, which are suspended within the carrier-gas.Surface (prereaction) which takes place between the two precursors in bound M–R fragments react with gas-phase H2E leading to the cold zone of the growth cell, upstream of the heated the elimination of RH and formation of new E–H surface substrate, can cause adverse eVects on the properties of the bound fragments.These E–H sites in turn react with gasgrown epilayers. It has been shown that the introduction of a phase R2M leading to the formation of new M–R sites on the s-donor compound to the reaction system can lead to a continuously growing particles. Termination of particle growth dramatic reduction of the prereaction. Compounds which have may occur by pyridine binding to the reactive surface M–R been investigated as prereaction suppressants include: 1,4- sites, preventing their reaction with H2E by a blocking dioxane,2–4 thioxane,4 triethylamine,5,6 1,3,5-trimethylhexahy- mechanism.10c dro-1,3,5-triazine,6 N,N,N¾,N¾-tetramethylethane-1,2-diamine,7 One step of this growth mechanism involves the reaction of N,N,N¾,N¾-tetramethyldiaminomethane8 and pyridine.9 surface bound E–H fragments with R2M to eliminate RH.While investigating the nature of the prereaction, recent Thus, one other approach that can be used in an eVort to studies of ours10 have shown that the gas-phase reactions that eliminate prereactions in the growth of II–VI semiconductors occur between H2S/H2Se and Me2Cd/Me2Zn result in the is to reduce or eliminate E–H bonds in the group 16 precursors.formation of chalcogenide deposits (ZnS, ZnSe, CdS and A number of research groups have, in their eVorts to control CdSe), with the deposits consisting of poorly formed nanocrys- II–VI epitaxial growth, used a range of dialkyl sulfides/ talline material of the hexagonal phase within the size range selenides and thiols/selenols as the group 16 precursor.The 10–100 nm. The addition of small amounts of pyridine to the most promising alternative group 16 precursors are those reaction system greatly improves the crystallinity exhibited which are bound to tert-butyl groups i.e. tBu2S12 and tBu2Se,13 by the particles, while addition of larger quantities of pyrialong with tBuSH14,15 and tBuSeH.16 dine retains the improved crystal quality whilst also lead- Although successful II–VI growth has been achieved with ing to a decrease in particle size.The ability of pyridine to tBuSH in combination with Me2Zn, Et2Zn and Me2Cd,14,15 it influence the particle size decreases in the order ZnS> has been suggested that the prereaction is only totally elimin- CdS>CdSe>ZnSe. These results, along with the work of ated when Et3N is added to the growth system.17 Herein, we report that in the gas-phase tBuEH and R2M (R=Et, Me; M=Zn, Cd; E=S, Se) react to form zinc/cadmium chalcogen- †Current address: School of Chemistry and Biochemistry, and School ide cluster complexes of the type [RMEtBu]n.The formation, of Materials Science and Engineering and Molecular Design Institute, Georgia Institute of Technology, Atlanta, Georgia 30332–0400, USA.characterisation and solid-state thermal decomposition of these J. Mater. Chem., 1998, 8, 2769–2776 2769clusters is described and discussed, with the goal of advancing fractional distillation at normal pressure through a high eYciency (30 theoretical plates) SpaltrohrA (Fisher Scientific the understanding of the epitaxial growth process using both conventional and single-source precursors.18 UK) distillation column.Bp 79–80 °C. Yield after purification 48 g (70%). 13C{1H} NMR (C6D6), d 36.06 (s, SeC(CH3)3) and 38.77 (s, SeC). 1H NMR (C6D6), d 0.10 (s, 1H, SeH) and Experimental 1.40 [s, 9H, SeC(CH3)3]. General (C2D5)2Zn. A Grignard solution of CD3CD2MgBr in Microanalytical data were obtained at the University of St. 100 cm3 of Et2O was prepared from 23 g of CD3CD2Br Andrews. Analysis of samples by powder X-ray diVraction (0.20 mol) and 6 g of Mg turnings (0.25 mol). This Grignard (PXRD) were carried out on a Sto�e STADI/P diVractometer solution was added dropwise to a rapidly stirred solution of using Cu-Ka radiation with data collected in the transmission 13.0 g of ZnCl2 (0.095 mol) in 100 cm3 of Et2O at a rate mode.Transmission electron micrographs (TEM) were suYcient to maintain a steady reflux (ca. 30 min). The resulting obtained at the University of St. Andrews on a Phillips EM suspension was externally heated to reflux for 1 h after com- 301 microscope at 80 keV and at Glasgow University on a plete addition of the Grignard reagent.All volatiles were JEOL 1200 EX operated at 120 keV (point resolution of collected by trap-trap distillation in vacuo into a -196 °C cold 0.3 nm) or on a ABT 002B operated at 200 keV (point trap and the (C2D5)2Zn subsequently purified by fractional resolution of 0.18 nm). Particle sizes were determined by direct distillation. Bp 118 °C. Yield 11.4 g, 90%. 13C{1H} NMR measurement of individual particles from the transmission (C6D6), d 5.74 (qnt, JD–C=18.5 Hz, ZnCD2CD3), 9.24 (spt, electron micrographs. 13C{1H} solid-state NMR spectra were JD–C=19.0 Hz, ZnCD2CD3). 2H NMR (C6H6–C6D6), d 0.27 obtained on a Bru�ker MSL 500 spectrometer using CPMAS (br s, 2 2H, ZCD3), 1.20 (br s, 3 2H, ZnCD2CD3). accumulation techniques with chemical shifts referenced to the CH2 resonance (d 38.56) of an external adamantane sample.[MeZnStBu]5. This compound was prepared according to Solution NMR data were recorded on a Bru�ker Associates the literature procedure,20 from the low temperature (-78 °C) AM300 spectrometer operating in the Fourier transform mode reaction between Me2Zn and tBuSH in light petroleum. with (for 13C) noise proton decoupling.The 13C{1H} and 1H Yielding, after work-up, a white powder of the product. NMR spectra were run in deuteriated solvents for the lock Found: C, 35.29; H, 7.65; C25H60S5Zn5 requires C, 35.41; H, signal, as given for each spectrum, with chemical shifts in ppm 7.13%. Solid-state 13C{1H} NMR, d-8.20 (s, ZnCH3),-6.76 to high frequency of tetramethylsilane (TMS) as the internal (s, ZnCH3), -5.78 (s, 2ZnCH3), -0.84 (s, ZnCH3), 35.39 [s, reference.The 2H NMR spectra were run in non-deuteriated SC(CH3)3], 36.09 [s, SC(CH3)3], 36.43 [s, SC(CH3)3], 36.50 [s, solvent, as given for each spectrum, with a few drops of the SC(CH3)3], 36.80 [s, SC(CH3)3], 46.64 (s, SC), 50.83 (s, SC), deuteriated analogue to act as both the lock signal and internal 51.59 (s, SC), 52.45 (s, SC), 52.74 (s, SC).reference—the 2H NMR signals were then back-referenced to TMS by using the known shift of the solvent relative to TMS. [MeZnStBu(py)]2. This compound was prepared according Simultaneous thermal analyses [STA, thermogravimetric to the literature procedure,20 from the low temperature analysis—diVerential temperature analysis (TGA–DTA)] were (-78 °C) reaction between Me2Zn, tBuSH and pyridine in performed on a ‘TA Instruments SDT 2960 Simultaneous light petroleum.Yielding, after work-up, a white powder of DTA-TGA’ instrument. the product. Found: C, 47.84; H, 6.83; N, 4.99; C20H34S2N2Zn2 Dry oxygen free nitrogen, helium and argon (BOC), purified requires C, 48.30; H, 6.89; N, 5.63%. Solid-state 13C{1H} by passing through two consecutive columns (2.5 ×80 cm) NMR, d -7.29 (s, 2ZnCH3), 36.77 [s, SC(CH3)3], 37.92 [s, packed with Cr2+ on silica, were used as the carrier gases and SC(CH3)3], 43.42 (s, SC), 43.91 (s, SC), 124.14 (s, pyC), as the inert atmospheres under which all preparations and 124.91 (s, pyC), 138.69 (s, pyC), 149.64 (br s, 2pyC).manipulations were carried out. Greaseless joints and taps were employed and manipulations were carried out using Gas-phase preparation of Zn and Cd chalcogenide cluster standard Schlenk line and catheter tubing techniques.complexes Dimethylcadmium, dimethylzinc and diethylzinc were prepared and purified as described previously.19 ZnCl2, Mg Prereaction experiments were conducted at just above turnings, tBuMgCl (2.0 mol dm-3 in Et2O), C2D5OH, pyridine atmospheric pressure (101.350 Pa) using an experimental setand tBuSH were purchased from Aldrich and the pyridine and up as schematically illustrated in Fig. 1. Carrier-gas containing tBuSH distilled from CaH2 prior to use. Amorphous Se powder specific gas-phase concentrations of reactants [Me2Cd, Me2Zn, (mesh size <325) was purchased from Johnson Matthey.Et2Zn and (C2D5)2Zn together with tBuSeH or tBuSH] were C2D5OH was transformed into C2D5Br by a standard method allowed to meet at the same point along a horizontal quartz on treatment with HBr. Light petroleum (bp 40–60 °C) and tube. Total gas flow through the apparatus was in the range diethyl ether were dried by distillation from sodium 300–450 cm3 min-1 and the gas-phase concentration of indidiphenylketyl and degassed prior to use.vidual reactants in the range 1×10-3–1×10-2 mol dm-3. EZuent gases containing unused reactants and any volatile Preparation of precursors products were allowed to pass through neutralising solutions followed by an industrial scrubber before being released into tBuSeH. A standard solution of tBuMgCl in Et2O (250 cm3, a fumehood.For specific experiments, eZuent gas containing 2.0 mol dm-3, 0.5 mol) was further diluted by the addition of any volatile products was diverted via a gas sampling port to 750 cm3 Et2O. Selenium powder (39.5 g, 0.50 mol ) was added a gas chromatograph (HP 5890)—mass spectrometer (HP in small batches over 2 h to the rapidly stirred Grignard 5972 series mass selective detector) for analysis (GC-MS).solution to form a voluminous white precipitate of tBuSeMgCl. The suspension was stirred overnight before being cooled in an ice bath and hydrolysed by the cautious addition of aqueous [EtZnStBu]5. Using the experimental set-up as described above, the gas-phase mixing of stoichiometric quantities of HCl (600 cm3, 1moldm-3, 0.6 mol). The clear pale-yellow organic layer was separated and the lower colourless aqueous Et2Zn and tBuSH produced a fine white powder along the entire length of the quartz tube.Attempts to obtain a satisfac- layer extracted with 3×100 cm3 of Et2O. The combined organic extracts were dried over MgSO4 followed by CaH2. tory elemental analysis were hampered by the sample constantly losing weight when exposed to air owing to hydrolysis.Finally, the tBuSeH, a clear, colourless and slightly light sensitive liquid, was separated from Et2O and purified by Solid-state 13C{1H} NMR, d 5.85 (s, ZnCH2CH3), 6.30 2770 J. Mater. Chem., 1998, 8, 2769–2776Fig. 1 Schematic diagram of the experimental set-up used both in the preparation and solid-state thermolysis of metal chalcogenides cluster complexes.(s, ZnCH2CH3), 7.19 (s, ZnCH2CH3), 7.69 (s, ZnCH2CH3), before thermolysis, 1.53 g; weight of residue, 0.32 g; 79.1% weight loss. [EtZnStBu]5: weight before thermolysis, 3.15 g; 8.63 (s, ZnCH2CH3), 14.13 (s, ZnCH2CH3), 14.61 (br s, 2ZnCH2CH3), 14.77 (br s, 2ZnCH2CH3), 35.94 [s, SC(CH3)3], weight of residue, 1.45 g; 54.0% weight loss. 36.21 [s, SC(CH3)3], 36.43 [s, SC(CH3)3], 36.75 [s, SC(CH3)3], 36.92 [s, SC(CH3)3], 45.47 (s, SC), 50.18 (s, SC), 50.46 (s, Formation of [Zn(StBu)2]n from [MeZnStBu(py)]2 SC), 50.93 (s, SC), 51.34 (s, SC).Heating [MeZnStBu(py)]2 to 100 °C in vacuo (ca. 1 h) liberated a colourless liquid, which was collected in a cold trap. This [MeZnSetBu]n. The gas-phase mixing of stoichiometric left a dirty white solid. The colourless liquid distillate was quantities of Me2Zn and tBuSeH produced a white deposit.shown to be [Me2Zn(py)2]: 1H NMR, d -0.70 (s, ZnCH3), Attempts to obtain a satisfactory elemental analysis were again 7.32 (m, pyH), 7.72 (m, pyH), 8.58 (m, pyH). 13C{1H} NMR, hampered by the sample constantly losing weight when exposed d -12.34 (s, ZnCH3), 124.07 (m, pyC), 136.69 (m, pyC), to air owing to hydrolysis.The solid-state 13C{1H} NMR 149.44 (m, pyC). The white solid was identified as impure spectrum consisted of broad peaks which merged into one [Zn(StBu)2]n: Found: C, 37.55; H, 7.47; N, 0.24; C8H18S2Zn another, but which had a similar appearance to those of requires C, 39.41; H, 7.44%. Solid-state 13C{1H} NMR, d 36.96 [MeZnStBu]5. 1H NMR (CD2Cl2), d -0.40 (s, ZnCH3), 1.70 [s, SC(CH3)3], 49.20 (s, SC), 50.08 (s, SC).[s, SC(CH3)3]. [MeCdStBu]n. The gas-phase mixing of stoichiometric Results and discussion quantities of Me2Cd and tBuSH produced a white deposit. The solid-state 13C{1H} NMR again gave broad peaks which Synthesis and characterisation of Zn and Cd chalcogenide cluster complexes merged into one another, but which had a similar appearance to those of [MeZnStBu]5.Solid-state 13C{1H} NMR, d -9.53 Using the system described in the experimental section, the (br, CdCH3), -5.81 (br, CdCH3), 37.44–39.31 [br, gas-phase mixing at room temperature of stoichiometric SC(CH3)3], 48.96–50.29 (br, SC). amounts of Et2Zn and tBuSH (He carrier gas) resulted in the formation of a fine white powder deposit. GC-MS analysis of Solid-state thermal decomposition of Zn chalcogenide cluster the eZuent carrier-gas confirmed that ethane is the only complexes gaseous by-product of this reaction.When Et2Zn was replaced by the deuteriated analogue, (C2D5)2Zn, the ethane produced A similar experimental set-up to that used above in the preparation and collection of prereaction deposits was was pure D5H, formed by protonation of (C2D5)2Zn by the acidic S–H of tBuSH.Upon exposure to air, the white powder employed (Fig. 1), with known amounts of metal chalcogenide cluster complexes contained inside a ceramic crucible placed gave oV a strong smell of thiol, presumably due to hydrolysis rather than from any remaining tBuSH. Without further inside the quartz tube and at the centre of the furnace.A carrier gas (Ar), at just above atmospheric pressure (101.350 purification, solid-state 13C{1H} NMR analysis of the powder gave a spectrum [Fig. 2(a)] which exhibits five diVerent sets of Pa), was passed along the quartz tube at a flow rate of 200 cm3 min-1 while the furnace was heated to a maximum ethyl resonances along with five diVerent tBu resonances. These results confirmed the presence of at least five diVerent temperature of 600 °C.As in the case of the preparation of prereaction deposits, eZuent gas containing unused reactants environments for both Zn and S within the deposit and the probable identity of the complex as the thiolate pentamer, and any volatile products was allowed to pass through neutralising solutions followed by an industrial scrubber before [EtZnStBu]5.A non-quaternary 13C{1H} suppression spectrum (which greatly reduces the intensity of CH and CH2 being released into a fumehood. For specific experiments, eZuent gas containing any volatile products was diverted via resonances) [Fig. 2(b)] shows the methyl signals of the ethyl groups, ZnCH2CH3, to be down field of the methylene, a gas sampling port to the GC-MS.[MeZnStBu]5: weight before thermolysis, 0.52 g; weight of ZnCH2CH3, signals due to the shielding eVect of the zinc. Attempts to establish the structure of the complex by X-ray residue, 0.16 g; 69.2% weight loss. [MeZnStBu(py)]2: weight J. Mater. Chem., 1998, 8, 2769–2776 2771E M E M R E M E M R E M tBu R tBu R tBu R tBu tBu Fig. 3 The basic structure of [RMEtBu]5 (M=Zn, Cd; E=S, Se; R=Me, Et).of the tert-butyl groups to be significantly shifted to higher magnetic field strength relative to the shifts of the other four. This unique resonance probably arises from the tert-butyl group in the ‘handle’ section of the molecule and the shift to higher field is probably due to the sulfur being bonded to two rather than three adjacent zinc atoms (resulting in a higher shielding eVect).Similarly, in [MeZnStBu]5, one of the ZnCH3 carbon resonances is clearly distinct from the others, with a shift this time to lower field, and again, this is probably the resonance of the ZnCH3 group in the ‘handle’ section of the molecule, which is deshielded due to the zinc atom being bonded to only two adjacent sulfur atoms rather than three as for all the other zinc atoms.White powders were also obtained from the gas-phase reactions of Me2Cd with tBuSH and from Me2Zn with tBuSeH. Although both compounds aVord poorly resolved solid-state 13C{1H} NMR spectra, the over-all shapes of the spectra are again similar to that of [MeZnStBu]5, suggesting that cluster complexes of the type [MeMEtBu]n are formed. Although the value of n cannot be established from the spectra alone, it has been proposed from NMR data that n=4 for [MeCdStBu]n when prepared from solution.21,22 As discussed in the Introduction, for MOVPE growth systems employing R2M and H2E (M=Zn, Cd; R=Me, Et; E=S, Se), pyridine can suppress the growth of particulate material by binding to the surface metal atom sites of particles growing within the gas phase.By a similar process, O’Brien and coworkers have reported that the reaction between Me2Zn, tBuSH and pyridine in solution (benzene) aVords the dinuclear Fig. 2 Solid-state 13C{1H} NMR spectra of (a) [EtZnStBu]5, (b) non- pyridine adduct complex, [MeZnStBu(py)]2, rather than the quaternary suppression spectrum of [EtZnStBu]5 and (c) [MeZnStBu]5. pentamer, [MeZnStBu]5, which forms in the absence of pyridine.11 These results confirm that for all MOVPE growth systems crystallography were thwarted due to our inability to grow X- which involve the gas-phase mixing of thiols/selenols with ray quality crystals. However, the solid-state 13C{1H} NMR dialkylzinc/cadmium compounds, a prereaction (detected or data of the ethyl complex suggested the complex to be isostruc- otherwise) will occur in the cold zone of the reactor.This will tural with the methyl pentamer, [MeZnStBu]5, first synthesised result in the depletion (or complete elimination) of the original by Coates and Ridley,20 and more recently structurally charac- precursors from the gas phase before the gas stream carrying terised by O’Brien and coworkers.11 Thus, in an attempt to the precursors reaches the heated substrate.Moreover, confirm the likely structure of the ethyl complex [EtZnStBu]5, Lovergine et al. have reported17 that the addition of triethyl- [MeZnStBu]5 was prepared by the literature method20 and amine to the growth system, when forming ZnS epilayers from characterised by solid-state 13C{1H} NMR spectroscopy. tBuSH and Me2Zn, although avoiding the formation of solid The solid-state 13C{1H} NMR spectrum of [MeZnStBu]5 prereaction deposits, still leads to non-homogeneous epitaxial does indeed correlate well with that of the ethyl complex, growth owing to some form of depletion of the two precursors essentially exhibiting five diVerent tBu resonances and four from the gas phase (other than through the epitaxial growth methyl resonances in a 1525151 ratio [Fig. 2(c)]. Interestingly, process itself ). This observation may be explained by the [MeZnStBu]5 is fluxional in solution (C6D6) with the 1H NMR formation of [MeZnStBu(Et3N)]2 dimers which then either spectrum exhibiting single resonances for all five ZnCH3 precipitate onto the substrate, or initiate gas-phase particulate groups and for all five tBu groups while the 13C{1H} solution growth, with the particulate matter then being precipitated spectrum of [MeZnStBu]5 also exhibits only one ZnCH3 onto the substrate.It is also possible that such dimers are not resonance and one tert-butyl resonance.11 This fluxionality is formed and that Et3N behaves like pyridine in terminating quenched in the solid-state, at least up to 25 °C.Thus, the particulate growth at an early enough stage that prereaction NMR data strongly support the view that in the solid-state deposits are not visible.10 [EtZnStBu]5 and [MeZnStBu]5 are isostructural, as schematically shown in Fig. 3. The structure consists of a cubic Thermogravimetric analysis and solid-state thermal arrangement of zinc and sulfur atoms with one edge of the decomposition of zinc chalcogenide cluster complexes cube broken open by the addition of the extra Zn–S unit, resembling a supermarket trolley.The solid-state 13C{1H} Because of the current interest in the use of metal chalcogenide complexes as ‘single-source’ precursors, both in MOVPE NMR data of both the methyl and ethyl complexes (Fig. 2) show the resonance for one of the quaternary carbon atoms growth and nanoparticle material preparation,18 as well as the 2772 J.Mater. Chem., 1998, 8, 2769–2776Table 1 Thermal analytical and weight loss data for zinc sulfide cluster complexes Step 1a Step 2a Residue (%)b Weight loss (%)b Weight loss (%)b Complex Tonset/ °C Obs. Calc.c Tonset/ °C Obs. Calc.d TGA Thermolysis Calc.e [MeZnStBu]5 110 24 28 230 45 43 31 31 29 [MeZnStBu(py)]5 70 46 51 230 29 29 26 21 20 [EtZnStBu]5 48 46 27f aFrom TGA (thermogravimetric analysis).bWeight losses and residual weights all relative to initial weight of complex. cLoss of half the Zn content as Me2Zn. dLoss of half the S content as tBu2S. eLoss of organics+half the ZnS content. fLoss of organics only, gives a calculated residue of 53%.role that zinc/cadmium chalcogenide cluster complexes might Me2Zn or [Me2Zn(py)2], and a second sharp endotherm with an onset temperature of ca. 230 °C corresponding to the loss play in MOVPE growth whilst employing separate group 12 and 16 precursors, it is important to understand the thermal of ‘tBu2S’. The residue consists of ZnS with a mass corresponding to one half of the total ZnS mass present in the original properties of complexes of the type: [EtZnStBu]5, [MeZnStBu]5 and [MeZnStBu(py)]2.Simultaneous thermal analyses sample. Macroscopic thermolysis also gave residual weights consistent with this sequence of reactions (Table 1). (STA), thermogravimetric analysis–diVerential temperature analysis (TGA–DTA) carried out on [MeZnStBu]5 and GC-MS analysis of the gas-phase products from the macroscopic thermolysis of [MeZnStBu]5, under Ar flow, [MeZnStBu(py)]2, indicates that these complexes demonstrate similar thermal behaviour [Table 1 and Fig. 4(a) and (b)]. confirmed that Me2Zn is released but that tBu2S is not the only S containing volatile decomposition product, 2-methyl- There is a broad endotherm at ca. 100 °C with a weight loss corresponding to the loss of all the zinc bound Me groups, as propene and tBuSH are also produced.Thus, at 100 °C, a trace amount of methane was seen while at 150 °C, traces of methane and 2-methylpropene along with large quantities of Me2Zn were detected. At 200 °C, large quantities of 2-methylpropene and tBuSH were detected along with traces of tBu2S. For the complex [MeZnStBu(py)]2, at 100 °C, a large quantity of pyridine was detected along with trace amounts of methane and toluene (resulting from the preparative method), while at 150 °C, pyridine, methane and 2-methylpropene (Me2Zn was not detected, see below) were observed, and above 200 °C, 2- methylpropene along with trace amounts of tBu2S and methane were seen.These results for [MeZnStBu(py)]2 do not demonstrate the production of Me2Zn.We propose that this is due to the presence of pyridine which condenses on the walls of the stainless steel tubing used to connect the thermolysis cell to the gas-sampling valve of the GC-MS. This pyridine will complex with any Me2Zn released to the gas-phase, forming the solid adduct complex, [Me2Zn(py)2]. At room temperature, dissociation and revapourisation of [Me2Zn(py)2] is at such a slow rate that gas-phase concentrations of Me2Zn are too low to detect.To confirm that [MeZnStBu(py)]2 decomposes by a similar mechanistic pathway to that of [MeZnStBu]5, [MeZnStBu(py)]2 was heated in the solid state under vacuum to 100 °C. The volatile products were collected in a cold trap and at room temperature gave a viscous liquid, which was shown by 1H NMR to contain Me2Zn and pyridine.Elemental analysis on the remaining solid, without washing or further purification, suggested it to be the complex [Zn(StBu)2]n. This was confirmed by solid-state 13C{1H} NMR as shown in Fig. 5(a). The 13C{1H} NMR of the original complex, [MeZnStBu(py)]2, is shown in Fig. 5(b) for comparison. The complex, [Zn(StBu)2)]n, shows two peaks of almost equal intensity due to the SCtBu carbons showing that both syn and anti orientations of the tBu groups occur with equal probability.Many complexes of this general formula, M(ER)2, including [Zn(StBu)2)]n,23f have previously been prepared and used as potential ‘single-source’ precursors for II–VI materials.23 It has been proposed by Steigerwald and Sprinkle that the use of Me2Cd and R2Te under MOVPE growth conditions results in the in situ formation Cd(TeR)2 which then decomposes to give CdTe.23h This could also be occurring in the growth of ZnS from Me2Zn and tBuSH.Although a detailed mechanistic pathway cannot be fully Fig. 4 Simultaneous thermal analyses (STA, thermogravimetric established it appears, from these results, that the overall analysis—diVerential temperature analysis (TGA–DTA)) of (a) [MeZnStBu]5, (b) [MeZnStBu(py)]2 and (c) [EtZnStBu]5. pathway for the thermal decomposition of [MeZnStBu]5 is as J.Mater. Chem., 1998, 8, 2769–2776 2773Fig. 6 X-Ray powder diVraction pattern of ZnS formed from the thermolysis of (a) [EtZnStBu]5, (b) [MeZnStBu]5 and (c) [MeZnStBu(py)]2. Standard X-ray powder diVraction patterns of (d) hexagonal (wurtzite) and (e) cubic (sphalerite) phases of ZnS are also shown.The peaks at 2h 21 and 24° arise from the Vaseline support. Fig. 5 Solid-state 13C{1H} NMR spectra of (a) the residue remaining after heating [MeZnStBu(py)]2 to 100 °C and (b) [MeZnStBu(py)]2 (the peaks in the range d 55–85 and d 190–220 are spinning side abstraction from the tert-butyl group to form 2-methylpropene bands of the main pyridine resonances at d 120–155).and [EtZnSH] with reductive elimination from the latter aVording the observed ethane and ZnS. Compounds similar to [EtZnSH] are known to give alkane and ZnS upon thermal shown in Scheme 1. This involves a rearrangement of the decomposition. Why this process should occur for [EtZnStBu]5 methyl radicals to liberate the volatile compound Me2Zn, rather than for [MeZnStBu]5 is not clear but demonstrates leaving the zinc-bis(thiolate), [Zn(StBu)2]n, as a residue.how a small diVerence within the precursor can radically alter Further heating leads to the loss of tert-butyl groups, mainly the pathway by which decomposition takes place. as 2-methylpropene and tBuSH along with far smaller amounts of tBu2S.H2S might be expected from the decomposition of Analysis of the solid thermolysis residues tBuSH, but we have shown that the extent of decomposition The macroscopic thermolysis experiments described above of tBuSH is very low below 300 °C.24 For [MeZnStBu(py)]2, produced grey or white residues, which were scraped from the a similar process occurs initially to give [Me2Zn(py)2] and pyrolysis tube and handled in air before being analysed by [Zn(StBu)2]n.powder X-ray diVraction (PXRD) and transmission electron The STA analysis of [EtZnStBu]5 is rather diVerent from microscopy (TEM). PXRD analyses of all the residues show those of the other compounds, since there is no evidence for broad peaks (Fig. 6) and although on first inspection the formation of [Zn(StBu)2]n (no sharp exotherm near 230 °C) diVraction patterns appear to resemble the PXRD of cubic and the decomposition occurs in three stages [Fig. 4(c)]. The (sphalerite) ZnS, previous studies of ours10,25 along with the residual mass obtained from both STA and macroscopic calculations of others26 suggest the residues are more likely to thermolysis (ca. 47%) is between those expected for the simple consist mainly of nanoparticles (responsible for the line broad- loss of the organic moieties (53.1%) and for the type of ening) of distorted hexagonal (wurtzite) phase ZnS (having a mechanism described above for [MeZnStBu]5 (26.6%). In the high number of stacking faults and leading to peak dependent thermolysis of [EtZnStBu]5, below 100 °C, trace amounts of line broadening).However, the observation of a peak at 2h= ethane were detected, at 150 °C, ethane along with larger 33° does suggest that some cubic phase is also present. The amounts of 2-methylpropene were seen, while above 200 °C, diVraction pattern of the residue from [MeZnStBu]5 [Fig. 6(b)] large amounts of 2-methylpropene along with trace amounts is somewhat broader than the others suggesting that the of tBu2S were detected.Although trace amounts of compounds average particle size is smaller. containing sulfur were detected, the mechanistic pathway of TEM images for the residue obtained from [EtZnStBu]5 thermal decomposition involves no major loss of zinc or sulfur. [Fig. 7(a)] show the ZnS to be in the form of circular agglomer- Thus, the decomposition pathway for [EtZnStBu]5 is rather ates 0.5–2 mm in diameter made up of individual ZnS nanopart- diVerent from that for [MeZnStBu]5 in that Et2Zn is not icles in the size range of 5-10 nm, whilst those from observed as a product, but rather ethane and 2-methylpropene [MeZnStBu]5 appear as laths [Fig. 7(b)], which probably arise are the major products. One possible decomposition route as a function of sample preparation.Individual crystals which explains the volatile product formation involves b-H within the aggregates are in the size range 5–20 nm. Although high resolution TEM images of the residue formed from [MeZnStBu]5 [Fig. 7(c)] reveal many of the nanocrystals to be twinned (or to be composed of a number of subcrystals), there are also a number of perfect single crystals within the residue.The particles produced on thermolysis by all three complexes are somewhat diVerent from those formed during gas-phase reactions between R2M and H2S (M=Zn, Cd) where in [MeZnStBu]5 2.5 [Zn(StBu)2] n + 2.5 Me2Zn T > 250 °C 2.5 ZnS + 2.5 tBuSH + 2.5 T = 100 °C general, and especially upon the addition of pyridine, more Scheme 1 Schematic representation of the processes occurring in the solid-state thermolysis of [MeZnStBu]5.homogeneous nanoparticles are formed.10 The particles 2774 J. Mater. Chem., 1998, 8, 2769–2776conditions used, the decomposition of the involatile chalcogenide cluster complexes proceeds at higher temperatures (>200 °C) to give agglomerates of nanometer sized particles of metal chalcogenides.However, the pathway by which decomposition occurs seems to be highly dependent upon the alkyl group within the chalcogenide cluster complex. Acknowledgements We thank the EPSRC for financial support (N.L.P., via a ROPA award, S.L. and D.F.F.) and Ciba Specialty Chemicals for a studentship (W.G.T.). References 1 (a) S. Fujiita, Y. Matsuda and A. Sasaki, J.Cryst. Growth, 1984, 68, 231; (b) A. Yoshikawa, S. Muto, S. Yamaga and H. Kasai, J. Cryst. Growth, 1988, 86, 279; (c) S. Yamaga, A. Yoshikawa and H. Kasai, J. Cryst. Growth, 1988, 86, 252. 2 P. J. Wright, B. Cockayne, A. J. Williams, A. C. Jones and E. D. Orrell, J. Cryst. Growth, 1987, 84, 552. 3 M. J. Almond, M. P. Beer, M. G. B. Drew and D. A. Rice, J. Organomet. Chem., 1991, 421, 129. 4 B. Cockayne, P. J.Wright, A. J. Armstrong, A. C. Jones and E. D. Orrell, J. Cryst. Growth, 1988, 91, 57. 5 P. J. Wright, P. J. Parbrook, B. Cockayne, A. C. Jones, E. D. Orrell, K. P. O’Donnell and B. Henderson, J. Cryst. Growth, 1989, 94, 441. 6 P. J.Wright, B. Cockayne, P. J. Parbrook, A. C. Jones, P. O’Brien and J. R.Walsh, J. Cryst. Growth, 1990, 104, 601. 7 M. J. Almond, M.P. Beer, K. Hagen, D. A. Rice and P. J.Wright, J. Mater. Chem., 1991, 1, 1065. 8 O. Briot, M. DiBlasio, T. Cloitre, N. Briot, P. Bigenwald, B. Gil, M. Averous, R. L. Aulombard, L. M. Smith, S. A. Rushworth and A. C. Jones, Mater. Res. Soc. Symp. Proc., 1994, 340, 515. 9 P. J. Wright, B. Cockayne, P. J. Parbrook, P. E. Oliver and A. C. Jones, J. Cryst. Growth, 1991, 108, 525. 10 (a) N. L. Pickett, D. F. Foster and D. J. Cole-Hamilton, J. Mater. Chem., 1996, 6, 507; (b) N. L. Pickett, D. F. Foster and D. J. Cole- Hamilton, J. Cryst. Growth, 1997, 170, 476; (c) N. L. Pickett, F. G. Riddell, D. F. Foster, D. J. Cole-Hamilton and J. R. Fryer, J. Mater. Chem., 1997, 7, 1855. 11 M. A. Malik, M. Motevalli, J. R. Walsh, P. O’Brien and A. C. Jones, J. Mater. Chem., 1995, 5, 731. 12 T. Obinata, K. Uesugi, G. Sato, I. Suemune, H. Machida and N. Shimoyama, Jpn. J. Appl. Phys., 1995, 34, 4143. 13 (a) W. Kuhn, A. Naumov, H. Stanzl, S. Bauer, K. Wolf, H. P. Wagner, W. Gebhardt, U. W. Pohl, A. Krost, W. Richter, U. Dumichen and K. H. Thiele, J. Cryst. Growth, 1992, 123, 605. 14 (a) D. N. Armitage, H. M. Yates, J. O. Williams, D. J. Cole- Hamilton and I.L. J. Patterson, Adv. Mater. Opt. Electron., 1992, 1, 43; (b) D. F. Foster, I. L. J. Patterson, L. D. James, D. J. Cole- Hamilton, D. N. Armitage, H. M. Yates, A. C. Wright and J. O. Williams, Adv. Mater. Opt. Electron., 1994, 3, 163. 15 K. Nishimura, K. Sakai, Y. Nagao and T. Ezaki, J. Cryst. Growth, 1992, 117, 119. 16 K. Nishimura, Y. Nagao and K. Sakai, J. Cryst. Growth, 1993, Fig. 7 TEM images of (a) ZnS from the thermolysis of [EtZnStBu]5, 134, 293. (b) ZnS from the thermolysis of [MeZnStBu]5 and (c) as (b) but at 17 N. Lovergine, M. Longo, C. Gerardi, D. Manno, A. M. Mancini higher magnification. and L. Vasanelli, J. Cryst. Growth, 1995, 156, 45. 18 (a) P. O’Brien and R. Nomura, J. Mater. Chem., 1995, 5, 1761; (b) P. O’Brien, Precursors for Electronic Materials, in Inorganic produced from these chalcogenide cluster complexes are similar Materials, ed.D. W. Bruce and D. O’Hare, John Wiley, London, to, but larger than, those reported by Steigerwald and 1992, p. 499; (c) M. Bochmann, Chem. Vap. Deposition, 1996, 2, 85. coworkers from the solid-state thermolysis of Zn(SPh)2 and 19 D. F. Foster and D. J. Cole-Hamilton, Inorg. Synth., 1997, 31, 29. [Zn(SPh)2(depe)] [depe=1,2-bis(diethylphosphino)ethane].23i 20 G. E. Coates and D. Ridley, J. Chem. Soc., 1965, 1870. 21 G. E. Coates and A. Lauder, J. Chem. Soc. A, 1966, 264. 22 J. D. Kennedy and W. McFarlane, J. Chem. Soc., Perkin Trans. 2, Conclusion 1977, 1187. Gas-phase mixing of thiols/selenols (REH) with dialkylzinc/ 23 (a) M. Bochmann, K. J. Webb, M. Harman and M. B. Hursthouse, Angew. Chem., Int. Ed. Engl., 1990, 29, 638; cadmium compounds (R2M) results in the precipitation of (b)M. Bochmann, K. J.Webb, J. E. Hails and D.Wolverson, Eur. zinc/cadmium chalcogenide cluster complexes of general J. Solid State Inorg. Chem., 1992, 29, 155; (c) G. Krau� ter and formula [tBuEMR]n. This demonstrates the importance of W. S. Rees, Jr., J. Mater. Chem., 1995, 5, 1265; (d)M. Bochmann, avoiding the low temperature premixing of such precursor A. P. Coleman and A. K. Powell, Polyhedron, 1992, 11, 507; combinations in the growth of II–VI materials by MOVPE, (e) J. G. Brennan, T. Siegrist, P. J. Carroll, S. M. Stuczynski, as this can lead to gas-phase depletion of precursors and non- L. E. Brus and M. L. Steigerwald, J. Am. Chem. Soc., 1989, 111, 4141; ( f )W. S. Rees, Jr. and G. Krau�ter, J. Mater. Res., 1996, 11, uniformity in the epitaxial growth. Under the reaction J. Mater. Chem., 1998, 8, 2769–2776 27753005; (g) R. D. Schluter, G. Krau� ter and W. S. Rees, Jr., J. Cluster 25 S. W. Haggata, X. Li, D. J. Cole-Hamiltion and J. R. Fryer, J. Mater. Chem., 1996, 6, 1771. Sci., 1997, 8, 123; (h)M. L. Steigerwald and C. R. Sprinkle, J. Am. Chem. Soc., 1987, 109, 7200; (i) J. G. Brennan, T. Siegrist, 26 M. G. Bawendi, A. R. Kortan, M. L. Steigerwald and L. E. Brus, J. Chem. Phys., 1989, 91, 7282. P. J. Carroll, S. M. Stuczynski, P. Reynders, L. E. Brus and M. L. Steigerwald, Chem.Mater., 1990, 2, 403. 24 N. L. Pickett, D. F. Foster, D. Ellis and D. J. Cole-Hamilton, J. Mater. Chem., to be submitted. Paper 8/06421K 2776 J. Mater. Chem., 1998, 8, 2769&ndas

ISSN:0959-9428    

DOI:10.1039/a806421k

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Lithium insertion in two tetragonal tungsten bronze type phases, M8W9O47 (M=Nb and Ta) Sagrario M. Montemayor, A. Alvarez Mendez, A. Martý�nez-de la Cruz, Antonio F. Fuentes* and Leticia M. Torres-Martý�nez Facultad de Ciencias Quý�micas, Divisio�n de Estudios Superiores, Universidad Auto�noma de Nuevo Leo�n, Apartado Postal 1625, Monterrey, Nuevo Leo�n, Mexico.E-mail: afernand@ccr.dsi.uanl.mx Received 10th June 1998, Accepted 21st September 1998 A study of lithium insertion in two tetragonal tungsten bronze (TTB) type phases of general formula M8W9O47 (M=Nb and Ta), is presented. The electrochemical insertion of up to 20 lithium atoms per formula unit in Nb8W9O47 (Li/SM=1.2) proceeds through a reversible reaction with several single phase and one two-phase domains, while in Ta8W9O47 the reversibility of lithium insertion is limited to 15 atoms (Li/SM=0.9).Structural changes on Nb8W9O47 as a function of the number of lithium atoms inserted have been studied by X-ray powder diVraction. dal coordination are preferentially occupied by niobium while Introduction in the corner sharing metal–oxygen octahedra, metal atoms The study of the chemistry of tungsten bronzes and related are on average (0.4 Nb+0.6 W).phases has attracted considerable attention because of their While only one polymorph for Nb8W9O47 has been reported potential use as electrodes, catalysts and in optical displays. to exist, two have been found for Ta8W9O47, a low (tetragonal ) Additionally, these phases were found to present interesting and a high temperature form (orthorhombic), both exhibiting optical and ferroelectric properties.In this context and continu- TTB type structures, although only in the high temperature ing with work started recently in our research group on one (1400 °C) is it possible to observe threefold TTB-type insertion reactions in some niobium–tungsten mixed oxides,1–3 superstructure.8 A similar cation distribution to that found a study of lithium insertion in two tetragonal tungsten bronze for the niobium–tungsten mixed phases was also proposed (TTB) type phases of general formula M8W9O47 (M=Nb and here.It is important to point out that although M8W9O47 is Ta), has been carried out. the ideal composition, a small range of non-stoichiometry may From the solid state chemistry point of view, the exist in both niobium and tantalum TTB type phases, due to Nb2O5–WO3 and Ta2O5–WO3 systems behave in a fairly disorder in the tunnels occupancy.similar way in the region rich in WO3. Phases containing Compounds presenting structures containing PCs as main between 50 and 80 wt.% of WO3, present a structure similar building units, such as W18O49, have been already studied as to the so-called tetragonal tungsten bronzes (TTB).These host materials for reversible lithium insertion reactions.9 In phases consist in a framework of MO6 octahedra sharing this work a study of the electrochemical lithium insertion in corners, linked in such a way that three, four- and five-sided M8W9O47 has been carried out.In order to evaluate the tunnels are formed. Decreasing oxygen-to-transition metal influence of lithium insertion on the structure of the parent ratio (3 in WO3 vs. 2.76 in M8W9O47) is achieved in these oxide, some lithiated phases were synthesised by indirect TTB type phases by filling a certain number of five-sided chemical reaction and characterised by X-ray powder diVractunnels with oxygen and metal atoms thus forming the so- tion experiments.As the main diVerence between both mixed called pentagonal columns (PCs for short):4aMO7 pentagonal tungsten oxides is the presence of Nb5+ instead of Ta5+, it bipyramid sharing equatorial edges with five MO6 octahedra. would be interesting to compare their behaviour versus lithium In Nb8W9O47 (Fig. 1) one third of the pentagonal tunnels are insertion reactions. No attempt is made in this work on testing M8W9O47 as cathode materials in lithium ion batteries.occupied in this way.5–7 Cation sites with pentagonal bipyrami- Fig. 1 Idealised structure of Nb8W9O47. When joining the diVerent polyhedra, five-, four- and three-sided tunnels are formed where additional ions can be inserted.J. Mater. Chem., 1998, 8, 2777–2781 2777the time of observation, intercalated samples were covered by Experimental a polyethylene film while recording data on the diVractometer Preparation and characterisation of pristine materials, to prevent sample oxidation. M8W9O47 The tantalum–tungsten and niobium–tungsten mixed oxides Results and discussion used in this study were synthesised by solid state reaction.The Electrochemical lithium insertion in Nb8W9O47 starting materials, WO3 (Aldrich Chem. Co, 99+%), Nb2O5 (Alfa Products 99.5%) and Ta2O5 (Aldrich Chem. Co, 99.99%), Electrochemical lithium insertion in Nb8W9O47 was carried were weighed in the appropriate stoichiometric ratio (4 out by discharging a cell with the following configuration: M2O559 WO3) and thoroughly mixed by grinding in an agate Li|1 mol dm-3 LiClO4 in (50% DEE+50% EC)|Nb8W9O47 mortar using Analar grade acetone.Powders were then pressed into 10 mm diameter pellets, placed in a platinum crucible and Results obtained from SPECS experiments run down to 1.1 V fired at 1250 °C in an electrical furnace. During firing samples vs. Li+/Li using ±10 mV (2 h)-1 potential steps, are shown were periodically extracted from the furnace and ground to in Fig. 2 as an E vs.composition (x in LixNb8W9O47) plot. favour reaction before being finally quenched at room tempera- As can be seen in this graph, Nb8W9O47 can reversibly ture. Phase identification was carried out by X-ray powder incorporate up to 20 lithium atoms per formula unit which diVraction in a Siemens D-5000 diVractometer using Cu-Ka correspond to about 1.2 Li per metal atom.Almost all lithium radiation (l=1.5418 A° ). A typical diVraction experiment for atoms inserted are removed after completing a charge– determining cell parameters was run with a step size of discharge cycle. The main features of this plot are two plateaux, 0.12° min-1 using KCl as internal standard.A and B, of approximately constant E values around 2.1 and 1.7 V vs. Li+/Li separating three regions where a continuous variation of E with composition is observed. In order to Electrochemical lithium insertion determine the existence of continuous transformations or to Electrochemical experiments were carried out with a multi- multiphase regions, potentiostatic experiments were analysed channel potentiostatic-galvanostatic system MacPile II,10 using in more detail.Representing these results as DQ/m vs. E, Fig. 3, a SwagelokTM type cell11 with metal lithium acting simul- it is possible to see that the electrochemical lithium insertion taneously as negative and reference electrode. Positive elec- in this material goes really through three very well defined trodes were prepared by mixing the phase being tested, reduction steps labelled here as A, B and C.The electrochemical Nb8W9O47 or Ta8W9O47, with carbon black and a binder, lithium extraction from LixNb8W9O47 follows a similar mech- (0.5% ethylene–propylene–diene terpolymer, EPDT, in cyclo- anism and three analogous steps, labelled as A¾, B¾ and C¾, are hexane) either in a 8951051 ratio (wt.%) for niobium contain- also present.ing samples or in a 5954051 ratio for the tantalum compound It was observed that the nature of the first step ( labelled as (this material is not a good electronic conductor so a larger A and A¾) was always better defined on oxidation and second proportion of carbon has to be used). The electrolyte used was a 1 mol dm-3 solution of LiClO4 in a previously dried 50550 mixture of ethylene carbonate (EC) and diethoxyethane (DEE).Cell assemblage was carried out in a MBraun glove box under an argon atmosphere with continuous purge of water vapour and oxygen ensuring an inside concentration for both compounds of <1.5 ppm. Two diVerent electrochemical experiments were carried out on these in a current or in a potential-controlled mode.Potentiodynamic titrations were carried out by a stepwise technique also known as step potential electrochemical spectroscopy, SPECS.12 In this technique, the potential is stepwise increased or decreased while recording charge increments vs. time at each potential level and therefore allowing the study of insertion reaction kinetics.The simultaneous determination of incremental capacities and Fig. 2 Evolution of cell voltage versus composition (x in observation of the kinetics of the redox process helps in LixNb8W9O47) obtained during a SPECS experiment run at determining the succession of single phase and two phase ±10 mV (2 h)-1 potential steps showing a complete charge– domains which might take place on insertion/deinsertion.discharge cycle. Typical experimental conditions were set at ±10 mV (2 h)-1 potential steps with a charge recording resolution of 5 mA h and a strict temperature control. Lithium chemical insertion Lithium insertion was also carried out by chemical reaction of Nb8W9O47 with an appropriate reducing agent, n-butyllithium in n-hexane. A given volume of this reagent (1.6 mol dm-3) was added to a known weight of host material previously placed in a glass reaction vessel.The mixture was stirred for 7 days. Reaction products were thoroughly washed with n-hexane, dried and stored in the glove box. Lithium content in LixNb8W9O47 was analysed by AAS using a Varian SpectrAA5 spectrometer after extraction of the inserted ions using concentrated nitric acid.Phase identification was carried Fig. 3 DQ/m vs. E plot obtained from a SPECS experiment run at out by X-ray powder diVraction as described above. Although ±10 mV (2 h)-1 potential steps showing the diVerent reduction and oxidation steps observed on lithium insertion and extraction. lithium inserted Nb8W9O47 proved to be stable at least during 2778 J. Mater. Chem., 1998, 8, 2777–2781Fig. 6 I/m and E vs.time plot obtained when crossing at Fig. 4 Evolution of cell current (,,) and voltage (———) with time during a SPECS experiment run using+10 mV (2 h)-1 potential steps. -10 mV (12 h)-1 potential steps the voltage region where peak B was observed. reduction than on the first discharge. Therefore, we shall start of the amount of lithium atoms inserted.14 A similar behaviour examining a chronoamperogram obtained when crossing at of the current decay versus time is observed on the third +10 mV (2 h)-1 potential steps, the potential region where the reduction peak labelled in Fig. 3 as C making for a second oxidation peak A¾ was observed, Fig. 4. The I vs. time plot continuous transformation between 1.3 and 1.6 V vs. Li+/Li shows a profile obviously not governed by a simple diVusion on lithium insertion in Nb8W9O47.process which would give a monotonic tendency towards I= A confirmation of these observations is presented in Fig. 7 0, but it is typical of a first order transition. In this type of which shows a comparison between two voltamperograms process, the response to a potential step will depend on the obtained when running SPECS experiments ±10 mV (2 h)-1 mobility of the interface between the two phases, relative to (filled circles) and ±10 mV (12 h)-1 (open squares) potential lithium diVusivity in both phases and on the kinetics of lithium steps.While steps labelled as A and A¾ show hysteresis in both transfer at the interface with the electrolyte assuming that experiments (initial slopes can be aligned) independently of there is no electronic conductivity limitation.13 Thus, it could the voltage scanning rate used, confirming our previous be assumed that the first reduction step A which corresponds assumption of dealing with a first order phase transition at to the first plateaux (at higher E values) also labelled as A in this potential range, a diVerent situation is observed for steps Fig. 2, can be associated with a multiphase domain separating labelled B and C. In these latter cases and for the experiment two solid solution regions, I and II. The phase I–phase II run at ±10 mV (12 h)-1, hardly any hysteresis eVect is noticed equilibrium potential can be given as the intercept at the slope between reduction and oxidation peaks. On the contrary, extrapolations at zero current, i.e. 2.12 V. Fig. 5 shows the peaks B¾ and C¾ are now found located on top of peaks B and evolution of the cell voltage versus composition obtained when C supporting the idea of crossing two consecutive continuous cycling a similar cell to the one described elsewhere in this phase transitions between 1.8 and 1.3 V vs. Li+/Li. Therefore, work, between 3 and 1.75 V vs.Li+/Li. As can be seen in this a complete phase diagram Li–Nb8W9O47 can be now proposed graph, phase I transformation to phase II is a reversible process as follows: above 2.12 V vs. Li+/Li, a solid solution domain with all the atoms initially incorporated being extracted after (I) from x=0 to x#2. At 2.12 V vs. Li+/Li, a two phase completing a charge–discharge cycle.Since potentiostatic equilibrium between phase I and phase II. Below 2.12 V vs. experiments run at ±10 mV (2 h)-1 potential steps did not Li+/Li, a solid solution domain (II) from x=6. Around 1.7 V reveal unambiguously the exact nature of reduction steps vs. Li+/Li, an incremental capacity peak probably due to a labelled in Fig. 3 as B and C and their corresponding oxidation tendency to local ordering in phase II around 11<x<12.steps B¾ and C¾, it was necessary to run experiments at much Below 1.7 V vs. Li+/Li, continuation of phase II with another lower scan rates [±10 mV (12 h)-1]. Time dependence of the possible tendency to local ordering giving rise to an increase current when crossing at -10 mV (12 h)-1 potential steps, the of the incremental capacity of the cell around 1.4–1.5 V vs. voltage region where the reduction peak labelled in Fig. 3 as Li+/Li. B was found, is shown in Fig. 6. In this case, the monotonic tendency towards I=0 observed in the current decay vs. time plot, can be associated with a continuous transformation. In homogeneous solutions and assuming that there is no ion–ion interaction, the system behaves in a similar way independently Fig. 7 Comparison between two DQ/m vs. E plots obtained Fig. 5 E (V vs. Li+/Li) vs. lithium content (x) in LixNb8W9O47 plot when running SPECS experiments at ±10 mV (2 h)-1 and ±10 mV (12 h)-1 potential steps. obtained when cycling a cell within a limited potential window. J. Mater. Chem., 1998, 8, 2777–2781 2779Fig. 8 E vs. lithium content (x) in Lix Nb8W9O47 plot obtained after completing three charge–discharge cycles.Fig. 10 E vs. lithium content (x) in LixTa8W9O47 plot obtained from Lithium insertion reaction reversibility in this material can a SPECS experiment run at ±10 mV (2 h)-1 potential steps showing be better appreciated in Fig. 8 where results obtained after two complete charge–discharge cycles within limited potential completing three charge–discharge cycles in a cell configured windows.as mentioned above are presented. Cycling is carried out with minimal capacity losses. Data shown were obtained from a galvanostatic experiment run at a cycling rate greater than plateaux indicating multiphase regions. As can be seen in this C/76 (C/76 corresponds to a charge or a discharge within 76 h).plot, Ta8W9O47 incorporates a larger number of lithium atoms According to the structure shown in Fig. 1, a total of 26 per formula unit (27, Li/SM=1.6) than its niobium analogue diVerent tunnels are found in a unit cell of this phase which (20 Li/formula unit) between 3.1 and 1.1 V. However, almost will also contain two formula units of Nb8W9O47. These 17% of those atoms remained in the structure after completing cavities are distributed as follows: 8 five-sided tunnels; 6 four- a charge–discharge cycle showing therefore the existence of sided and 12 three-sided channels.In order to be able to larger structural changes on insertion. Fig. 10 shows the evolaccommodate 40 lithium atoms (20 per formula unit) in 26 ution of the cell voltage versus composition for a similar cell tunnels of a unit cell, multiple lithium occupancy has to be to the one described above but cycled up to 1.85 (open dots) assumed in some cavities (ratio of lithium atoms intercalated and 1.35 volts (solid dots) showing this time that the insertion in a unit cell to available cavities =1.54). Based on tunnel of up to 15 lithium atoms in Ta8W9O47 (Li/SM=0.9), is a size, it is most likely that only one lithium atom would enter reversible reaction. into three-sided tunnels.This assumption would leave 28 Both ions Nb5+ and Ta5+ have similar ionic radii (0.64 A ° , lithium atoms to be distributed in 14 tunnels which makes a six-coordinate),15 but contain a diVerent number of electrons lithium atoms to available cavities ratio of 2.Thus, It is (36 and 68 respectively). Therefore, Ta5+ has a larger electronic reasonable to suggest that two lithium atoms would enter in density than Nb5+ which makes lithium diVusion more diYcult each of the four (coordination number 12) and five-sided in Ta8W9O47 relative to that in Nb8W9O47 and proving that tunnels (coordination number 15). With the available data it these metals play an important role in the insertion reaction.is not possible to give a filling sequence of these cavities. An attempt was made to prepare lithiated phases by chemical reaction with n-butyllithium but these proved to be extremely Electrochemical lithium insertion in Ta8W9O47 unstable upon exposure to the atmosphere. A preliminary study of the electrochemical lithium insertion Chemical lithium insertion in Nb8W9O47 in Ta8W9O47, an isostructural phase with Nb8W9O47 containing tantalum instead of niobium, was carried out by discharg- In order to study the influence of lithium insertion on the ing a similar cell to the one described above. Results obtained structure of the parent oxide, diVerent inserted compositions from a SPECS experiment run at ±10 mV (2 h)-1 potential included in the solid solution regions detected during the steps, are shown in Fig. 9 as the evolution of cell voltage electrochemical study, were prepared by chemical reaction of versus composition (x in LixTa8W9O47). Only slight slope Nb8W9O47 and n-butyllithium. Inserted materials prepared changes are observed with apparently no existence of voltage were characterised as previously described and resulted in the following compositions: Li1.8Nb8W9O47, Li7.3Nb8W9O47 and Li18.8Nb8W9O47. Phase characterisation was carried out by Xray powder diVraction using KCl as internal standard.Powder patterns, which are shown in Fig. 11, were indexed using the same orthorhombic cell described in the literature for the starting material by using a least-squares cell refinement program.Results are shown in Table 1. Powder patterns showed a gradual change in position and intensity of some reflections as the insertion reaction proceeded which can be related with small displacements of the atoms in Nb8W9O47 in order to accommodate the lithium ions. The most important changes are observed between diVraction patterns b and c which correspond to phase I transition to phase II.A colour change for these materials was observed as the number of lithium atoms incorporated increased, going from pale yellow for the pristine phase to blue developing to Fig. 9 Evolution of cell voltage vs. composition (x in LixTa8W9O47) black for phase II which corresponds to an increase in the obtained during a SPECS experiment run at±10 mV (2 h)-1 potential steps showing a complete charge–discharge cycle.electronic conductivity of Nb8W9O47 as the insertion reaction 2780 J. Mater. Chem., 1998, 8, 2777–2781(11<x<12) and between 1.40 and 1.50 V probably indicate the existence of a tendency to local ordering in phase II. These results are supported by X-ray diVraction data of inserted materials which show important changes on the transition from phase I to II with additional variations as the number of lithium atoms inserted increased.In order to accommodate such a number of lithium atoms and taking into account the number of tunnels available per unit cell, multiple lithium occupancy has to be considered in some cavities. Electrochemical lithium insertion was also studied in Ta8W9O47 where the reaction seems to follow a diVerent mechanism confirming that the second transition metal in these TTB type structures (Nb and Ta) plays an important role in the insertion reaction.Lithium insertion reversibility in Ta8W9O47 was observed for a smaller number of atoms than its niobium analogue (Li/SM=0.9). Acknowledgements Fig. 11 X-Ray powder diVraction patterns of (a) Nb8W9O47, (b) Li1.8Nb8W9O47, (c) Li7.3Nb8W9O47 and (d) Li18.8Nb8W9O47. Financial support from CONACYT (Project 3862P-A9607) is gratefully acknowledged. The authors are also especially grate- Table 1 Cell parameters and cell volume values obtained from X-ray ful to Dr.Yves Chabre for useful discussions and his critical powder diVraction data for LixNb8W9O47 remarks on the original manuscript.Compound a/A° b/A° c/A° Cell volume/A° 3 References W9Nb8O47 a 36.692 12.191 3.945 1764.65 1 A. F. Fuentes, A. Martý�nez de la Cruz and L. M. Torres-Martinez, W9Nb8O47 36.675(8) 12.186(2) 3.9447(8) 1763.0±0.7 Solid State Ionics, 1996, 92, 103. Li1.8Nb8W9O47 36.71(1) 12.201(4) 3.942(1) 1766.1±0.9 2 A. F. Fuentes, E. Briones Garza, A. Martý�nez de la Cruz and Li7.3Nb8W9O47 36.9(1) 12.29(3) 3.93(1) 1787.8±9.0 L.M. Torres-Martý�nez, Solid State Ionics, 1997, 93, 245. Li18.8Nb8W9O47 36.9(1) 12.28(5) 3.94(1) 1790.0±12.7 3 A. F. Fuentes, A. Martý�nez de la Cruz and L. M. Torres-Martinez, aRef. 5. Mater. Res. Soc. Symp. Proc., 1997, 453, 659. 4 M. Lundberg, Chem. Commun. Univ. Stockholm, 1971, No. XII. 5 R. S. Roth and J. L. Waring, J. Res. Nat. Bur.Stand., Sect. A, 1966, 70, 281. proceeds. Powder patterns obtained after Li removal were 6 A. W. Sleight, Acta Chem. Scand., 1966, 20, 1102. very similar to those of the pristine phase confirming the 7 D. C. Craig and N. C. Stephenson, Acta Crystallogr., Sect. B, reversibility of this reaction. 1969, 25, 2071. 8 F. Krumeich and T. Geipel, J. Solid State Chem., 1996, 124, 58. 9 A. Martý�nez de la Cruz, F. Garcý�a-Alvarado, E. Mora�n, Conclusions M. A. Alario-Franco and L. M. Torres-Martý�nez, J. Mater. Chem., 1995, 5, 513. A study of lithium insertion in M8W9O47 (M=Nb and Ta), a 10 C. Mouget and Y. Chabre, Multichannel Potentiostatic and tetragonal tungsten bronze type phase (TTB) has been carried Galvanostatic System MacPile, Licensed from CNRS and UJF out. The electrochemical insertion of up to 20 lithium atoms Grenoble to Bio-Logic Corp., 1 Av. de l’Europe, F-38640, Claix, France. per formula unit of Nb8W9O47 (Li/SM=1.2) is a reversible 11 J. M. Tarascon, J. Electrochem. Soc., 1985, 132, 2089. reaction which proceeds through three reduction steps: a first 12 Y. Chabre, J. Electrochem. Soc., 1991, 138, 329. order transition at around 2.12 V and two continuous trans- 13 Y. Chabre and J. Pannetier, Prog. Solid State Chem., 1995, 23, 1. formations at around 1.70 and between 1.40 and 1.50 V vs. 14 C. J. Wen, B. A. Boukamp, R. A. Huggins and W. J. Weppner, Li+/Li. These transformations originated at least two solid J. Electrochem. Soc., 1979, 126, 2258. solution regions of general formula LixNb8W9O47 with the 15 R. D. Shannon, Acta Crystallogr., Sect. A, 1976, 32, 751. following approximate composition limits: I: 0x2; II: 6x20. Two incremental capacity peaks observed at 1.70 V Paper 8/04410D J. Mater. Chem., 1998, 8, 2777

ISSN:0959-9428    

DOI:10.1039/a804410d

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Lithium loss kinetics from polycrystalline LixNi1-xO at high temperatures Ermete Antolini ENEA C.R. Casaccia, Via Aguillarese 301, I-00060 Santa Maria di Galeria, Roma, Italy Received 29th July 1997, Accepted 21st September 1998 Lithium loss kinetics from polycrystalline LixNi1-xO has been investigated in the temperature range 900–1500 °C by measuring the time and temperature dependence of the weight and lattice constant change of the samples.At 900 °C the rate of Li2O evaporation was controlled by lithium ion diVusion in LixNi1-xO. An initial region of fast diVusion followed by a region of slower diVusion was observed in the thermogravimetric measurements. This observation can be interpreted as the rapid diVusion of lithium ions along the grain boundaries and subsequent diVusion into the bulk of the grain.A plot of the logarithm of lithium loss following 2 h of thermal treatment at diVerent temperatures vs. the reciprocal of absolute temperature consisted of two straight lines, the slope depending on the activation energy of the process and the change of slope occurred at 1300 °C. This behaviour suggests that up to 1300 °C lithium ion diVusion from the bulk to the surface of LixNi1-xO particles was the rate-determining step.Above 1300 °C, instead, the evaporation process depended on the demixing reaction of Li2O(g) and NiO(s) at the surface of the particles. Pure, stoichiometric nickel oxide is a compound with cubic detect, and for x>0.30 the change of crystal structure during the evaporation process would complicate the interpretation NaCl-type structure.Reaction of Li2O and NiO in the presence of the results. of oxygen gives rise to the formation of LixNi2+1-2xNi3+xO solid solution, where the oxidation state of Ni partially changes from +2 to +3.1,2 The unit cell slightly decreases with increasing lithium content as a consequence of the diVerence Experimental in the ionic radius between Ni2+ and Ni3+ ions.For lithium Lithium nickel oxide solid solution was prepared by solid state atomic fraction x>0.31 a rhombohedral distortion of the cubic reaction of Ni and Li2CO3. Nickel powder (INCO 255) and structure by lithium nickel ordering on alternate (111) planes lithium carbonate (Merck 5671) in the nominal lithium atomic takes place, giving rise to a hexagonal simmetry.2–4 LixNi1-xO fraction x=0.30 were used as starting materials.This powder solid solutions are used as a cathode material in molten mixture was submitted to the following thermal treatments in carbonate fuel cells.5 Moreover LixNi1-xO has been studied air: (i) heating from room temperature to 900 °C at with the aim of developing humidity ceramic sensors.6 Studies 1.0 K min-1; (ii) isothermal treatment at 900 °C up to 40h; on the synthesis of LixNi1-xO with low (x0.23)7,8 and high (iii) heating at 3.0 K min-1 from 900 °C to the maximum (x0.30)9 lithium content by solid state reaction of a Ni and temperature in the range 1000–1500 °C; (iv) isothermal treat- Li2CO3 powder mixture indicated that solid solution occurs ment at the maximum temperature for 2 h and (v) cooling to in two steps: (i) first, formation of LixNi1-xO with x higher room temperature at 10 K min-1.Moreover two samples were than the nominal value at the grain surface, then (ii) diVusion treated for diVerent times at 1200 and 1400 °C, respectively. of lithium ions from the surface to the bulk of the particle.Thermogravimetric measurements were performed using a Few studies have dealt with Li2O evaporation from Du Pont 2000 thermal analysis system equipped with a 951 LixNi1-xO. A recent work of Sata10 on the vaporization of TGA module. lithium oxide from LixNi1-xO solid solution at temperatures XRD patterns were collected at room temperature after up to 700 °C stated that Li2O2 formation and its diVusion rate quenching of the samples on a Philips PW 1729 powder in the specimen might be related to the rate-determining step diVractometer equipped with a 1771 vertical goniometer using in the vaporization process.Above 1000 °C, Iida found that filtered Cu-Ka radiation. the evaporation of lithium oxide from the solid solution is diVusion-controlled and governed by a parabolic law.11 The distribution of remaining lithium ions in the solid solution Results and discussion following the loss of some Li2O was also investigated.Sata LixNi1-xO was formed during the dynamic step up to 900 °C found that the lithium concentration in the specimen decreased of the thermal treatment. As lithium loss takes place during linearly from the surface to the interior along the specimen the process of formation of LixNi1-xO from Ni/Li2CO3 mix- thickness.10 Azzoni et al.12 observed the coexistence of substitures, 14 we determined the lithium atomic fraction at the tutional solid solution NiO type and ordered solid solution beginning of the isothermal step at 900 °C from both X-ray LiNiO2 type structures with diVerent lithium content following diVraction (XRD) and thermogravimetric measurements.Li2O evaporation from ordered solid solutions (x>0.30). From XRD measurements, the following relation between Moreover, Berbenni et al.13 revealed structural and microlithium atomic fraction x and lattice constant a/A° was used:15 structural changes of LixNi1-xO solid solution during Li2O evaporation at 800 °C. x=(4.1748-a)/0.17756 (1) The aim of this work is to better understand the process of Li2O evaporation from the solid solution at high temperatures From the weight change, we utilised the following relation:16 in the range 900 to 1500 °C.The nominal lithium content of the composition investigated was x=0.30, as for x<0.30 the x=[A(1+Dm/m0)-MNiO]/[A(1+Dm/m0)-(MNi-MLi)] (2) final amount of lithium would be very low and diYcult to J.Mater. Chem., 1998, 8, 2783–2786 2783with A=MNi+xn MLi2CO3/2(1-xn) (3) where Dm/m0 is the weight change of the samples, xn is nominal lithium atomic fraction, and MNiO, MNi, MLi and MLi2CO3 are the molecular weights of the compounds. The value of x obtained was 0.260 from XRD measurements (a= 4.1286 A° ) and 0.265 from thermogravimetric measurements (Dm/m0=0.113), i.e.the same within experimental error. The Li2O evaporation reaction is given by eqn. (4). LixNi1-xOAa LiyNi1-yO+b Li2O+(b/2)O2 (4) where a=(1-x)/(1-y) and b=(x-y)/2(1-y) for x>y. This reaction takes place in two steps: (i) lithium ion diVusion to the grain surface and (ii) a demixing reaction of Fig. 2 Log–log plot of fractional lithium loss from the solid solution vs. thermal treatment time at 900 °C.Li2O(g) and NiO(s) at the surface of the particles. The basic phenomenological mass transport relation, governing a solid state diVusion process like lithium ion diVusion in LixNi1-xO, grain from the bulk to the lithium-poor boundaries. DiVusion is that of Fick: in polycrystalline solids is known to occur along grain bound- J=-D grad C (5) aries more rapidly than through the interior of the crystals.Atkinson found that Ni diVusion in NiO was enhanced at where J is the number of atoms crossing a unit area in unit grain boundaries with respect to lattice diVusion and that the time, C is the concentration of the mobile species, and the faster diVusion pathway had the lower activation energies.19 constant D is the chemical diVusion coeYcient. In solids atoms The fast grain boundary diVusion is caused by the segregation adopt reasonably well defined positions.Mass transport occurs of point defects to the core region where they have higher by atoms making transitions between these positions in such concentration and higher mobility than in the lattice.20 As a way that the time of transit is much less than the residence long as the rate of grain-boundary diVusion is greater than time at any particular position.Thus, diVusion can be thought that of lattice diVusion at all temperatures, we can assume of as occurring by particle hopping in a random way on a that in the diVusion controlled region not only at 900 °C, lattice of sites distributed in space.17 The hopping event but at all temperatures, lattice diVusion is initiated after between sites involves the particle crossing an energy barrier, the boundaries are lithium-poor owing to lithium loss by the necessary energy coming from thermal fluctuations with a grain-boundary diVusion. probability described by the Boltzmann distribution.Hence, XRD measurements confirm the results of thermogravi- the diVusion process is thermally activated and the diVusion metric measurements.Fig. 3 shows the 204 reflections of the coeYcient has the Arrhenius form: solid solution following diVerent thermal treatment times at D=D0 exp (-Ea/RT) (6) 900 °C. The reflection shifts towards lower angles with time, owing to lithium loss. From 0 to 3 h, we detect a broadening where Ea is the activation energy of diVusion.The evaporation of the peak, attributed to a lithium concentration gradient of kinetics can be expressed as: outer and inner parts of the grain, as a consequence of fast Cev=ktn (7) lithium loss at the boundaries. The result of the Rietveld refinement procedure indicated the presence of three solid where Cev is the amount of evaporated species, k=k0exp solutions with x=0.26, 0.24 and 0.22, respectively.A sharpen- (-Ea/RT) is the rate constant, t is the thermal treatment time ing of the reflection then occurs between 3 and 20 h, related and the exponent n is related to the reaction mechanism. to homogenisation of the solid solution owing to the diVusion DiVusion-controlled evaporation is described by n=0.5 to of lithium ions from the bulk to the boundaries of the particles. 0.8.18 Fig. 1 and 2 show normal and log–log plots of fractional The transition point between grain-boundary and lattice lithium loss from the solid solution vs. thermal treatment time diVusion indicates that about 30% of lithium in the solid at 900 °C, respectively. As can be seen in Fig. 2, the log–log solution is present at the grain boundaries or at layers adjacent plot breaks into two straight lines with diVerent slopes.The to the grain boundaries. slopes of these lines were 0.75 and 0.48, respectively. In both After 40 h of isothermal treatment at 900 °C, the lithium cases, a diVusion-controlled process is occurring. This result atomic fraction x remaining in LixNi1-xO is 0.123. Then, these can be interpreted in terms of rapid lithium diVusion along specimens were thermally treated at various temperatures for the grain boundaries, followed by lithium diVusion into the Fig. 3 DiVractometric traces of the 204 reflection of LixNi1-xO Fig. 1 Dependence of fractional lithium loss from the solid solution on thermal treatment time at 900 °C. following diVerent times of thermal treatment at 900 °C. 2784 J. Mater. Chem., 1998, 8, 2783–2786Fig. 6 Log–log plot of lithium loss a from the solid solution vs. Fig. 4 The logarithm of lithium loss a from the solid solution as a function of the reciprocal of absolute temperature. thermal treatment time at 1200 and 1400 °C. 2 h. From the relation: Cev=k0 exp(-Ea/RT) tn (7) plotting ln Cev vs. 1/T at constant time, we can determine the activation energy of the process.Fig. 4 shows the dependence of the logarithm of lithium loss a=mLiev/m0Li, (where mLiev is the amount of evaporated lithium, in this case after 2 h of thermal treatment, and m0Li is total lithium amount at the beginning of isothermal treatment) on the reciprocal of absolute temperature. A change in the slope of the plot is observed at 1300 °C.This feature is related to the change of the ratedetermining process. It can be inferred that at high temperatures the rate-determining step is the demixing reaction at the surface of the grain. To confirm this result we have evaluated the time-dependence of lithium loss at 1200 °C (diVusioncontrolled region) and 1400 °C (demixing reaction-controlled Fig. 7 Dependence of the logarithm of remaining lithium (1-a) in region), respectively.Fig. 5 shows a plot of lithium loss a(= the solid solution on thermal treatment time at 1200 and 1400 °C. mLiev/m0Li) vs. time of samples thermally treated at 1200 and 1400 °C. There are two pathways by which the rate of respectively. We have determined the correlation parameters evaporation of Li2O can occur at high temperature.(r2) for the experimental data and theoretical equations, the (a) DiVusion control: the relation of lithium loss to time is slopes of the plots shown in Fig. 6 (the slope is the value of given by eqn. (8): n) and 7 (the slope is the value of k) and their estimated a=ktn (8) standard errors. As can be inferred from the values reported in Table 1, at 1200 °C the best fit is shown according to the (b) Surface reaction: Li2O evaporation at the grain surface diVusion process, whereas at 1400 °C the best fit is shown follows first order kinetics [eqn.(9)]: according to the demixing reaction at the grain surface. d(1-a)/dt=k (1-a) (9) The activation energy for the diVusion process was 179 kJ mol-1. From the slope of ln[-ln(1-a)] vs. 1/T, the Thus, the relation between lithium loss a and isothermal heat of evaporation of Li2O from LixNi1-xO was evaluated treatment time fits a first-order rate law [eqn.(10)]: as 92 kJ mol-1. a=1 -exp(-kt ) (10) Fig. 6 and 7 show log–log plots of lithium loss (a) vs. time Conclusions (diVusion control ) and the plot of the logarithm of remaining In the present study it was observed that lithium loss from lithium (1-a) in the solid solution vs.time (surface reaction), LixNi1-xO with x=0.26 at 900 °C initially occurs by grainboundary diVusion of lithium ions in the solid solution. Then, for long times and temperatures up to 1300 °C, the ratedetermining step was lattice diVusion of lithium ions into LixNi1-xO. Above 1300 °C, a demixing reaction of Li2O(g) and NiO(s) at the grain surfaces was the rate-controlling step.Table 1 Correlation parameters (r2) of experimental data and theoretical equations, k and n values, and their estimated standard errors Temperature/ °C log a=log k+n log t log (1-a)=-kt 1200 r2=0.991 r2=0.985 n=0.65 Dn/n=0.063 k=0.014 h-1 Dk/k=0.071 1400 r2=0.978 r2=0.996 Fig. 5 Dependence of lithium loss a from the solid solution on thermal n=0.60 Dn/n=0.103 k=0.179 h-1 Dk/k=0.042 treatment time at 1200 ($) and 1400 °C (().J. Mater. Chem., 1998, 8, 2783–2786 278510 T. Sata, Ceram. Int., 1998, 24, 53. References 11 Y. Iida, J. Am. Ceram. Soc., 1960, 43, 117. 12 C. B. Azzoni, A. Paleari, V. Massarotti, M. Bini and D. Capsoni, 1 S. Van Houten, J. Phys. Chem. Solids, 1960, 17, 7. Phys. Rev. B, 1996, 53, 703. 2 J. B. Goodenough, D. G. Wickham and W. J. Croft, J. Phys. 13 V. Berbenni, V. Massarotti, D. Capsoni, R. Riccardi, A. Marini Chem. Solids, 1958, 5, 107. and E. Antolini, Solid State Ionics, 1991, 48, 101. 3 I. J. Pickering, J. T. Lewandowski, A. J. Jacobson and 14 E. Antolini, Mater. Lett., 1993, 16, 286. J. A. Goldstone, Solid State Ionics, 1992, 53–56, 405. 15 E. Antolini, M. Leonini, V. Massarotti, A. Marini, V. Berbenni 4 R. Stoyanova and E. Zhecheva, J. Solid State Chem., 1994, 108, and D. Capsoni, Solid State Ionics, 1990, 39, 251. 211. 16 A. Marini, V. Massarotti, V. Berbenni, D. Capsoni, R. Riccardi, 5 Fuel Cells Handbook, ed. A. J. Appleby and F. R. Foulkes, Van E. Antolini and B. Passalacqua, Solid State Ionics, 1991, 45, 143. Nostrand Reinhold, New York, 1990. 17 A. Atkinson, Mater. Sci. Technol., 1994, 11, 303. 6 T. Sato, C. Hsien-Chang, T. Endo and M. Shimada, J. Mater. Sci. 18 T. Sata and T. Yokoyama, Yogyo Kyokaishi, 1973, 81, 170. Lett., 1986, 5, 552. 19 A. Atkinson, Adv. Ceram., 1987, 23, 3. 7 E. Antolini, J. Mater. Sci. Lett., 1993, 12, 1947. 20 D. M. DuVy, J. Phys. C, 1986, 19, 4393. 8 E. Antolini, A. Marini, V. Berbenni, V. Massarotti, D. Capsoni and R. Riccardi, Solid State Ionics, 1992, 57, 217. 9 E. Antolini and M. Ferretti, Mater. Lett., 1997, 30, 59. Paper 8/05948I 2786 J. Mater. Chem., 1998, 8, 2783–2786

ISSN:0959-9428    

DOI:10.1039/a805948i

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Chemical interactions between strontium-doped praseodymium manganite and 3 mol% yttria-zirconia Jin-Ping Zhang,*a San-Ping Jiang,b Jonathan G. Love,a Karl Fogera and Sukhvinder P. S. Badwalb aCeramic Fuel Cells Limited, 170 Browns Road, Noble Park, Vic. 3174, Australia bCSIRO, Manufacturing Science & Technology, Private Bag 33, Clayton South MDC, Clayton, Vic. 3169, Australia Received 27th July 1998, Accepted 14th September 1998 The interfacial reaction between (Pr0.8Sr0.2)yMnO3 (PSM, y=0.9, 1.0, 1.05) film and 3 mol% yttria tetragonal zirconia (TZ3Y) substrate has been studied at 1200 and 1400 °C in air. A diVusion layer of Pr and Mn in zirconia, which was identified to be a cubic phase of zirconia, was detected in all specimens.When the solubility limit of Pr ions in the cubic zirconia was reached, a pyrochlore phase, Pr2Zr2O7, was formed. A delay for the formation of pyrochlore phase was observed for the A-site sub-stoichiometric and stoichiometric PSM at 1200 °C. For the A-site over-stoichiometric PSM, a Pr-rich (Pr,Zr)Ox phase was detected at the interface besides the pyrochlore phase. At 1400 °C, a relatively thick layer of pyrochlore phase was formed after 24 hour heat treatment in all specimens.The amount of the pyrochlore phase formed at the interface depends on the A-site stoichiometry of perovskite in the initial stage. The growth of the pyrochlore layer after the initial stage, however, appears to be determined by contact area between PSM and the substrate.changes between porous (Pr0.8Sr0.2)yMnO3 ( y=0.9, 1.0, 1.05) 1 Introduction film and 3 mol% Y2O3-ZrO2 (TZ3Y) electrolyte have been Doped perovskite oxides and yttria-zirconia (YSZ) are investigated in air at 1200 and 1400 °C. Although not much commonly used as cathode and electrolyte materials respect- information is reported on interactions between TZ3Y and ively in solid oxide fuel cells (SOFCs) operating at tempera- LSM, this electrolyte composition, despite its low conductivity, tures of around 900–1000 °C.1,2 Interfacial reactions between is of considerable interest to many SOFC technology develthe cathode [especially in Sr-doped lanthanum manganite opers, because its mechanical strength is high and it is easy to (LSM)] and YSZ, have been studied extensively at high fabricate very thin (60–70 mm) sheets of this material (as temperatures, and the reaction products have been well charac- opposed to 150 mm for the 8 mol% Y2O3-ZrO2). terised and documented.3–7 The stoichiometric LSM reacts with YSZ extensively at temperatures above 1200 °C,6–10 form- 2 Experimental ing lanthanum zirconate La2Zr2O7 and/or strontium zirconate SrZrO3 phases at the interface depending on the La/Sr ratio The PSM powders with composition (Pr0.8Sr0.2)yMnO3 ( y= at the A-site.The interfacial reactions have also been reported 0.9, 1.0 and 1.05, coded hereafter PSM-A, PSM-B and PSMat lower temperatures by some authors (1150, 11004,11 and C respectively) were prepared by co-precipitation followed by 1000 °C12). It is generally known that an A-site deficient LSM calcination at 1000 °C for 4 h in air.The value of y is used to suppresses the formation of La2Zr2O7.4,6,11,13 The formation indicate the stoichiometry of the perovskite phase for conof zirconates at the interface is detrimental to the performance venience only and no assumption has been made that it is a of a solid oxide fuel cell system, causing substantial increase single phase material.The electrolyte substrate TZ3Y, 20 mm in the overpotential and resistivity at the cathode/electrolyte in diameter and 150 mm in thickness, was prepared from interface.14 3 mol% Y2O3-ZrO2 (Tosoh Corporation, Japan) by tape cast- Owing to severe corrosion of stack components, high cost ing and sintering at 1500 °C. The PSM was screen-printed (ca.and degradation of stack performance, it is necessary to lower 40 mm thick) on the TZ3Y substrate and sintered at 1200 °C the operating temperature of SOFCs from 900–1000 °C to the for 4, 24 or 168 h. Another batch was sintered at 1400 °C for intermediate range 700–800 °C.2 One of the requirements for 24 h. The heat treatment temperature of 1200 °C has been lowering the operating temperature is to develop a new cathode commonly used in the LSM/YSZ system for the investigation material which has reasonably low overpotential losses in the of the interfacial reaction.The 1400 °C temperature was used 700–800 °C temperature range. Ishihara et al.15 studied the to accelerate the solid state reaction for a more conspicuous electrochemical behaviour of the Sr-doped praseodymium observation of the interactions between PSM and TZ3Y.manganite (PSM) and found that the overpotential losses for PSM powders after calcination were characterised by X-ray PSM were significantly lower than those of other Sr-doped diVraction (XRD) for phase analysis and by scanning electron lanthanide manganites at intermediate operating temperatures. microscopy (SEM) for powder morphology.After heat treat- Therefore the material is a potential cathode for intermediate ments, the reaction couples were carefully fractured and in temperature SOFC operation. However, information on inter- some cases polished cross-sections were prepared. Both fracactions between PSM and YSZ is scarce. Wen et al.16 sintered tured face and polished cross-sections were examined with a PSM and YSZ powder mixture at 1000 °C for 100 h and did SEM.X-Ray energy dispersive spectroscopy (EDS) was used not detect any interfacial reactions. In order to examine to study the elemental distribution in the PSM/TZ3Y interface interfacial reactions between PSM and YSZ, a higher heat region. In some samples, PSM was carefully removed from the treatment temperature (1200 °C), similar to those used for TZ3Y substrate and the exposed TZ3Y surface was examined the LSM/YSZ system in most investigations, may be required. with XRD and SEM/EDS.A Siemens D500 X-ray diVractometer (Siemens, Germany) with Cu-Ka radiation and In the present study, the microchemical and microstructural J. Mater. Chem., 1998, 8, 2787–2794 2787Fig. 1 XRD patterns of the powders calcined at 1000 °C showing that PSM-A and PSM-B are relatively pure perovskite, while PSM-C contains the major phase perovskite and a small amount of praseodymium oxide. a Leica 360 field emission SEM (Cambridge, UK) equipped with an Oxford Link EDS system were used for specimen characterisation. Identification of phases from XRD patterns was based on the JCPDS-ICDD database. The Rietveld method,17 a technique for crystal structure refinement from powder diVraction data, was used in the current study to analyse zirconia phases at the PSM/TZ3Y interface.The Rietveld refinement was performed using the program LHPM1.18 3 Results XRD traces of powders calcined at 1000 °C are displayed in Fig. 1. The diVraction peaks marked ‘P’ in Fig. 1 are of perovskite, and those marked ‘O’ belong to praseodymium oxide Pr6O11. The stoichiometric PSM-B and the A-site substoichiometric PSM-A powders are relatively pure perovskite while A-site over-stoichiometric PSM-C powder contains a small amount of praseodymium oxide besides the major perovskite phase. A few minor unidentified reflections are also noticed in all the traces. The powder morphology of PSM-A, PSM-B and PSM-C is about the same after calcination at 1000 °C for 4 h.The PSM particle size was in the range 0.1–0.2 mm. Fig. 2 The backscattered electron micrographs of the polished cross (1) Reaction products after heat treatment at 1400 °C section of the interfaces after 24 h at 1400 °C: (a) PSM-A/TZ3Y; (b) PSM-B/TZ3Y; (c) PSM-C/TZ3Y. ‘L1’ and ‘L2’ refer to the Two reaction layers were identified between TZ3Y and PSM.pyrochlore layer and the Pr- and Mn-diVused zirconia layer Fig. 2 displays the backscattered electron micrographs taken respectively. from a polished cross section of the PSM/TZ3Y interface after heat treatment at 1400 °C for 24 h. The first reaction layer (marked L1 in Fig. 2) can be clearly observed in these micrographs.Pr and Zr were identified by EDS analysis as the major elements in the reaction layer. To identify the phase of the reaction layer, PSM was removed carefully by scraping and the exposed surface of the reaction layer was examined by XRD. A representative XRD trace from the reaction layer on TZ3Y substrate which was in contact with PSM-C is presented in Fig. 3 with the identification of each reflection.The major reflections (marked ‘X’ in Fig. 3) match with those of the pyrochlore phase, Pr2Zr2O7 (ICDD file No.: 19-1021), except that the intensity of the reflection (400) (2h=33.45°) is about twice that reported in 19-1021. This probably is due to Fig. 3 A representative XRD pattern from the substrate originally in contact with PSM-C showing that the pyrochlore phase was formed the preferred orientation. Thus both XRD and EDS analysis after 24 h at 1400 °C.results confirm that the major phase in the reaction layer ‘L1’ is Pr2Zr2O7, the pyrochlore phase. Some minor reflections in Fig. 3 arise from the substrate TZ3Y, and the perovskite when the layer was formed with PSM-A, 3.1 mm with PSM-B and 12.0 mm with PSM-C after heat treatment at 1400 °C powder left-over after scraping.The average thickness of the praseodymium zirconate layer for 24 h. In addition to the pyrochlore phase, a second reaction layer determined from backscattered electron micrographs varied with the compositions of PSM. It was about 9.5 mm thick was also observed from the fractured surface. Fig. 4 shows a 2788 J. Mater. Chem., 1998, 8, 2787–2794representative micrograph of the fractured surface of PSMA/ TZ3Y.In this micrograph a 28 mm thick distinct layer is obvious between PSM-A and TZ3Y. It contains the pyrochlore layer ‘L1’ that is only 9.5 mm thick, and another reaction layer ‘L2’. The technique of energy dispersive X-ray mapping was used to identify elements in the layer ‘L2’. Fig. 5 displays Xray maps recorded from the polished cross section of PSMA/ TZ3Y showing the distribution of related elements Zr, Y, Sr, Pr and Mn around the interface. It should be noted that the Mn Ka peak partly overlapped with Pr Lb2, the intensity of which has been subtracted from Mn Ka in Fig. 5. In order to show clearly the element distribution in layer ‘L2’, X-ray maps (g) and (h) do not include the PSM layer because the contrast between the Mn concentration in PSM and in ‘L2’ is so high that the Mn distribution in ‘L2’ can not be seen from Fig. 5(f ) when the PSM layer is included.From Fig. 5, it can Fig. 4 SEM micrograph of the fractured surface of PSM-A/TZ3Y be seen that the layer ‘L2’ consists of Pr, Mn and the elements showing a ca. 28 mm thick distinct reaction layer formed between of TZ3Y.The Pr ions and some Mn ions appeared to have PSM-A and TZ3Y after 24 h at 1400 °C, which consists actually of entered the TZ3Y lattice forming a solid solution. More results two layers of products, ‘L1’ (pyrochlore) and ‘L2’ (Pr- and Mn-diVused zirconia). about the nature of this diVusion layer will be presented in the following sections. The XRD, SEM and EDS observations thus far for Fig. 5 The EDS X-ray maps recorded from the polished cross section of PSM-A/TZ3Y after 24 h at 1400 °C showing the distribution of related elements in PSM and reaction layers L1 and L2: (a) backscattered electron image; (b) Zr La1; (c) Y Ka; (d) Sr Ka; (e) Pr La1; (f )Mn Ka. The Xray maps (g) Pr La1 and (h) Mn Ka do not include the PSM layer in order to show clearly the distribution of Pr and Mn in the diVusion layer ‘L2’, which can not be seen clearly from (e) and (f ) in which the PSM layer is included.J. Mater. Chem., 1998, 8, 2787–2794 2789Table 1 The thickness of the zirconate layer and the diVusion distance It was also noticed from SEM examination that the PSM-A of Pr and Mn in TZ3Y after 24 h at 1400 °C and PSM-C coatings were highly sintered compared with the PSM-B coating.Thickness of Pr diVusion Mn diVusion Specimen zirconate/mm distancea/mm distancea/mm (2) Reaction products after heat treatment at 1200 °C PSM-A/TZ3Y 9.5 28 55 Fig. 7 displays the backscattered electron micrographs of PSM-B/TZ3Y 3.1 6 14 polished cross sections of PSM-A/TZ3Y, PSM-B/TZ3Y and PSM-C/TZ3Y 12.0 30 70 PSM-C/TZ3Y after heat treatment at 1200 °C for 4, 24 and aThe distance was measured from the PSM/pyrochlore phase interface. 168 h respectively. The top part in each micrograph shows PSM, and the bottom section the substrate. Pyrochlore (PZ) formed at the interface is marked on the micrographs. From PSM-A/TZ3Y interface heat treated at 1400 °C can be summarthese micrographs it can be seen that the reaction products ised as below.between PSM and TZ3Y vary with the A-site stoichiometry 1 Pr ions have diVused into TZ3Y. The diVusion distance of PSM and the time of heat treatment. Substantial amounts is about 28 mm (from the PSM/pyrochlore interface), correof praseodymium zirconate were detected after heat treat- sponding to the thickness of the dense layer viewed from the ment of PSM-A/TZ3Y at 1200 °C for 168 h, PSM-B/TZ3Y for fractured surface. 24 h, and PSM-C/TZ3Y for 4 h. In PSM-B/TZ3Y and PSM- 2 Mn ions have also diVused into TZ3Y. The diVusion C/TZ3Y, praseodymium zirconate formed a continuous layer distance of Mn ions (ca. 55 mm from PSM) is much larger than at the interface, whereas in PSM-A/TZ3Y the zirconate formed that of Pr. However, Mn was not detected in the praseodymium islands at contact points between PSM-A and the substrate.zirconate layer. The thickness of the reaction layer grew with the time of heat 3 A small amount of Sr, estimated to be less than a few wt.%, treatment. was also detected in the zirconate layer (L1), but no strontium The inset in the micrograph of PSM-C/TZ3Y/24 h in Fig. 7 zirconate phase was formed.is an enlargement of the reaction layer, showing clearly two 4 Zr or Y was not found in the PSM phase. XRD study distinct layers of products between the PSM-C coating and showed that the perovskite phase of the PSM layer did not the substrate. Fig. 8 presents the EDS spectra (a) for the top change its structure after the reaction. layer (the brightest in contrast in the micrograph) and (b) for 5 The yttrium was detected in both reaction layers (L1 and the bottom layer, showing that both layers contain Pr and Zr L2) at about the same level as in the TZ3Y bulk phase (i.e. but with diVerent atomic ratios.The experimental conditions 3 mol% Y2O3). for the EDS X-ray analysis were kept the same in all cases so The same microstructure of the interface was also found in that a semi-quantitative comparison of elemental concentration the other two specimens PSM-B/TZ3Y and PSM-C/TZ3Y in diVerent specimens could be carried out.Trace (b) in Fig. 8 heat-treated at 1400 °C for 24 h. The thickness of the zirconate is a typical EDS spectrum of praseodymium zirconate. From and the diVusion distance (measured from the PSM/pyrochlore comparison of the two EDS traces it is known that the atomic interface) of the major elements Pr and Mn in TZ3Y in three ratio of Pr/Zr of the top layer is higher than that of the specimens are summarised in Table 1.It can be seen from pyrochlore layer which is about 151. The top layer, therefore, Table 1 that the degree of the reaction for PSM-A/TZ3Y is consists of a (Pr,Zr)Ox phase with the atomic ratio of Pr/Zr>1.slightly lower than that for PSM-C/TZ3Y interface whereas A very thin layer of (Pr,Zr)Ox phase was also detected in PSM-B/TZ3Y interface showed relatively higher stability. PSM-C/TZ3Y after 4 h at 1200 °C when it was examined at a The fact that Zr and Y were not detected in the PSM layer, higher magnification. After 168 h at 1200 °C the amount of clearly indicates that the growth of the zirconate layer is in (Pr,Zr)Ox was much less than that after 24 h, and did not form the direction of the abutting electrolyte.This was also obvious a distinct layer, as shown in Fig. 7. from Fig. 6, a micrograph taken from the cross section of After removal of most of the PSM-C coating from the PSM-A/TZ3Y after 24 h at 1400 °C. In the area where there specimen sintered for 24 h at 1200 °C, XRD analysis was was no PSM-A, the pyrochlore phase was not formed.It can carried out on the substrate and part of the trace is displayed be seen from the micrograph that the TZ3Y substrate and the in Fig. 9. There are a few extra peaks (marked ‘O’) in the XRD praseodymium zirconate top surfaces are almost level, indicattrace besides those of expected phases pyrochlore, TZ3Y and ing that the zirconate phase has grown into the TZ3Y substrate.perovskite ( left over from scraping). These extra reflections are The absence of PSM-A in such areas probably arose from the probably of the (Pr, Zr)Ox solid solution phase because they shrinkage of the coating at high temperature. match with the reflections of Pr6O11 with a systematic peak position shift (to larger angle) that is not a zero point error. The Pr and Mn diVusion layer observed in the specimens sintered at 1400 °C was also detected in all specimens heated at 1200 °C.Some representative EDS spectra of the diVusion layer are displayed in Fig. 10. Spectra (a) and (b) were recorded from the diVusion layers formed in PSM-A/TZ3Y after heat treatment for 4 and 24 h respectively at 1200 °C, showing clearly the presence of Pr and Mn in TZ3Y.For comparison the EDS spectrum of praseodymium zirconate formed in PSMA/ TZ3Y after 168 h is also displayed in Fig. 10(c). It can be seen from Fig. 10 that the concentration of Pr in TZ3Y increases with the time of heat treatment. The diVusion layer is not obvious from the contrast of SEM micrographs taken from polished cross sections (Fig. 7). However, the change of the substrate microstructure near the interface due to diVusion of Pr and Mn and the formation of the pyrochlore phase can be seen on the substrate surface by Fig. 6 SEM micrograph from the cross section of PSM-A/TZ3Y after removing the PSM coating carefully from the substrate. This 24 h at 1400 °C showing that the TZ3Y substrate and the praseodymis illustrated in the secondary electron micrographs shown in ium zirconate top surfaces are almost leveled, indicating that the zirconate has grown into the TZ3Y substrate.Fig. 11. Fig. 11(a) was taken from the unreacted TZ3Y for 2790 J. Mater. Chem., 1998, 8, 2787–2794Fig. 7 The backscattered electron micrographs of the polished cross sections of PSM-A/TZ3Y, PSM-B/TZ3Y and PSM-C/TZ3Y after heat treatment at 1200 °C for 4, 24 and 168 h respectively.The top part in each micrograph shows PSM, and bottom part the substrate. Pyrochlore (PZ) formed at the interface is marked in the micrographs. The inset in the micrograph of PSM-C/TZ3Y/24h is an enlargement of the reaction layer, showing clearly two distinct layers of products between the PSM-C coating and the substrate.Fig. 10 The EDS spectra recorded from the substrate near the interface of PSM-A/TZ3Y after sintering at 1200 °C for 4 h (a), 24 h (b) and Fig. 8 The EDS spectra recorded from the top layer (a) and the 168 h (c). The electron beam was located on the zirconate for bottom layer (b) shown as inset in the micrograph PSM-C/TZ3Y/24 h spectrum (c).in Fig. 7, indicating that both layers contain Pr and Zr but with diVerent atomic ratios. comparison; Fig. 11(b) from the Pr- and Mn-diVused zirconia in PSM-A/TZ3Y showing a dramatic increase of the grain size of zirconia near the interface after 24 h at 1200 °C; Fig. 11(c) from the same specimen as in Fig. 11(b) but in a diVerent area showing the trace of some contact points with PSM; Fig. 11(d) from the praseodymium zirconate surface formed in PSMA/ TZ3Y after heat treatment at 1200 °C for 168 h; and Fig. 11(e) from the substrate surface near the edge of the reaction layer in PSM-B/TZ3Y after heat treatment at 1200 °C for 168 h showing the microstructure of diVerent layers including the pyrochlore layer ‘L1’, the Pr- and Mn-diVused zirconia layer ‘L2’ and the unreacted TZ3Y.EDS analysis on the trace of the contact points shown in Fig. 11(c) showed a typical composition of the pyrochlore phase. This indicates that some crystal nuclei of praseodymium zirconate had been formed on the surface of Pr- and Mn-diVused zirconia in PSM-A/TZ3Y Fig. 9 The XRD pattern recorded from the substrate surface of after 24 h at 1200 °C.The crystal nuclei of praseodymium PSM-C/TZ3Y after 24 h at 1200 °C showing the formation of the zirconate were not observed from the cross-section in Fig. 7 praseodymia solid solution (Pr,Zr)Ox besides the expected phases pyrochlore, TZ3Y and perovskite. probably because the scale is too small. J. Mater. Chem., 1998, 8, 2787–2794 2791Fig. 11 SEM micrographs showing the morphology of the substrate surface: (a) unreacted TZ3Y; (b) the Pr- and Mn-diVused zirconia in PSM-A/TZ3Y showing a dramatic grain growth of zirconia at the Fig. 12 (a) The XRD patterns recorded from the substrate surface: interface after 24 h at 1200 °C; (c) the same specimen as (b) but in a (i) unreacted TZ3Y, (ii) reacted with PSM-A at 1200 °C for 24 h and diVerent area showing the formation of the pyrochlore crystal nuclei; (iii) reacted for 168 h.(b) Enlarged part of (a) showing that the (d) the praseodymium zirconate surface formed in PSM-A/TZ3Y after reflection marked * in (ii) consists of two peaks. sintering at 1200 °C for 168 h; and (e) the substrate surface near the edge of the reaction layer in PSM-B/TZ3Y after 168 h at 1200 °C showing the microstructure of diVerent layers including the pyrochlore layer ‘L1’, the Pr- and Mn-diVused zirconia layer ‘L2’ and the unreacted TZ3Y.Fig. 12(a) displays the XRD patterns recorded from the substrate surface: (i) unreacted TZ3Y, (ii ) reacted with PSMA at 1200 °C for 24 h and (iii ) reacted for 168 h. The TZ3Y has a tetragonal form of crystal structure, and the trace (i) in Fig. 12(a) is a typical XRD pattern of a tetragonal zirconia.A comparison between the trace (ii) and the trace (i) shows that the relative intensities of a number of reflections (marked with an asterisk) in trace (ii) are much higher than in trace (i). In fact, when the traces were enlarged, it was observed that each of the enhanced reflections consisted of two peaks, and one example is shown in Fig. 12(b). This suggests that another phase has been formed with the diVusion of Pr and/or Mn ions into the zirconia. Trace (iii ) in Fig. 12(a) is composed of the diVraction peaks belonging to the Pr- and Mn-diVused zirconia shown in the trace (ii) and those of pyrochlore phase. All XRD results are consistent with SEM observations. In order to characterise the phase of the Pr- and Mn diVused zirconia, the XRD data of trace (ii) in Fig. 12(a) were analysed using the Rietveld method. During the Rietveld refinement the unit cell parameters, zero point, scale factors, peak Fig. 13 The output from the Rietveld refinements of the XRD pattern of the Pr- and Mn-diVused TZ3Y in PSM-A/TZ3Y after 24 h at width/shape/asymmetry parameters and background 1200 °C using (a) tetragonal lattice, and (b) cubic as well as tetragonal coeYcients were refined simultaneously to convergence.The lattice. The observed data are indicated by crosses and the calculated atomic position parameters were fixed as reported in the by a continuous line overlying them, and the diVerence profile is the literature.19 When the structure parameters of tetragonal zir- lower curve in each figure.The short vertical lines show the positions conia were tested in the Rietveld refinement, the agreement of all possible Bragg reflections. index Rwp was 0.132. When the cubic as well as the tetragonal lattice of zirconia was used in the refinement, Rwp dropped significantly to 0.068. The output from the Rietveld refinements was improved significantly when the cubic lattice was included in the refinement, as shown in Fig. 13(b). This suggests that using (a) the tetragonal lattice, and (b) the cubic as well as the tetragonal lattice is displayed in Fig. 13. Large diVerence (lower the Pr- and Mn-diVused zirconia region has a cubic structure and is consistent with the observation of grain growth in the curve in each figure) in reflection intensities between the observed (+ markers) and the calculated (continuous line) diVusion layer because usually dopant-stabilised cubic zirconia has much larger grain size than tetragonal TZ3Y.XRD profiles was observed from Fig. 13(a) when only a tetragonal lattice was used for the refinement. The fit, however, The alteration of the crystal structure from tetragonal to 2792 J. Mater.Chem., 1998, 8, 2787–2794Table 2 A summary of reaction products between PSM and TZ3Y at 1200 °C Products Specimen 4 h 24 h 168 h PSM-A/TZ3Y DiVusion layera Pr2Zr2O7 nuclei Pr2Zr2O7 islands DiVusion layer DiVusion layer PSM-B/TZ3Y DiVusion layer Pr2Zr2O7 layer Pr2Zr2O7 layer DiVusion layer DiVusion layer PSM-C/TZ3Y (Pr,Zr)Ox layer (Pr,Zr)Ox layer (Pr,Zr)Ox islands Pr2Zr2O7 layer Pr2Zr2O7 layer Pr2Zr2O7 layer DiVusion layer DiVusion layer DiVusion layer aDiVusion layer refers to Pr- and Mn-diVused zirconia layer.cubic indicates that at least part of the Pr and Mn ions have In PSM-A and PSM-B, there was no free praseodymia in the coating. The atomic ratio of Pr/Zr>1 at the interface with entered the lattice of zirconia. The interfacial reaction products between TZ3Y and PSM TZ3Y was unlikely to occur, and therefore, the Pr-rich phase (Pr,Zr)Ox was not formed.of three compositions at 1200 °C are summarised in Table 2. The formation of pyrochlore phase was delayed in PSM-A/TZ3Y and PSM-B/TZ3Y. It has been reported before 4 Discussion that A-site deficiency in the perovskite may suppress or delay the formation of pyrochlore phase in the LSM/YSZ The above results can be summarised as follows. 1 A diVusion layer of Pr and Mn in zirconia was formed system.4,6,11,13 This was explained by the hypothesis that the diVusion of Mn into YSZ produced chemically active La2O3, in all specimens regardless of the A-site stoichiometry and the heat treatment temperature. The Pr- and Mn-diVused zirconia which formed pyrochlore phase with YSZ.Therefore, extra manganese in the perovskite should suppress the formation of is cubic and its grain size is much larger than that in TZ3Y. 2 The pyrochlore phase was formed at the interface of PSM free La2O3, and hence the pyrochlore phase. However, this hypothesis cannot explain why the formation of pyrochlore with Pr- and Mn-diVused zirconia after heat treatment at 1200 °C for diVerent times for all compositions of PSM studied.phase was delayed in PSM-B/TZ3Y with no excess Mn in the perovskite. Furthermore, the results of the current study Significant amounts of pyrochlore phase were detected after 168 h in PSM-A/TZ3Y, after 24 h in PSM-B/TZ3Y and after showed clearly that the Pr and Mn ions had diVused into zirconia, regardless of the stoichiometry of PSM.The chemi- 4 h in PSM-C/TZ3Y. 3 Besides the pyrochlore phase, a Pr-rich (Pr,Zr)Ox phase cally active praseodymia was unlikely to have been produced during the interaction. It seems that a diVerent explanation is was detected at the interface of the A-site over-stoichiometric PSM-C/TZ3Y. The (Pr,Zr)Ox phase formed a distinct layer required to understand the delay for the formation of the pyrochlore phase in PSM-A/TZ3Y and PSM-B/TZ3Y. after 24 h of heat treatment at 1200 °C, and was much less conspicuous after 168 h.From the phase relations, it is known that the pyrochlore phase is formed only when the local Pr concentration has 4 Relatively thick layers of pyrochlore phase were formed after 24 h at 1400 °C for all compositions of PSM in contact reached the solubility limit in cubic zirconia.When there was no extra praseodymia in the PSM coating as in PSM-A/TZ3Y with TZ3Y. The degree of the reaction for PSM-B/TZ3Y was much lower than that for PSM-A/TZ3Y and PSM-B/TZ3Y. and PSM-B/TZ3Y, the atomic ratio of Pr/Zr was initially very low in zirconia in the area near the interface. With time more The dissolution of Pr ions in TZ3Y can be understood from the phase relations between zirconia and praseodymia. Zirconia Pr ions diVused into the region near the interface.Some of it diVused out into an area in zirconia further away from the can react with praseodymia, forming solid solutions varying from tetragonal zirconia, cubic zirconia, pyrochlore phase to interface due to the gradation of Pr concentration. It would therefore take some time at 1200 °C for the Pr ions to reach praseodymia with the increase of the praseodymia content at 1600 °C.20 The solid state reactions between zirconia and the solubility limit in cubic zirconia in the area near the interface.Accordingly, the formation of the pyrochlore phase praseodymia at 900 and 1100 °C have also been reported.21 The current study has shown that Pr ions can enter the lattice was delayed in PSM-A/TZ3Y and PSM-B/TZ3Y. As the Pr content in PSM-A was lower than in PSM-B, the migration of zirconia at 1200 and 1400 °C and form a cubic phase of zirconia solid solution. The Pr concentration increased in the of Pr ions from PSM-A to TZ3Y is therefore expected to be slower than from the stoichiometric PSM-B during the initial cubic zirconia with the heating time.When the solubility limit of Pr ions in cubic zirconia was reached, the pyrochlore phase stage of heat treatment. Therefore the delay for forming pyrochlore phase with PSM-A was more pronounced than crystallised out. In the A-site over-stoichiometric PSM-C coating, a that with PSM-B at 1200 °C. At 1400 °C the whole process was accelerated.A complete considerable amount of free praseodymium oxide is present, which reacted with TZ3Y when heat treated at high tempera- layer of pyrochlore phase was formed for all specimens after 24 h. It needs to be noted that diVerent from the observations ture. A gradation of Pr concentration from the Pr6O11/TZ3Y interface (high) to the Pr-diVused zirconia (low) is expected.at 1200 °C, the amount of pyrochlore phase formed with PSMA at 1400 °C is much more than with PSM-B. This may be Therefore, various layers ranging from (Pr,Zr)Ox with atomic ratio of Pr/Zr>1, the pyrochlore with Pr/Zr$1, to the explained from the change of the migration rate of Pr ions at various stages of heat treatment and in particular, the diVerence diVusion layer with Pr/Zr<1, were formed between PSM and TZ3Y at 1200 °C.However, when the free Pr6O11 was con- in PSM/substrate contact areas between the two specimens. As the Mn content in PSM-A was higher than in PSM-B, the sumed, no further (Pr,Zr)Ox phase was formed. The Pr ions from the (Pr,Zr)Ox solid solution, which contains the highest migration of Mn ions from PSM-A to zirconia is expected to be faster than from PSM-B during the initial stage of heat Pr concentration in the specimen, would diVuse through the pyrochlore layer and react with more zirconia.Therefore, the treatment. On the other hand, the higher Pr content in PSMB resulted in a further migration of Pr ions from PSM-B to amount of (Pr,Zr)Ox solid solution after 168 h was less than that after 24 h at 1200 °C.It is expected to eventually disappear zirconia. This means that more Mn ions would leave from PSM-A than from PSM-B and more Pr ions would leave from completely. J. Mater. Chem., 1998, 8, 2787–2794 2793PSM-B than from PSM-A. It is quite likely that after some MnO3/TZ3Y were reached after much longer time compared with (Pr0.8Sr0.2)1.05MnO3/TZ3Y. The formation of the pyroch- time of heat treatment (the initial stage), the chemical composilore phase, therefore, was delayed in these two specimens at tions for PSM-A and PSM-B may become somewhat similar. 1200 °C. The amount of the pyrochlore phase formed at the After this initial stage, the amount of the pyrochlore formed interface was determined by the A-site stoichiometry of PSM between PSM and the substrate will be primarily determined in the initial stages of heat treatment.The growth of the by the contact area between the two materials. A large contact pyrochlore layer after the initial stage, however, appeared to area, as present in PSM-A (Fig. 2), will lead to enhanced Pr be controlled by the contact area between PSM and the migration from PSM to the substrate.This mechanism would substrate. It is believed that the initial stage is complete after explain the fact that the pyrochlore layer formed with PSM-A 24 h at 1400 °C but not even after 168 h at 1200 °C for all the after 24 h at 1400 °C is thicker than with PSM-B. compositions of PSM. The A-site deficient (Pr0.8Sr0.2)0.9MnO3 It appears that at the initial stage of heat treatment, the in contact with TZ3Y, following long term heat treatment, will amount of Pr ions migrating into TZ3Y was determined mainly form more pyrochlore phase than the stoichiometric by the A-site stoichiometry of the perovskite.The composition (Pr0.8Sr0.2)MnO3 due to the large contact area between PSM- diVerence between the diVerent perovskites may disappear A and TZ3Y. after the initial stage of the heat treatment is complete.The amount of Pr ions diVusing into the substrate was then determined mainly by the contact area between PSM and the Acknowledgements substrate. It seems that the initial stage is complete after 24 h The authors are grateful to Dr. R. Ratnaraj and Mr. K. at 1400 °C but not yet after 168 h at 1200 °C for all the Wilshier for reviewing this paper, Mr. D.Milosevic for compositions of PSM. synthesising the PSM powders, Mr. R. Donelson for supplying The relatively large contact area between PSM-A and the the TZ3Y substrates, Mr. H. Jaeger for assistance with substrate may be attributed to the high sinterability for the A- microscopy, Natasha Rockelmann with XRD, and Kristine site deficient perovskite. This has been observed previously by Giampietro and Kylie Chapman with specimen and microother authors.13 It was assumed that the vacancy within the graph preparation.structure of the A-site deficient perovskite enhanced the cation diVusion in perovskite itself and hence improved the References sinterability of the perovskite. 1 N. Q. Minh, J. Am. Ceram. Soc., 1993, 76, 563. The highest reaction degree for PSM-C at 1400 °C among 2 S.P. S. Badwal and K. Foger, Mater. Forum, 1997, 21, 183. all the diVusion couples can be attributed to the presence of 3 L. Kindermann, D. Das, D. Dahadur, R. Weiß, H. Nickel and free Pr6O11 at the initial stage as well as to the high contact K. Hilpert, J. Am. Ceram. Soc., 1997, 80, 909. area between PSM-C and the substrate. 4 A. Mitterdorfer, M.Cantoni and L. J. Gauckler, in Second European Solid Oxide Fuel Cells Forum, ed. B. Thorstensen, Mn was always detected in the diVusion layer in the current Druckerei J. Kinzel, Gottingen, 1996, p. 373. study. It appears that both Pr and Mn migrate out of the 5 D. Kuscer, J. Holc, M. Hrovat, S. Bernik, Z. Samardzija and PSM electrode (for all compositions) into the electrolyte.D. Kolar, Solid State Ionics, 1995, 78, 79. However, some Pr is trapped by the formation of the pyro- 6 G. Stchniol, E. Syskakis and A. Naoumidis, J. Am. Ceram. Soc., chlore phase, while Mn keeps on diVusing into zirconia. This 1995, 78, 929. may explain much deeper diVusion zone for Mn than that for 7 H. Taimatsu, K. Wada and H. Kaneko, J. Am. Ceram. Soc, 1992, 75, 401.Pr even after longer heat treatment times. There is no consist- 8 C. Brugnoni, U. Ducati and M. Scagliotti, Solid State Ionics, 1995, ency in the literature about the manganese diVusion in yttria 76, 177. stabilised zirconia. Some authors have reported manganese 9 G. Stochniol, H. Grubmeier, A. Naoumidis and H. Nickel, in diVusion in zirconia,3,6,22,23 whereas others did not detect any Proceedings of the Fourth International Symposium on Solid Oxide Mn in zirconia.4,8,12 Waller et al.recently studied the manga- Fuel Cells (SOFC-IV), ed. M. Dokiya, O. Yamamoto, H. Tagawa and S. C. Singhal, The Electrochemical Society, Pennington, NJ, nese diVusion in single crystal and polycrystalline yttria stabil- 1995, p. 995. ised zirconia using XRD and dynamic secondary ion mass 10 C.Clausen, C. Bagger, J. B. Bilde-Sorensen and A. Horsewell, spectrometry techniques, and showed convincingly that Mn Solid State Ionics, 1994, 70/71, 59. ions have diVused into both single crystal and polycrystalline 11 T. Tsepin and S. A. Barnett, Solid State Ionics, 1997, 93, 207. zirconia.24 The diVusion coeYcient of Mn in polycrystalline 12 J. A. M. van Roosmalen and E.H. P. Cordfunke, Solid State Ionics, zirconia was significantly higher than that in the single crystal. 1992, 52, 303. 13 J. W. Stevenson, T. R. Armstrong and W. J.Weber, in ref. 9, p. 454. They concluded that in polycrystalline zirconia, the grain 14 H. Y. Lee and S. M. Oh, Solid State Ionics, 1996, 90, 133. boundary diVusion of Mn dominates the mechanism for Mn 15 T. Ishihara, T.Kudo, H. Matsuda and Y. Takita, J. Am. Ceram. migration into zirconia. The substrate used in the current Soc, 1994, 77, 1682. study is polycrystalline TZ3Y that has a much higher density 16 T. L. Wen, H. Tu, Z. Xu and O. Yamamoto, in Extended Abstracts of grain boundaries to provide diVusion paths for manganese of 11th International Conference on Solid State Ionics, Honolulu, Hawaii, Nov. 16–21, 1997, p. 188. than the 7.5 mol% yttria stabilised zirconia used by Waller 17 H. M. Rietveld, J. Appl. Crystallogr., 1969, 2, 65. et al.24 Therefore, a considerable amount of manganese 18 R. J. Hill and C. J. Howard, Report No. M112, Australian Atomic diVusion into TZ3Y at both 1200 and 1400 °C is expected as Energy Commission (now ANSTO), Lucas Heights Research indeed was the case in the present study. Laboratories, New SouthWales, Australia, 1986. 19 C. J. Howard, R. J. Hill and B. E. Reichert, Acta Crystallogr., Sect. B, 1988, 44, 116. 5 Conclusions 20 R. L. Withers, J. G. Thompson and P. J. Barlow, J. Solid State Chem., 1991, 94, 89. The interfacial reactions between (Pr0.8Sr0.2)yMnO3 electrode, 21 M. K. Nasakar and D. Ganguli, J.Mater. Sci., 1996, 31, 6263. with diVerent A-site stoichiometry, and TZ3Y electrolyte at 22 A. Khandkar, S. Elangovan and M. Liu, Solid State Ionics, 1992, 1200 and 1400 °C in air have been studied. It has been found 52, 57. that both Pr and Mn ions diVused into TZ3Y forming a cubic 23 S. K. Lau and S. C. Singhal, in Corrosion 85, The National zirconia solid solution. When the solubility limit of Pr ions in Association of Corrosion Engineers (NACE) Meeting, Boston, MA, March 25–29, 1985, p. 345/1. the cubic zirconia was reached, a pyrochlore phase was formed. 24 D.Waller, J. D. Sirman and J. A. Kilner, in Proceedings of the Fifth When there was free praseodymia in the PSM coating, a International Symposium on Solid Oxide Fuel Cells (SOFC-V), ed. praseodymia–zirconia solid solution (Pr,Zr)Ox was formed, U. Stimming, S. C. Singhal, H. Tagawa and W. Lehnert, The which occurred only in the case of the (Pr0.8Sr0.2)1.05- Electrochemical Society, Pennington, NJ, 1997, p. 1140. MnO3/TZ3Y couple. The solubility limit of Pr in zirconia for specimens (Pr0.8Sr0.2)MnO3/TZ3Y and (Pr0.8Sr0.2)0.9- Paper 8/05835K 2794 J. Mater. Chem., 1998, 8, 2787–2794

ISSN:0959-9428    

DOI:10.1039/a805835k

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Oriented growth of hydroxyapatite on (0001) textured titanium with functionalized self-assembled silane monolayer as template Chuanbin Mao,* Hengde Li, Fuzhai Cui, Qinglin Feng, Hao Wang and Chunlai Ma Department of Materials Science and Engineering, Tsinghua University, Beijing 100084, People’s Republic of China Received 18th February 1998, Accepted 19th August 1998 A highly (0001) textured hydroxyapatite [Ca10(PO4)6(OH)2, HA] coating on polycrystalline titanium plate is successfully synthesized by a biomimetic process mimicking biomineralization.To simulate the first stage of biomineralization, that is, supramolecular preorganization, a template surface with highly organized arrangement of carboxyl (KCOOH) and alcoholic hydroxyl (KOH) groups is prepared through self-assembly of vinyltriethoxysilane [(C2H5O)3SiCHLCH2, VTS] on hydroxylated titanium with strong (0001) texture, followed by oxidation of the vinyl groups (KCHLCH2) with dilute KMnO4 solution into alcoholic hydroxyl and then into carboxyl groups.The functionalized substrate can induce oriented nucleation and growth of HA with (0001) planes parallel to the substrate surface from supersaturated HA solution through interfacial molecular recognition.The mechanisms of molecular recognition are also discussed. is to highly texture the substrate before functionalization 1 Introduction through self-assembly, resulting in a large quantity of domains Biomineralization has been regarded as an excellent archetype with nearly the same arrangement of surface atoms.Then a for inorganic materials synthesis by materials chemists.1 Four surface with a large quantity of domains with nearly the same sequential stages are involved in the biomineralization includ- arrangement of functional groups may be prepared after ing supramolecular preorganization, interfacial molecular rec- chemisorption and self-assembly of functional molecules. That ognition, vectorial regulation and cellular processing.1 In is, an organized SAM can be fixed on a highly textured recent years, with these four stages as archetypes, a new substrate through a self-assembly process.synthetic strategy, termed biomimetic synthesis or template In this paper, we hope to demonstrate the above synthesis, has been developed to synthesize a variety of inor- consideration through biomimetic synthesis of an HA coating ganic materials such as nanoparticles,2 thin films,3 coatings,3 on polycrystalline titanium. The titanium is textured with porous materials4 and materials with complex forms.5 (0001) preferably parallel to the surface through rolling defor- Currently, biomimetic synthesis of inorganic materials through mation since the (0001) texture is a typical kind of rolling a process mimicking biomineralization has been a very deformation induced texture for hexagonal metals.8 The funcpromising approach to prepare materials of low cost.5 tionalized surface is prepared through initial hydroxylation To perform a successful biomimetic synthesis, it is most with H2O2 and subsequent self-assembly of vinyltriethoxysilane important to mimic the first stage of biomineralization (i.e., (VTS), followed by oxidation of vinyl groups with dilute supramolecular preorganization) where a highly organized KMnO4 solution into alcoholic hydroxyl groups and then into reaction template is built through self-assembly of an organic carboxyl groups.The principle of such a biomimetic process matrix.For synthesis of films or coatings on inorganic sub- is illustrated in Fig. 1. strates, a surface with organized functional groups is often prepared to act as reaction template through chemisorption 2 Experimental and self-assembly of functional organic molecules on the substrates.3,6,7 The functional molecules bond with the surface 2.1 Biomimetic synthesis atoms which are usually coordinatively unsaturated. Therefore, A supersaturated solution of HA ([Ca2+]=4mM) was prepared the arrangement of the functional organic molecules is often according to the procedure in ref. 9. A 2 mm thick titanium determined by the surface lattice. For a single crystal substrate, plate was prepared by rolling deformation to introduce the the surface to be functionalized is always a single plane (i.e., (0001) texture and was cut into small square pieces no grain boundaries and little variation in atomic arrangement).For a polycrystalline substrate, the arrangement of atoms usually diVers significantly between two exposed crystal grains unless the substrate surface is highly textured with one plane preferably parallel to the surface, therefore the arrangement of chemisorbed and self-assembled functional molecules is usually not uniform and organized across the whole surface.Hence, the successful biomimetic growth of inorganic films or coatings is usually realized on self-assembled monolayer (SAM) covered single crystals.6,7 However, practical films or coatings are usually on polycrystalline substrates, for example, the hydroxyapatite (HA) coating on titanium as biomedical bone implant.Therefore, it is necessary to find a way to obtain a highly organized functionalized surface on polycrystalline substrates. We consider that one eYcient way to increase the degree of Fig. 1 Principle of biomimetic synthesis of HA coating on titanium through a process mimicking biomineralization. organization of the arrangement of the functional molecules J.Mater. Chem., 1998, 8, 2795–2801 2795(1 cm×1 cm). The plates were metallographically polished with SiC emery paper to remove the oxide surface layer. The final polishing was performed with No. 800 paper and the final thickness of the polished plates was about 1 mm. The plates were ultrasonically washed in acetone for ca. 10 min and rinsed in deionized water for 1 min, and then placed in a 10% (by volume) solution of VTS in benzene for 2 weeks.After being rinsed in deionized water for several seconds, they were further aged in 5 mass% aqueous alkaline KMnO4 solution first at 0 °C for one week and then in 5 mass% aqueous KMnO4 solution containing sodium periodate (NaIO4) at room temperature for one week. The resulting plates are denoted as VTS-Ti.One VTS-Ti plate was put into 50 ml HA solution and aged for up to 2 months. For comparison, a 1 cm×1 cm×2 mm control titanium plate with poor (0001) texture was prepared by a diVerent rolling procedure,8 and subjected to the subsequent experiment as above. 2.2 Characterization Fig. 3 Time dependences of pH in VTS-Ti solution and control The changes of pH during aging were measured in situ with a solution.numeric acidity meter. The phase composition and orientation of the substrate and coating were determined by X-ray diVracnaked eye upon aging both titanium plates in HA solutions. tion (XRD) on a Rigaku D/max RD diVractometer with a In HA solution containing VTS-Ti (VTS-Ti solution), crys- Cu target.To observe the morphologies of the coating by tallization only occurs on the surface of VTS-Ti to form a scanning electron microscopy (SEM), a knife was used to cut continuous white layer, indicating that heterogeneous between the coating and the substrate, and then the coatings nucleation and growth of calcium phosphates is favored on were fractured to expose a fresh fractured cross section, VTS-Ti.However, in the HA solution containing control VTS- followed by coating with gold by ion beam sputtering. SEM Ti (control solution), crystallization occurs not only on the images were observed by a Hitachi S-450 scanning electron surface of control VTS-Ti to form several deposits of island- microscope coupled with an energy dispersive X-ray analysis like white precipitates, and at the air/solution and solution/ (EDAX) system.Functional groups were detected by X-ray beaker interfaces to form island-like white precipitates, but photoelectron spectroscopy (XPS) on an ESCALAB 220i-XL also in the bulk solution to form floating white precipitates. system with Mg-Ka source. In VTS-Ti solution, the white coating can be observed by the naked eye after about 2 days and grows thicker with aging for 3 Results about one week.Crystallization elsewhere was not observed for up to 2 months. In the control solution, the island-like 3.1 Coating formation precipitates did not form a continuous layer for up to 2 The XRD results in Fig. 2 indicate that the degree of (0001) months, which is consistent with the results of Hanawa and texture of VTS-Ti is much higher than that of control VTS- Ota.10 Ti through the comparison of the relative intensity ratio Crystallization of HA can be written in terms of eqn.(1) between (0001) and (10191) diVraction peaks. Much diVerent 10Ca2++6PO43-+2OH-=Ca10(PO4)6(OH)2 (1) crystallization of calcium phosphate was observed by the and reduction of pH can be an eVective indicator of HA crystallization on titanium or elsewhere in the bulk solution.Fig. 3 depicts the time dependences of pH values of the VTSTi solution and control solution. The pH of the VTS-Ti solution decreases after an incubation period due to the crystallization on VTS-Ti and levels oV at about 6.6 after about one week indicative of a typical nucleation and growth process. For the control solution, the pH begins to decrease a little later and then decreases more rapidly.It also levels oV at about 6.6 after ca. 3 days. 3.2 Coating characterization The XRD spectrum for VTS-Ti aged for 2 days as shown in Fig. 2(c) suggests that the HA crystals have nucleated and grown with (0001) preferably parallel to the substrate. After VTS-Ti has been aged for 7 days, i.e., when the coating on VTS-Ti no longer grows as judged from a constant solution pH (Fig. 3) the coating is composed of highly (0001) textured HA as shown in Fig. 2(d). Since only several small deposits of island-like precipitates occur on the control VTS-Ti, they can not be characterized by XRD and thus SEM-EDAX is adopted. Fig. 4 shows the SEM morphologies of the fractured cross-sections of the coating on VTS-Ti at diVerent stages.After being aged for 2 days [Fig. 4(a)], the coating on VTSTi is ca. 10 mm thick. Previous studies have shown that HA Fig. 2 XRD patterns of (a) control VTS-Ti; (b) VTS-Ti; (c) coating on VTS-Ti after 2 days and (d) coating on VTS-Ti after 7 days. grains always tend to grow along the [0001] direction, which 2796 J. Mater. Chem., 1998, 8, 2795–2801of plate-like HA grains which are ca. 10 mm wide. Since it is diYcult to make a fractured cross-section of one island, the length of the HA grains can not be determined. However, analogous to the grains on VTS-Ti, the length of HA grains may be ca. 40 mm. Most of the ball-like islands are much thicker than the coating on VTS-Ti. This is because there are several HA grains along the normal of the substrate.From the SEM images in Fig. 5(b), the plate-like HA grains have grown toward the solution preferably with [0001] parallel to the normal of the substrate. These results demonstrate our considerations in section 1, that is, a functionalized surface on the textured substrate can be an eVective template for biomimetic growth. 4 Discussion 4.1 Formation of an organized and hydroxylated surface After the old oxide surface layer is removed through polishing, a fresh oxide layer will quickly form on the plate. Fig. 6 shows the XPS Ti 2p spectrum of a titanium surface before treatment with H2O2. Curve-fitting revealed that the spectrum can be resolved into two peaks centered at 463.89 and 458.27 eV, corresponding to Ti 2p1/2 and Ti 2p3/2 in TiO2, respectively.12,13 It is thus evident that the surface oxide layer of titanium is Fig. 4 SEM morphologies of the fractured cross sections of coatings TiO2. Many researchers also found that TiO2 is the natural on VTS-Ti, after 2 (a), 4 (b) and 7 days (c). A top-view SEM oxide layer on the titanium at room temperature.13–15 Since morphology of the coating on VTS-Ti after 7 days is shown in (d).Regions A in (a) denote the crystallization area that is not on the thermodynamically stable form of TiO2 at low temperature SAM-(110)TiO2 domains. Note that the incident electron beam was is anatase (tetragonal, a=0.3783 nm, c=0.951 nm, JCPDS 4- tilted by 45° in (a), (b) and (c). 477),16 the surface oxide layer can be regarded as the anatase form.17 Recently, Azoulay et al.18 made a systematic investiresults in a plate habit, and the long axes of the plate-like gation on the interactions of oxygen with polycrystalline grains are parallel to the crystallographic [0001] direction.11 titanium surfaces at a very low pressure of O2.They found Thus, one can judge the type of preferred orientation by the that the initial accumulation of oxygen follows an island fashion of the grain arrangements shown in the SEM images.formation model, resulting in a patch-like pattern of oxide From Fig. 4(a), one can observe that the HA crystals have surface layer containing diVerent titanium valence states, and been oriented with [0001] perpendicular to the substrate, the oxide with better lattice match with underlying parent consistent with the XRD result in Fig. 2(c). After 4 days titanium structure will be favored to form first. Ti2O (hexag- [Fig. 4(b)], the coating is ca. 30 mm thick and the highly onal, a=0.29593 nm, c=0.4845 nm, JCPDS 11-218) shows (0001) textured state is very obvious in the SEM figure. After the best lattice match with titanium (a=0.2950 nm, c= 7 days [Fig. 4(c)], the coating is ca. 40 mm thick and consists 0.4686 nm, JCPDS 5-682) among all the titanium oxides and of highly textured plate-like HA grains, consistent with the can be directly transformed from titanium through the ordering XRD result in Fig. 2(d). From the top-view SEM morphology of the dissolved oxygen atoms (cf. notes on JCPDS 11-218). of the coating on VTS-Ti [Fig. 4(d)], one can see that the Electron microscopy also revealed that Ti2O can grow out of coating is nearly uniform across the whole surface.The above titanium through coherent transformation.19 In addition, Ti2O results suggest that HA crystals nucleate on VTS-Ti with has the lowest content of oxygen among all the titanium (0001) preferably parallel to the substrate and grow toward oxides.20 Consequently, the fresh surface oxide layer, which is the solution, giving rise to a textured coating with the thickness too thin to be detected by XRD, may be formed from Ti equal to the length of one HA crystal along the [0001] directly to TiO2, or through intermediate oxides with low direction. Fig. 5 shows the top-view SEM morphologies of the oxygen content, i.e., low chemical valence of titanium, such island-like precipitates on the control VTS-Ti after 7 days.as Ti2O, and finally to TiO2. The intermediate Ti2O will be This reveals that the precipitates consist of islands, which are also (0001) textured due to coherent transformation from significantly diVerent from the coating on VTS-Ti as shown (0001) textured titanium. Since there is a good lattice matching in Figs. 4(c) and (d).The islands are ball-like in shape and their size ranges between 50 and 150 mm. They are composed Fig. 5 Top-view SEM morphologies of island-like precipitates on Fig. 6 Curve-fitting of the Ti 2p XPS spectrum for titanium before control substrate after aging for one week: (a) whole morphology and (b) one island at higher magnification. H2O2-treatment. J. Mater. Chem., 1998, 8, 2795–2801 2797Fig. 7 Schematic illustration of the lattice match between (0001)Ti and Fig. 8 Schematic illustration of formation and functionalization of a (110)anatase. The dashed and bold line frames show the outlines of the self-assembled silane monolayer. unit cells of Ti along [0001] and TiO2 along [110], respectively. To demonstrate the formation of functionalized SAM, VTS- between (0001)Ti or (0001)Ti2O and (110)anatase (for structure Ti was subjected to XPS characterizations.Fig. 9 shows the O of anatase see ref. 16) as illustrated in Fig. 7, the transformed 1s XPS spectrum of VTS-Ti. Curve-fitting revealed three peaks TiO2 will be (110) textured due to the (0001) texture of with binding energies of 533.04, 531.32 and 529.65 eV, corre- underlying titanium and may be why the texture of substrate sponding to O 1s electrons in KOKSi(O)KOK before SAM-functionalization plays an important role in biom- Si(O)KOK( like SiO2), KCOOH and KOH, and surface oxide imetic mineralization of HA.It should be stated that the lattice, respectively.12 It is worthwhile to point out that due process of surface oxidation on titanium in air has scarcely to adsorption of some other phases containing oxygen such been studied perhaps because the oxidation from Ti to TiO2 as H2O and KMnO4, their contribution in O 1s photoemission is very quick and TiO2 is always the main detected oxide.13–15 peaks may be overlapped and cannot be resolved from the During subsequent aging in H2O2 solution, a surface interabove species.For example, the O 1s binding energy in KMnO4 action will occur between the fresh oxide layer formed on the is very similar to that in KCOOH.12 Therefore, the C 1s XPS titanium and H2O2 to form active hydroxyl groups.21,22 The spectrum (Fig. 10) may be more suitable to characterize the peroxide ion (O22-) can be regarded as an active state of O2.functionalized SAM. Curve-fitting showed three peaks with Its electron configuration is sls2, s*ls2, s2s2, s*2s2, s2px2, binding energies of 288.24, 286.02 and 284.63 eV, which corre- p2py2, p2pz2, p*2py2, p*2pz2, s*2px0 where s and p refer to spond to carbon in KCOOH, KCH(OH)CH2OH, KCHLCH2 bond-forming orbitals, and s* and p* to antibonding orbitals, respectively and there is an empty s*2px orbital in O22-.21,22 The coordinatively unsaturated surface (c.u.s.) Ti4+ ions have unbonded valence electrons and will interact with O22- in H2O2.The coordinatively saturated bulk Ti4+ ions show an outermost electron configuration of 3d04s0. Some c.u.s. titanium sites may occur as Ti3+ (3d14s0).23,24 During the surface reaction, 3d electrons will be favored to enter empty s*2px orbitals, resulting in cleavage of the OKO bond in O22- and thereby the formation of hydroxyl groups (KOH).21,22 The fresh KOH will be thermodynamically favored to be chemisorbed on the surface oxide layer through bonding with c.u.s.Ti4+ ions.17,25 In addition, chemisorbed hydroxyl groups can also be formed through dissociation of H2O on titanium. However, the concentration of such chemisorbed hydroxyl groups is very low because the major part of H2O is physisorbed. 17,25 Hence, a highly hydroxylated surface will be Fig. 9 Curve-fitting of the O 1s XPS spectrum for VTS-Ti. realized after H2O2 treatment. The chemisorbed hydroxyl groups on the c.u.s. Ti4+ ions of the fresh oxide layer can be denoted as TiKOH. Moreover, the texturing of the substrate before chemisorption of hydroxyl groups makes a majority of TiO2 crystals with (110) planes exposed, that is, a majority of domains with same c.u.s. Ti4+ arrangement on surface and thus a majority of domains with the same arrangement of TiKOH.Consequently, the OH groups on the textured substrate are more organized than a nontextured substrate, and may be more suitable to act as a template for biomimetic growth of coatings after further functionalization. 4.2 Formation of a self-assembled monolayer (SAM) According to known organic processes,26 vinyl groups will be oxidized by dilute KMnO4 solution first to give alcoholic hydroxyl groups at low temperature and then into carboxyl groups at room temperature and the formation of VTS-SAM Fig. 10 Curve-fitting of the C 1s XPS spectrum for VTS-Ti. and further functionalization of SAM is illustrated in Fig. 8.27 2798 J. Mater. Chem., 1998, 8, 2795–2801and contaminated carbon inherent to the XPS spectrometer, respectively.12,28 These results suggest that the oxidation of KCHLCH2 into KCH(OH)CH2OH and further into KCOOH is not complete, which is characteristic for this organic reaction. Therefore, the SAM on VTS-Ti is a mixture of OH-terminated and COOH-terminated silane moleculars fixed on the titanium plate, i.e., a mixture of states as shown in Fig. 8(b)–(d).Since we are dealing with a self-assembly process and the hydroxylated surface on highly (0001) textured titanium is highly organized, the mixture may be uniform, i.e., the arrangement of KCHLCH2, OH and COOH groups in the functionalized SAM on highly textured titanium is highly organized.Among these groups, OH and COOH can induce biomimetic mineralization6,7 and may be the reason why control VTS-Ti can not induce coating formation; this will be discussed later. 4.3 Mechanism of oriented coating formation In biomineralization, oriented nucleation and growth result from the mediation by preorganized supramolecules through interfacial molecular recognition including the complementarities of lattice geometry, electrostatic potential, polarity, stereochemistry, space symmetry and topography.1 Similarly, in this biomimetic synthesis, the oriented nucleation and growth of HA originate from the regulation by OH- and COOHfunctionalized SAM fixed on textured titanium through interfacial molecular recognition.The interfacial molecular recognition involves at least four aspects: (1) Crystal lattice matching.The arrangement of OH and COOH groups in SAM fixed on the (110) oriented TiO2 layer along [001]anatase shows excellent one-dimensional (1-D) coherent matching with the (0001) plane of HA along [01190]HA as illustrated in Fig. 11.29 (2) Hydrogen bonding interaction. OH and COOH on SAM can form hydrogen bonds not only with OPO33- on (0001) planes of HA, i.e., OH,OPO3 and Fig. 11 Schematic illustration of the one-dimensional lattice match relation between (0001) planes of HA and SAM treated TiO2. COOH,OPO3,29,30 but also with OH- on the (0001) planes (a) Idealized crystal structure of HA[30] (note that the atomic arrange- of HA, i.e., OH,OKH and COOH,OKH, respectively. ment of PO4 and OH is not resolved for clarity), (b) (0001) plane of (3) Electrostatic potential interaction.The negatively charged HA at z=3/4 and (c) lattice overlapping between the functionalized COO- on SAM can attract the positively charged Ca2+ in (110) plane of TiO2 and (0001) plane of HA. The dashed and bold HA solution to substrate surface. (4) Stereochemistry match. line frames show the outlines of the unit cells of HA along [0001] and The directions of the valence bonds of OH and OH in COOH TiO2 along [110], respectively. tend to be parallel to that of OKH on (0001) planes of HA (parallel to the c-axis of HA).29 The recognition elements (1) and (4) explain why (0001) Thus finally a continuous textured coating can be formed on VTS-Ti [Fig. 12(d)]. However, on control VTS-Ti, few isolated textured HA is favored to form.Now one can understand why VTS-Ti can induce formation of (0001) textured HA domains are exposed with (110) planes of TiO2, and then the eVective molecular recognition, especially the 1-D lattice coating while control VTS-Ti can only induce formation of island-like precipitates, and moreover, both the coating on matching, can only occur on such few isolated domains.Consequently, for control VTS-Ti, after HA crystals nucleate VTS-Ti and the island-like precipitates on the control substrate exhibit order on a length scale larger than the expected size and grow on these few SAM-(110)TiO2 domains as the case on VTS-Ti [Fig. 12(a¾) and (b¾)], the degree of supersaturation is of SAM-(110)TiO2 domains as shown in Fig. 4 and 5. This can be explained by the formation process of HA on the substrate relatively higher for HA to nucleate in head-to-head fashion, through a 2-D lattice match between (0001) planes of the new as illustrated in Fig. 12. On VTS-Ti, many domains are exposed with (110) planes of TiO2 and functionalized with OH or HA nuclei and the underlying HA crystals31 [Fig. 12(b¾) and (c¾)] and also in other places such as the beaker/solution COOH groups and thus there will be many domains in which all the above molecular recognition processes are active and and air/solution interface.The distance between two SAM-(110)TiO2 domains is much longer for the control sub- induce oriented nucleation and growth of HA with (0001) planes parallel to VTS-Ti. Therefore, HA crystals nucleate strate than for VTS-Ti.For control VTS-Ti, before the crystallization can cover the exposed area between preferably on many SAM-(110)TiO2 domains with (0001) preferably parallel to the substrate on such domains [Fig. 12(a)]. SAM-(110)TiO2 domains, the solution has been saturated. Therefore, there will be exposed areas where no HA crystallizes As these crystals are growing toward solution along the [0001] direction, HA crystals will also begin to nucleate and grow on and island-like precipitates exist on the substrate [Fig. 12(d¾)].One can thereby understand why the height of islands on other SAM domains in the nearby pores between SAM-(110)TiO2 domains close to the growing crystals in a side- control substrate is much larger than that of the coating on VTS-Ti, and both the coating and the island-like precipitates by-side fashion, through the above recognition processes except lattice matching between SAM and nuclei and 2-D show order on a length scale larger than the expected size of SAM-(110)TiO2 domains.It is possible that the final coating lattice matching along substrate normal to the growing HA crystals and the new nuclei31 [Fig. 12(b) and (c)].The crystals on VTS-Ti will tend to be uniform as shown in Fig. 4 because there will be a larger driving force for growth for smaller grown according to this fashion are labelled A in Fig. 4(a). J. Mater. Chem., 1998, 8, 2795–2801 2799HA will begin earlier on VTS-Ti than on control VTS-Ti, and the crystallization always occurs on VTS-Ti before elsewhere in the VTS-Ti solution while relatively few crystals crystallize on control VTS-Ti before crystallizing elsewhere in the control solution.Thus the pH decreases at a earlier stage in the VTSTi solution than in the control solution, i.e., the incubation period is shorter in the VTS-Ti solution than in the control solution. After the incubation period, the crystallization on VTS-Ti is controlled by the reaction template through the above recognition in VTS-Ti solution.However, in the control solution, crystallization occurs elsewhere such as at the air/ solution and beaker/solution interface besides the control substrate due to the existence of a much smaller area of the eVective reaction template. Thus the pH value decreases more rapidly and levels oV more quickly in the control solution than in the VTS-Ti solution, i.e., the nucleation and growth period is longer in the VTS-Ti solution than in the control one. 4.4 Importance of texturing HA coatings So far, there have been no reports concerning the successful synthesis of highly (0001) oriented HA coatings on titanium or other substrates by biomimetic or other processes.3,32–36 In natural bone, HA is highly oriented with (0001) perpendicular to the collagen fibrils.37 Therefore, as a bone implant material, it may be necessary for HA coating to be (0001) oriented to help growth of collagen fibrils in surrounding bone tissue into the implant to enhance biointegration.35 Fig. 12 Schematic illustration of the biomimetic mineralization process on (A) VTS-Ti and (B) control VTS-Ti.Note that the thicker and 5 Conclusion thinner lines correspond to the SAM domains fixed on (110)TiO2 and those on TiO2 crystal surfaces except (110). A highly oriented HA coating is successfully prepared by a biomimetic process on a textured titanium substrate. The oriented growth of HA results from the mediation by a crystals (i.e., HA crystals which nucleate later on SAM functionalized self-assembled monolayer with organized domains fixed on TiO2 surfaces except (110)].It seems that arrangement of OH and COOH groups fixed on the textured the geometry matching between the arrangement of functional titanium plate through interfacial molecular recognition such groups and 1-D or 2-D crystal lattice of inorganic nuclei plays as crystal lattice matching, hydrogen bonding interaction, an important role in biomimetic mineralization.Further inves- electrostatic potential interaction and stereochemistry match. tigation of this point is now under way in our laboratory. Another interesting result to be explained is the diVerent Acknowledgements time dependences of the pH values between VTS-Ti solution and control VTS-Ti solution.From reaction (1), it can be This project was granted financial support from China deduced that the more rapidly the HA crystals nucleate and Postdoctoral Science Foundation. grow, the more quickly the solution pH decreases. Therefore, the diVerent time dependences of pH values reflect the diVerent rates of HA crystallization. The pH-time curves in Fig. 3 can References be resolved into three stages.The initial stage is the incubation 1 S. Mann, J. Mater. Chem., 1995, 5, 935 and references therein. period during which the pH value remains unchanged and no 2 F. J. Fendler and F. C. Meldrum, Adv. Mater., 1995, 7, 607. HA crystals are formed. The subsequent stage is the nucleation 3 B. C. Bunker, P. C. Rieke, B. J. Tarasevich, A. H.Campbell, and growth period during which the pH decreases because G. E. Fryxell, G. L. Graft, L. Song, J. Liu, J. W. Virden and G. L. McLay, Science, 1994, 264, 48. OH- ions are constantly incorporated into HA structure to 4 Q. Huo, H. I. Margolese and G. D. Stucky, Chem. Mater., 1996, make HA nucleate and grow together with PO43- and Ca2+ 8, 1147. ions. At this stage, HA crystals begin to nucleate and further 5 S.Mann and G. A. Ozin, Nature, 1996, 382, 313. grow into plate-like grains. The last stage is an equilibrium 6 H. Shin, R. J. Collins, M. R. De Guire, A. H. Heuer and period when the pH value no longer changes because the C. N. Sukenik, J. Mater. Res., 1995, 10, 692. solution has reached a saturated state. Such patterns are 7 H. Shin, R. J. Collins, M.R. De Guire, A. H. Heuer and C. N. Sukenik, J. Mater. Res., 1995, 10, 699. characteristic for crystallization which undergoes changes from 8 J. H. Keeler, W. R. Hibbard and B. F. Decker, Trans. AIME, the incubation period through nucleation and growth period 1953, 197, 932. to saturation. The incubation period for VTS-Ti solution 9 H. B. Lu, C. L. Ma, H. Cui, L. F. Zhou, R.Z.Wang and F. Z. Cui, (0–8 h) is shorter than that for control solution (0–12 h), J. Cryst. Growth, 1995, 155, 120. whereas the nucleation and growth period for the former 10 T. Hanawa and M. Ota, Appl. Surf. Sci., 1992, 55, 269. (8–180 h) is longer than that for the latter (12–70 h). During 11 M. Kukura, L. C. Bell, A. M. Posner and J. P. Quirk, J. Phys. Chem., 1972, 76, 900. the incubation period, the ionic species (PO43-, OH-, Ca2+) 12 C.D. Wagner, W. M. Riggs, L. E. Davies, J. F. Moulder and are bonded to the functionalized surface through molecular G. E. Muilenberg, Handbook of X-ray photoelectron spectroscopy, recognition, which will favor the subsequent nucleation and Perkin-Elmer Corporation, Eden Prairie, MN, 1979, pp. 58–70. growth of HA, although HA crystals do not crystallize at this 13 C.N. Sayer and N. R. Armstrong, Surf. Sci., 1978, 77, 301. time. Because there are many more SAM-(110)TiO2 domains 14 E. Bertel, R. Stockbauer and T. E. Mady, Surf. Sci., 1984, 141, as reaction template for the above recognitions on VTS-Ti 355. 15 D. E. Eastman, Solid State Commun., 1972, 10, 933. than on the control substrate, the nucleation and growth of 2800 J.Mater. Chem., 1998, 8, 2795–280116 K. I. Hadjiivanov and D. G. Klissurski, Chem. Soc. Rev., 1996, 28 G. Beamson and D. Briggs, High resolution XPS of organic poly- 25, 61. mers, the Scienta ESCA300 database, John Wiley & Sons, New 17 T. Hanawa, M. Kon, H. Ukai, K. Murakami, Y. Miyamoto and York, 1992, pp. 66, 116. K. Asaska, J. Biomed. Mater. Res., 1997, 34, 273. 29 T. Hanazawa, Inorganic Phosphate Materials, Elsevier, 18 A. Azoulay, N. Shamir, E. Fromm and M. H. Mintz, Surf. Sci., Amsterdam, 1989, pp. 24–29, 91–92. 1997, 370, 1. 30 P. F. Gonzalez-Diaz and M. Santos, J. Solid State Chem., 1997, 19 K. Huelse and K. H. Kramer, in Titanium Science and Technology, 22, 193. ed. G. Luetjering, U. Zwicker and W. Bunk, Deutsche 31 P. G. Koutsoukos and G. H. Nancollas, J. Cryst. Growth, 1981, Gesellschaft fuer Metallkunde e. V., Germany, 1984, vol. 2, 53, 10. p. 1072. 32 C. Ohtsuki, H. Iida, S. Hayakawa and A. Osaka, J. Biomed. 20 P. G. Wahlbeck and P. W. Gilles, J. Am. Ceram. Soc., 1966, Mater. Res., 1997, 35, 39. 49, 180. 33 S. Li, Q. Liu, J. De Wijn, K. De Groot and B. L. Zhou, J. Mater. 21 J. Weis, Adv. Catal., 1952, 4, 343. Sci. Lett., 1996, 15, 1882. 22 J. C. Bailar, Comprehensive Inorganic Chemistry, Pergamon Press, 34 T. KoKubo, E. Miyaji, H. M. Kim and T. Nakamura, J. Am. 1973, vol. 2, p. 685. Ceram. Soc., 1996, 79, 1127. 23 K. Hadjiivanov, A. Davydev and D. Klissurski, Kinet. Katal., 35 K. A. Khor, P. Cheang and Y. Wang, J. Miner. Met. Mater. S, 1988, 29, 161. 1997, 2, 51. 24 K. Hadjiivanov, O. Saur, J. Lavnotte and J. C. Lavalley, Z. Phys. 36 C. Ohsuki, H. Iida, S. Hayakawa and A. Osaka, J. Miomed. Chem. (Munich), 1994, 187, 1281. Mater. Res., 1997, 35, 39. 25 C. Morterra, J. Chem. Soc., Faraday Trans., 1988, 84, 1617. 37 A. Veis, The Chemistry and Biology of Mineralized Connective 26 R. T. Morrison and R. N. Boyd, Organic Chemistry, Prentice Hall, Tissues, Elsevier, Amsterdam, 1981, p. 617. New Jersey, 1992, 6th edn., pp. 357–360. 27 T. B. McPherson, H. S. Shim and K. Park, J. Biomed. Mater. Res., 1997, 38, 289. Paper 8/01384E J. Mater. Chem., 1998, 8, 2795–2801 2801

ISSN:0959-9428    

DOI:10.1039/a801384e

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Porous hydroxyapatite monoliths from gypsum waste Sachiko Furuta,a Hiroaki Katsuki*a and Sridhar Komarnenib aSaga Ceramics Research Laboratory, 3037-7, Arita-machi, Saga, 844-0024, Japan. E-mail: katsuki@emon.scrl.pref.saga.jp bIntercollege Materials Research Laboratory, and Department of Agronomy, The Pennsylvania State University, University Park, PA 16802-4801, USA Received 25th August 1998, Accepted 8th October 1998 Porous hydroxyapatite (HAp) monoliths were synthesized from gypsum mold waste with diammonium hydrogen phosphate solution by hydrothermal treatment at 50–190 °C and their properties studied.Gypsum waste samples could be completely converted to HAp with 0.5 mol dm-3 (NH4)2HPO4 at 50 °C in 15 days and 100 °C in 2 days.This paper describes the eVect of hydrothermal reaction conditions such as reaction temperature, time and concentration of (NH4)2HPO4 on the formation of new HAp monoliths and their porous properties. vessel of a stainless steel hydrothermal reactor. The initial pH 1 Introduction of the (NH4)2HPO4 solution was ca. 7.5. The hydrothermal Hydroxyapatite [Ca10(PO4)6(OH)2], which is well known as reactor was introduced into a drying oven which was heated a main component of bone and teeth, has been used widely in in the range 50–100 °C for 1–15 days.After hydrothermal many industrial applications and in the medical field. For treatment, reacted samples were washed with distilled water example, HAp has been found to have cation-exchange proper- to remove residual ions such as SO42-, NH4+ and PO43- and ties where Ca2+ ions of HAp can be exchanged with poisonous were dried.The drying temperature was kept below 50 °C so heavy metal ions.1,2 HAp also has the ability to adsorb organic that unreacted gypsum crystals would not be decomposed chemicals.3,4 Generally, HAp powder is synthesized by sev- during the drying treatment.To promote the formation rate eral routes such as precipitation from solution containing of HAp, 1.0 mol dm-3 (NH4)2HPO4 was used as the starting Ca2+ and PO43- ions5,6 and hydrolysis of CaHPO4 or solution and the reaction temperature was increased to 190 °C. CaHPO4·2H2O at room temperature7 and hydrothermal syn- The conversion of gypsum to HAp was estimated from the thesis of Ca3(PO4)2 at 200 °C.8 One of the problems with these ratio of X-ray intensities of the peaks of gypsum (hkl=020, processes is that HAp monoliths such as bone or tooth, etc. 2h=11.62°) and HAp (hkl=211, 2h=31.77°) by powder Xcan not be obtained from these chemicals by an in situ method. ray diVraction (XRD, Model RAD-2B, Rigaku Co., Japan). Therefore, synthetic HAp powder is molded and then sintered The HAp samples were dissolved in Na2CO3–H3BO3 solution, at 1000–1250 °C.9 and the chemical composition of HAp was analyzed by ICP A large amount of gypsum plaster (CaSO4·2H2O) is used emission spectrometry (ICPS-2000, Shimadzu Co., Japan).as molds for slip or pressure casting in the ceramic industry. The crystal morphology of the synthetic samples was observed These gypsum molds wear out after repeated use (ca. 100 by scanning electron microscopy (SEM, Model JXA-840, times) and are then discarded as waste.Gypsum is a sparingly JEOL Co., Japan) and transmission electron microscopy soluble calcium salt, and has never been used previously as (TEM, Model 2010, JEOL Co., Japan). The pore diameter, the source material to prepare HAp. Here we develop a novel porosity and pore size distribution of products were investiprocess for preparing porous HAp monoliths directly from gated by mercury porosimetry (Pore Sizer 9310, Micromeritics gypsum waste by in situ crystallization using the following Co., USA) and the surface area was measured by the BET chemical reaction.method with N2 gas adsorption (Autosorb 1, Quantachrome Co., USA). 10CaSO4·2H2O+6(NH4)2HPO4�Ca10(PO4)6(OH)2 +6(NH4)2SO4+4H2SO4+18H2O 3 Results and discussion This concept will alleviate the environmental problem from the following two standpoints: (1) eVective recycling of indus- 3.1 Synthesis of HAp from gypsum waste trial waste materials, and (2) use of this new porous HAp as Gypsum waste used in this study was composed of fine needle- a purification system of waste water containing heavy metal like crystals of length 5–10 mm and 1–2 mm thickness as shown ions or organic chemicals. in Fig. 1. This gypsum has a porous structure with 2.4 mm In this study, porous HAp monoliths were produced by in median pore diameter, 60% porosity and 1.6 m2 g-1 surface situ crystallization from gypsum mold waste by a conventional area. The eVect of reaction temperature on the formation of hydrothermal treatment and its characteristics were HAp was investigated with 0.5 mol dm-3 (NH4)2HPO4.Fig. 2 investigated. shows XRD patterns of the as-received gypsum and the samples synthesized for 1, 2 and 3 days at 100 °C, and Fig. 3 2 Experimental shows XRD patterns of the samples synthesized at 50, 75 and 100 °C for 6–15 days. Fig. 4 summarizes the conversion rate Gypsum waste mold was washed and cut into rectangular pieces of 5×10×20 mm. The surface layer of gypsum mold from gypsum to HAp based on the calculation from XRD intensity ratios of gypsum and HAp. At 100 °C, ca. was removed since small amounts of some impurities such as Na and Si components from the ceramic raw materials are 58–59 mass% gypsum was hydrothermally converted into HAp after 1 day, and the conversion of the gypsum to HAp reached deposited on the surface of the repeatedly used gypsum mold.The gypsum and 0.5 mol dm-3 diammonium hydrogen phos- 100 mass% after 2 days. HAp monoliths could be hydrothermally prepared by in situ crystal growth from gypsum waste phate [(NH4)2HPO4] solution were placed in a 50 ml Teflon J. Mater.Chem., 1998, 8, 2803–2806 2803Fig. 1 Morphology of gypsum waste. Fig. 4 The conversion rate of gypsum to HAp based on the calculation from the X-ray intensity ratio under the hydrothermal reaction with 0.5 mol dm-3 (NH4)2HPO4. &: 50°C, $: 75°C and +: 100 °C. involves slight leaching of Ca2+ ions from gypsum crystals and these ions then react with PO43- and OH- ions in solution to form HAp until all the gypsum crystals are converted.These results show that the novel HAp monoliths can be synthesized by in situ crystallization of gypsum waste with (NH4)2HPO4 solution under conventional hydrothermal treatment at 50–100 °C. 3.2 Properties of new HAp crystals Fig. 2 XRD patterns of (a) gypsum waste and samples synthesized by To characterize the chemical composition of synthesized HAp hydrothermal reaction with 0.5 mol dm-3 (NH4)2HPO4 at 100 °C for crystals, the Ca/P molar ratio of the crystals was analyzed by (b) 1 day, (c) 2 days and (d) 3 days.ICP. The Ca/P molar ratios of crystals obtained at 50 °C for 15 days and 100 °C for 6 days were calculated to be 1.25 and 1.30, respectively. These new HAp crystals were nonstoichiometric compared to theoretical hydroxyapatite [Ca10(PO4)6(OH)2] whose Ca/P ratio was 1.67. HAp is known to be a stable substance even though its composition is nonstoichiometric. There are several previous reports11–13 on explaining why the Ca/P ratio of apatite can be varied over a wide range: (1) other crystal phases such as CaH(PO4)3·2H2O and Ca4H(PO4)3·3H2O can co-exist,11 (2) phosphoric acid can be adsorbed on the surface of HAp,12 (3) there may be a deficiency of Ca2+ in the crystal lattice of apatite.13 In our study, only the HAp phase was observed by XRD after complete conversion.Thus it is considered that residual phosphoric acid is adsorbed on the surface of HAp because the concentration of (NH4)2HPO4 in the starting solution is stoichiometrically higher than that of gypsum.The crystallinity Fig. 3 XRD patterns of HAp synthesized from gypsum waste and of the HAp crystals increased with reaction temperature as 0.5 mol dm-3 (NH4)2HPO4 under the following hydrothermal reaction can be deduced from peak widths of the XRD patterns conditions: (a) 100 °C6 days; (b) 75 °C, 7 days; (c) 50 °C, 15 days. in Fig. 3. Fig. 5 and 6 show the morphology of HAp synthesized at 50, 75 and 100 °C.After reaction at 50 °C for 7 days, unreacted and (NH4)2HPO4 at 100 °C in 2 days. To prepare smaller HAp crystals, the growth of HAp was examined at lower gypsum crystals were still present within the body of the monolith, but the surface was coated with fine HAp crystals temperatures. The conversion of gypsum to HAp was 100 mass% at 75 °C after treatment for 7 days and after 15 of diameter 10–30 nm as shown in Fig. 5(b). Almost all gypsum crystals were completely converted into the HAp days at 50 °C. Monoliths of HAp resulted under all reaction conditions, phase after 15 days leading to an HAp monolith which was composed of crystals of 80–120 nm. With increasing reaction and the size of the HAp monoliths were the same size as the starting gypsum piece in all cases.During this preparation temperature, the crystal growth of HAp was enhanced and the crystals grew to 3–8 mm in length at 100 °C for 3 days. process, HAp is produced in a solution under neutral or mildly alkaline condition with the solubility product To promote the formation rate of HAp, the gypsum waste was reacted with 1.0 mol dm-3 (NH4)2HPO4 at 190 °C.[Ca2+]5[PO43-]3[OH] exceeding that of HAp.6 The solubility of CaSO4·2H2O in water in the present reaction is ca. However, under these conditions the gypsum could not be completely converted into the HAp phase because the surface 0.16–0.21% at 50–100 °C.10 The mechanism of HAp formation 2804 J. Mater. Chem., 1998, 8, 2803–2806Fig. 6 Morphology of HAp crystals synthesized from gypsum waste Fig. 5 Morphology of HAp crystals synthesized from gypsum waste and 0.5 mol dm-3 (NH4)2HPO4 by hydrothermal reaction. (a) and and 0.5 mol dm-3 (NH4)2HPO4 by hydrothermal reaction. (a) 75 °C, 3 days; (b) 75 °C, 7 days; (c) 100 °C, 1 day; (d) 100 °C, 3 days. (b): 50 °C, 7 days; (c) and (d): 50 °C, 15 days. were thus not carried out using this concentration of (NH4)2HPO4.of gypsum rapidly reacted with (NH4)2HPO4, and a tight HAp layer of about 1 mm thickness was formed on the surface 3.3 Porous properties of HAp monoliths after reaction for 3 days and prevented further conversion of gypsum crystals to HAp. Therefore, unreacted gypsum crystals In this study, novel HAp monoliths were easily synthesized from the reaction of gypsum waste with (NH4)2HPO4 and it remained within the body of the monolith.Further reactions J. Mater. Chem., 1998, 8, 2803–2806 2805at 100 °C for 3 days. The surface area strongly depends on the morphology of HAp, and decreased with increased crystal size. The porosities of the HAp samples which were synthesized using 0.5 mol dm-3 (NH4)2HPO4 at 75 °C for 1–10 days were in the range 75–78% while the porosity of starting gypsum waste was 60%.Pore size distributions of the starting gypsum and products at 75 °C are shown in Fig. 8. The bimodal pore distribution in Fig. 8(b) and (c) is attributed to the diVerent pore sizes of gypsum waste and HAp. Pore diameters of products obtained after treatment for 1–7 days centered at 2–3 mm were apparently from unreacted gypsum, and the pore distribution of HAp aggregates was between 0.01 and 1 mm with a median pore diameter of around 0.3 mm.The pore size distribution of gypsum centered at 2–3 mm remained in the sample treated up to 7 days at 75 °C, and the sample had double-pore structures, one from gypsum and the other from HAp. A single pore distribution derived from HAp crystals was obtained after 10 days at 75 °C.While it was not possible to obtain a single pore structure at 50 °C over short times, the sample prepared at 100 °C for 2 days showed only the pore size distribution of HAp. Fig. 7 The surface area of synthesized HAp with 0.5 mol dm-3 of (NH4)2HPO4. &: 50°C, $: 75°C and +: 100 °C. Conclusion In this study, porous HAp monoliths were synthesized and characterized from gypsum mold waste and (NH4)2HPO4 solution by hydrothermal reaction.The main results are as follows. (1) The conversion of gypsum to HAp was aVected by reaction time and temperature taking about 2 weeks for complete conversion at 50 °C, but only 2 days at 100 °C when 0.5 mol dm-3 (NH4)2HPO4 was used. The crystallinity of the synthesized HAp was enhanced by increasing the reaction temperature.SEM observation revealed that HAp monoliths formed at 50 °C in 15 days and were composed of fine crystals of HAp of 80–120 nm diameter and 0.5–1 mm length while HAp formed at 100 °C for 3 days showed larger crystals of size 3–8 mm. (2) The surface area of synthesized HAp monoliths ranged from 20–45 m2 g-1 at 50–100 °C and was related to the morphology of the HAp crystals.The pore size distribution of gypsum was ca. 2–3 mm and decreased in intensity during the hydrothermal reaction while the pore size distribution of HAp, which was between 0.01 and 1 mm, increased. Under appropriate conditions of time and temperature of reaction, porous HAp monoliths with single pore structure resulted from the reaction of gypsum waste and (NH4)2HPO4.Fig. 8 The pore size distributions of (a) starting gypsum and HAp synthesized with 0.5 mol dm-3 (NH4)2HPO4 at 75 °C for (b) 1 day, (c) 7 days and (d) 10 days. References 1 T. Suzuki, T. Hatsushika and Y. Hayakawa, J. Chem. Soc., Faraday Trans. 1, 1981, 77, 1059. was composed of fibrous or needle-like crystals. To determine 2 S. Suzuki, T. Fuzita, T. Maruyama and M.Takahashi, J. Am. the value of such HAp monoliths for applications, porous Ceram. Soc., 1993, 76, 1638. properties such as porosity, pore diameter and the surface 3 E. C. Moreno, M. Kresak and A. GaVar, J. Colloid Interface Sci., 1994, 168, 173. area are important and therefore, these were investigated. 4 T. Akazawa, M. Kobayashi, T. Kanno and K. Kodaira, J. Mater. Fig. 7 shows the eVect of reaction temperature and time on Sci., 1998, 33, 1927. the surface area of HAp monoliths. The surface area of 5 A. L. Bosky and A. S. Posner, J. Phys. Chem., 1976, 80, 40. products at 75 °C for 2–10 days and at 100 °C for 1–6 days 6 H. Mcdowell, T. M. Gregory and W. E. Brown, J. Res. Natl. Bur. ranged from 26–30 and 19–23 m2 g-1, respectively. However, Stand., Sect.A, 1977, 81, 273. 7 E.J.DuV, J. Chem. Soc. A, 1971, 917. these values increased to 35–45 m2 g-1 at 50 °C when the 8 M. Yoshimura, H. Suda, K. Okamoto and K. Ioku, J. Mater. Sci., gypsum was reacted for 7–15 days due to the formation of 1994, 29, 3399. smaller HAp crystals as can be seen in Fig. 5. Upon in situ 9 J. G. Peelen, B. V. Rejda and K. De Groot, Ceram. Int., 1978, synthesis of HAp monoliths from gypsum waste, the surface 4, 71. area increased from 20 to 45 m2 g-1 upon lowering the tem- 10 H. F. W. Tayler, The Chemistry of Cements, 1964, vol. 1, p. 321. 11 D. McConnel, Arch. Oral Biol., 1965, 10, 421. perature from 100 to 50 °C and this is related to the mor- 12 A. S. Posner and S. R. Stephenson, J. Dent. Res., 1952, 31, 371. phology of HAp crystals. From the observation of HAp 13 G. Kuhl and W. E. Nebergall, Z. Anorg. Allg. Chem., 1963, 24, crystals by TEM and SEM (Fig. 5 and 6), the crystals were 313. of length 100–300 nm long and thickness 10–30 nm at 50 °C for 7 days, 2–5 mm long at 75 °C for 3 days, and 3–8 mm long Paper 8/06659K 2806 J. Mater. Chem., 1998, 8, 2803–2806

ISSN:0959-9428    

DOI:10.1039/a806659k

出版商:RSC    

年代:1998

数据来源: RSC

摘要:

J O U R N A L O F C H E M I S T R Y Materials Growth of calcium phosphate onto coagulated silica prepared by using modified simulated body fluids Juan Coren�o,a Rogelio Rodrý�guez,*a,b Miguel A. Araizac and Victor M. Castan�ob aDepartamento de Fý�sica, Universidad Auto�noma Metropolitana-Iztapalapa, Apdo. Postal 55-534, Me�xico, D.F. 09340 bInstituto de Fý�sica, Universidad Nacional Auto�noma de Me�xico, Apdo.Postal 1-1010, Quere�taro, Qro. Me�xico 76001 cFacultad de Odontologý�a, Universidad Nacional Auto�noma de Me�xico, Ciudad Universitaria, Me�xico, D.F. 04510 Received 11th May 1998, Accepted 25th September 1998 Silica sols prepared by the alkaline hydrolysis of tetraethylorthosilicate were coagulated by adding an excess of a CaCl2 aqueous solution. The aggregates were immersed into three diVerent modified simulated body fluids at 90 °C to allow the growth of a calcium phosphate phase onto the silica aggregates.The apatite phase grew faster compared to previous studies. Also, the amount of the crystalline apatite yield was higher when the simulated body fluid employed had the largest Ca/P ratio, as measured by X-ray diVraction. The relevance of these findings is discussed in terms of the current and future trends in biomaterials research and development.A hydrated silica gel layer is formed on the glass surface Introduction and provides favorable sites for apatite nucleation. It has The preparation of synthetic biomaterials, aimed to be been suggested in the literature11 that, in contrast to dense employed for prosthestic applications in living organisms, has silica, this silica layer produced on the bioglasses, is flexible attracted growing interest in the last few decades.For instance, enough to provide the oxide–oxide spatial requirements to apatite coatings have been applied onto diVerent substrates to match the bone lattice, thus providing epitaxial sites for bone produce a number of interesting materials with potential growth.biomedical applications.1 The development of novel nano- The potential role of hydrated silica species on biological composites opens a broad new field of research since, in mineralization9 also points out the relevance of producing principle, it is possible to control, to a great extent, the materials with high surface areas, controlled porosity and corresponding morphology, which in turn is related to the appropriate chemical groups available for the desired interbiomedical and physical properties.2 actions.One possible approach is the coating of diVerent A number of important requirements, in addition to the substrates with biocompatible apatite layers, by using various obvious biocompatibility, ought to be fulfilled if synthetic techniques, which range from standard vacuum evaporamaterials are to be used for human applications. In fact, one tion to chemical vapor deposition (CVD) technologies. of the most important characteristics is the corresponding Unfortunately, long-term animal studies suggest that the HAp porous structure, since the extra-cellular fluids must be allowed coatings may degrade or come oV.9 to flow through the inner structure of the biomedical device One interesting method to produce calcium phosphate to allow an adequate osteoconduction.3,4 apatites at low temperature is through the immersion of silica Among the materials that are known to have an appropriate gels into simulated body fluid (SBF).12–17 Besides the promisbiological activity when implanted in living organisms calcium ing potential applications of these experiments, from a practical phosphates deserve special attention.5 Historically speaking, standpoint, the available reports in the literature reveal a very low reaction kinetics and a limitation on the number of the modern use of these materials for bone repair was probably available surface chemical groups for the reaction with the pioneered by Albee in the 1920s6 which are, nowadays, widely surrounding SBF, since the silica gels reportedly employed employed for low-loading hard tissue repair and augmentation only react on the areas exposed to the SBF.Accordingly, an of living bone.4 The excellent histological behaviour of these alternative method, by starting from silica sols rather than materials, either synthetic or natural, is attributed to their gels, which are coagulated in situ by adding controlled amounts chemical similarity to the mineral phase of natural bone, of calcium ions, and by using three diVerent modified SBFs, hydroxyapatite [Ca10(PO4)6(OH)2] or HAp, which is the is reported here.The higher temperature of reaction, namely classical example.In fact, the crystallographic and chemical 90 °C, makes the available reacting chemical groups more properties of HAp closely resemble those of bone and tooth reactive; additionally, due to the use of nanometer-sized sols minerals.7 and to the coagulation process, one can obtain faster reaction Bioactive glasses have also been used for implants because times, higher yields and more controllable conditions these materials bond to bone through an intervening bonecompared with previous studies.like apatite layer formed on their surfaces in a living bodylike environment.5,8 Interestingly this apatite layer is calciumdeficient compared with hydroxyapatite, being formed by small Experimental crystallites and containing small amounts of carbonate.9 In CaO–SiO2-based bioactive glasses,10 it has been found that Coagulated silica was obtained by preparing pure silica sol by silica forms a low solubility matrix in which the network of the sol–gel method under alkaline conditions of the hydrolysis silicate chains acts as a framework for ionic species (Ca2+, reaction, followed by a coagulation procedure of the silica PO43-, Na+, etc.), whose role is to stimulate the biochemical particles through the addition of an excess of calcium chloride.It is important to point out that, unlike previous reports in environment surrounding the bioactive glass. J. Mater. Chem., 1998, 8, 2807–2812 2807the literature, in the present studies, coagulated silica sols were 9 days. After this period the solid was washed and dried at 100 °C for 2 h.employed instead of the standard silica gels. SBF and mSBFs solutions were prepared by dissolving, at diVerent concentrations, reagent grade NaCl, NaHCO3, KCl, Sol preparation Na2HPO4·12H2O, MgCl2·H2O, CaCl2 or Na2SO4, buVered A solution of 4 moles of distilled water in 6 moles of ethanol at pH=7.4 by using tris(hydroxymethyl )aminomethane (reactive grade)(Baker Co.) was added to another solution of [(CH2OH)3CNH2] and hydrochloric acid. 1 mole of tetraethylorthosilicate (TEOS) (Aldrich Chem. Co.) in 6 moles of ethanol, under vigorous stirring. The starting Characterization techniques pH of water was adjusted to 12 by using NH4OH. The resulting mixture was poured into a round double-necked The dynamic light scattering (DLS) apparatus used to measure the particle size was a Brookhaven Instrument with a digital flask, and heated to reach reflux conditions.A profile of the silica particle size as a function of the reaction time was correlator model 9000; in all cases the scattering angle was set to 90°, the measurements were done at room temperature and obtained by sampling at regular time intervals by using dynamic light scattering techniques.Once a constant particle the light source was an argon ion laser operating at 488 nm. The phosphate concentration was measured by the molyb- diameter was obtained, the sol–gel reaction was stopped by diluting with ethanol and cooling the system to prevent denum blue method18 using a UV–VIS absorption spectrophotometer (Perkin Elmer Lambda 5) at 690 nm.The X-ray gelation. Silica sols prepared under these (basic) conditions are electrically stabilized; consequently, they can be coagulated diVractograms (XRD) of the samples were obtained in a Phillips diVractometer in the range 20–50° with a scanning at will by the addition of a suitable salt. rate of 2° min-1. Scanning electron microscopy (SEM) was carried out in a Zeiss model DSMSol coagulation ground and carbon-coated.m-Raman characterization was An aqueous 0.1 M CaCl2 solution was added dropwise to a carried out in a DILOR apparatus model Labram equipped fixed volume of colloidal silica suspension under stirring; the with a He–Ne laser as a light source and a confocal optical final concentration of calcium ions was 0.01 M.After stirring microscope; the wavenumber range of the scattered light was for one hour, the flocculate was allowed to settle under the varied from 100 to 1200 cm-1. 31P NMR spectra were taken influence of gravity; the liquid was decanted and the solid on a Bruker ASX300 NMR spectrometer using a 4 mm CP centrifuged. The solid part was washed twice by resuspending MAS probe at 5 kHz.H3PO4 was used as reference. it in distilled water and stirring gently for 20 min. The coagulates were dried at 400 °C for 2 days and ground by using an Results and discussion agate mortar. Fig. 1 shows the X-ray diVraction pattern corresponding to Calcium–phosphates growth pure HAp where the plane indices of the main reflections are shown. XRD for pure silica, prior to the immersion into the The growth of calcium phosphate crystals was achieved by mSBF and with the same thermal treatment as all samples, resuspending the silica coagulates in an aqueous solution was also obtained (not shown); for this sample no crystalline containing calcium and phosphate ions.Four diVerent ionic reflections were observed but only a broad peak around 23° solutions were used: due to the amorphous phase. 1) SBF at 37 °C XRD for samples containing calcium phosphate grown on their surfaces at diVerent immersion times in diVerent mSBFs 2) mSBF-1.0: modified SBF with Ca/P=1.00 at 90 °C are shown in Fig. 2–4. Fig. 2 corresponds to the diVractograms 3) mSBF-1.7: modified SBF with Ca/P=1.67 at 90 °C 4) mSBF-2.5: modified SBF with Ca/P=2.50 at 90 °C Table 1 summarizes the ionic composition for the four SBF solutions; human plasma is also reported for comparison purposes.Each experiment using mSBF was carried out in three round flasks. The first flask contains only 20 ml of mSBF, to rule out a possible spontaneous crystallization. The second flask contains 0.35 g of 0.125–0.25 mm ground silica in 30 ml of mSBF and the third contains 0.20 g of coagulated silica ground to sizes smaller than 0.125 mm in 20 ml of mSBF. The third flask was sampled on a regular basis to measure phosphate concentration in the solution as a function of time and to obtain X-ray diVractograms of the corresponding products.The experiment at 37 °C was taken as reference. In this case, the silica aggregates were used without previous drying.They Fig. 1 X-Ray diVractogram of pure crystalline hydroxyapatite. were added to the SBF and kept at 37 °C under agitation for Table 1 Ion concentration of standard blood plasma, SBF, and the modified SBF used for calcium phosphates growth onto silica aggregates Ion concentration/mM Ca/P Na+ K+ Ca2+ Mg2+ Cl- HCO3- HPO42- SO42- Blood plasma 2.5 142.0 5.0 2.5 1.5 103.0 27.0 1.0 0.5 SBF 2.5 142.0 5.0 2.5 1.5 147.0 4.2 1.0 0.5 Modified SBF, No. 1 1.67 142.4 — 2.0 — 144.0 — 1.2 — Modified SBF, No. 2 1.0 143.0 — 1.5 — 143.0 — 1.5 — Modified SBF, No. 3 2.5 142.0 — 2.5 — 145.0 — 1.0 — 2808 J. Mater. Chem., 1998, 8, 2807–2812Fig. 2 X-Ray diVractograms of samples immersed in mSBF with Ca/P=1.0 at 90 °C and diVerent immersion times: (a) 10 h, (b) 23 h, Fig. 5 X-Ray diVractogram of a sample prepared at 37 °C immersed (c) 33 h and (d) 48 h. in SBF for 7 days. equivalent to Fig. 2 but using mSBF-1.7 and mSBF-2.5, respectively. For comparison purposes, the diVractogram for the experiment at 37 °C is shown in Fig. 5. The diVractograms of Fig. 2–4 show, in addition to the contribution of the amorphous silica, the crystalline reflection corresponding to the apatite phase. As the Ca/P ratio in the mSBFs was increased, the intensity of the reflection was also increased.This is clearly observed in Fig. 6 where the intensity of the characteristic apatite peak is plotted as a function of the immersion time for the three diVerent mSBFs. This plot shows that, at the beginning of the crystal growth on the silica surface, the amount of the crystalline phase is small, corresponding to an induction period where the silica surface oVers multi-nucleating sites where the apatite can grow.Up to 33 h the amount of crystals remains practically constant as revealed Fig. 3 As for Fig. 2 but for samples immersed in mSBF with by the intensity of the reflection, which is nearly the same as Ca/P=1.67. after induction.However, after this time, for Ca/P ratios of 1.67 and 2.5, their heights have increased producing sharper and stronger X-ray reflections. Since the height of the diVraction peak is known to be proportional to the amount of crystals which contribute to the peak, there is a 2.2 fold increase in this crystal content for mSBF-2.5 compared with mSBF-1.0. It is important to mention that these growth times are considerably shorter than typical HAp formation periods reported previously.19 m-Raman spectra were obtained for all samples and pure HAp.The Raman spectrum for pure HAp is shown in Fig. 7(a), while the sample with Ca/P 2.5 is shown in Fig. 7(b) where it is possible to observe a small signal corresponding to the apatite superimposed on a strong fluorescence signal produced by the silica substrate.As before, only the strongest apatite band at 962 cm-1 due to the phosphate symmetric stretching vibration20 is observed because this phase represents only a small percentage of the resulting material. Similar Fig. 4 As for Fig. 2 but for samples immersed in mSBF with Ca/P=2.5. of the samples immersed in mSBF-1.0 at 90 °C for 10, 23, 33 and 48 h, respectively.The whole set of reflections for the crystalline phase is diYcult to obtain since it is the minor component of the resulting material (ca. 10%). It is worth mentioning that, since the apatite growth takes place on the surface of small coagulated silica particles, the silanol groups on their surfaces and the open structure produced by the coagulating process allow the coating to be both internal and external, unlike the methods which use silica gels; this means that the reflections may also contain contributions from the internal coating.Since the reaction conditions favor apatite formation,10 the characteristic reflection for the crystalline phase is located around 32°; this peak is assigned to an overlap of the diVraction bands of three crystalline spacings: (211), Fig. 6 Plot of the 32° reflection intensity as a function of the immersion time for samples immersed in the three mSBFs. (112) and (300) according to the literature.19 Fig. 3 and 4 are J. Mater. Chem., 1998, 8, 2807–2812 2809Fig. 10 Consumption of phosphorus as a function of time for all Fig. 7 m-Raman spectra of pure HAp (a) and for the sample Ca/P 2.5 experiments.immersed in mSBF for 48 h. spectra (not shown) were obtained for the remainder of apatites with Ca/P ranging from 1.14 to 1.66 (d 2.8±0.2) the samples. which is clearly diVerent from those corresponding to other Fig. 8 and 9 show the 31P MAS NMR spectra for pure HAp calcium phosphates.21 This result confirms the identification and for the sample immersed in mSBF-2.5 for 48 h, respect- of the crystalline phase formed as apatite.ively. In Fig. 9 only one 31P resonance with weak sidebands is Fig. 10 shows the change in P concentration for the four observed. The isotropic chemical shifts obtained are d 2.710 experiments. Both the rate and the amount of P consumption for pure HAp and d 2.924 for the sample. These values are in are similar for all the three experiments at 90 °C.The decrease accord with the chemical shift reported for calcium phosphate in P concentration for these experiments ends approximately 23 h after the beginning of the reaction, whereas for the experiment at 37 °C, it starts approximately at this point. It is interesting that, even for the experiment at 37 °C, which corresponds to the slowest apatite growth, an apatite growth rate faster than those reported for gels dried at 400 °C is observed.19 There are a number of factors which explain this behavior: (a) the present experiments were carried out using coagulated silica sols instead of the standard gels, which increases the active surface area available for the chemical reaction, (b) the higher temperature of reaction makes the available chemical groups more reactive so reducing the reaction time.Additionally, the Ca ions adsorbed on the silica sol may favour the growth of the apatite phase by increasing the ionic activity of the surrounding solution near the flocs. It is important to mention that, for the series of experiments without silica sols (i.e., containing only mSBF), there was no P consumption during the same time intervals reported here.SEM micrographs for the whole set of samples are shown in Fig. 11–15. Fig. 11 reveals a smooth surface corresponding to pure silica. The micrographs shown in Fig. 12–14 correspond to experiments where the apatite grew onto Fig. 8 31P MAS NMR spectrum of pure hydroxyapatite. Fig. 9 31P MAS NMR spectrum of the sample Ca/P 2.5 immersed in Fig. 11 SEM micrograph of pure silica. mSBF for 48 h. 2810 J. Mater. Chem., 1998, 8, 2807–2812Fig. 14 SEM micrograph from a sample with Ca/P=1.0 for an Fig. 12 SEM micrograph from a sample with Ca/P=1.67 for an immersion time of 48 h. immersion time of 48 h. Fig. 13 SEM micrograph from a sample with Ca/P=2.5 for an immersion time of 48 h. Fig. 15 SEM micrograph from a sample at 37 °C and 7 days immersion in SBF. 0.125–0.25 mm ground silica particles at 90 °C during 48 h. These do not show the typical spherical morphology that has been observed in similar work at 37 °C.14,19 modified by the presence of some ions. In this work it is observed that the Ca/P ratio aVects both the crystal size and Fig. 12 shows the micrograph corresponding to samples immersed in mSBF-1.7; this shows smaller and more uniform the way they cluster together, but not significantly the crystal morphology.It seems that the eVect of Ca/P ratio on the crystallite sizes (ca. 12 mm) with respect to the other two experiments at diVerent Ca/P ratios. Fig. 13 shows a micro- crystal morphology is diminished at the higher temperature employed for the apatite growth.graph corresponding to Ca/P =2.5 where the apatite crystals form a loose cluster with a wider particle size distribution It is interesting to note the dependence of the apatite morphology on the nature of the substrate. In this work, the compared with Fig. 12. These characteristics are not observed in Fig. 14 (Ca/P=1.0) where the crystals in the aggregate are sample prepared in SBF at 37 °C does not show the flake-like crystals reported when fired silica was used,22 but rather a tightly distributed. For comparison, the micrograph for the experiment at 37 °C uniformly distributed crystal layer without a defined morphology.Also, this layer was grown more rapidly than that is shown in Fig. 15. Here, it can be seen that this rough surface is quite diVerent from the other SEM micrographs.The apatite previously reported.19 These observations can be explained by the greater silanol density present in a wet coagulated silica phase shows smaller and poorly defined crystallites, compared with samples at 90 °C. This sample was kept in the SBF until sol, compared with fired silica, since it has been proposed that silanol groups are responsible for apatite nucleation.11 XRD characteristic peaks were observed (after 7 days).These results demonstrate the strong influence of temperature on the All the above results demonstrate that the silica is covered with a layer of apatite whose morphology depends on tempera- growth of the crystalline phase. For apatite growth onto silica fired at 400 °C22 it has been ture, immersion time and the type of substrate.The morphology of this nano-composite can also be modified, either reported that the morphology is Ca/P dependent, and can be J. Mater. Chem., 1998, 8, 2807–2812 28113 L. L. Hench and J. Wilson, Science, 1984, 226, 630. by changing the Si5Ca/P ratio, or by modifying the route of 4 F. H. Albee, Ann. Surg., 1920, 71, 32. adding the constituents. 5 R.Z. LeGeros, Calcium Phosphates in Oral Biology and Medicine, One important advantage of the present approach over the Karger, Basel, Switzerland, 1991. previous reports is that, since the silica sol was coagulated, 6 K. de Groot, in Contemporary Biomaterials, ed. J. W. Boretos and the active sites suitable for apatite growth are evenly distributed M. Eden, Noyes Publications, Park Ridge, NJ, 1984, pp. 477–492. over the whole sample surface, consequently it is expected that 7 R. H. Doremus, J. Mater. Sci., 1992, 27, 285. the crystalline phase will grow both inwardly and outwardly 8 M. M. Pereira, A. E. Clark and L. L. Hench, J. Am. Chem. Soc., 1995, 18, 2463. on the silica flocs. This shows the clear advantage of using 9 L. L. Hench, J. Am. Ceram. Soc., 1991, 74, 1487.sols rather than gels to allow the presence of more active 10 L. L. Hench and A. E. Clark, Biocompatibiity of Orthopedic groups available for the corresponding reaction. Additionally, Implants, ed. E. F. Williams, CRC Press, Boca Raton, FL, 1982, by choosing the appropriate starting silica particle size it is vol. 2, pp. 129–170. possible to control the porosity and the interstitial volume of 11 L.L. Hench and E. C. Ethridge, Biomaterials: An Interfacial the material. The immersion time, and, in turn, the thickness Approach, Biophysics and Bioengineering Series, ed. of the HAp layers, also slightly modifies the average pore size. A. Noordergraaf, Academic Press, New York, 1982, vol. 4, p. 139. 12 A. Ravaglioli and A. Krajewski, Bioceramics: Materials, Properties and Applications, Chapman & Hall, London, 1992, Conclusions pp. 140 and 175. 13 T. J. Kokubo, J. Non-Cryst. Solids, 1990, 120, 138. A novel composite material with potential biomedical 14 P. Li, C. Ohtsuki, T. Kokubo, K. Nakanishi, N. Soga, applications prepared by coagulating silica sols with calcium T. Nakamura and T. Yamamuro, J. Am. Ceram. Soc., 1992, 75, ions, is reported. With this procedure, a silica substrate covered 2094.inwardly and outwardly with a crystalline apatite phase can 15 T. Kokubo, Biomaterials, 1991, 12, 1155. 16 R. Fresa, A, Constantini, A. Buri and F. Branda, J. Non-Cryst. be produced. The role of the liquid phase (SBF) was analyzed Solids, 1995, 16, 1249. by using several modified simulated body fluids. The kinetics 17 M. Tanahashi, T. Kokubo, T. Nakamura, Y. Katsura and of each process is certainly an area worth further study, not M. Nagano, J. Non-Cryst. Solids, 1996, 17, 47. only because of its relevance for the production of novel 18 C. Sung-Baek, N. Kazuki, T. Kokubo, N. Soga, C. Ohtsuki, biomaterials, but also because it opens the possibility of T. Nakamura, T. Kitsugi and T. Yamamuro, J. Am. Ceram. Soc., preparing other ceramic-like systems at low temperature. 1995, 78, 1769. 19 P. Li, C. Ohtsuki, T. Kokubo, K. Nakanishi, N. Soga, T. Nakamura and T. Yamamuro, J. Mater. Sci. Mater. Med., 1993, Acknowledgments 4, 127. 20 M. A.Walters, Y. C. Leung, N. C. Blumenthal, R. Z. Legeros and The authors are indebted to Ing. Francisco Rodrý�guez K. A. Konsker, J. Inorg. Biochem., 1990, 39, 193. Melgarejo from CINVESTAV, Quere�taro and to Dra. 21 J. W. P. Rothwell, J. S.Waugh and J. P. Yesinowski, J. Am. Chem. Antonieta Mondrago�n from IFUNAM, for their valuable Soc., 1980, 102, 2637. assistance in the Raman measurements. 22 P. Li, K. Nakanishi, T. Kokubo and K. de Groot, Biomaterials, 1993, 14, 963. References 1 K. A. Khor, P. Cheang and Y. Wang, JOM, 1997, 49, 51. Paper 8/03503B 2 R. Rodrý�guez, J. Coren�o and V. Castan� o, Adv. Comp. Lett., 1996, 5, 25. 2812 J. Mater. Chem., 1998, 8, 280

ISSN:0959-9428    

DOI:10.1039/a803503b

出版商:RSC    

年代:1998

数据来源: RSC