摘要:

Spirooxazine- and spiropyran-doped hybrid organic–inorganic matrices with very fast photochromic responses Barbara Schaudel,a Ce�line Guermeur,a Cle�ment Sanchez,*a Keitaro Nakatanib and Jacques A. Delaireb aL aboratoire de Chimie de la Matie`re Condense�e, URA CNRS 1466, Universite� Pierre etMarie Curie, 4 Place Jussieu, 75252 Paris, France bL aboratoire de Photophysique et Photochimie Supramole�culaires etMacromole�culaires, URA CNRS 1906, ENS CACHAN, 61 Avenue du Pre�sident Wilson, 94235 Cachan Cedex, France Both spirooxazine and spiropyran dyes have been embedded into two different hybrid matrices, which were formed from hydrolysis and cocondensation between diethoxydimethylsilane and zirconium propoxide and between methyldiethoxysilane (DH) and triethoxysilane (TH) respectively.The nature and the kinetics of the photochromic response depend strongly on the hydrophobic/hydrophilic balance (HHB) of the hybrid material. The HHB controls the competition between direct and reverse photochromism. The photochromic behaviour of the strongly hydrophobic spirooxazine-doped DH/TH coatings is direct, highly efficient (DA>1), reversible and extremely fast (thermal bleaching time constant, k=0.2 s-1).The photochromic kinetics of this hybrid material are, to the best of our knowledge, much faster than those reported for spirooxazine in any other solid matrix. The mild characteristics offered by the sol–gel process allow interactions on the kinetics of colouration and thermal fading. As far as photochromic devices are concerned, tuning between the introduction of organic molecules within an inorganic a strong and fast photochromic colouration (large DA) and a network.1 Inorganic and organic components can then be very fast thermal fading is needed.Usually spiropyran- or mixed at the nanometric scale in virtually any ratio, leading spirooxazine-doped sol–gel matrices or even spirooxazine- to the so-called hybrid organic–inorganic nanocomposites.2–4 doped polymeric matrices exhibit slow thermal fading (at least These hybrids are extremely versatile in their composition, several minutes).10–17 processing, and optical and mechanical properties.5 Organic This article addresses the photochromic behaviour of a molecules play an important role in optics; many hybrid optical systems such as luminescent solar concentrators, solidstate dye lasers, optical sensors, photochromic and NLO devices have been developed in the past few years.6–9 Spiropyrans and spirooxazines are two of the fascinating families of molecules exhibiting photochromic properties. Upon irradiation, the colourless spiropyran or spirooxazine undergo a heterolytic CMO ring cleavage, producing coloured forms of merocyanines (Fig. 1).The merocyanines may interact with their environment, i.e. solvent, matrix etc., leading to different photochromic responses. Levy and co-workers10,11 first demonstrated the important role played by the dye–matrix interactions in the photochromic response of spiropyrans. They studied the photochromism of spiropyrans trapped in sol–gel matrices synthesized via polymerization of Si(OCH3)4 or RSi(OEt)3 (R=ethyl, methyl, etc.) precursors, and observed two types of photochromic behaviour.When the photochromic dye is trapped within a hydrophilic domain of the matrix (domain containing residual SiMOH groups), the open zwitterionic coloured forms are probably stabilized through hydrogen bonding with the acidic silanol groups present at the pore surface.The result of this stabilization is the observation of the coloured forms before irradiation. These coloured forms can be bleached by irradiation in the visible range. This has been termed ‘reverse photochromism’. On the other hand, spiropyran dyes embedded in a more hydrophobic hybrid network made by hydrolysis of RSi(OEt)3 exhibit direct photochromism, i.e.the colourless form is stable without irradiation. Such photochromic behaviour has been reported for many spiropyran- or spirooxazine-doped sol–gel matrices.12–17 Moreover, for hybrid organic–inorganic matrices containing different chemical environments (hydrophilic and hydrophobic domains) a competition between direct and reverse photochromisms can be observed.17 However, many fundamental questions still need to be considered.Little is Fig. 1 Representation of the two photochromic dyes SP and SO and their open form known concerning the role of the photochromic dye–matrix J. Mater. Chem., 1997, 7(1), 61–65 61spiropyran, SP (6-nitro-1¾,3¾,3¾-trimethylspiro-2H-1-benzopyran- 2,2¾-indoline) and a spirooxazine, SO [1,3,3-trimethylspiroindoline- 2,3¾-(3H)-naph(2,1-b)-(1,4)oxazine] (Fig. 1) embedded within two new hybrid matrices. The tuning of dye–matrix interactions allows us to obtain spirooxazine-doped hybrid coatings which exhibit a strong and very fast photochromic response. Two kinds of hybrid matrices have been synthesized by using organically modified silicon alkoxide precursors [R¾xSi(OEt)4-x] (R¾=CH3, H), eventually cocondensed with zirconium alkoxide, Zr(OPrn)4.Fig. 2 Experimental setup used for photochromic behaviour studies Experimental Synthesis of the samples by measuring the light intensity transmitted The SP and SO dyes were purchased from Aldrich. through the sample during and after irradiation. The wave- The first SP- or SO-doped hybrid matrix was prepared as length of this probe beam was selected by a set of two follows.(CH3)2Si(OC2H5)2 (D; Fluka), absolute ethanol and monochromators. Its intensity was attenuated strongly comwater in a y50.55y molar ratio were mixed for three minutes pared to the irradiation beam. The transmitted light was under magnetic stirring. The pH of the water was adjusted to detected by a photomultiplier linked to a computer-driven 2 by addition of hydrochloric acid.The appropriate amounts digital multimeter (Keithley 2000). The incidence was close to of Zr(OPrn)4 (Fluka) were added to the solutions in order to normal for both beams. produce Zr5Si (x5y) molar ratios ranging from 10590 to 30570. Kinetics of bleaching were studied by following the fading After ageing for 1 h, the photochromic dye solution (10-2 mol of the absorbance (A) at 490 nm for SP and 610 nm for SO, dm-3 in ethanol) was added to the sol.Samples will be labelled which are the absorption maxima in the visible region for the D/Zrx, where Zr stands for the zirconium, x for the amount doped D/Zrx matrices. The thermal bleachings were fitted by of zirconium (Zr5Si, x5y). using mono [A=Bexp(-kt)+C] or bi-exponential [A= The second SP- or SO-doped matrix was prepared from the Bexp(-k1t)+Cexp(-k2t)+D] equations.hydrolysis and cocondensation of (CH3)HSi(OC2H5)2 (DH; ABCR) and HSi(OC2H5)3 (TH; Fluka), precursors. The Results and Discussion DH5TH5EtOH5H2O (pH=7) molar ratios were 0.750.350.551. The dye solution (10-2 mol dm-3 in ethanol) Materials was added after a few minutes. Samples are labelled The D/Zrx matrices have been characterized already by 13C DH70/TH30.MAS, 29Si MAS and CP MAS NMR studies.18,19 These data Bulk samples and coatings a few mm thick were prepared revealed that these D/Zrx systems are hybrid nanocomposites easily from both doped D/Zrx and DH70/TH30 sols. made from polydimethylsiloxane chains and zirconium oxopolymers. Moreover, FTIR and DTA show that the zirconium NMR experiments oxopolymers are hydrophilic domains that still contain The MAS NMR experiments were realised on a Bruker MSL hydroxo groups coming from residual ethanol or ZrMOH 300 spectrometer using a Bruker 7 mm rotor.Spectra were ligands.20 The size and the spacing between the ZrO2-based recorded with 1 ms pulses, a 0.1 s delay and a 5 kHz spinning domains is about a few nm, as indicated by SAXS.20 However speed for 17O, with 2 ms pulses, a 10 s delay and a 4 kHz the nature of the interface between the PDMS chains and the spinning speed for 29Si and with 3 ms pulses, a 10 s delay and zirconium oxo-based domains was not defined in these hybrid a 4 kHz spinning speed for 1H spectra.The positions of the materials. 17O MAS re therefore carried out NMR resonances were located taking Me4Si (29Si and 1H) and to clarify the nature of the interface.water (17O) as d 0 references. The 17O MAS NMR spectrum of the D/Zr20 matrix (Fig. 3) The low natural abundance of the 17O nucleus shows large resonances located at d 400 and 290 which [(3.7×10-2)%] and its quadrupole moment renders its detec- represent OZr3 and OZr4 respectively.21 The assignement of tion difficult. However, the use of 10% 17O-enriched water for the sharper resonance located at d 336 is not obvious at the the hydrolysis of precursors lead to a specific labelling of moment.It may be due to some residual molecular OZr4 SiMO*H, SiMO*MSi and SiMO*MM groups, and thus species. The main peak at d 73 is due to bridging OSi2 and greatly enhances their detectability compared to ROH or the broad signal around d 160 to SiMOMZr bonds.22 The SiMOR groups.FTIR spectroscopy IR spectra were recorded on powdered samples with the conventional KBr pellet technique using a 550 Magna Nicolet FTIR spectrometer. Optical experiments The photochromic behaviour of the samples was studied using the experimental setup described in Fig. 2. A xenon mercury arc lamp (450W), providing light in the UV–VIS spectrum, was used to irradiate the sample. The appropriate irradiation wavelength was chosen by means of a narrow-band (10 nm) interference filter, and commutation of a shutter allowed us to make irradiation cycles. A beam from another light source, a Fig. 3 17O MAS NMR spectrum of the D/Zr20 matrix xenon lamp (150 W), was used to follow the absorbance change 62 J.Mater. Chem., 1997, 7(1), 61–65peaks assigned to homocondensation are then the major signals. These data revealed clearly that these D/Zrx systems can be better described as a nanocomposite because homocondensation OZr3, OZr4, OSi2 species have been identified clearly. This composite is built from hydrophobic polydimethylsiloxane chains covalently linked through ZrMOMSi bonds to hydrophilic domains made of zirconium oxopolymers (Fig. 4). The DH70/TH30 matrix was characterized by 1H MAS, 29Si MAS NMR and FTIR spectroscopies. The 29Si MAS NMR spectrum (Fig. 5) exhibited only two pairs of doublets located at d 32.5, -36.8 and d 82, -87.4. These resonances are due to fully condensed DHand THunits, respectively.The doublets are due to J coupling between Si and H via SiMH bonds. They are observed in the NMR solid-state spectrum because the 29Si resonances are particularly narrow suggesting a quasiliquid behaviour of the DH and TH units. The ratio between these resonances is 70530, as in the initial mixture, showing that upon hydrolysis and cocondensation reactions the SiMH Fig. 6 1H MAS NMR spectrum of the DH70/TH30 matrix bonds of DH and TH precursors have not been cleaved.24 The 1H MAS NMR spectrum (Fig. 6) exhibited one peak at according to the integration of the peaks, are in a 0.151 ratio d 0.4 due to the methyl protons of the DH units and two with silicium. The network is then almost fully condensed. peaks at d 4.5 and 4.9 due to SiH in the TH and DH units In the FTIR spectrum of the DH70/TH30 coatings (Fig. 7) respectively. The peaks at d 1.4 and 4.0 which correspond to the presence of strong IR bands located at 2237 and 2176 cm-1 CH3 and CH2 groups, respectively, are due to residual species: confirmed that the SiMH bonds of the DH and TH precursors ethoxy groups and asmall proportion of ethanol. These species, have not been cleaved.These bands correspond to nSiMH in TH and DH units respectively. The DH70/TH30 coatings also exhibit strong IR bands located at 1000–1100 cm-1, indicating the formation of SiMOMSi linkages. Moreover, the 2300–4000 cm-1 frequency range (the nOMH region) is absolutely flat, suggesting that these matrices have an extremely low hydroxy group content. In agreement with data reported previously23 the strongly hydrophobic DH/TH network can be described as a copolymer formed by short chains of DH units crosslinked by TH units.Both D/Zrx and DH/TH exhibit glass-transition temperatures at about -100 °C23,24 and their specific areas measured by nitrogen adsorption porosimetry are extremely low (<5 m2 g-1). These two matrices are very flexible and do not present any open porosity under nitrogen probing.At room temperature, both SP and SO dyes embedded in these hybrid matrices exhibit good stability. However, the photostability of these materials is currently under investigation. Photochromic properties D/Zrx matrices doped with SP or SO are lightly coloured (pink with SP or blue with SO) before irradiation. However, Fig. 4 Schematic representation of the D/Zr matrix the absorbance in the visible region is weak in comparison with the total amount of embedded photochromic dyes.The amount of coloured form depends on x. Fig. 8 shows the photochromic behaviour of SP-doped D/Zrx gels for three x Fig. 5 29Si MAS NMR spectrum of the DH70/TH30 matrix Fig. 7 IR spectrum of the DH70/TH30 matrix J. Mater. Chem., 1997, 7(1), 61–65 63pyran- and spirooxazine-doped sol–gel matrices and polymers are also given.The kinetic data of the SP- or SO-doped D/Zr20 samples are similar to those reported for other modified sol–gel matrices or in organic polymers.17,25 As in organic polymers, the bleaching follows a biexponential equation which can be explained by an inhomogeneous distribution of free volumes in the gel.Moreover, the presence of different stereoisomers (cis or trans) could also account for this behaviour. The different isomer–matrix interactions could explain the different kinetics observed for SO and SP. The thermal fading is longer for SP-doped hybrids than for SO-doped ones. This phenomenom can be correlated to the fact that SP open forms are known for their tendency to form zwitterionic species, while non-charged quinonic species are usually favoured for open SO molecules.Zwitterionic species Fig. 8 Photocolouration (lirr=320 nm) and photodecolouration (lirr= can be stabilized markedly by hydrogen bonding with the 547 nm) for SP-doped D/Zr bulks at 490 nm, (a) D/Zr10, (b) D/Zr20, matrix, thus lowering the decay times of thermal fading.(c) D/Zr30 The SO or SP DH70/TH30 doped matrices exhibit normal photochromism. All the samples are colourless before values. When the amount of zirconium increases, the irradiation. This is probably due to the strong hydrophobic absorbance variation due to the colouration decreases while character of this matrix. For the two photochromic dyes, the that due to decolouration increases: there are more open forms thermal fading can be fitted, with excellent agreement, to a in the gel.The amount of coloured form increases pro- monoexponential equation. This may be related to the quasi- portionally with x and is much higher for D/Zr30 than for liquid mobility observed by NMR for this matrix. D/Zr10 samples. The rate constants obtained for the two dyes embedded in This indicates that before irradiation the SO and SP dyes the DH70/TH30 matrix are also reported in Table 1.The are split roughly into two populations. The coloured merocya- thermal fading of SP in the DH70/TH30 matrix is faster than nine open forms of SO and SP are stabilized by hydrogen those reported for other sol–gel matrices10,11 or for PMMA.25 bonding within the hydrophilic regions of the zirconium oxopo- The time dependence of the absorption upon repeated lymers, while the closed SO and SP forms are probably located irradiation with 365 nm light for SO-doped DH70/TH30 coat- in the environment of the hydrophobic polydimethylsiloxane ings is reported in Fig. 9. The photochromic behaviour is chains. Therefore, for these D/Zrx matrices the photochromism reversible, extremely fast (k=0.2 s-1) and corresponds to a is partially reversible and can be balanced by tuning the very high absorption jump (DA=1.2).The photochromic D/Zr ratio. kinetics of this SO-doped material are, to the best of our The thermal bleaching behaviour of the D/Zr20 samples knowledge, much faster than those reported for SO in any were fitted with a biexponential equation. The SP-doped other matrix (sol–gel matrices, organic polymers, alcohols, materials exhibited a very long bleaching time (ca. 24 h) while etc.).14,15,17,25,27 for the SO-doped D/Zr20 materials the thermal fading was It is interesting to note that when embedded within the much faster. The rate constants for SO- and SP-doped D/Zr20 same DH70/TH30 matrix, SP shows a much longer thermal samples are reported in Table 1.For comparison, data from fading rate than SO. Moreover, a substantial part of the merocyanine form does not revert back to the initial form, the literature concerning the photochromic properties of spiro- Table 1 Photochromic behaviour of different spiropyrans and spirooxazines in sol–gel matrices, PMMA and ethanol (R=3-glycidoxypropyl, R¾=3-aminopropyl) chromophore matrix effecta characteristics ref.SP SiO2 D�R D:k=6.7×10-5 s-1 R: k=1.7×10-5 s-1 10 EtSiO1.5 D k=1.7×10-4 s-1 11 SiO2–Me2SiO D�R D:k=6.7×10-5 s-1 R: k=5×10-5 s-1 Me2SiO–ZrO2 D�R this work MeHSiO–HSiO1.5 D k=5×10-3 s-1 this work PMMA D k1=7×10-4 s-1 k2=10-4 s-1 25 ethanol D k=3.7×10-4 s-1 26 spiropyran SiO2 D/R t0.5=2.3×105 s MeSiO1.5 D 13 SO MeSiO1.5 D k1=1.15×10-2 s-1 17 k2=1.4×10-3 s-1 RSiO1.5–EtSiO1.5–EtSiO1.5 t0.5=2 s 14 MeSiO1.5–RSiO1.5–R¾SiO1.5–SiO2–Me2SiO D t0.5=2 s 15 Me2SiO–ZrO2 D�R k1=3.1×10-2 s-1 this work k2=2×10-3 s-1 MeHSiO–HSiO1.5 D k=0.2 s-1 this work PMMA D k1=4×10-2 s-1 25 k2=4×10-3 s-1 ethanol D k=0.2 s-1 27 spirooxazine SiO2 D/R k=1.6×10-3 s-1 13 MeSiO1.5 D k=1.2×10-2 s-1 aD, direct photochromism; R, reverse photochromism. 64 J. Mater. Chem., 1997, 7(1), 61–65doped DH/TH hybrid coatings exhibit, after a strong colouration (DA>1.2) a very fast thermal bleaching, which is, to the best of our knowledge, the fastest thermal bleaching reported for spirooxazine-doped inorganic or organic materials. References 1 H. Schmidt and B. Seiferling, Mater. Res. Soc. Symp. Proc., 1986, 73, 739. 2 C. J. Brinker and G.Scherrer, Sol–Gel Science, the Physics and Chemistry of Sol–Gel Processing, Academic Press, San Diego, 1989. 3 B.M. Novak, Adv. Mater., 1993, 5, 422. 4 C. Sanchez and F. Ribot, New J. Chem., 1994, 18, 1007. 5 Sol–Gel Optics, Processing and Applications, ed. L. C. Klein, Kluwer Academic, Boston, 1993. 6 Sol–Gel Optics I, ed. J. D. Mackenzie and D. R. Ulrich, Proc. SPIE, vol. 1328,Washington, DC, 1990. Fig. 9 Photocolouration (lirr=365 nm) and thermal bleaching for the 7 Sol–Gel Optics II, ed. J. D. Mackenzie, Proc. SPIE, vol. 1758, SO-doped DH70/TH30 film at 610 nm Washington, DC, 1992. 8 Sol–Gel Optics III, ed. J. D. Mackenzie, Proc. SPIE, vol. 2288, Washington, DC, 1994. even after two months. This phenomenom is probably due to 9 B. Dunn and J. I. Zink, J.Mater. Chem., 1991, 1, 903. the favoured zwitterionic form of the merocyanine SP which 10 D. Levy and D. Avnir, J. Phys. Chem., 1988, 92, 734. should interact slightly with the weakly polar d+SiMHd- 11 D. Levy, S. Einhorn and D. Avnir, J. Non-Cryst. Solids, 1989, 113, 137. bonds28 of the DH/TH matrix. As a consequence, the return 12 D. Preston, J. C. Pouxviel, T. Novinson, W.C. Kaska, B. Dunn of the SP dye to the closed form is slowed by these interactions. and J. I. Zink, J. Phys. Chem., 1990, 94, 4167. 13 H. Nakazumi, R. Nagashiro, S. Matsumoto and K. Isagawa, SPIE Proc. Vol 2288, Sol–Gel Optics III, Proc. SPIE, vol. 2288, San Conclusions Diego, 1994. 14 L. Hou, B. Hoffmann, M. Menning and H. Schmidt, J. Sol–Gel Sci. Photochromic hybrid materials were prepared by using SO T echnol., 1994, 2, 635.and SP dyes in two different hydrid matrices (D/Zr and 15 L. Hou, M. Menning and H. Schmidt, J. Sol–Gel Sci. T echnol., DH/TH). These experiments show the very high sensitivity of 1996, in press. the photochromic behaviour of SO and SP dyes to dye–matrix 16 L. Hou, M. Menning and H. Schmidt, Proc Eurogel’92, 1992, 173. interactions.The sol–gel materials, in particular the hybrid 17 J. Biteau, F. Chaput and J. P. Boilot, J. Phys. Chem., 1996, 100, ones, allow tuning of these interactions, which are of para- 9024. 18 S. Dire�, F. Babonneau, C. Sanchez and J. Livage, J. Mater. Chem., mount importance for control of the kinetics. The first matrix, 1992, 2, 239. D/Zr, shows reverse and direct photochromism. The ratio 19 S.Dire�, F. Babonneau, G. Carturan and J. Livage, J. Non-Cryst. between reverse and direct photochromism increases with the Solids, 1992, 147 & 148, 62. amount of hydrophilic zirconium oxo polymers present in the 20 C. Guermeur and C. Sanchez, to be published. material and emphasises that the open forms of these dyes are 21 T. J. Bastow, M. E. Smith and H. J. Whitfield, J.Mater. Chem., trapped via their interaction with residual MMOH groups. As 1992, 2, 989. 22 F. Babonneau, J. Maquet and J. Livage, Ceram. T rans., 1995, 55, in matrices presenting an inhomogeneous distribution of free 53. volume, the photodynamics of these SO- and SP-doped D/Zr 23 G. D. Soraru, G. D’Andrea, R. Campostrini and F. Babonneau, materials are not first order. J. Mater. Chem., 1995, 5, 1363. In contrast, the photodynamics of SO- and SP-doped hybrid 24 F. Babonneau, L. Bois, J. Livage and S. Dire�, Mater. Res. Soc. matrices made from hydrolysis of methyldiethoxysilane and Symp. Proc., 1993, 286, 289. triethoxysilane (DH/TH) materials are first order, in agreement 25 Y. Atassi, Thesis, Ecole Nationale Supe�rieure de Cachan France, 1996. with the pseudo-liquid behaviour observed by solid-state NMR 26 J. B. Flannery, J. Chem. Soc. A, 1968, 5660. studies. SP dyes are very sensitive to weak dye–matrix inter- 27 Applied photochromic polymer systems, ed. C. B. McArdle, New actions and are able to probe the weak polarity of SiMHbonds. York, 1991. The negative partial charge carried by the hydrogen atoms 28 A. P. Altshuller and L. Rosenblum, J. Am. Chem. Soc., 1955, 77, of the SiMH bonds makes these DH/TH matrices strongly 272. hydrophobic. This hydrophobicity is responsible for the direct and very fast photochromic behaviour observed. The SO- Paper 6/06859F; Received 7th October, 1996 J. Mater. Chem., 1997, 7(1)

ISSN:0959-9428    

DOI:10.1039/a606859f

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Unsupported SiO2-based organic–inorganic membranes Part 1.—Synthesis and structural characterization Sandra Dire`,*a Eva Pagani,a Florence Babonneau,b Riccardo Ceccatoa and Giovanni Carturana aDipartimento di Ingegneria dei Materiali, Universita` di T rento, v.Mesiano 77, 38050 T rento, Italy bChimie de la Matie`re Condense�e, Universite� Pierre etMarie Curie/CNRS, 4, place Jussieu, Paris, France Tetraethoxysilane (TEOS) and methyltriethoxysilane (MTES) have been used to prepare hybrid SiO2-based membranes.These self-supported materials were obtained from controlled polymerization reactions for various TEOS/MTES molar ratios ensuring the achievement of crack-free disks 8 cm in diameter and 10–40 mm in thickness. The rheological behaviour of precursor solutions was studied and gelling times were determined.The whole process, from starting solution to xerogel, was followed by FTIR spectroscopy, viscosity measurements and multinuclear solid-state NMR, and is discussed in terms of the hydrolysis–condensation kinetics of tetrafunctional and trifunctional silicon alkoxides. Density, shrinkage, elastic modulus (E), modulus of rupture (MOR) and elongation at break were all determined and related to preferential structural arrangements of networks according to the TEOS/MTES ratio.Hybrid organic–inorganic materials are currently under inves- solution at pH=1.5. Various water/alkoxide ratios were calcutigation as potential multifunctional and high-performance lated for the different compositions, so that H2O/SiOEt=0.5. materials: organic modified ceramics (ORMOCERs) properties No solvent was used except for the pure Si(OEt)4 gel (ethanol).are exploited in the fields of adhesives, sealing and modified Details concerning the labelling of samples, compositions and glass surface materials,1–3 sensors4 and artificial membranes.5–8 preparation conditions are reported in Table 1. As regards membrane technology, the extension of appli- Solutions were stirred at room temperature, using storage cation fields to more and more sophisticated processes parallels times determined by solution rheology, affected by the the development of the new class of inorganic membranes, TEOS/MTES ratio.Each solution (2.4 ml) poured into poly- which can also compete with organic polymers in terms of styrene Petri dishes (diameter 8 cm) and covered by polymer chemical and thermal durability and separation performances. foil afforded clear, homogeneous gel membranes in ca. 10 days, Sensors and membranes are involved with mass transport during which time several pinholes on the foil were made. phenomena due to concentration, pressure or electrical field Samples of thickness 10–40 mm were allowed to dry in air at gradients: any chemical modification causing changes of net- constant temperature and humidity (Fig. 1). work structure and surface features affects detection or separation properties so that hybrid materials are expected to improve the development of inorganic membranes. Characterization techniques The sol–gel process can lead easily to silica-based hybrid The rheological properties of the gelling solutions were deter- materials containing SiMO and SiMC bonds: intrinsic molecu- mined at 24.8±0.2°C with a cone/plate system using a MC- lar gels with both organic and inorganic moieties are obtained 10 Physica viscosimeter.Values of shear stress, shear rate, by using R¾nSi(OR)4-n as precursors.9–11 Note that the con- temperature and viscosity were collected for the various solu- ditions chosen for hydrolysis–condensation of these precursors tions during ageing.affect the structural features and physical properties of the Solid-state NMR measurements were performed on a MSL final products.12,13 Obviously, with a constant CH3/Si/O ratio, 400 Bruker spectrometer. Solid samples were spun at 4 kHz. the ordinary synthesis of hybrid SiO2 materials, from mixtures For the 29Si single-pulse experiments (SPE), pulse width (2 ms: of Si(OR)4, as cross-linking agent and di- or tri-functional h=30°) and relaxation delays (30 s) were chosen to take into methylsilanes, may not lead to the same product as that account the long T1 relaxation times. 29Si MAS NMR spectra obtained from pure CH3Si(OR)3 ; in particular, the different were also recorded using cross polarization (CP) techniques spatial location of CH3MSi bonds may result in different and variable contact times.The spectra were analysed using porosities and pore size distributions. This apparent complithe WINNMR and WINFIT programs.14 13C CP MAS NMR cation may modify mechanical properties: in particular, concentration and distribution of CH3MSi groups may influence spectra were recorded using 2 ms contact time.Recycle delays the achievement of self-carrying items. of 10 s were used for all CP spectra. In this initial work, the sol–gel process was used to prepare FTIR spectra were recorded on a Nicolet 5DXC instrument self-supported, hybrid organic–inorganic membranes from mix- equipped with both horizontal attenuated total reflectance tures of silicon tetraalkoxide and CH3Si(OEt)3: results on the (HATR) and diffuse reflectance (DR) accessories. In the HATR development of the gel structure from these precursors and on mode, a known volume of sol was poured directly onto the the characterization of the resulting xerogels with various surface of the ZnSe crystal, closed with an appropriate cap; 64 Si(OEt)4/CH3Si(OEt)3 ratios, will be presented in this paper, scans at 45° to the IR incident radiation were collected.For while gas-separation properties and surface characterization DR measurements, xerogels were powdered and dispersed in will be reported in a future paper. KBr, accumulating 64 scans for each spectrum. Densities were measured in water and hexane at 25°C by Archimedes’ method; linear shrinkage was calculated by meas- Experimental uring the xerogels’ diameters and expressing them as percent- Synthesis of Si(OEt)4/MeSi(OEt)3 gels ages referring to the Petri dish diameter.MOR, Young’s modulus (E) and maximum elongation at break (eR ) were Various mixtures of Si(OEt)4 (TEOS) and CH3Si(OEt)3 (MTES) were hydrolysed at room temperature with an HCl obtained with a three-point bending test, using an Instrom J. Mater.Chem., 1997, 7(1), 67–73 67Table 1 Sample labels and compositions Sample label, composition, H2O/Si ethanol/Si TxMy molar ratio molar ratio molar ratio [Si]0/mol l-1 T100 100 TEOS 2 0.5 3.47 T70M30 70 TEOS/30 MTES 1.85 — 4.01 T50M50 50 TEOS/50 MTES 1.75 — 4.12 T30M70 30 TEOS/70 MTES 1.65 — 4.24 M100 100 MTES 1.5 — 4.42 Fig. 1 T70M30 xerogel disk testing machine with a load cell of 100 N and a testing rate of 1 mm min-1. Results The hydrolysis and condensation process of MTES and TEOS precursors at various ratios was studied using different techniques, in order to compare the advance of the gelling process and of the gel network structure. IR spectra Fig. 2(a) shows FTIR spectra recorded in the HATR mode in the interval 4000–650 cm-1 for T100, T50M50 and M100 samples, 15 min after preparation of the solutions.In the highfrequency field, CMH stretching vibrations at 3000–2850 cm-1 and a wide band centred at 3350 cm-1 due to OMH stretching were present. The bending vibrations of the Et and Me groups were observed in the interval 1500–1300 cm-1; for compositions containing MTES, a precise signal due to SiMCH3 stretching was recorded at 1276 cm-1.The 1200–1000 cm-1 region showed overlapping signals of OMSiMO bonds in different species. The signal due to SiMOEt bonds at 954 cm-1 was recognized clearly in the T100 spectrum, absorption due to ethanol being observed at 880 cm-1; the M100 spectrum displays peaks at 916 cm-1 attributed to SiMOH and SiMO- bonds,15 and that of ethanol at 880 cm-1.The signal at 955 cm-1 was virtually lost, suggesting that the hydrolysis of ethoxide groups in this sample was almost completed. In the T50M50 spectrum, an intermediate situation appeared with peaks at955, 920 and 880 cm-1.Spectra recordedafter 120 min are shown ig. 2(b). In the T100 sample, the band at 3350 cm-1 revealed large amounts of water and ethanol, which decreased as the MTES content increased, in agreement with a decrease of the signal at 1637 cm-1 [d(OH)].The M100 sample was characterized by almost complete hydrolysis, the SiMCH3 signal at 1276 cm-1 and the SiMOH one at 916 cm-1 still being visible. Instead, a remarkable concentration of SiMOEt was observed in T100, while T50M50 showed intermediate behaviour. The M100 spectrum recorded between Fig. 2 HATR-FTIR spectra recorded (a) 15 min and (b) 120 min after 120 min and gelling time (220 min) remained virtually preparation of the T100, T50M50 and M100 sols unchanged, whereas slow, progressive evolution was found for 68 J. Mater. Chem., 1997, 7(1), 67–73T100 in the interval 120–330 min, i.e. T100 gelling time.This between the viscosity increase and composition; samples with higher concentrations of MTES displayed a viscosity increase supports the fact that hydrolysis–condensation processes still occur for T100 when the MTES sample had already gelled. at 238–250 min ageing time, whereas TEOS-rich compositions showed a net increase after 327–342 min, the slope of the curve Fig. 3 reports 4000–400 cm-1 spectra recorded in the DR mode, which revealed the spectral window between 650 and being appreciably smaller. A trend discontinuity was observed for T30M70, whose transition occurred at 226 min. 400 cm-1. TxMy films, stored at the same temperature and humidity, show that adsorbed water increased as %TEOS Shear stress vs. shear rate and viscosity vs. shear rate plots are shown in Fig. 5. M100 lost Newtonian flow behaviour increased (bands at 3300 and 1630 cm-1). In the 3000–2800 cm-1 interval, CMH stretching vibrations of after 200 min [Fig. 5(a)], while T100 showed it up to 305 min [Fig. 5(b)]; in both cases, further ageing produced yield behav- residual OEt and Me groups were present; the intensity of the peak at 1272 cm-1 [n(SiMCH3 )] decreased as the MTES iour [Fig. 5(c) and (d)].17 TxMy compositions displayed intermediate behaviour. concentration decreased. Other features were the intensity inversion of peaks in the 1200–1000 cm-1 interval and the lowering of the peak at 940 cm-1 (SiMOH and SiMOEt) as NMR results the amount of MTES increased. The position of the peak Solid-state NMR (MAS NMR) can provide direct information corresponding to the angular deformation d(SiMOMSi) and on the local environment of different structural units, and thus related to the silica units arrangement,16 at 466 cm-1 in the on the degree of condensation of the network.Fig. 6 shows T100 spectrum, was shifted to lower frequencies as %MTES the 29Si MAS NMR spectra of T100, T50M50 and M100 increased. recorded as single-pulse experiments (SPE); Table 2 shows the percentages of various units obtained after simulation of SPE Viscosity measurements spectra, as well as the degree of condensation of the network.Fig. 4 reports the evolution of solution viscosities of the TxMy It confirms what was published previously,18 that the degree samples. Viscosity vs. time diagrams indicate a relationship of condensation decreases with the average functionality of the precursors. It is interesting to note that the degree of condensation of the T50M50 sample is intermediate between those of the T100 and M100 samples.Series of 29Si CP MAS NMR spectra with variable contact times were also recorded on the three samples. Analysis of the variation of magnetization vs. contact time was carried out according to the simplest model describing CP between two spin reservoirs, one for dilute spins, S, and one for abundant spins, I, using the well known formula:19 MS(tc)=cI cS M0S 1 1-lG1-exp[-(1-l) tc TIS ]Hexp(- tc T 1rI) with l=TIS/T1rI.M0S is the magnetization at the equilibrium in the static field B0, TIS is the CP standard time which is related to the strength of the I–S dipolar coupling and T1rI is the relaxation time of the abundant spins in the rotating frame, which will cause a loss of magnetization for long contact time.cI and cS are the magnetogyric ratios for spins I and S, respectively. This formula applies for S spins which are partially decoupled from the protons. The TSiH and T1rH fitted values are reported in Table 2 as well as the percentages of the various units detected by CP, and estimated from the M0S value. For T100 and M100 samples, the agreement factor for the fitting procedure, R, is > 0.99, showing a good agreement between the experimental behaviour and the theoretical model.In the case of the T50M50 CP MAS spectrum, a lower agreement factor is obtained (R>0.97) which may be related to the presence of several local environments for the different T and Q sites.Fig. 3 DRIFT spectra of the TxMy xerogels Two interesting features can be pointed out from this study. First, a comparison between the quantitative analyses of the various sites carried out from the SPE and CP spectra shows perfect agreement for the M100 and T50M50 samples, and a discrepancy for the T100 sample.In this last sample, the number of Q4 units is underestimated in the CP technique, suggesting that the analysis of the CP dynamics measures only the amount of Q4 units close to Q3 or Q2 units. In contrast, in the T50M50 system, the Q4 units should be in close proximity to T units, which allows their detection via the CP technique. A second interesting point is the difference in the TSiH values corresponding to the Q units, in the T100 and T50M50 samples.Their decrease in the T50M50 sample indicates stronger 1H–29Si dipolar coupling for these units and could be related to the close proximity of the T units. These results strongly suggest a good chemical homogeneity of the Fig. 4 Viscosity vs. time diagrams of the TxMy samples: 1, M100; #, T100; %, T50M50; +, T30M70; ×, T70M30 T and Q units within the T50M50 sample.J. Mater. Chem., 1997, 7(1), 67–73 69Fig. 5 Viscosity vs. shear rate [(a) and (b)] and shear stress vs. shear rate [(c) and (d)] plots of M100 [(a) and (c)] and T100 [(b) and (d)] samples present in the 1H MAS NMR spectrum of the T50M50 xerogel [Fig. 7(b)]. Physical and mechanical results The density, shrinkage and mechanical features of the xerogel membranes were also studied.Densities are reported in Table 3; a linear decrease as the percentage of MTES increased was observed. Shrinkages, calculated as described in the Experimental section, displayed a linear relationship with MTES content (Fig. 8), proving the absence of linear shrinkage for M100. Mechanical measurements performed on thin bars of the xerogel were obtained by a three-point bending test; the results are reported in Table 4.The elastic modulus, E, is an intrinsic property depending on the bond density and related to the material structure. The E value of T100 agrees with reported values of silica xerogels prepared under acidic conditions, i.e. 5–10 GPa.20 The decrease in elastic modulus as the organic load increased accounted for the lower cross-linking consequent upon the introduction of a trifunctional precursor.The consequent reduction in SiMO bond density also affected the maximum elongation at break eR , which increased with MTES concentration. Fracture surfaces were featureless and disks appeared dense and macroscopically homogeneous, as shown in Fig. 9. Discussion For a two-step acid–base process, van Bommel et al.21 reported that, under acidic and neutral conditions, the hydrolysis rate Fig. 6 29Si MAS NMR spectra of M100, T50M50 and T100 xerogels of alkyl-substituted silicon alkoxides is faster than for TEOS; moreover, condensation already occurs in the acid step. These observations have been confirmed recently:22a after the addition 13C CP MAS NMR and 1H MAS NMR spectra (Fig. 7), recorded on M100 and T50M50 samples, show the presence of water, under acidic conditions, the hydrolysis of alkylsubstitutedalkoxides with short alkyl chains is almost complete of SiMMe groups with signals at d -3.7 and 0.4, respectively. Moreover, these spectra indicate an incomplete hydrolysis– within the first few minutes and the condensation degree increases quickly during the first 2 h and then slows. condensation process.As a matter of fact, the signals due to residual SiMOEt groups are present in the 13C CP MAS NMR Results obtained for TxMy samples are consistent with these works, based on viscosity and NMR results:21,22 the HATR- spectra (d 18.4 and 58.1 for M100, d 17.6 and 59.9 for T50M50), and a signal attributed to SiMOH terminal units (d 4.1) is FTIR spectra indicate that MTES is hydrolysed faster than 70 J.Mater. Chem., 1997, 7(1), 67–73Table 2 29Si solid-state NMR data units obtained (%) degree of sample sitea d TSiH/ms T1r/ms SPE CP condensation M100 T3 -65.2 1.7 52 88 86 0.96 T2 -57.6 1.1 48 12 14 T100 Q4 -110.9 8.7 82 49 25 Q3 -101.4 4.2 27 44 64 0.85 Q2 -92.1 3.6 21 7 11 T50M50 T3 -63.1 2.2 123 45 42 T2 -55.3 1.6 70 5 10 Q4 -108.9 5.1 177 21 20 0.89 Q3 -100.7 2.2 113 25 25 Q2 -91.7 1.5 65 4 3 a Tn and Qn: n=number of bridging oxygens.Fig. 7 1H MAS NMR and 13C CP MAS NMR spectra of (a) M100 and (b) T50M50 xerogels Table 3 TxMy density results Table 4 Mechanical characterization of TxMy samples sample density/g cm-3 sample E/GPa eR (%) MOR (MPa) T100 1.84±0.04 T100 5.2±0.6 7 40±10 T70M30 1.61±0.03 T50M50 2.6±0.3 19 50±8 T50M50 1.48±0.02 M100 0.7±0.1 43 30±10 T30M70 1.34±0.04 M100 1.29±0.05 Fig. 9 Fracture surface of the T50M50 bar (SEM image) Fig. 8 Linear shrinkage vs. %MTES of TxMy disks J. Mater. Chem., 1997, 7(1), 67–73 71TEOS, since the concentration of ethoxy groups for M100 is gels (Table 2) and the lower agreement factor between the experimental and simulated spectra.The decrease in TSiH of Q reduced markedly after 15 min; moreover, hydrolysis seems to be complete in 120 min, and extensive condensation of silanols units in T50M50, compared also with results reported for silica gels prepared from TEOS in acid conditions,24 could be may be deduced from the negligible evolution of the M100 spectra from 120 min to gelling time.The different behaviour ascribed to the formation of an intimate mixture between silica gel and organic modified moieties.25 TSiH values of 1 ms have in the reactivities of M100 and T100 under acidic conditions may be ascribed to the inductive effect of the CH3 group, been reported for T units in which the Si atom is surrounded by three bridging oxygens.26 The slight increase in TSiH of T leading to activated species 1 which is lower in energy than 2, ultimately resulting in faster hydrolysis–condensation reactions units in T50M50, owing to a low effectiveness of magnetization transfer for short contact times, could be related to a different for MTES.23 mobility of T units in the network.Peeters et al.27 recently studied the functionality of hybrid gels obtained from mixtures of TEOS and various trifunctional organosilanes and concluded that, in spite of the enhanced condensation of Q and T units, the total number of network bonds decreases with increasing substitution level. This general This interpretation is consistent with the short time required behaviour may be extended to our results on the physical and for M100 viscosity increase.The fast condensation reaction mechanical properties of the gels. Density values decrease from leads to extended oligomers, which collapse to the gel network 1.84 (T100) to 1.29 g cm-3 (M100), as a result of a less well by condensation of few residual SiMOH bonds: indeed, the interconnected network. The same fact is observed for the E gelling time of M100 is shorter than that of T100, in which and eR values: these parameters are related to bond density the viscosity increase takes much longer, suggesting slower, and reflect increased network mobility as %MTES continuous SiMOH condensation.17 The gelling kinetics for increases.13,28 Linear shrinkage is maximum for T100 and M100 and T100 solutions having intermediate compositions decreases with increasing MTES content.Shrinkage may be do not lead to behaviours corresponding to the sum of M100 ascribed to byproduct release and condensation between and T100 behaviours multiplied by the molar fractions of residual SiMOH groups after gel formation; the latter process relevant precursors: for instance, the viscosity trend of T50M50 is prominent in our case, owing to the experimental conditions is very close to that of M100 (Fig. 4). This observation deserves employed. The possibility of affording a siloxane bond depends some specific comments. If gelling time is considered as a on the concentration of terminal SiMOH groups. Solid-state kinetic parameter which substantiates the occurrence of a NMR spectra show that the M100 xerogel displays slow closed SiO2 network, i.e.the occurrence of a structure where availability of SiMOH, since the T2 units (OH and OEt most oxygens bridge two Si atoms, parameter t, defined as: terminal units) only amount to 12% and terminal silanols may be not present, in agreement with our 1H MAS NMR and t=Si functionality gelling time [Si]0 FTIR results. These facts and the hydrophobicity of sample M100,29 preventing hydrolysis of residual SiMOR groups by where [Si]0=silicon concentration in the starting solutions adsorbed water, accounts for the absence of linear shrinkage (Table 1), Si functionality=number of SiMOR in the starting observed in M100.In contrast, the shrinkage of TxMy xerogels precursor, quantifies the rate required for engaging the SiMO increases with %TEOS, owing to the greater concentration of bonds of the precursors in SiMOMSi units per volume unit.SiMOH groups and to the possible involvement of moisture Values of t for each sample are shown in Fig. 10 vs. MTES to complete SiMOR hydrolysis. mol%. Clearly, MTES accelerates the occurrence of the gel; in fact t increases above the value expected if the process resulted Provincia Autonoma di Trento is greatly acknowledged for from the sum of the individual gelling processes of TEOS and financial support.MTES (Fig. 10, dotted line). From the viewpoint of reaction mechanism, this fact implies that terminal SiMOH groups References derived from TEOS are involved more rapidly in condensation to SiMOMSi upon reaction with SiMOH groups derived from 1 H.Scholze, J. Non-Cryst. Solids, 1985, 73, 669. MTES. Indeed, partial cocondensation between T and Q units 2 H. Schmidt, H. Scholze and G. Tunker, J. Non-Cryst. Solids, 1986, may be presumed in the T50M50 gel, considering the variations 80, 557. 3 J. Wen, V. J. Vasudevan and G. L. Wilkes, J. Sol–Gel Sci. T echnol., in TSiH of T and Q units compared to T100 (Q) and M100 (T) 1995, 5, 115. 4 P. Lacan, P. Le Gall, I. Rigola, C. Lurin, D. Wettling, C. Guizard and L. Cot, Sol–Gel Optics II, Proc. SPIE Vol. 1758, ed. J. D. Mackenzie, SPIE, Washington, DC, 1992, p. 464. 5 A. Kaiser, H. Schmidt and H. Bottner, J.Membr. Sci., 1985, 22, 257. 6 C. Guizard, N. Ajaka, M. P. Besland, A. Larbot and L. Cot, in Polyimides and Other High T emperature Polymers, ed. M. K. M. Abadie and B.Sillion, Elsevier, Amsterdam, 1991, p. 537. 7 C. Guizard and P. Lacan, in Proc. 1st Eur. Workshop on Hybrid Organic Inorganic Materials, Bierville, November 8–10, 1993, ed. C. Sanchez and F. Ribot, CNRS, Paris, 1993, p. 153. 8 T. Okui, Y. Saito, T. Okubo and M. Sadakata, J. Sol–Gel Sci. T echnol., 1995, 5, 127. 9 S. Dire`, F. Babonneau, C. Sanchez and J. Livage, J. Mater. Chem., 1992, 2, 239. 10 Z. Zhang, Y. Tanigami, R. Terai and H. Wakabayashi, J. Non- Cryst. Solids, 1995, 189, 212. 11 F. Babonneau, L. Bois, J. Maquet and J. Livage, in Eurogel 91, ed. S. Vilminot, R. Nass and H. Schmidt, Elsevier, Amsterdam, 1992, p. 319. 12 W. G. Fahrenholtz and D. M. Smith, Mater. Res. Soc. Symp. Proc., 1992, 271, 705. Fig. 10 Parameter t vs. MTES mol(%) (t=Si functionality 13 H.H. Huang, B. Orler and G. Wilkes, Macromolecules, 1987, 20, 1322. ×[Si]0/gelling time) [$, experimental; #, calculated (mixture rule)] 72 J. Mater. Chem., 1997, 7(1), 67–7314 Programs from Bruker Spectrospin, Wissembourg, France. 1994, 346, 365; (b) S. Prabakar, R. A. Assink, N. K. Raman and C. J. Brinker,Mater. Res. Soc. Symp. Proc., 1994, 346, 979. 15 R. M. Almeida, T. A. Guitton and C. G. Pantano, J. Non-Cryst. Solids, 1990, 121, 193. 23 C. J. Brinker and G. W. Scherer, Sol–Gel Science, Academic Press, New York, 1990, ch. 3. 16 F. Babonneau, K. Thorne and J. D. Mackenzie, Chem. Mater., 1989, 1, 554. 24 K. L.Walther, A.Wokaun and A. Baiker,Mol. Phys., 1990, 71, 769. 25 F. Babonneau, Mater. Res. Soc. Symp. Proc., 1994, 346, 949. 17 M. D. Sacks and R. Sheu, J. Non-Cryst. Solids, 1987, 92, 383. 18 R. H. Glaser, G. L. Wilkes and C. E. Bronnimann, J. Non-Cryst. 26 G. S. Carajava, D. E. Leyden, G. R. Quinting and G. E. Maciel, Anal. Chem., 1988, 60, 1776. Solids, 1989, 113, 73. 19 M. Mehring, Principles of High Resolution NMR in Solids, 27 M. P. J. Peeters, W. J. J. Wakelkamp and A. P. M. Kentgens, J. Non-Cryst. Solids, 1995, 189, 77. Springer-Verlag, Berlin, 1983, p. 129. 20 M. J. Muratagh, E. K. Graham and C. G. Pantano, J. Am. Ceram. 28 R. H. Glaser and G. L. Wilkes, Polym. Bull., 1988, 19, 51. 29 C. Della Volpe, S. Dire` and E. Pagani, J. Non-Cryst. Solids, 1996, Soc., 1986, 69, 775. 21 M. J. van Bommel, T. N. M. Bernards and A. H. Boonstra, J. Non- in press. Cryst. Solids, 1991, 128, 231. 22 (a) L. Delattre and F. Babonneau, Mater. Res. Soc. Symp. Proc., Paper 6/03554J; Received 21stMay, 1996 J. Mater. Chem., 1997, 7(1), 67–73 73

ISSN:0959-9428    

DOI:10.1039/a603554j

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Molecular epitaxy of perfluoroicosane on PTFE tribological transfer films studied by XPS and RAIRS Graham Beamson,* David T. Clark, Neil W. Hayes and Daniel S-L. Law Research Unit for Surfaces, T ransforms and Interfaces, Daresbury L aboratory, Warrington, Cheshire, UK WA4 4AD The molecular orientation of perfluoroicosane (PFI), n-C20F42, vacuum-deposited onto plain substrates (silicon and gold) and onto poly(tetrafluoroethylene) (PTFE) tribological transfer films, has been studied by angle-dependent X-ray photoelectron spectroscopy (XPS) and reflection–absorption IR spectroscopy (RAIRS).On the plain substrates PFI forms an ordered film with the molecular chains oriented perpendicular to the substrate. However, when deposited onto a PTFE tribological transfer film, epitaxial growth occurs such that the PFI chains align with the PTFE chains, parallel to the substrate.For PFI on silicon, XPS reveals a surface chemical shift for the uppermost CF3 groups of the sample. Tribological transfer films of poly(tetrafluoroethylene) (PTFE) substrate.25 However, because the crystal lattice parameters were first studied in the 1960s and 1970s,1,2 but their ability and helix pitch of PFI and PTFE are very closely to induce epitaxy effects in deposited overlayers has been matched,24,26,27 epitaxial growth of PFI on a PTFE transfer recognised only recently.3 The films are formed when PTFE is film might be expected to occur such that the PFI chains align drawn, under load, across a smooth substrate at temperatures parallel with the substrate.Here we present angle-resolved below the polymer melting point (327 °C). Several techniques XPS and reflection–absorption IR spectroscopy (RAIRS) data have been used to characterise PTFE tribological transfer which clearly demonstrate that this is the case. films, including X-ray photoelectron spectroscopy (XPS),4 Fourier-transform IR spectroscopy (FTIR),5,6 near-edge X-ray absorption fine structure (NEXAFS),7 electron microscopy1,2 and atomic force microscopy (AFM).8–10 The films consist of Experimental PTFE ribbons, typically 200–1000 nm wide by 5–10 nm high PTFE tribological transfer films were prepared by drawing a and many hundreds of micrometres long, aligned with the PTFE blade, under load, across a heated substrate at a film-draw direction and separated by channels of bare sub- controlled speed.4 The PTFE blade (ca. 13 mm wide×0.5 mm strate. The thickness of the ribbons and the coverage of the thick) was first cleaned and conditioned by drawing several substrate increases with deposition temperature and load.4 times across a clean silicon wafer at 270°C, under a load of Within the ribbons the helical PTFE chains are also aligned 750 g and at 0.6 mm s-1. The silicon wafer was then discarded with the draw direction.4,7–10 and replaced with either a fresh wafer or a gold coated glass A wide range of materials, including small organic molecules microscope slide.The PTFE transfer film was deposited in a and polymers, form ordered overlayers when deposited onto single sweep under identical conditions to those used for PTFE transfer films.3,11–21 Deposition techniques include from cleaning.Silicon wafer substrates (75 mm diameter, 100 the vapour phase, from the melt and from solution.3 This is orientation) were used as received from the supplier. The gold- an example of the general phenomenon of polymer on polymer coated slides were prepared by argon ion cleaning and vacuum epitaxy,22 which is thought to arise from a combination of deposition of ca. 1 nm of chromium to promote adhesion, classical epitaxy, i.e. the matching of specific crystal lattice followed by ca. 200 nm of gold. parameters in the substrate with those in the overlayer, and Films of PFI were deposited onto ca. 12×12 mm2 pieces graphoepitaxy, molecular alignment in the overlayer induced of the four different substrates (silicon, gold/glass, PTFE/ by the gross aligned morphology of the substrate.Polymer on silicon and PTFE/gold/glass) by heating a small quantity of polymer epitaxy already has technological applications in the the powder (Aldrich) in an electric crucible in vacuum (ca. field of liquid crystal displays and future applications may 10-5 mbar). Under these conditions, PFI sublimes below its include the molecular alignment of conducting polymer films melting point (164 °C).The substrates were positioned for electronic devices. Several publications describethe oriented ca. 30 mm above the crucible and held at room temperature. growth of conducting polymer films on a PTFE transfer Films of PFI prepared in this way were sufficiently thick to film.3,13–15 give a metallic blue interference colour. Linear oligomers of PTFE, CnF2n+2, are known to form PFI/gold/glass and PFI/PTFE/gold/glass samples were ordered overlayers on PTFE transfer films3 but details of studied by RAIRS using a Biorad FTS60 spectrometer with a molecular orientation in the overlayers have not been reported.wide-band MCT (mercury cadmium telluride) detector and a The deposition of perfluorotetracosane, n-C24F50, onto copper Spectratech FT80 specular reflectance accessory.The IR beam and rubbed low molecular mass PTFE has been investigated was incident at 80° to the normal of the sample plane and by NEXAFS.23 This shows that on copper the n-C24F50 spectra were collected in the range 4000–400 cm-1 at a reso- molecular chains orient perpendicular to the substrate whereas lution of 4 cm-1. Two hundred scans were coadded and ratioed on rubbed PTFE they align parallel with the PTFE surface.against the spectrum of the relevant substrate as reference. In it’s ‘natural’ state, perfluoroicosane (PFI), n-C20F42,forms PFI/silicon and PFI/PTFE/silicon samples were analysed lamellar crystals with stacked layers of helical molecular chains.by XPS using a Scienta ESCA300 spectrometer. This employs Within each layer the molecules are arranged on a hexagonal a high-power rotating anode and monochromated Al-Ka X- lattice with their long axes perpendicular to the plane of the ray source (hn=1486.6 eV), high transmission electron optics layer.24 When vapour deposited onto a plain substrate, PFI orients with the molecular chains perpendicular to the and a multichannel detector.28,29 The geometry of the X-ray J.Mater. Chem., 1997, 7(1), 75–78 75source, sample manipulator and electrostatic lens is such that considerably more intense than the A2 species. However, in RAIRS the metal surface selection rule allows only those at low electron take-off angle (h, defined relative to the sample surface) the X-ray beam strikes the sample at grazing incidence. vibrations with a transition dipole moment perpendicular to the surface to be observed.33,34 Hence, the occurrence of This gives enhanced sensitivity at low h and is ideal for the study of thin films.29 The samples were mounted on standard strong parallel polarised vibrations and weak perpendicular polarised vibrations implies that the PFI molecules are stubs with double-sided tape and analysed at ambient temperature.The pressure in the analysis chamber of the spectrometer arranged with the molecular chain axis perpendicular to the substrate. The observation of weak E1 bands may be due to was ca. 10-9 mbar, and the X-ray source power was 2.8 kW. Charge compensation was achieved using a low-energy electron slight deviations of the molecular chains from perpendicular or to slight misalignment of the optical pathway of the flood gun (Scienta FG300) with settings adjusted to give minimum peak widths.Survey and C 1s, F 1s region spectra spectrometer. The RAIRS spectrum of PFI/PTFE/gold/glass is quite were recorded for each sample at h=90° and 10°.The entrance slit to the hemispherical analyser was 0.5 mm wide and a pass different from that of PFI/gold/glass. The strong parallel polarised vibrations have almost disappeared and new bands energy of 150 eV was used, giving an overall instrument resolution of ca. 0.32 eV.30 Quantification was achieved using have appeared at 1296 (sh), 1283 (sh), 1264, 1230 (sh), 1166 and 556 cm-1.The 1264, 1230, 1166 and 556 cm-1 vibrations C 1s and F 1s sensitivity factors measured using clean PTFE tape. Total acquisition times for survey and region spectra at are assigned to the perpendicular fundamental bands E1(1), E1(2), E1(3) and E1(4), and the 1296 and 1283 cm-1 the two values of h were ca. 40 min per sample. X-Ray-induced sample degradation during this time is judged to be insignifi- vibrations to the perpendicular progression bands n2(7) and n2(8), respectively.The implication is that the molecular chain cant. Some sublimation of the PFI films occurred during exposure to the vacuum system, detected by a lightening of axis now lies parallel to the substrate surface. Again, the observation of weak parallel polarised bands may be due to the film interference colour after ca. 60 min in vacuum. However, the films were sufficiently thick that the sublimation slight deviations of the molecular chains from parallel to the substrate or to slight misalignment of the spectrometer. did not affect the XPS data. The E1(1), E1(2) and E1(3) bands of PFI shift slightly on going from PFI/gold/glass to PFI/PTFE/gold/glass, i.e. from Results and Discussion 1254, 1219 and 1150 cm-1 to 1264, 1230 and 1166 cm-1, respectively.This may represent a small change in the structure RAIRS of the molecule in going from ‘naturally’ grown material on a The IR spectrum of perfluoroicosane in the range plain substrate to epitaxially grown material on a PTFE 4000–400 cm-1 has been well documented6,31,32 and is known transfer film substrate.to consist of degenerate fundamental vibrations (E1 species), non-degenerate fundamentals (A2 species) and progression XPS bands arising from interactions with the chain ends. The E1 species are polarised perpendicular to the molecular chain XPS survey spectra of PFI/silicon and PFI/PTFE/silicon were identical, showing only C and F (see Fig. 2) Fig. 3 shows C 1s axis, the A2 species parallel, and the progression bands may be either perpendicular or parallel depending on the partic- spectra of the samples at h=90° and 10°.The C 1s component at ca. 292.5 eV binding energy represents the CF2 carbon ular vibration. The E1 species occur at 1260–1150 and 560–550 cm-1, whereas the A2 species occur at 650–630 and atoms of the PFI chain and that at ca. 294.6 eV represents the CF3 end groups.35 At h=90° XPS samples to a depth of ca. 535–520 cm-1. The progression bands are spread across the entire spectrum. 5 nm, whereas at h=10° it is more surface-sensitive, sampling to ca. 1 nm. The PFI molecular chain is ca. 2.5 nm long.24 For RAIRS spectra of PFI/gold/glass and PFI/PTFE/gold/glass are shown in Fig. 1. The spectrum of PFI/gold/glass shows PFI/silicon the CF35CF2 intensity ratio increases on going from h=90° to 10°, consistent with PFI molecules oriented numerous vibrations, the strongest of which occur at 1373, 1340, 647, 545 and 531 cm-1.The 1373, 1340 and 545 cm-1 perpendicular to the substrate. This effect has been seen in other materials where the terminal CF3 group of a perfluori- vibrations are assigned to the parallel progression bands n2(2), n2(4) and n5(4), and the 647 and 531 cm-1 vibrations to the nated alkyl chain is located at the uppermost surface of the sample.36,37 For PFI/PTFE/silicon the CF35CF2 ratio remains parallel fundamental bands A2(2) and A2(3), respectively.The E1 species appear as relatively weak bands at 1254, 1219 and unchanged as h is reduced, consistent with PFI molecules aligned parallel to the substrate. 1150 cm-1. They would normally dominate the transmission IR spectrum of a randomly oriented PFI sample, being Table 1 shows quantification and curve-fit data for the samples. For PFI/PTFE/silicon the C5F atomic ratio is very close to the theoretical value for C20F42 at both h=90° and Fig. 1 RAIRS spectra of (a) PFI/gold/glass and (b) PFI/PTFE/gold/ Fig. 2 XPS survey spectrum of PFI/silicon at h=90° glass 76 J. Mater. Chem., 1997, 7(1), 75–78Fig. 4 C 1s curve-fit for PFI/silicon at h=10° value for PTFE.38 The spectra shown in Fig. 3 and 4 are also binding energy referenced in this way. For PFI/PTFE/silicon the C 1s curve-fit ratios for CF35CF2(2)5CF2(1) are very close to 10510580 at both h= 90° and 10°, as expected for PFI molecules lying parallel to the substrate.For PFI/silicon the CF35CF2(2)5CF2(1) ratio changes from close to 10510580 at h=90° to 16.4516.4567.1 at h=10°. Calculation of electron attenuation predicts that for PFI chains oriented perpendicular to the substrate the CF3: total CF2 intensity ratio is given by: CF3/CF2= [1+exp(-19a)] exp a 1-exp(-18a) where a=Dz/l sin h, Dz is the displacement between successive CF2 units along the chain axis and l is the C 1s inelastic mean free path.With Dz=0.130 nm24 and l=3.0 nm39 the CF3/CF2 Fig. 3 XPS C 1s spectra at h=90° and 10° of (a) PFI/silicon and (b) intensity ratio is predicted to be 10.5589.5 at h=90° and PFI/PTFE/silicon 22.4577.6 at h=10°. Hence the agreement at h=90o is good. The discrepency at h=10° could be due to several factors, such as roughness and disorder of the sample surface on an 10°, consistent with PFI molecules aligned parallel to the substrate.For PFI/silicon the C5F atomic ratio decreases on atomic scale, the finite solid angle of collection of the spectrometer electron lens and small errors in setting h.40 going from h=90° to 10°, consistent with the perpendicular molecular orientation. However, it is surprising that at h=90° Note that for PFI/silicon the C 1s CF3 binding energy increases by ca. 0.10 eV, relative to CF2(1), on going from h= the C5F ratio shows an excess of carbon over the theoretical value. This may be due to a dependence of the C 1s and F 1s 90° to 10°. The F 1s binding energy also increases by ca. 0.06 eV. This must be due to the different chemical environment sensitivity factors (via the photoelectron inelastic mean free paths) on alignment between the electron take-off direction of the surface CF3 groups compared with the same groups deeper in the sample.Calculation shows that for PFI/silicon and the channels between the molecular chains. Components representing CF3, CF2 adjacent to CF3 at h=90° the surface CF3 group contributes <50% of the total C 1s CF3 intensity, the remainder being due to CF3 [CF2(2)] and the remainder of the CF2 units of the molecular chain [CF2 (1)] were fitted to the C 1s envelopes.The greater groups at the interface between the first and second molecular layers. However, at h=10° the C 1s CF3 intensity is due binding energy of CF2(2) relative to CF2(1) is a secondary chemical shift effect due to the additional b-position fluorine entirely to the surface CF3 group.Surface chemical shifts in polymer systems have been reported previously41 and are due substituent.38 A single component was fitted to the F 1s envelope. The C 1s curve fits were constrained to give equal to reduced intermolecular polarisation relaxation at the surface, resulting in slightly higher binding energies.42 The C 1s curve- areas for the CF3 and CF2(2) components; a typical curve fit is shown in Fig. 4. After curve fitting the spectra were referenced fit binding energy for CF2(2) decreases by ca. 0.17 eV, relative to CF2(1), on going from h=90° to 10°. However, as CF2 (2) to a binding energy of 292.48 eV for CF2(1), a recent literature Table 1 XPS quantification and curve-fit data at h=90° and 10° for PFI/silicon and PFI/PTFE/silicon C 1s curve-fit data EB/eV atom% F 1s curve-fit sample h C (atom%) F (atom%) CF2(1) CF2 (2) CF3 CF2 (1) CF2(2) CF3 EB/eV theory (C20F42) 32.3 67.7 80 10 10 PFI/PTFE/Si 90 32.4 67.6 292.48 293.13 294.56 79.1 10.4 10.4 689.69 10 32.7 67.3 292.48 293.11 294.58 79.7 10.2 10.2 689.70 PFI/Si 90 35.5 64.5 292.48 293.14 294.60 80.8 9.6 9.6 689.71 10 32.8 67.2 292.48 292.97 294.70 67.1 16.4 16.4 689.77 J.Mater. Chem., 1997, 7(1), 75–78 779 P. Dietz, P. K. Hansma, K. J. Ihn, F. Motamedi and P. Smith, lies within the overall CF2 envelope and as the constraint of J. Mater. Sci., 1993, 28, 1372. equal CF3 and CF2(2) areas is less valid at h=10° than at 90°, 10 P.Bodo and M. Schott, Synth.Met., 1994, 67, 647. this observation is not as reliable as for the CF3 group. 11 N. W. Hayes, G. Beamson, D. T. Clark, D. T. Clarke and For PFI/PTFE/silicon no change is observed in the C 1s D. S-L. Law, Polym. Commun., 1996, 37, 523. CF3 or F 1s binding energies, relative to C 1s CF2(1), as h is 12 F. Motamedi, K. J. Ihn, D. Fenwick, J. C.Wittmann and P. Smith, J. Polym. Sci., Part B Polym. Phys., 1994, 32, 453. reduced, and the values are very similar to those for PFI/ 13 S. Meyer, P. Smith and J. C. Wittmann, J. Appl. Phys., 1995, silicon at h=90°. This is not surprising given the parallel 77, 5655. orientation of the PFI molecules at the surface. In this arrange- 14 K. Pichler, R. H. Friend, P. L. Burns and A. B.Holmes, Synth. ment the uppermost PFI molecules may show a surface Met., 1993, 55, 454. chemical shift relative to those deeper in the sample, but 15 M. Fahlman, J. Rasmusson, K. Kaeriyama, D. T. Clark, ‘internal’ surface chemical shifts would not be expected. G. Beamson and W. R. Salaneck, Synth.Met., 1994, 66, 123. 16 P. M. McNellis, C. Mathis, B. Francois, S. Meyer, J. C. Wittmann, C.Godon and S. Lefrant, Synth.Met., 1994, 66, 185. Conclusion 17 P. Damman, M. Dosiere, P. Smith and J. C. Wittmann, J. Am. Chem. Soc., 1995, 117, 1117. Both XPS and RAIRS clearly demonstrate a difference in 18 D. Fenwick, K. Pakbaz and P. Smith, J.Mater. Sci., 1996, 31, 915. molecular orientation for PFI vapour deposited onto plain 19 C. Y. Yang, Y. Yang and S. Hotta, Mol.Cryst. L iq. Cryst., 1995, substrates (silicon and gold) and onto PTFE tribological 270, 113. 20 C. Y. Yang, Y. Yang and S. Hotta, Synth. Met., 1995, 69, 303. transfer films. In PFI/silicon and PFI/gold/glass the PFI 21 Y. Ueda, T. Kuriyama, T. Hari and M. Ashida, J. Electron molecular chains orient perpendicular to the substrate, whereas Microsc., 1994, 43, 99. in PFI/PTFE/silicon and PFI/PTFE/gold/glass they take up 22 J.C. Wittman and B. Lotz, Prog. Polym. Sci., 1990, 15, 909. a parallel orientation. Because the structural parameters of 23 K. Nagayama, M. Sei, R. Mitsumoto, E. Ito, T. Araki, H. Ishii, bulk PFI and PTFE are very closely matched, this is probably Y. Ouchi, K. Seki and K. Kondo, J. Electron Spectrosc. Relat. a case of classical epitaxy.However, some authors9 have Phenom., 1996, 78, 375. 24 H. Schwickert, G. Strobl and M. Kimmig, J. Chem. Phys., 1991, suggested that the structure of PTFE chains at the surface of 94, 2800. a transfer film may be different from those in the bulk, in 25 T. Ohta, K. Seki, T. Yokoyama, I. Morisada and K. Edamatsu, which case the graphoepitaxy effect of the aligned PTFE Phys. Scr., 1990, 41, 150.ribbons would be more important. We note a small frequency 26 J. J.Weeks, E. S. Clark and R. K. Eby, Polymer, 1981, 22, 1480. shift for the E1 vibrations of PFI on going from PFI/gold/glass 27 B. L. Farmer and R. K. Eby, Polymer, 1985, 26, 1944. to PFI/PTFE/gold/glass, which may represent a small differ- 28 G. Beamson, D. Briggs, S. F. Davies, I. W. Fletcher, D. T. Clark, J.Howard, U. Gelius, B. Wannberg and P. Balzer, Surf. Interface ence in shape of the molecule in the two situations. For PFI/ Anal., 1990, 15, 541. silicon, XPS reveals a surface chemical shift for the uppermost 29 U. Gelius, B. Wannberg, P. Baltzer, H. Fellner-Felldeg, CF3 groups of the sample. G. Carlsson, C-G. Johansson, J. Larsson, P. Munger and G. Vegerfors, J. Electron Spectrosc.Relat. Phenom., 1990, 52, 747. The Engineering and Physical Sciences Research Council 30 ESCA300 instrument manual, Scienta Instruments AB, Uppsala, 1988. (EPSRC) is thanked for financial support of the RUSTI 31 M. Kobayashi and T. Adachi, J. Phys. Chem., 1995, 99, 4609. core program. 32 S. L. Hsu, N. Reynolds, S. P. Bohan, H. L. Strauss and R. G. Snyder, Macromolecules, 1990, 23, 4565. 33 R. G. Greenler, J. Chem. Phys., 1996, 44, 310. References 34 H. A. Pearce and N. Sheppard, Surf. Sci., 1976, 59, 205. 1 K. R. Makinson and D. Tabor, Proc. R. Soc. L ondon, A, 1964, 35 A. Dilks, in Electron Spectroscopy : T heory, T echniques and 281, 49. Applications, ed. A. D. Baker and C. R. Brundle, Academic Press, London, 1981, p. 277. 2 C. M. Pooley and D. Tabor, Proc. R. Soc. L ondon, A, 1972, 329, 36 G. Beamson and D. Briggs, unpublished work. 251. 37 S. D. Evans, T. M. Flynn, A. Ulman and G. Beamson, Surf. 3 J. C. Wittmann and P. Smith, Nature (L ondon), 1991, 352, 414. Interface Anal., 1996, 24, 187. 4 G. Beamson, D. T. Clark, D. E. Deegan, N. W.Hayes, D. S-L. Law, 38 G. Beamson and D. Briggs, High Resolution XPS of Organic J. R. Rasmusson and W. R. Salaneck, Surf. Interface Anal., 1996, Polymers. T he Scienta ESCA300 Database, John Wiley, 24, 204. Chichester, 1992. 5 M. Schott, Synth.Met., 1994, 67, 55. 39 H. Zhuang, K. Gribbon Marra, T. Ho, T. M. Chapman and 6 M. Kobayashi, M. Sakashita,T. Adachi and M. Kobayashi, J. A. Gardella, Macromolecules, 1996, 29, 1660. Macromolecules, 1995, 28, 316. 40 C. S. Fadley, Prog. Solid State Chem., 1976, 11, 265. 7 Ch.Ziegler, Th.Schedel-Niedrig, G. Beamson, D. T. Clark, W. R. 41 G. Beamson and D. Briggs, Mol. Phys., 1992, 76, 919. Salaneck, H. Sotobayashi and A. M. Bradshaw, L angmuir, 1994, 42 C. B. Duke, Surf. Sci., 1978, 70, 674. 10, 4399. 8 H. Hansma, F. Motamedi, P. Smith, P. Hansma and J. C. Wittmann, Polymer, 1992, 33, 647. Paper 6/05114F; Received 23rd July, 1996 78 J. Mater. Chem., 1997, 7(1), 75–78

ISSN:0959-9428    

DOI:10.1039/a605114f

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Characterisation, conductivity and mechanical properties of the oxygen-ion conductor La0.9Sr0.1Ga0.8Mg0.2O3-x John Drennan,a† Viktor Zelizko,a David Hay,a Fabio T. Ciacchi,a S. Rajendranb and Sukhvinder P. S. Badwal*b aCSIRO, Division ofMaterials Science and T echnology, Private Bag 33, RosebankMDC, Clayton, V ictoria 3169, Australia bCeramic Fuel Cells L imited, 710 Blackburn Road, Clayton, V ictoria 3168, Australia The new oxygen-ion conductor La0.9Sr0.1Ga0.8Mg0.2O3-x has been prepared by conventional solid-state reaction at high temperatures and characterised by X-ray diffraction, scanning and transmission electron microscopy, and conductivity (four-probe dc and impedance) measurements.The room-temperature structure is orthorhombic, space group Pnma (no. 62), with a= 5.5391(7) A° , b=7.8236(12) A° , c=5.5224(7) A° .The material undergoes a phase transition at 445 K to a rhombohedral structure. Mechanical property measurements at room temperature and at 1173 K give average strength measurements of 162±14 MPa and 55±11 MPa respectively. Conductivity and ionic transport number measurements confirm predominantly ionic conduction. The contribution from the grain boundary conductivity is extremely small at temperatures below 673 K.At 1073 K, an ionic conductivity value of 0.12 S cm-1 was recorded in air. It has been reported recently1–4 that the LaGaO3 perovskite temperature. The required amounts of lanthanum oxide substituted at the A and B sites shows good oxygen-ion (>99%, calcined at 1000 °C for 2 h before use), gallium oxide conducting properties at elevated temperature.The most pro- (99.99%), magnesium oxycarbonate (>99%) and strontium mising candidate materials have been shown to be those carbonate (>99%) were mixed and milled together in isopropyl substituted at the A site with Sr and at the B site with Mg. alcohol for 24 h followed by calcination at 1423 K for 4 h. The The highest conductivity has been reported for the composition calcined powder was grey and X-ray diffraction pattern showed La0.8Sr0.2Ga1-yMgyO3-x (y=0.10–0.15).4 The electrolyte is it to be single phase.The oxide powder was milled again in commonly referred to as LSGM. The material is reported to isopropyl alcohol for 24 h, dried and pressed into bar shapes be stable in both reducing and oxidising atmospheres up to for four-probe dc and impedance measurements, or disc shapes 1223 K and shows a reported ionic conductivity of for mechanical strength measurements, and sintered at > 0.1 S cm-1 at 1073 Kand an ionic transport number of close 1723 K for 15 h (heating and cooling rates of 300°C h-1).The to unity. Other materials with high ionic conductivity, such as sintered discs were darkish grey but had a density doped ceria and bismuth oxide, are unstable in reducing (6.58 g cm-3)>98.5% of the theoretical.environments and develop substantial electronic conductivity. The sintered and polished specimens were examined with a The thermalexpansion coefficient and the oxygen ionic conduc- scanning electron microscope. Detailed characterisation of tivity domain (temperature and oxygen partial pressure range) sintered samples was undertaken using a combination of are close to those of stabilised zirconias.The perovskite-based analytical electron microscopy (ATEM) and X-ray diffraction materials have often been discussed5,6 as possible ionic conduc- techniques. ATEM was carried out on both crushed and tors with the tantalising prospect of being able to engineer ion beam thinned specimens using a Philips CM30 series substitution of aliovalent cations onto both the A and B sites electron microscope.Energy dispersive X-ray spectra were (ABO3) with a view to introducing a variety of vacancy recorded using an EDAX 9900 system and all micrographs, schemes and enhanced oxygen-ion conductivity.Until the diffraction patterns and spectra were recorded with the micro- recent report by Ishihara et al.1 on LSGM, the results have scope operating at an accelerating voltage of 300 keV. In the been disappointing. With this discovery the renewed interest case of the ion beam thinned specimens it was found necessary in this class of materials may prove to be a fruitful area to coat the sample with a thin layer of carbon to avoid of research.charging problems. A Siemens D500 diffractometer was used Apart from ionic conductivity measurements and some to collect X-ray diffraction patterns using graphite-monochro- information on the thermal expansion behaviour (thermal mated Cu-Ka radiation. XRD data were also collected at expansion coefficient a=10-5 K-1 3) only limited physical 483 K using a locally constructed temperature stage for the property data are available for LSGM.Moreover, there is D500 diffractometer. Refinement of structural parameters was some confusion over the crystallographic characterisation of carried out by Rietveld methods using the program the material. In this paper we attempt to address these WYRIET 3.7 deficiencies by reporting mechanical property data, conduc- Differential thermal analysis (DTA) measurements were tivity measurements (both four-probe dc and impedance) and made with a Stanton Redcroft Thermochemical Analyser TMA crystallographic characterisation which was obtained using a series 793.The heating rate used was 10°C min-1. combination of analytical electron microscopy and X-ray Four-probe dc conductivity measurements were performed diffraction techniques. in air as a function of temperature (673–1273 K) at 10–25 K intervals during both heating and cooling cycles and as a Experimental function of time at 1123 and 1273 K on different specimens.The specimens were ca. 20–22 mm long with linear conduction The powder of composition La0.9Sr0.1Ga0.8Mg0.2O2.85 was areas of ca. 0.21 cm2. The details of the experimental set-up prepared by conventional solid-state reaction at a high have been described in a previous publication.8 For impedance measurements, specimens were cut from bars and had dimen- † Present Address: Centre for Microscopy and Microanalysis, The University of Queensland, Brisbane, QLD 4072, Australia. sions of a=5.3 mm, b=7.5 mm and thickness t=3.9 mm.J. Mater. Chem., 1997, 7(1), 79–83 79Table 1 Room-temperature strength of LSGM samplesa sample tb/mm wc/mm load/N strength/MPa 1 1.484 16.32 279 138 2 1.510 16.31 339 162 3 1.495 16.31 335 164 4 1.263 15.60 265 183 5 1.248 15.61 233 165 6 1.257 15.64 228 159 aAverage strength=162±14 MPa. bt=thickness. cw=diameter. Table 2 High-temperature strength of LSGM samplesa sample t/mm w/mm load/N strength/MPa 1 1.478 15.13 137 67 2 1.513 14.98 125 58 3 1.496 14.97 126 60 Fig. 1 Optical micrograph of an LSGM sample 4 1.548 14.98 112 50 5 1.540 15.06 143 64 6 1.471 15.02 72 35 7 1.484 15.00 104 50 aAverage strength=55±11 MPa at 1173 K. Platinum paste electrodes sandwiched between platinum mesh were used for current collection. Impedance measurements were performed over the temperature range 573–723 K and the frequency range 5 Hz–5 MHz (with a Hewlett Packard 4192A impedance analyser) in air at 25 K intervals. Transport number was measured by constructing a galvanic cell with platinum paste electrodes and exposing one side of the cell to air and the other side to controlled H2–H2O gas mixtures at different temperatures.Flexural strength measurements were made at room temperature and at 1173 K. Disc-shaped sintered specimens (dimensions are presented in Tables 1 and 2) were used. Prior Fig. 2 Rietveld pattern refinement of LSGM as the orthorhombic to testing, both surfaces of the discs were ground using a (Pnma) structure. The observed XRD trace is shown, together with diamond wheel and one of the surfaces was polished to a 1 mm the theoretical trace using refined parameters, and a difference trace finish using diamond paste.This is to ensure that the tensile is shown below. Expected peak positions from the model structure surface is free from any pre-existing flaws. The loading con- are shown as bars. figuration for the test was similar to that described in the ASTM F394-78 9 publication.In the room-temperature case, b=7.8236(12) A° and c=5.5224(7) A° . Refinement converged to the three support balls were equally spaced on a circle of Bragg R=5.1%. Refined atomic coordinates are given in 12.098 mm diameter with the loading ram of 3.946 mm diam- Table 3. eter. At 1173 K, support pins with hemispherical ends and The room-temperature symmetry was further confirmed by having diameters of 3 mm were used.The pins were spaced analytical transmission electron microscopy (ATEM). equally on a circle of diameter 11.493 mm and the loading ram Convergent-beam electron diffraction patterns recorded from was 3.95 mm in diameter. individual grains of the LSGM showed extra reflections which are not consistent with a primitive cubic cell.Fig. 3 shows an Results and Discussion indexed orthorhombic pattern ([100]) based on the unit cell described above. To further confirm that the material was in The material was easily fabricated into dense bodies. The fact homogeneous, energy dispersive X-ray spectra were microstructure of a polished and etched specimen is shown in recorded from a number of grains and an example is shown Fig. 1. In general, very few but uniformly distributed pores in Fig. 4. Clearly a solid-state reaction has taken place and were observed in the microstructure. Clearly a bimodal grain La, Sr, Ga and Mg were detected; composition variations size distribution was observed in the microstructure. Small between grains were insignificant. Quantitative analysis grains were in the 3–5 mm and large grains in the 15–25 mm determined using the peak intensities gave values close to the size range.The room-temperature X-ray diffraction patterns nominal composition but difficulties with overlapping peaks of the prepared samples used in this study showed a basic perovskite-related structure. However, on closer examination Table 3 Refined fractional atomic coordinates with estimated standard it is clear that some reflections showed shoulders and splitting deviations in parentheses representing a reduction in symmetry.Goodenough and Feng3 reported that the La0.9Sr0.1Ga0.8Mg0.2O2.85 phase was cubic atom 103x/a 103y/b 103z/c perovskite, whilst Ishihara et al.1 have reported an orthorhombic structure for this material. In this study the powder patterns La 9998(14) 2500 -18(36) Sr 9998(14) 2500 -18(36) were indexed using the orthorhombic cell and Rietveld refine- Ga 0 0 5000 ments carried out on this basis.The results of Rietveld full pat- Mg 0 0 5000 tern refinement are shown in Fig. 2. The atomic model which O1 4380(119) 2500 584(129) gave the best fit to the data was orthorhombic, space group O2 2409(127) 504(32) 7522(174) Pnma (no. 62), with refined cell dimensions a=5.5391(7) A° , 80 J. Mater. Chem., 1997, 7(1), 79–83Fig. 3 The [100] diffraction pattern of an LSGM specimen Fig. 5 Bright-field micrograph recorded from part of a grain of LSGM showing the strain and massive twinning which is a feature of this material. Inset in the micrograph are microdiffraction patterns recorded from the regions directly underneath the patterns.Clearly a twin relationship exists. in the high-temperature trace, part of which is shown in Fig. 6, Fig. 4 A portion of the energy dispersive X-ray spectrum showing the where peaks appear as a multiplet. This is characteristic of a presence of La, Sr, Ga and Mg in the grain of an LSGM specimen change to rhombohedral (R3c) symmetry, where these peaks appear as a triplet, as observed for the phase transition at and absorption effects introduced some uncertainties into these 418 K in LaGaO3.12 It should be noted that Petric et al.4 calculations.observed extra reflections for LSGM in the electron diffraction A further observation of the microscopy revealed that the pattern. They described these reflections as superstructure grains of LSGM at room temperature contained contrast reflections resulting from the formation of microdomains.In features consistent with twinning. Microdiffraction patterns the work reported in this study, although such reflections were confirmed that these domains are associated with twinning observed, these could be fully indexed using the orthorhombic and this is illustrated in Fig. 5. This phenomenon has been cell structure. observed in electron microscopy studies of the related perov- The consequence that this phase transition has on the skite LaGaO3.10 It should be noted that extensive crystallo- suitability of LSGM as an electrolyte material is unknown but graphic studies have been made of LaGaO3 and the material the amount of strain evident in the micrographs (Fig. 5) has an orthorhombic structure at room temperature with lattice parameters a=5.5232(5) A° , b=7.776(2) A° , c= 5.4925(7) A° and the space group Pnma (no. 62) satisfies the diffraction evidence. In addition, LaGaO3 is known11 to undergo a phase transition at 418 K in which the system transforms from orthorhombic symmetry to rhombohedral symmetry and the consequent readjustment gives rise to elimination of twins.DTA results obtained from the LSGM specimens examined in this work revealed that an endothermic peak was detected at 445 K, slightly higher than that reported for pure LaGaO3 (418 K), but certainly indicating a phase transition. It is a logical assumption, therefore, that the twin boundaries observed in LSGM are a result of a possible phase transition similar to that observed in LaGaO3.To further confirm the DTA results, high-temperature X-ray diffraction was employed to examine for any evidence of this phase transition. Indeed, evidence of a phase change is seen in the XRD data at 483 K. The group of peaks at 2h#32° is seen as a doublet in the orthorhombic (Pnma) room-temperature struc- Fig. 6 Expansion of the peak structure near 32° showing the roomtemperature (a) and 483 K (b) XRD traces ture.However, further splitting of this group of peaks is seen J. Mater. Chem., 1997, 7(1), 79–83 81Table 5 Activation energy data for LSGM, 9 mol% Sc2O3–ZrO2 and suggests that thermal cycling through the transition point may 9 mol% Y2O3-ZrO2 be deleterious to mechanical properties. The results of strength measurements on LSGM samples at (a) From four-probe dc conductivity data: room temperature and at 1173 K are presented in Tables 1 Ea/kJ mol-1 and 2.The average strength determined was 162±14 MPa at room temperature. Specimens dimensions, fracture loads and 400–450°C 850–1000°C determined strength for 1173 K strength tests are given in Table 2. At this temperature the average strength determined La0.9Sr0.1Ga0.8Mg0.2O3-x 109±3 64±2 9 mol% Sc2O3–ZrO2 130±2 72±3 for LSGM was 55±11 MPa.The flexural strength of these 9 mol% Y2O3–ZrO2 107±2 81±3 materials is significantly lower in comparison to the zirconiabased electrolytes presently in use. It is worth mentioning here (b) From impedance dataa (300–450 °C): that the mechanical strength of fully stabilised zirconia is in the vicinity of 300 MPa at room temperature and 120 MPa at Ea/kJ mol-1 1273 K.13 For yttria–tetragonal zirconia ceramics, flexural strengths of around 1000 MPa and 350–400 MPa have been La0.9Sr0.1Ga0.8Mg0.2O3-y Rv Rgb Rtotal measured at room temperature and 1273 K respectively.13 before annealing 107±1 97±2 105±1 Fig. 7 shows Arrhenius plots for the conductivity data of after 1000 °C anneal, 5000 min 105±1 96±2 104±1 LSGM, 9 mol% Sc2O3–ZrO2 and 9 mol% Y2O3–ZrO2.Clearly after 850 °C anneal, 5000 min 106±1 98±2 105±1 the conductivity of LSGM is higher than that of 9 mol% Y2O3–ZrO2 , but over the temperature range 1073–1273 K the aRv=volume (or lattice) resistivity, Rgb=grain boundary resistivity, ionic conductivity values for both materials are comparable.Rtotal=Rv+Rgb . Table 4 gives ionic conductivity values for several high-conductivity electrolyte materials. At 1073 K, the ionic conductivity of Sm2O3-doped ceria is slightly lower than that of LSGM. The activation energy values for conduction are given in Table 5. At high temperatures (1123–1273 K), the activation energy for LSGM was lower than for both 9 mol% Sc2O3–ZrO2 and 9 mol% Y2O3–ZrO2 materials.Conductivity data as a function of time at 1123 and 1273 K for LSGM are displayed in Fig. 8. Only a slight decrease in the conductivity was observed with time at 1273 K. At 1123 K, the effect of time on conductivity was insignificant. This absence of any ageing process clearly indicates that the phase assemblage or the structure of the material is stable.This was further confirmed by transmission electron microscopy of the annealed (1273 K, 5000 min) specimen. Impedance measurements on LSGM specimens over the 573–723 K range show a very low contribution from the grain Fig. 8 Dc conductivity for LSGM as a function of time at 1123 (a) and 1273 K (b) boundary resistivity (Fig. 9). A slight decrease in the volume resistivity as a result of annealing at 1273 K could not be explained by the conductivity or activation energy data or by the detailed microstructural analysis.In order to determine ionic transport number, small fuel cells were constructed with Pt air and fuel electrodes and LSGM as the electrolyte. Fig. 10 shows the results of open circuit voltage measurements at several temperatures for air vs.H2–H2O mixture. In all cases the measured voltage was close to the theoretical (ionic transport number close to unity) indicating that the material is mainly an ionic conductor. Fig. 7 Arrhenius plots (four-probe dc data) for LSGM (#), 9 mol% Sc2O3–ZrO2 (%) and 9 mol% Y2O3–ZrO2 (') Table 4 Conductivity data for some high-conductivity oxygen-ion conducting electrolytes s/S cm-1 s/S cm-1 system (1073 K) (1273 K) ref.LSGM 0.121 0.316 (ZrO2)0.91 (Y2O3)0.09 0.046 0.166 14 (ZrO2)0.91 (Sc2O3 )0.09 0.109 0.306 14 (CeO2)0.82 (Gd2O3)0.18 — 0.235 15 (CeO2)0.80 (SmO1.5)0.20 0.096 0.25 16 (Bi2O3)0.80 (Er2O3)0.20 0.37a — 17 (Bi2O3)0.80 (Nb2O5)0.20 0.19a — 18 Fig. 9 Impedance plots at 623 K in air for LSGM before (a) and after (b) annealing at 1273 K for 5000 min aAt 973 K. 82 J. Mater. Chem., 1997, 7(1), 79–83specimen preparation and to Dr. K. Foger for reviewing this manuscript. References 1 T. Ishhihara, H. Matsuda and Y. Takita, J. Am. Chem. Soc., 1994, 116, 3801. 2 T. Ishihara, H. Matsuda and Y. Takita, Solid State Ionics, 1995, 79, 147. 3 M. Feng and J. B. Goodenough, Eur. J. Solid State Inorg. Chem., 1994, 31, 663. 4 A.Petric, P. Huang and A. Skowron, Proc. 2nd Eur. SOFC Forum, ed. B. Thorstensen, Druckerei J Kinzel, Go�ttingen, Germany, 1996, Fig. 10 Open-circuit voltage [measured (#), calculated (%)] for Pt, pp. 751–760. air|LSGM|Pt, H2–2%H2O cell as a function of time. The solid line 5 J. B. Goodenough, A. Manthiram and J-F. Kuo, Mater. Chem. represents the calculated cell voltage. Phys., 1993, 35, 221. 6 A. F. Sammells, R. L. Cook, J. H. White, J. J. Osborne and R. C. MacDuff, Solid State Ionics, 1992, 52, 111. Conclusions 7 M. Schneider, Program WYRIET 3, version 3, D-8134 Pocking, West Germany, 1992. At 1273 K, the ionic conductivity of La0.9Sr0.1Ga0.8Mg0.2O3-x 8 S. P. S. Badwal, F. T. Ciacchi and D. V. Ho, J. Appl. Electrochem., is higher by a factor of two compared to ZrO2 doped with 8 1991, 21, 721.mol% Y2O3 ; however, it was similar to that of ZrO2 doped 9 ASTM Designation: F 394–78. with 9 mol% Sc2O3. The major drawback of these materials 10 M. Sundberg, P-E. Werner, M. Westdahl and K. Mazur, Mater. as potential electrolytes for use in solid oxide fuel cells or other Sci. Forum, 1994, 166–169, 795. 11 H. M. O’Bryan, P. K. Gallagher, G. W.Berkstresser and similar applications is the high cost of gallium compounds, C. D. Brandle, J.Mater. Res., 1990, 5, 183. related to its scarcity, and the low mechanical strength of 12 Y. Wang, X. Liu, G.-D. Yao, R. C. Lieberman and M. Dudley, LSGM, especially at the fuel cell operating temperatures. It is Mater. Sci. Eng. A, 1991, 132, 13. unlikely that such materials can be used in electrolyte-sup- 13 V. Zelizko, unpublished work. ported designs for solid oxide fuel cells or in applications 14 S. P. S. Badwal, F. T. Ciacchi, J. Drennan and S. Rajendran, to where they must play some role in the structural design of a be published. 15 T. Kudo and H. Obayashi, J. Electrochem. Soc., 1976, 123, 415. device. A more likely application would be in their use as thin 16 K. Eguchi, T. Setoguchi, T. Inoue and H. Arai, Solid State Ionics, coatings on an electrode substrate. Alternatively these materials 1992, 52, 165. may find use in sensors where the current-carrying capacity of 17 M. J. Verkerk, K. Keizer and A. J. Burggraaf, J. Appl. Electrochem., the device is not important and the materials are used in small 1980, 10, 81. quantities. 18 T. Takahashi, H. Iwahara and T. Esaka, J. Electrochem. Soc., 1977, 124, 1563. The authors are thankful to Dr. S. P. Jiang for determining the ionic transport number, Miss Kristine Giampietro for Paper 6/04563D; Received 1st July, 1996 J. Mater. Chem., 1997, 7(1), 79–83

ISSN:0959-9428    

DOI:10.1039/a604563d

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Thermoelectric properties of Al-doped ZnO as a promising oxide material for hightemperature thermoelectric conversion Toshiki Tsubota, Michitaka Ohtaki, Koichi Eguchi and Hiromichi Arai* Department of Materials Science and T echnology, Graduate School of Engineering Science, Kyushu University, 6-1 Kasugakoen, Kasuga, Fukuoka 816, Japan The thermoelectric properties of a mixed oxide (Zn1-xAlx)O (x=0, 0.005, 0.01, 0.02, 0.05) are investigated in terms of materials for high-temperature thermoelectric conversion.The electrical conductivity, s, of the oxide increases on Al-doping by more than three orders of magnitude up to ca. 103 S cm-1 at room temperature, showing metallic behaviour. The Seebeck coefficient, S, of (Zn1-xAlx)O (x>0) shows a general trend in which the absolute value increases gradually from ca.-100 mVK-1 at room temperature to ca.-200 mVK-1 at 1000°C.As a consequence, the power factor, S2s, reaches ca. 15×10-4 Wm-1 K-2, the largest value of all reported oxide materials. The thermal conductivity, k, of the oxide decreases with increasing temperature, owing to a decrease in the lattice thermal conductivity which is revealed to be dominant in the overall k.In spite of the considerably large values of k, the figure of merit, Z=S2s/k, reaches 0.24×10-3 K-1 for (Zn0.98Al0.02)O at 1000 °C. The extremely large power factor of (Zn1-xAlx)O compared to other metal oxides can be attributed to the high carrier mobility revealed by the Hall measurements, presumably resulting from a relatively covalent character of the ZnMO bond owing to a fairly small difference of the electronegativities of Zn and O.The dimensionless figure of merit, ZT , of 0.30 attained by (Zn0.98Al0.02)O at 1000 °C demonstrates the potential usefulness of the oxide. Thermoelectric power generation converts thermal energy V-1 s-1 (which is the same as the drift mobility of copper) directly to electrical energy via the Seebeck effect induced by reported for ReO3,13 which shows completely metallic behava temperature difference in solid materials.The energy balance iour and the highest electrical conductivity of all the known of thermoelectric conversion gives the figure of merit, Z= oxides in the normal state. With respect to these advantages S2s/k, where s is the electrical conductivity, S is the Seebeck of oxide materials, we have begun to investigate the thermocoefficient and k is the thermal conductivity of the material.electric properties of oxides, and have reported already that Practical applications generally require Z1×10-3 K-1. On In2O3-based mixed oxides14 and CaMnO3-based perovskite- the other hand, the Carnot efficiency of thermoelectric conver- type oxides15 have potential for new thermoelectric materials sion improves with an increasing temperature difference over at high temperatures.which the thermoelectric device operates. A product of Z and Zinc oxide is well known as a conductive oxide with a wide the absolute temperature, the dimensionless figure of merit, forbidden bandgap of 3.5 eV. The electrical properties of ZnO ZT , is thereby employed as the most inclusive criterion for have attracted much interest in many fields related to chemistry evaluating thermoelectric materials.Efficient materials have and physics, and are applied to electronic devices such as gas ZT values of nearly unity or greater. ZT=1 corresponds to a sensors16 and varistors.17 Although Zn is sometimes regarded conversion efficiency>10%; no materials have achieved ZT= as a member of the 3d transition metals, ZnO is actually not 2 to date.a transition metal oxide because the 3d orbital of Zn2+ is In the last three decades, Si–Ge alloys,1 several metal filled. In general, the valence band of a metal oxide semicon- chalcogenides,2,3 transition-metal disilicides4–6 and some ductor consists mainly of the 2p orbital of the O2- anions, boron compounds7,8 have been developed as materials for while an empty or half-filled orbital with the lowest energy in high-temperature thermoelectric power generation.However, the metal cations dominates the characteristics of the conduc- practical utilization has been limited because many of these tion band. The conduction electrons in the 3d transition-metal materials require costly surface protection to prevent oxi- oxides are therefore largely influenced by a localized character dation or vaporization, and some others have inherent of the 3d orbitals of the metal cations.However, the conduction temperature limits owing to phase transitions at high tempera- band of ZnO is constituted mainly of the lowest unoccupied tures.With respect to high-temperature operation in air, the 4s and 4p orbitals of Zn2+, and thus the conduction electrons advantages of metal oxides in their common oxidation state would be more mobile. The fairly large electronegativity of are apparent because of their excellent stability to heat. Zn, and its relatively small ionicityas well as astrong preference Moreover, a number of metal oxides have high electrical for the sp3 hybridization of the ZnMO bond also appear to conductivities, and also the Seebeck effect in oxide materials be very promising, since they imply that the ZnMO bond has has been studied widely.Nonetheless, reports on the appli- a rather covalent character, which would lead to a high carrier cation of oxides to thermoelectric materials are surprisingly mobility in the oxide.In the present study, we have investigated scarce. While high-temperature superconducting cuprates the thermoelectric properties of ZnO-based mixed oxides in were once proposed for thermoelectric refrigerators,9 they terms of thermoelectric power generation at high temperatures, were revealed to have inherently poor performances10,11 revealing the highly promising thermoelectric figure of merit mainly due to their extremely low carrier mobilities (as low of Al-doped ZnO for the first time as an oxide material.18 Here as 0.1 cm2 V-1 s-1).Metal oxides are generally considered to we discuss the thermoelectric properties in terms of the electri- be low-mobility materials, mainly because of their highly cal and thermal transport properties, analytical and spectro- ionic characters.In fact, however, some oxides show considerscopic information, and we consider the microstructures of ably high carrier mobilities, e.g. a reported value of the Hall the oxide as a potential material for high-temperature mobility of SnO2 single crystals is 150–240 cm2 V-1 s-1 at 300 K,12 which is much larger than the value of 30 cm2 thermoelectrics. J.Mater. Chem., 1997, 7(1), 85–90 85Experimental Preparation and characterization of samples The mixed oxides (Zn1-xAlx)O (x=0, 0.005, 0.01, 0.02, 0.05) were prepared from ZnO and Al2O3 powders of guaranteed grade. The powders were mixed and pulverized in a Nylonlined ball mill for 24 h. The powder mixture was pressed into a pellet and sintered at 1400 °C for 10 h in air. The heating and cooling rates were 200 °C h-1.The crystal phases in the samples thus obtained were examined by a powder X-ray diffraction (XRD) study on a Rigaku RINT-1400 diffractometer using Cu-Ka radiation. Scanning electron microscopy was carried out on aJEOL JSM-T330A instrument equipped with an EDX spectrometer. Auger microprobe analysis was performed on a field emission type JEOL JAMP-7800F spectrometer operated at a probe energy of 10 keV. 27Al solidstate NMR spectra were collected on a Varian VXR-400S Fig. 1 Arrhenius plots for the electrical conductivities of (Zn1-xAlx)O spectrometer at an oscillation frequency of 104.21 MHz using for x=0 (#), 0.005 (&), 0.01 (%), 0.02 ($), 0.05 (') the MAS technique at room temperature.An AlCl3 aqueous solution was employed as an external reference (d -0.1). A single pulse sequencewith a delay time (D1 ) of 2.0 s was applied. density >99%, density difference between the samples cannot explain the variation in s. The incorporation of Al into ZnO should therefore be responsible for the noticeable changes of Measurement of thermoelectric properties the value and the behaviour of s observed here.The samples for electrical measurements were cut from the The Seebeck coefficients, S, of (Zn1-xAlx)O (x=0, 0.005, sintered pellets as rectangular bars of ca. 15×5×3 mm3, and 0.01, 0.02, 0.05) are shown in Fig. 2. The S values of all the polished with SiC emery papers. The relative densities of all samples are negative within the whole temperature range the samples were measured by Archimedes’ method.The examined, indicating n-type conduction. The Seebeck measurements of the electrical conductivity and the Seebeck coefficient of undoped ZnO has large and negative values of coefficient were carried out simultaneously in air from room ca.-300 to -400 mV K-1 from room temperature to 1000°C, temperature to 1000 °C.The measurement procedures have and showed no particular dependence on temperature. On the been described elsewhere in detail.14 Briefly, the s values were other hand, (Zn1-xAlx)O (x>0) gave smaller but still moderate measured by the dc four-probe technique. The S values were S values, with a general trend in which the absolute value obtained from the least-squares regressions of the thermoelec- increased gradually from ca.-100 mV K-1 at room tempera- tromotive force as a function of the temperature difference ture to ca. -200 mV K-1 at 1000 °C. <5 K applied by a heater at each measurement temperature. The power factor, S2s, which represents the electrical contri- All the measurements were carried out after attaining the bution to the overall thermoelectric performance, is calculated steady-state temperature at each step.The carrier concen- from the results obtained above and depicted in Fig. 3 as a tration, n, and the Hall mobility, mH, were determined from function of temperature. The S2s value of undoped ZnO the Hall measurements carried out at room temperature by increases with increasing temperature owing to a large increase the van der Pauw method for sliced samples at an applied in s.The maximum S2s value of undoped ZnO is magnetic field of 0.8 T. The thermal conductivity was deter- 3.68×10-4W m-1K-2 at 1000 °C; even this value appears mined from the thermal diffusivity and the specific heat to be surprisingly large compared with those shown by oxide capacity obtained by the laser flash measurement on an materials.All the Al-doped samples, nonetheless, attain exceed- ULVAC TC-7000 instrument from room temperature to ingly large values of S2s, i.e. 8–15×10-4W m-1 K-2 over a 1000 °C for sample disks 10 mm in diameter and 1–2 mm in wide temperature range from room temperature to 1000°C. thickness. The k data were calibrated with a standard sample These values are the largest ever reported on oxide materials, of a sapphire single crystal.and also surpass those shown by b-FeSi2 and b-SiC,19 which have been proposed as non-oxide candidates for high-temperature thermoelectric materials. The unexpectedly large power Results and Discussion factor of (Zn1-xAlx)O denotes the superiority of the electrical Electrical transport properties The temperature dependence of the electrical conductivity, s, of (Zn1-xAlx)O (x=0, 0.005, 0.01, 0.02, 0.05) is shown in Fig. 1. The s value of undoped ZnO increased sharply with increasing temperature to 700 °C, and appeared rather constant between 700 and 1000 °C. The Arrhenius plot of s for undoped ZnO indicates that the sample has an impurity region up to 700 °C, showing no intrinsic conduction up to 1000 °C.With the addition of a small amount of Al2O3 (x=0.005), the behaviour of s changed from semiconducting to metallic, and the values became higher than that for undoped ZnO by more than three orders of magnitude at room temperature. The s values of (Zn1-xAlx)O showed a monotonic increase with increasing amounts of added Al2O3 up to x=0.02, and then decreased slightly on further Al-doping, as seen for x=0.05. Consequently, (Zn0.98Al0.02)O gave the highest s value of all the samples. Because all the sintered samples, including the Fig. 2 The Seebeck coefficients of (Zn1-xAlx)O as a function of temperature for x=0 (#), 0.005 (&), 0.01 (%), 0.02 ($), 0.05 (') undoped one, have a dense microstructure with a relative 86 J. Mater. Chem., 1997, 7(1), 85–90Fig. 3 The power factors of (Zn1-xAlx)O as a function of temperature Fig. 4 Powder X-ray diffraction patterns of (Zn1-xAlx)O for x=0 (a), for x=0 (#), 0.005 (&), 0.01 (%), 0.02 ($), 0.05 (') 0.005 (b), 0.01 (c), 0.02 (d), 0.05 (e). #, ZnO; ', ZnAl2O4; %, a-Al2O3. even with a fraction smaller than the nominal compositions, a transport properties of the present oxide in contrast to a solid solution of Al2O3 in ZnO occurs.The ZnAl2O4 phase general view of metal oxides as ionic solids. with high resistivity would be, however, responsible for the We also carried out Hall measurements to investigate the decrease in s at x=0.05. carrier concentration and the mobility of each sample. The We have also carried out EDX and Auger microprobe carrier concentrations, n, of (Zn1-xAlx)O (x=0, 0.005, 0.01, analyses for cross-sections of the samples, and detected no 0.02, 0.05) atroom temperature are shown in Table 1.Assuming signals due to elements other than Zn, Al and O (except surface that all Al atoms occupy Zn sites in ZnO and that they act as carbon contamination). This ruled out any unintentional donors to provide one electron per Al, theoretical values of n doping of impurities other than Al as a reason for the increase can also be calculated from the amount of doped Al2O3.in s. Although the low doping levels caused some difficulties Whereas the observed n value for undoped ZnO is of the order in the determination of the precise concentrations of Al in the of 1023 m-3, all of the observed n values for the Al-doped samples, the intensity ratios of Zn/Al in the EDX spectra were samples are of the order of 1025 m-3.The dramatic increase in good agreement with the nominal compositions, within in the observed n values brought about by Al-doping is experimental error. Two-dimensional Auger microprobe map- consistent with the marked increase in s accompanying the ping for the cross-sectioned surface of (Zn0.98Al0.02)O over an semiconductor-to-metal transition seen in Fig. 1. However, the area of 70×70 mm2 with a 256×256 resolution (corresponding observed n value appears to be almost saturated at x#0.005, to a spatial resolution<300 nm) showed random distributions being smaller than the theoretical values even at x=0.005. of Zn and Al with no correlation, confirming a sufficient These results suggest that, at higher doping levels, only a homogeneity of the sample composition.Auger microprobe limited portion of the added Al might be effective as a dopant. spectroscopy (probe diameter 15 nm) detected no difference between the elemental compositions within the grains and at Phase, structural and spectroscopic considerations the grain boundaries of the sample, thus presenting no evidence Phase diagrams reported previously on the ZnO–Al2O3 system of secondary phase segregation at the grain boundary. tell us that there is no solid solution region near the ZnO end Furthermore, we performed 27Al solid-state NMR spec- member.A recent article on the phases in the system troscopy in order to clarify the site and chemical environment Al2O3–ZnO also concluded that Al2O3 does not dissolve in at which the doped Al atoms might be located.Because the ZnO.20 However, as mentioned above the addition of Al2O3 Zn atoms in ZnO are four-coordinated, it is expected that the to ZnO has been regarded as a typical example of n-type coordination number of the doped Al would be four if they doping or valence control of semiconductors.21 Studies on substitute successfully the Zn sites in ZnO.By contrast, in a- phase diagrams are generally carried out with compositional Al2O3, which is the thermodynamically stable phase of alu- changes at intervals of at least a few mol%, which appear to minium oxide at temperatures >1000 °C, the Al atoms are all be too large to examine the occurrence of solid solutions at six-coordinated, Al(6).It is well known that Al(4) and Al(6) very low doping levels. The results of an XRD study for can be distinguished easily by 27Al NMRspectroscopy, because (Zn1-xAlx)O (x=0, 0.005, 0.01, 0.02, 0.05) are shown in Fig. 4. in oxides the chemical shift of the former is in the range The diffraction patterns indicate that the samples at x0.02 d 50–77, while the latter appears at d 0–20.22 Although the Al contain a small amount of a ZnAl2O4 spinel phase which atoms in the ZnAl2O4 spinel are also six-coordinated, the increases with increasing x.NoAl2O3 single phase was detected crystallographic symmetry around the atoms should be differ- for all the compositions examined. A single-phase sample of ent from those in a-Al2O3. Accordingly, as shown in Fig. 5(d) ZnAl2O4 was also prepared, and the spinel was revealed to and (e), 27Al MAS NMR spectra of a-Al2O3 gave a completely have a high electrical resistivity. We therefore concluded that, symmetric single peak at d 12.5, while ZnAl2O4 showed a single peak split into an asymmetric doublet at d 7 and 14, ascribedto second-order nuclear quadropole interaction, which Table 1 The carrier concentrations of (Zn1-xAlx)O at room temperature can be observed well for nuclei with the nuclear spin |I|>1 and a large nuclear quadropole coupling constant.23 Both n/1025 m-3 spectra contain no signals due to Al(4).The Al-doped samples of (Zn1-xAlx)O at x0.02 showed two distinct features in x observed theoretical their NMR spectra, as seen in Fig. 5(b) and (c); a doublet at 0 0.052 — d 7 and 14, the same as for ZnAl2O4, and a broad resonance 0.005 6.5 18 in the region d 30–50 which is assigned to Al(4).However, no 0.01 7.7 37 peak was observed in the lower field region of d60, where 0.02 7.2 74 Al(4) of c-Al2O3 is reported to appear at d 68.1 and is actually 0.05 6.5 193 found at d 67 in our experiment. These results indicate clearly J.Mater. Chem., 1997, 7(1), 85–90 87Fig. 6 The Hall mobility of (Zn1-xAlx)O at room temperature as a function of x addition of Al, the samples at x=0.02 and 0.05 showed similar microstructures in which the grain sizes are comparable to Fig. 5 27Al MAS NMR spectra of (a) x=0.005, (b) x=0.02, (c) x=0.10 that of undoped ZnO and the grain interior appeared to be for (Zn1-xAlx )O and (d) ZnAl2O4, (e) a-Al2O3.Arrowheads indicate the spinning side bands. rather even. The variation in mH may be due to the changes in the microstructure observed here, although the reasons for these phenomena are still unclear. The changes in s as a whole the formation of the spinel phase at least at x0.02 and the may be explained as follows: at a low doping level, incorpor- absence of any Al2O3 single phase in all the samples, which is ation of Al atoms into the Zn sites dramatically increased the consistent with the XRD results.The Al(4)/Al(6) intensity carrier concentration but suppressed the mobility by changing ratio became larger for smaller amounts of added Al2O3. At the microstructure, and then further Al-doping caused an very low doping levels, a weak but very broad ‘bump’ due to improvement in the microstructure, resulting in recovery of the sample probe of the NMR spectrometer overlaps in the the mobility and an additional increase in s until the secondary region d 10–50, as observed for x=0.005 in Fig. 5(a); the spinel phase came to affect the conductivity. These results also sample at x=0.01 also gave virtually the same spectrum as imply that the electrical properties of the oxide might be Fig. 5(a). However, differential spectra, obtained by subtracting improved further by optimizing the microstructure. the background spectrum measured for a ZnO raw powder, revealed that Al(4) with a peak at d ca. 48 is dominant at x= Thermal transport properties 0.005. These findings confirm that, at least at the lower doping The thermal conductivity, k, is another fundamental parameter levels, a substantial proportion of the doped Al really exists as for the evaluation of the thermoelectric performance of mate- Al(4) which should be absent in ZnAl2O4 and a-Al2O3, and rials, and it may cancel out an advantage brought about by a that they have settled in an environment which is different large value of s as is in the case for metals.Actually, the k from those for Al(4) in c-Al2O3 ; this suggests strongly that the values of the present oxides are rather high, ca. 40 Wm-1 K-1 Al atoms occupy the Zn sites in ZnO. Even at this composition, at room temperature, whereas they decrease to ca. 5W however, there seems to be a small peak at d ca. 14, implying m-1 K-1 at 1000 °C with increasing temperature as shown in that the spinel phase coexists even at the lowest doping level Fig. 7. In general, the thermal conductivity of metal oxides in this study. These results are in good agreement with the becomes higher as the atomic mass ratio M/O (M is a metal results of the n measurements stated above, in which the n atom) approaches unity. Also, a smaller unit cell or a simpler values were practically unchanged on addition of increasing lattice structure results in improved thermal conduction due amounts of Al2O3 (x>0.005), also suggesting that the solubility to lattice vibrations or phonons.The lattice structure of ZnO, limit of Al into ZnO is actually x<0.005. However, s of the which forms the wurtzite structure with simple hexagonal Al-doped ZnO increased until x=0.02.This fact seems to symmetry, would hence largely be responsible for the high require an additional explanation. thermal conductivity values. Since thermoelectric materials in The Hall mobility, mH, of (Zn1-xAlx)O at room temperature practical use generally have k values of ca. 1–3W m-1 K-1, was determined by Hall measurements and the results are shown in Fig. 6 as a function of x. The mH value of undoped ZnO, 65 cm2 V-1 s-1, is quite reasonable compared to the values in the literature,21 and appears to be rather large for an oxide ceramic. However, the mH values for the samples at x=0.005 and 0.01 are about half of that for undoped ZnO, and the further addition of Al caused the mH values to recover. These phenomena may be associated with changes in the sample microstructures.We therefore examined the microstructures of (Zn1-xAlx)O (x=0, 0.005, 0.01, 0.02, 0.05) by SEM. The samples, polished to obtain mirror-like surfaces, were heated at 1200 °C for 20 min in air for heat etching. Undoped ZnO has a simple microstructure with large grains without noticeable interior structures. However, at x=0.005 the grains showed a fine-layered interior structure, and at x=0.01 the cross-sectional surface became roughened with a small pointed particle-like structure.We consider that these structures are responsible for the decrease in mH because they can provide Fig. 7 The thermal conductivities of (Zn1-xAlx)O as a function of temperature for x=0 (#), 0.02 ($), 0.05 (') many more scattering centres than undoped ZnO.On further 88 J. Mater. Chem., 1997, 7(1), 85–90the present k values for (Zn1-xAlx)O appear to be too high coordinate geometry requires a ratio >0.414, one can expect six-coordinate Zn2+ in ZnO if the ZnMO bond is sufficiently for thermoelectric applications. ionic. However, in reality ZnO crystallizes in the wurtzite structure in which Zn2+ coordinates with four oxygen atoms.Thermoelectric performance The sp3 hybrid character of Zn2+, in which the cations have In spite of the high k values, the figure of merit, Z=S2s/k, of the d10 closed shell and show the strongest four-coordination the present oxide is revealed to be still prospective, benefiting preference of the first-row transition metals, results in the from the considerably large power factors. The figures of merit strong directionality of the MMO bond required for the of (Zn1-xAlx)O (x=0, 0.02, 0.05) are shown in Fig. 8 as a wurtzite structure. Accordingly, Zn forms a chemical bond function of temperature. The Z values of x=0.02 and 0.05 with a rather covalent character even in its oxide; the ionicity, increased with increasing temperature, and Z attained the which is the proportion of an ionic character in an MMO largest value of 0.24×10-3 K-1 at 1000 °C for x=0.02.This bond, is reported to be 0.62 for ZnO in contrast to 0.84 for value is much larger than that of b-SiC,19 and is as large as MgO in the rocksalt structure.29 The unexpectedly large Z that of the best result reported on b-FeSi24–6 which shows a value of (Zn1-xAlx)O should be attributed to these facts, which maximum at ca. 600 °C. These facts suggest strongly that the would lead to a high carrier mobility. electrical properties of the present oxide would be sufficiently The thermal conductivity can be varied greatly by extrinsic advantageous to overcome the unfavourable thermal factors such as microstructures and impurities. The overall k properties.value of a solid is given as The three parameters defining the figure of merit, s, S and k=kel+kph k, are all functions of the carrier concentration. Assuming that the broad-band semiconductor model holds good, an increase where kel and kph are the electron and lattice thermal conduc- in the carrier concentration makes s increase while |S| tivities, respectively.It is well known that kph is proportional decreases. According to Ioffe,24 at a constant carrier concen- to T -1 above the Debye temperature (this is usually below tration the Z value is determined by a material factor, b, room temperature).30 The temperature dependence of k of defined as (Zn0.98Al0.02)O shows a good linearity vs. T -1, as shown in Fig. 9, suggesting a predominant contribution of kph to the b=T 5/2 m*3/2 m/kph overall k. This is further confirmed in Fig. 9 by a negligible where T is the absolute temperature, m* is the effective mass proportion of kel estimated from the Wiedemann–Franz of the carrier, m is the mobility, and kph is the lattice thermal relation, kel=L sT (L is the Lorentz number), for which the conductivity. This equation indicates that a large effective mass validity for several electronic-conducting oxides has also been and a high mobility, as well as a small lattice thermal conduc- confirmed.31 It is also known that for semiconductors the tivity, are desirable.Although higher carrier mobility generally electronic contribution in k is a few percent at most, and the requires a smaller effective mass, an increase in the mobility lattice contribution is at least 90% of the overall thermal overcomes the corresponding decrease in the effective mass conductivity, in good agreement with the results obtained here.when the dominant carrier scattering mechanism is impurity Asuppression of kph is therefore desirable for efficient reduction scattering rather than acoustic phonon scattering.25 The small of the thermal conductivity of (Zn1-xAlx)O without seriously difference of the electronegativity between the constituent affecting the electrical properties.This strategy is supported atoms generally increases the mobility. Moreover, heavy atoms by our experimental results for k in Fig. 7, in which the k and a large unit cell reduce the lattice thermal conductivity. value decreases with increasing amounts of added Al2O3, even Caillat et al.investigated the binary compounds on the basis with the considerable increase in s. of these concepts, and discovered skutterudite-related com- The temperature dependence of the dimensionless figure of pounds having extremely high mobilities, CoSb3,26 IrSb327 and merit, ZT , is shown in Fig. 10. The ZT values for x=0.02 and RuSb228 etc., as potential materials for the next generation of 0.05 are much larger than that of undoped ZnO over the thermoelectrics.whole temperature range. The sample at x=0.02 attains the As mentioned above, ZnO consists of rather light atoms largest ZT value of 0.30 at 1000 °C, in spite of its high thermal and has a simple crystal structure and, accordingly, the ZnO- conductivity. The thermoelectric performance of the oxide as based oxide showed a high thermal conductivity.However, ZT=0.30 at present is evaluated as approximately one-third the fairly small electronegativity difference between Zn and O of the standard requirement for thermoelectric materials in should lead to a more covalent character of the MMO bond practical use. However, it should be noted that the present compared to those in other metal oxides.The ratio of the ionic oxide attains such ZT values by overcoming a large disadvan- radii, r(cation)/r(anion), for Zn2+ and O2- is 0.53. Since six- tage due to the markedly high thermal conductivity with a predominant contribution from phonons. Because several Fig. 8 The figures of merit of (Zn1-xAlx)O as a function of temperature Fig. 9 The thermal conductivities of (Zn0.98Al0.02)O as a function of inverse temperature. $, k; #, kph=k-ke. for x=0 (#), 0.02 ($), 0.05 (') J. Mater. Chem., 1997, 7(1), 85–90 89to the overall k value of (Zn1-xAlx)O. It is therefore expected that reduction of the lattice thermal conductivity without seriously affecting the electrical properties, as has been proposed and proven to be effective in many cases, will further improve the ZT value of (Zn1-xAlx)O. The authors thank Mr.Yasuhiro Yamada of the Government Industrial Research Institute, Kyushu, for his kind cooperation on the laser flash measurement of k. We also acknowledge Professor Isao Mochida and Dr. Kinya Sakanishi of the Institute of Advanced Material Study of our University for the 27Al NMR measurements.References Fig. 10 The dimensionless figures of merit of (Zn1-xAlx )Oas a function 1 C. M. Bhandari and D. M. Rowe, Contemp. Phys., 1980, 21, 219. of temperature for x=0 (#), 0.02 ($), 0.05 (') 2 J. C. Bass and N. B. Elsner, in Proc. 3rd Int. Conf. T hermoelec. Energy Conv., ed. K. R. Rao, IEEE, New York, 1980, p. 8. 3 J. F. Nakahara, T.Takeshita, M. J. Tschetter, B. J. Beaudry and methodologies have been reported for effective reduction of K. A. Gschneidner Jr., J. Appl. Phys., 1988, 63, 2331. the lattice thermal conductivity, circumventing degradation in 4 I. Nishida, Phys. Rev. B, 1973, 7, 2710. the electrical properties, we can expect the ZT of (Zn1-xAlx)O 5 I. Nishida and T. Sakata, J. Phys. Chem. Solids, 1978, 39, 499.to improve further. 6 T. Kojima, Phys. Status Solidi A, 1989, 111, 233. 7 C. Wood and D. Emin, Phys. Rev. B, 1984, 29, 4582. 8 S. Yugo, T. Sato and T. Kimura, Appl. Phys. L ett., 1985, 46, 842. Conclusions 9 W. J. Macklin and P. T. Moseley,Mater. Sci. Eng. B, 1990, 7, 111. 10 T. O. Mason,Mater. Sci. Eng. B, 1991, 10, 257. The thermoelectric properties of (Zn1-xAlx)O were investi- 11 W.J. Macklin and P. T. Moseley, Mater. Sci. Eng. B, 1991, 10, 260. gated and revealed to be highly promising for a potential high- 12 C. G. Fonstad and R. H. Rediker, J. Appl. Phys., 1971, 42, 2911. temperature thermoelectric material. The s values of 13 T. P. Pearsall and C. A. Lee, Phys. Rev. B, 1974, 10, 2190. 14 M. Ohtaki, D. Ogura, K. Eguchi and H. Arai, J.Mater. Chem., (Zn1-xAlx)O (0<x0.05) were more than three orders of 1994, 4, 653. magnitude higher at room temperature than that of undoped 15 M. Ohtaki, H. Koga, T. Tokunaga, K. Eguchi and H. Arai, J. Solid ZnO, showing metallic behaviour. The absolute value of the State Chem., 1995, 120, 105. Seebeck coefficient of (Zn1-xAlx )O decreased accordingly, but 16 S. J. Jung, H.Ohsawa, Y. Nakamura, K. Hasumi and O. Okada, remained moderate with a general trend in which |S| increased J. Electrochem. Soc., 1994, 141, L53. gradually from ca. 100 mV K-1 at room temperature to ca. 17 S. N. Bai and T. Y. Tseng, J. Appl. Phys., 1993, 74, 695. 18 M. Ohtaki, T. Tsubota, K. Eguchi and H. Arai, J. Appl. Phys., 200 mVK-1 at 1000 °C. These results appeared to be consistent 1996, 79, 1816.with an increase in the carrier concentration by Al-doping, 19 K. Koumoto, M. Shimohigasi, S. Takeda and H. Yanagida, and were confirmed by the Hall measurements. However, we J. Mater. Sci. L ett., 1987, 6, 1453. concluded that the apparent solubility limit of Al in ZnO is 20 G. Heiland, E. Mollwo and F. Stuo�ckmann, Solid State Phys., less than x=0.005 from the results of the n measurements and 1959, 8, 191. 27Al NMR studies. The x dependence of the Hall mobility was 21 M. Nakamura, N. Kimizuka, T. Mohri and M. Isobe, J. Solid State Chem., 1993, 105, 535. considered to be lated to the grain structures. The power 22 J. W. Akitt, in Multinuclear NMR, ed. J. Mason, Plenum Press, factors of all the Al-doped samples attain markedly large New York, 1987, p. 259. values of (8–15)×10-4W m-1 K-2 over a range from room 23 B. C. Gerstein, Anal. Chem., 1983, 55, 781A. temperature to 1000 °C. These are the largest values of all 24 A. F. Ioffe, Semiconductor T hermoelements and T hermoelectric oxides reported, and surpass those of b-FeSi2 and b-SiC which Cooling, Infosearch Limited, London, 1957. have been studied extensively as high-temperature non-oxide 25 R. P. Chasmer and R. J. Stratton, J. Electron. Control, 1959, 7, 52. 26 T. Caillat, A. Borshchevsky and J-P. Fleurial, in Proc. 13th Int. candidates. The thermal conductivities of all the samples Conf. T hermoelectrics, 1994, ed. B. Mathiprakasem, ALP Press, decreased (from ca. 40W m-1 K-1 at room temperature to New York, p. 58. ca. 5W m-1 K-1 at 1000 °C) with increasing temperature, and 27 T. Caillat, A. Borshchevsky and J-P. Fleurial, in Proc. 13th Int. decreased with increasing amount of doped Al2O3. In spite of Conf. T hermoelectrics, ed. B. Mathiprakasem, ALP Press, New the high k value, the oxide at x=0.02 attained a figure of York, 1994, p. 31. merit of Z=0.24×10-3 K-1 and a dimensionless figure of 28 T. Caillat, A. Borshchevsky and J-P. Fleurial, in Proc. 13th Int. Conf. T hermoelectrics, 1994, ed. B. Mathiprakasem, ALP Press, merit of ZT=0.30 at 1000 °C. The unexpectedly high thermo- New York, p. 209. electric performance of the present oxide can be attributed to 29 J. C. Phillips, Rev. Mod. Phys., 1970, 42, 317. the rather covalent character of the metal-to-oxygen bond in 30 C. Kittel, Introduction to Solid State Physics, 6th edn., Wiley, New the oxide, resulting in a high carrier mobility. Although the k York, 1986. values of (Zn1-xAlx)O appeared to be considerably high 31 M. E. Fine and N. Hsieh, J. Am. Ceram. Soc., 1974, 57, 502. compared to other thermoelectric materials, the lattice thermal conductivity was revealed to have the dominant contribution Paper 6/02506D; Received 10th April, 1996 90 J. Mater. Chem., 1997, 7(1), 85–90

ISSN:0959-9428    

DOI:10.1039/a602506d

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Structural and dielectric study of the Na0.5Bi0.5TiO3–PbTiO3 and K0.5Bi0.5TiO3–PbTiO3 systems Omar Elkechai, Pascal Marchet, Philippe Thomas, Michel Manier and Jean-Pierre Mercurio* L aboratoire deMate�riaux Ce�ramiques et T raitements de Surface, URA CNRS no 320, Faculte� des Sciences, 123, Avenue Albert-T homas, 87060 L imoges Cedex, France A study of the Na0.5Bi0.5TiO3–PbTiO3 and K0.5Bi0.5TiO3-PbTiO3 systems has been carried out using X-ray diffraction, differential scanning calorimetry and dielectric measurements.The limits of the rhombohedral (Na0.5Bi0.5TiO3-rich side) and tetragonal (PbTiO3-rich side) solid solutions have been determined, as well as the evolution of their lattice parameters as a function of composition and temperature. Ceramic materials have been prepared by natural sintering (1090–1220 °C; 0.5 h) of powders obtained by solid-state reaction (900–1000°C; 20 h) of the corresponding oxides and carbonates. The dielectric permittivities of these materials have been measured in a wide frequency range for temperatures between 20 and 800 °C.The results showed that they all are ferroelectric at room temperature and exhibit either a diffuse, probably second-order phase transition, or a sharp, firstorder- like phase transition from the ferroelectric to the paraelectric state, depending on the composition.Lead titanate–zirconate ceramics (PZT) are currently the most The powders were pressed into discs of diameter 10 mm and thickness 1 mm and were sintered in air at 1090–1220 °C, used ferroelectric materials in the field of piezoelectric applications.A noticeable feature of these materials is the occurrence depending on the composition. The thermal cycle consisted of heating at 5°C min-1 to the highest temperature (dwelling of a morphotropic phase boundary (MPB) in the phase diagram which separates tetragonal and rhombohedral ferro- time 30 min) followed by natural cooling to room temperature in the oven. Samples with 90–95% of the theoretical density electric regions.Solid solutions with compositions close to the MPB present the best electromechanical properties. were obtained. After polishing the major faces, they were coated with a low-temperature silver or gold paste fired at In the research of new complex systems with interesting piezoelectric characteristics, care should be taken that thermal 600°C for 10 min, aged overnight at 100 °C and left for 24 hours at room temperature before the measurements.and time stability of the properties is connected with the value of the Curie temperature of the components. Previous works Low-frequency dielectric measurements were carried out between room temperature and 1000°C (at increasing and have shown that in most cases this corresponds to the situation where the MPB is located between tetragonal and rhombo- decreasing temperature) at chosen frequencies in the range 10 kHz–5 MHz using a HP 4194A impedance analyser.hedral phases.1 Up to now, results concerning systems other than PZTs have only been devoted to either Na0.5Bi0.5TiO3–(Sr,Ba)TiO3 or the low-lead range of Na0.5Bi0.5TiO3–PbTiO3.2–6 A study Results was therefore conducted on the whole Na0.5Bi0.5TiO3–PbTiO3 Structural characterisation (NBT–PT) and K0.5Bi0.5TiO3–PbTiO3 (KBT–PT) systems, which are both likely to present MPBs.Actually, at room At room temperature, NBT is rhombohedral (a=0.3891 nm, temperature NBT is rhombohedral (Tc#320 °C), KBT is tetra- a=89.6°) whereas KBT and PT are tetragonal with the gonal (Tc#380 °C) as is PT (Tc#490 °C).7 This paper presents following lattice parameters; KBT, a=0.3918 nm, c= the results concerning the preparation of some materials 0.3996 nm; PT, a=0.3899 nm, c=0.4155 nm.belonging to this family as well as their thermal and dielectric properties. NBT–PT system. XRD data showed two (Na0.5Bi0.5)1-xPbxTiO3 solid solution ranges; rhombohedral on the sodium-rich side and tetragonal on the lead-rich side, Experimental separated by a biphasic region for 0.10<x<0.18.This result is somewhat different from that obtained by Takenaka et al. Polycrystalline compounds were prepared by solid-state reaction of the corresponding oxides or carbonates. Stoichiometric and recently by Park and Hong, who did not find a biphasic range, but found a morphotropic phase boundary at x= mixtures of reagent grade TiO2, Bi2O3, Na2CO3, K2CO3 and PbO (or PbCO3) were mixed thoroughly and calcined in 0.13–0.14 between rhombohedral and tetragonal symmetry at room temperature.4,6 The observed discrepancy of the present alumina crucibles between 900 and 1000°C for 20 h.Further calcinations were necessary to achieve complete reaction.results with respect to other works could be attributed to the different preparation processes, especially the annealing con- Phase characterisation and phase boundary limits were determined by X-ray diffraction (XRD) with a Siemens D5000 ditions. Fig. 1 shows the evolution of the room-temperature lattice parameters as a function of the composition as well as diffractometer (graphite monochromator, Cu-Ka radiation). The lattice parameters were refined by a least-squares method.the c/a ratio within the tetragonal range. As expected, they increase with x according to the increase of the mean ionic The structural evolution of the compounds with temperature was observed using a high-temperature X-ray (HTXRD) radii which change from 0.134 nm (Na,Bi) to 0.147 nm (Pb).Nevertheless, in the tetragonal region, the c parameter increases attachment (Anton Parr HTK10) working between room temperature and 1000°C. Differential scanning calorimetry more strongly than a, leading to a change of the tetragonality from 1.02 (x=0.18) to 1.06 (x=1). This result is slightly (DSC) analyses were performed in air using a Netzsch STA 409 DSC device.different from the parent system NBT-KBT in which the J. Mater. Chem., 1997, 7(1), 91–97 91Fig. 3 shows the thermal evolution of the lattice parameters of tetragonal (Na0.5Bi0.5)1-xPbxTiO3 solid solutions with x= 0.30, 0.50, 0.70 and 1. In each case, increasing temperature causes the tetragonal lattice parameters to vary in opposite senses: i.e.up to the transition temperature a increases and c decreasesalmost monotonously. Nevertheless, the way in which the parameters change between room temperature and the transition temperature is dependent on the composition. On the left-hand side of the tetragonal domain, a and c reach almost the same value corresponding to the cubic lattice parameter at the transition temperature as shown in Fig. 3 (left-hand side). On the right-hand (PbTiO3-rich) side the thermal variation of the lattice parameters at the transition shows a strong discontinuity, which is clearly visible when the lead content is higher than 0.5. Fig. 1 NBT–PT system. Room-temperature lattice parameters vs. composition. KBT–PT system. As expected, XRD data showed that KBT and PT give rise to a full range solid solution with tetragonal symmetry.The evolution of the room-temperature lattice par- tetragonal distortion does not vary greatly with the composiameters of (Na0.5Bi0.5)1-xPbxTiO3 solid solutions as a function tion within the tetragonal domain.8 of composition is given in Fig. 4. As Pb is substituted for Variable-temperature XRD experiments were performed on (Na,Bi), the a parameter of the tetragonal cell decreases slightly several samples belonging to both single-phase domains but the c parameter increases strongly increases, both quasi- (rhombohedral and tetragonal) up to temperatures well above linearly, leading to a more and more distorted unit cell: the the Curie temperature.tetragonality, c/a, changes from 1.02 (KBT) to 1.06 (PT).The thermal evolution of lattice parameters and volume for The thermal evolution of the lattice parameters of some Na0.5Bi0.5TiO3 and (Na0.5Bi0.5)1-xPbxTiO3 (x=0.03, 0.05 and selected compositions of the system are given in Fig. 5. The 0.09 are presented in Fig. 2 as examples of rhombohedral solid behaviour is similar to that observed for the tetr solid solutions. In the temperature range studied, the parameters solutions of the NBT–PT system; the higher the lead content, show a very slight monotonic increase with increasing tempera- the steeper the discontinuity of the lattice parameters (cf.ture, and there are no anomalies at temperatures close to the Fig. 3). expected structural phase transitions. This is because the lowtemperature rhombohedral unit cell is not strongly distorted Dielectric properties with respect to the prototype high-temperature cubic cell.This result is similar to that observed previously in the NBT–KBT The dielectric properties of the ceramic materials were measured at several frequencies between 1 kHz and 10 MHz. In system.8 Fig. 2 NBT–PT system: rhombohedral range. Lattice parameters and cell volume vs.temperature for (Na0.5Bi0.5)1-xPbxTiO3 : x=0 (a), 0.03 (b), 0.05 (c) and 0.09 (d). 92 J. Mater. Chem., 1997, 7(1), 91–97Fig. 3 NBT–PT system: tetragonal range. Lattice parameters and cell volume vs. temperature for (Na0.5Bi0.5)1-xPbxTiO3: x=0.30 (a), 0.50 (b), 0.70 (c) and 1 (d). In the tetragonal range, the sharpness of the permittivity maximum becomes more and more pronounced as the lead content is increased.KBT–PT system. In contrast to the previous system, KBT–PT does not exhibit any lack of miscibility over the whole composition range. As a consequence, one expects that the dielectric properties, especially the thermal variations of the dielectric permittivities, will have continuous behaviour from KBT to PT. Fig. 7 shows the thermal variations of the dielectric permittivities and losses at 1 MHz of compositions within the KBT–PT system.As observed already for the tetragonal compositions of the NBT–PT system, the curves e(T ) show a progressive evolution from a diffuse maximum (KBT) to a very sharp one (PT). Fig. 4 KBT–PT system. Room-temperature lattice parameters vs. composition. Discussion general, the results showed no significant frequency dispersion of either dielectric permittivity or loss. For clarity, only the The dielectric permittivities of NBT and KBT exhibit broad maxima at 320 and 380 °C respectively, undoubtedly connected results obtained at 1 MHz will be given hereafter.with the transition towards the cubic paraelectric state.7 In contrast,PT is known to present asharp permittivity maximum NBT–PT system.The thermal variations of the dielectric permittivities and losses at 1 MHz of some compositions of at 490°C. As the compositions at the extremes of the systems under investigation show well established ferroelectric behav- the NBT–PT system are given in Fig. 6A (x=0, 0.03, 0.08 and 0.09—rhombohedral) and B (x=0.19, 0.20, 0.30 and 0.60— iour at room temperature, the discussion of the thermal evolution of the structural and dielectric characteristics over tetragonal).They clearly show different behaviour according to the symmetry: rhombohedral low-lead content materials the whole composition ranges will be conducted in terms of the phenomenological approach developed by Devonshire exhibit a diffuse character whereas a sharp evolution of the permittivity is present in tetragonal lead-rich materials.after the Ginzburg–Landau theory of phase transitions.9 A main consequence of this theory is that the thermodynamical The rhombohedral solid solutions exhibit the same behaviour as similar materials in the NBT–KBT system.8 The nature of the phase transition of a ferroelectric material can be derived qualitatively from the evolution of the reciprocal temperature of the maximum permittivity is lower than that for pure NBT, but is almost composition independent along permittivity below and above the transition temperature, Tt.For a so-called first-order phase transition, the ratio r of the the rhombohedral range. J. Mater. Chem., 1997, 7(1), 91–97 93Fig. 5 KBT–PT system. Lattice parameters vs.temperature for (K0.5Bi0.5)1-xPbxTiO3: x=0.05 (a), 0.20 (b), 0.40 (c) and 0.90 (d ). slopes of the reciprocal permittivities, e-1(T<Tt)/e-1(T>Tt), over the whole solubility range for NBT–PT and for x>0.20 should be ca. 4, whereas it should be only 2 for a second-order in the case of KBT–PT. The Curie constants calculated at high phase transition. In addition, for a second-order phase trans- temperature are in the range 2–2.5×105 K, as observed comition, the Curie temperature is very close to the temperature, monly for one-dimensional ferroelectrics (e.g. perovskites, tung- T0, extrapolated from the high temperature variations of e-1, sten bronzes).When the lead content increases to PT the whereas this extrapolation always leads to a T0 value less than thermal variations of the permittivities change gradually from Tc for a first-order phase transition.a diffuse to a non-diffuse character (Fig. 6B and 7), as do the The overall properties of the materials under investigation slope ratios r of the reciprocal permittivities which change are strongly dependent on their crystal structures, either from 2.5 to 4 (Fig. 8). After Devonshire, this behaviour should rhombohedral for the (Na,Bi)-rich side of NBT–PT, or tetra- correspond to a progressive evolution from a diffuse, second- gonal for both the lead-rich side of NBT–PT and the whole order to a sharp, first-order phase transition as the composi- composition range of KBT–PT. The discussion will be focused tions approaches PT. This assumption should be supported separately on the properties of materials with (i) rhombohedral by the thermal evolution of the lattice parameters.At low lead and (ii) tetragonal structures. contents, the transition from the tetragonal to the cubic symmetry occurs gradually by a mere convergence of the (i ) Rhombohedral solid solutions: (Na0.5Bi0.5)1-xPbxTiO3 lattice parameters a and c towards the value of the parameter (0<x0.09) of the cubic unit cell at the transition temperature.So, there is no volume change at the transition. However, for x>0.40, A careful examination of the thermal evolutions of the permitthere is a discontinuous change of symmetry at the transition tivities show a composition dependent weak shoulder on the temperature leading to an increase in cell volume.This change low-temperature side of the peak. In NBT, this shoulder was becomes larger as x increases. The volume variation, which assigned initally to a ferroelectric–antiferroelectric transition.2 Recently, Park and Hong refined this sequence and proposed changes from -0.07×10-3 (NBT–PT) to -0.4×10-3 nm3 a diffuse phase transition zone in the temperature range (PT), is characteristic of a first-order phase transition.In 240–280 °C for x<0.1.6 In the present work, the temperatures addition, DSC experiments carried out at increasing and at which this shoulder appears decreases from ca. 230 °C for decreasing temperatures have shown very well defined peaks NBT to ca. 150 °C for the upper limit of the rhombohedral corresponding to the transition temperatures.With a heating/ domain. This result is similar to that already observed for cooling rate of 20°C min-1, the DSC curves exhibit thermal NBT–ST solid solutions where the ferroelectric–antiferroelec- peaks at the transition temperatures with a thermal hysteresis tric transition temperature decreases monotonically as the DT#-15 °C, as observed already in pure PT.These peaks strontium content increases.5 As a consequence, Devonshire are even stronger and sharper when the lead content is calculations cannot be applied to the low-temperature side of increased. The energy involved in the phase transitions can be the permittivity peak. calculated accurately only for compositions where x>0.30. These calculations show that the energies increase linearly up (ii) Tetragonal solid solutions: (Na0.5Bi0.5 )1-xPbxTiO3 to x=0.70 and then increase more rapidly as x approaches ( 0.18<x1) and (K0.5Bi0.5)1-xPbxTiO3(0x1) unity.This behaviour is very similar to that of the increase in volume change at the transition calculated from the lattice These tetragonal solid solutions show similar behaviour. The phenomenological calculation can be carried out successfully parameters. 94 J. Mater. Chem., 1997, 7(1), 91–97Fig. 6 NBT–PT system. Permittivity and loss vs. temperature for (Na0.5Bi0.5)1-xPbxTiO3 . A, rhombohedral range: x=0 (a), 0.03 (b), 0.08(c) and 0.09 (d). B, etragonal range: x=0.19 (a), 0.20 (b), 0.30 (c) and 0.60 (d). In the tetragonal domain, the more striking feature is the of the Curie temperatures as observed, for example, in BT–ST or BT–PT.However, in both cases, the Curie temperatures evolution of the temperature of the phase transition from tetragonal to cubic symmetry. As the tetragonal compositions show a maximum value at ca. 70–80 mol% lead, in close agreement with HTXRD and DSC experiments (Fig. 9). This seem to behave as regular solid solutions (e.g.KBT–PT solid solutions follow Vegard’s law), one expects a linear variation unusual behaviour could be connected with the respective J. Mater. Chem., 1997, 7(1), 91–97 95Fig. 7 KBT–PT system. Permittivity and loss vs. temperature for (K0.5Bi0.5)0.51-xPbxTiO3: x=0 (a), 0.09 (b), 0.18 (c), 0.30 (d), 0.40 (e), 0.50 (f ), 0.60 (g) and 0.80 (h). polarizabilities of lead and (Na,Bi) or (K,Bi).Up to ca. 60 Conclusion mol% lead, the influence of increasing lead content dominates Ferroelectric to paraelectric phase transitions in the NBT–PT over the effect of decreasing the (Na,Bi) or (K,Bi) content and and KBT–PT systems were studied as a function of composi- the Curie temperature increases, whereas above 60 mol% lead, tion and temperature. The nature of the transition was dis- the decrease of (Na,Bi) or (K,Bi) content is not compensated cussed in terms of the phenomenological derivation of the by the increase in Pb content and the Curie temperature decreases to that of PT.Ginzburg–Landau theory. The diffuse character decreases 96 J. Mater. Chem., 1997, 7(1), 91–97Fig. 9 Transition temperatures vs. composition from HTXRD (%, increasing T ) and DSC (', increasing T ; #, decreasing T ) for (a) Fig. 8 Direct and reciprocal permittivities of two selected NBT–PT NBT–PT and (b) KBT–PT solid solutions showing (a) second-order (x=0.19) and (b) first-order (x=0.60) phase transitions 3 C. S. Tu, I. G. Siny and V. H. Schmidt, Phys. Rev. B, 1994, 49, strongly as the lead content increases in both systems. A study 11550. of the thermal evolution of the polarisation coupled with 4 T. Takenaka, K. Sakata and K. Toda, Ferroelectrics, 1990, 106, variable-temperature optical microscopy examinations of these 375. materials is now in progress in order to confirm the above 5 K. Sakata and Y. Masuda, Ferroelectrics, 1974, 7, 347. statements. 6 S-E. Park and K. S. Hong, J. Appl. Phys., 1996, 79, 383. 7 G. A. Smolenskii, V. A. Isupov, A. I. Agranovskaia and N. N. Krainik, Fiz. T vergodo T ela, 1960, 2, 2982. References 8 O. Elkechai, M. Manier and J. P. Mercurio, Phys. Status Solidi A, in the press. 1 B. Jaffe, W. R. Cook and H. Jaffe, Piezoelectric Ceramics, Academic 9 A. F. Devonshire, Philos. Mag., 1949, 40, 1040. Press, London, 1971. 2 B. Vakhrushev, V. A. Isupov, B. E. Kvyatkovsky, N. M. Okuneva, I. P. Pronin, G. A. Smolenskii and P. P. Syrnikov, Ferroelectrics, 1985, 63, 153. Paper 6/02148D; Received 27thMarch, 1996 J. Mater. Chem., 1997, 7(1), 91–97 97

ISSN:0959-9428    

DOI:10.1039/a602148d

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Local density functional calculations of the electronic structures of Ti2AlC and Ti3AlC Samir F. Matar, Yann Le Petitcorps and Jean Etourneau Institut de Chimie de laMatie`re Condense�e de Bordeaux–CNRS, Cha�teau Brivazac, Avenue du Docteur Schweitzer, F33600 Pessac, France Local density functional calculations are used to address the electronic structures and the properties of chemical bonding of two definite phases formed within the ternary system Ti, Al and C: Ti2AlC and Ti3AlC.From the analyses of the density of states and of the crystal orbital overlap populations of the respective phases within the ASW method the role of C is assessed. Moreover, the bonding within TiC is discussed concomitantly. These calculations are of interest in the composite field to understand the mechanisms of formation of new compounds at the matrix/reinforcement interface.Carbon fibre-reinforced titanium–aluminium intermetallic ably large atomic sphere radii leading to a large overlap in the ASA. composite materials are of interest for space and aeronautics applications. In recent years several works have been devoted In k-space, the Brillouin zone integration is achieved on a uniform mesh of points in the irreducible wedge of the relevant to their investigation, both experimentally1–3 and theoretically. 4,5,6 From the latter point of view, we modelled recently Bravais lattice. The matrix elements are constructed involving solutions of the Schro�dinger equation up to the secondary l the influence of substituted and inserted carbon within the alloy lattice of TiAl on its electronic structure.6 Based on quantum number, lmax+1, where lmax=2 for Ti and Al and 1 for C and the ES.The contributions associated with the lmax+1 quantitatively resolved chemical bonding criteria we proposed that carbon should substitute preferentially for aluminium higher angular momenta are relative to non-explicitly calculated terms in the limited ASW basis set7 but should always when it enters the TiAl lattice.This is supported by the actual occurrence of Ti-rich carbide compounds such as Ti2AlC in be lower than 0.1 electron in order to ensure a convergence of the immediate neighbourhood of the intermetallic matrix of TiAl. Moreover, the growth of such a carbide phase from this alloy leads to an enrichment of Al atoms at the intermetallic/ ternary compound interface.For the titanium-rich alloy Ti3Al, carbon is inserted to give Ti3AlC. In all cases TiC is formed in the vicinity of the carbon fibre (Fig. 1). Thus Ti–C interactions form TiC, whereas Ti2AlC and Ti3AlC are formed respectively with the TiAl–C and Ti3Al–C couples, according to a diffusion path determined by Clochefert.3 In this, the second part of our investigation of carboncontaining TiAl, we address the electronic properties of the titanium-rich carbides Ti2AlC and Ti3AlC actually forming in the {Ti–Al–C} phase diagram, with the objective of examining the influence of carbon on the chemical bonding within the alloy lattice.Method of calculation As in our earlier investigation,6 the electronic properties of all carbon-containing alloy systems were calculated using the ab initio self-consistent augmented spherical wave (ASW) method.7 The ASW method allows one to describe the electronic properties of a material starting from those of its atomic constituents.The calculations are based on the density functional theory in which the effects of exchange and correlation are treated in the local density approximation within the scheme of von Barth and Hedin, and Janak.8 The ASW method uses the atomic sphere approximation (ASA) where each atom is surrounded by a sphere.Within the atomic spheres the potential is assumed to be spherically symmetric. The ASA imposes a unit-cell volume equal to the total volume of the spheres, leading to their overlap. This is unproblematic for close-packed crystal structures, but for loosely packed ones the empty space must be represented by use of ‘empty spheres’ (ES), i.e., pseudo-atoms with Z=0 atomic number and no core Fig. 1 Electron micrographs of the chemical interaction between states. ES are introduced in order to account for the interstitial carbon and Ti–Al intermetallics. Upper: C–TiAl; lower: C–Ti3Al.(Reproduced with permission from ref. 3.) space in the lattice and to avoid the use of otherwise unreason- J. Mater. Chem., 1997, 7(1), 99–103 99the charges. The self-consistent cycle is carried out until the is an anion at face centres of the cube; space group Pm3m). Therefore Ti3AlC [Fig. 1(b)] can be described as an antiper- following convergence criteria are reached: DE=10-8 Ry [1 Ry (rydberg)=13.6 eV] for the total energy and DQ=10-8 ovskite because Ti atoms occupy the face centres, Al and C being at corner and cube-centre positions respectively.From for the charge difference between two successive cycles. In this work all calculations were carried out at the experimental the point of view of coordination polyhedra, the structure can be regarded as a three-dimensional array of Ti6C octahedra lattice constants obtained from ref. 3. Furthermore, in this work the chemical bonding features are discussed based sharing corners. It is hence a poorly packed structure because the midpoints of the edges, D 0 0, 0 D 0 and 0 0 D, are vacant on the so-called COOP (read CO-OP: crystal orbital overlap populations), of which a comprehensive account was given by sites where ES had to be introduced in the ASA. Before examining the electronic structures of these two Hoffmann from the quantum chemistry standpoint (extended Hu�ckel calculations).9 This allows for the DOS features to be carbides, it is relevant to consider the coordination polyhedra in TiC.We stress that this binary carbide is modelled here in discussed on bases of chemical bonding criteria by weighting them with the sign and magnitude of the overlap integral a 151 composition although it is known to be sub-stoichiometric in carbon, i.e.TiCx with 0.56<x<0.98. between the relevant orbitals. We recently implemented the COOP in the ASW method10 with the objective of obtaining TiC crystallizes in the NaCl-type structure (space group Fm3m) with four formula units per unit cell [Fig. 2(c)]. Ti and more precise information on the chemical bonding from first principles. C are at the origin and D D D positions, respectively. Ti and C are octahedrally coordinated with each other; consequently, Ti6C octahedra share edges. Crystal structures and setup of the unit cells for By comparing the three structures an interesting observation ASW calculations appears: dimensionality increases from Ti2AlC to Ti3AlC and TiC.In contrast to several Ti2AlM (M=Nb, V, Cr, Mn) compounds Ad hoc and non-unique choices of the atomic spheres radii which crystallize in a tetragonal structure,5 Ti2AlC is hexagonal in the ASA were such that: rTi/rES=1.26, rAl/rES=1.30 and with a large c axis and two formula units per cell.It crystallizes rC /rES=1.10. Such values were tested as one choice which in the Cr2AlC-type structure11 with the P63/mmc space group simultaneously minimizes the overlap between the spheres and and Ti at (4f ), Al at (2c) and C at (2a) Wyckoff positions. The yields in converged lmax+1 residual charges. structure is shown in Fig. 2(a). It can be regarded as an alternating stacking of triangular prisms and octahedral Ti polyhedra containing Al and C atoms, respectively, along the Calculations and Results c-axis.Ti6C octahedra share edges and form two-dimensional Partial charges layers perpendicular to the c axis. From this low dimensionality, the structure is poorly packed and in the ASA, ES had to Table 1 gives the partial charges for Ti2AlC and Ti3AlC (DQ be introduced between the layers at sites related to those of Ti designates the deviation from neutrality).The overall features general positions. of charge transfer are similar, in that it occurs from the two With one formula unit per unit cell, Ti3AlC has a structure metallic atoms towards the non-metal and the empty spheres, derived from the cubic perovskite ABX3 (where A and B are i.e.Ti,Al�C,ES. The averure from neutrality per large and small cations at corner and centre positions, and X metal is then ca. 0.73 in Ti2AlC and ca. 0.79 in Ti3AlC. In as far as carbon receives ca. 0.6 electrons in both carbides, this charge excess leads to the larger occupancy of ES in the latter. However, the differences which characterize the DQ values of Ti and Al in each compound should be addressed.They arise from the fact that Al exhibits a larger d character in Ti2AlC (represented by the higher d occupancy) than in Ti3AlC. This should indicate a larger hybridization between Ti and Al in Ti2AlC with respect to Ti3AlC. By virtue of this mixing there is an enhancement of the sp character of Ti which could be due to its interaction with Al and/or with C.This establishes a covalent character of the bonding in these materials, to be further illustrated in next section. Density of states (DOS) The upper panels (a) of Fig. 3 and 4 show the site-projected DOS of Ti2AlC and Ti3AlC. Energy reference along the x axis Table 1 Site and l-projected partial charges for Ti2AlC and Ti3AlC s p d f DQ Ti2AlC(ES)2:a Al 1.08 0.82 0.23 (0.02) -0.84 Ti 0.37 0.56 2.33 (0.08) -0.67 C 1.40 3.05 0.15 (0.02) 0.62 ES 0.54 0.20 (0.04) — 0.78 Ti3AlC(ES)3:b Al 0.88 1.25 0.15 (0.01) -0.69 Ti 0.32 0.51 2.28 (0.07) -0.82 C 1.29 3.07 0.21 (0.03) 0.60 ES 0.62 0.19 (0.05) — 0.85 Fig. 2 (a) Hexagonal structure of Ti2AlC. (b) Perovskite-derived structure of Ti3AlC. (c) NaCl-type structure of TiC. (Reproduced with aQ=2(-0.67)-0.84+0.62+2(0.78)=0 (neutrality).bQ=3(-0.82)- 0.69+0.60+3(0.85)=0 (neutrality). permission from ref. 3.) 100 J. Mater. Chem., 1997, 7(1), 99–103is taken with respect to the Fermi level (EF) within a reduced energy range (-8 to +8 eV), i.e. excluding the low lying C 2s states, to make the presentation clear. The y axis gives the DOS per atom and unit energy (atom-1 eV-1).In both carbides, the Fermi level crosses the lower part of the Ti 3d states centred above EF because of the nearly empty d band. These states show much larger structures towards the lower energies than in TiAl,6 where they interact solely with Al s,p states because of the extra interaction with carbon. This is shown by the peaks between -6 and -4 eV in Ti2AlC and -7 and -3 eV in Ti3AlC.From a preliminary crystal-field analysis, the peaks in Ti DOS at -2, 1 and 2 eV arise mainly from in-plane xy and x2-y2 d orbitals. However, it is difficult to separate totally the contributions of the five different d orbitals because of their hybridization and of the collective character of the electrons. The DOS at EF , n(EF), are dominated by Ti 3d and show a sharp peak in Ti2AlC, probably due to Ti–C interactions (see next section) whereas such a feature is absent in Ti3AlC where n(EF ) are three times lower. There is a larger contribution from Al states at and above EF in Ti2AlC which agrees with our discussion of the charges, leading to a mixing between Ti 3d and Al 3p.In Ti3AlC the DOS are dominated by carbon and Ti on one hand and Ti d on the other hand below and above EF, respectively.Al plays a less important role at EF in this carbide and its ‘sp’ DOS are seen in the energy windows -8 to -6 eV and -4 to -2 eV. A relevant feature is the broadness of the band over the energy range -8 to-2 eV as opposed to the sharp peaks in the same range in Ti2AlC. In both compounds the DOS of the ES closely follow those of the other species, which is consistent with charge transfer into them from the other sites.Fig. 3 Ti2AlC: (a) site projected densities of states in atom-1 eV-1 (solid line Ti; dashed line Al; dotted line C; dash-dotted line ES); (b) At this point the discussion of the mixing features solely partial COOP for pair interactions: Ti–C (solid line), Ti–Al (dashed from the partial DOS cannot give more information about the line), Al–C (dotted line) chemical bonding in the two compounds.A further step must be undertaken, by examining the COOP. Crystal orbital overlap populations (COOP) The features of chemical bonding can be assessed further by using the COOP. In the lower panels (b) of Fig. 3 and 4, the COOP are shown for the interactions between the different atoms in the two compounds plotted in the same energy range, i.e.for Ti–C, Ti–Al and Al–C. Along the y(COOP) axis, positive, negative and zero values point to bonding, antibonding and non-bonding states, respectively. In Ti2AlC, below EF, Ti–C interactions predominate, whereas in TiAl Ti–Al interactions are the driving interaction for the bonding.6 They exhibit largely bonding character in the range -6 to-2 eV, and follow exactly the Ti DOS in the same energy range in which C 2p states dominate.Thus the sp character introduced into Ti (cf. Table 1) mainly arises from its interaction with carbon. The antibonding counterpart can be seen in the conduction band (2 to 6 eV). The large separation between the bonding and antibonding peaks is indicative of a strong interaction assimilated with a s-like interaction.This somehow opposes the Ti–C interaction in Ti3AlC, where less localized bonding states are seen to extend over a wide band in a larger energy window (s- and p-like bonding). Interestingly, carbon is engaged not only in Ti–C interactions but also in Al–C ones, Ti–C and Al–C interactions having bonds in the same energy range.This is in contrast to Ti2AlC, where only Ti–C interactions are present in the valence band. This is supported experimentally, whereby the solubility of carbon is much more important in Ti3Al than in TiAl.11 However, the Al–C bond seems weaker because it is largely antibonding from -4 to -2 eV whereas Ti–C is bonding over a wider energy range. For the Ti–Al interaction, the distance between these two Fig. 4 Ti3AlC: (a) site projected densities of states in atom-1 eV-1 sites is 23% shorter in Ti2AlC than in Ti3AlC. This should (solid line Ti; dashed line Al; dotted line C; dash-dotted line ES); (b) explain the differences appearing between the two panels for partial COOP for pair interactions: Ti–C (solid line), Ti–Al (dashed line), Al–C (dotted line) the Ti–Al interaction and should assess the d character brought J.Mater. Chem., 1997, 7(1), 99–103 101into Al by its bonding with Ti (cf. Table 1). As a matter of fact, Ti–Al becomes important only around EF, i.e. in the DOS region where Ti d states predominate. In contrast, Ti–Al interactions in Ti3AlC are mainly seen in the energy range of Al 3p states, around -2 eV, in the same energy range as Ti–C and Al–C bonds with the largest bonding contribution at the top of the valence band, whereas they are nearly absent in the Ti–C energy range in Ti2AlC owing to the nearly two-dimensional array of Ti6C octahedra.Comparison with TiC At this point a comparison with the electronic structure of TiC is in order. Fig. 5(a) gives the site-projected DOS of TiC. They are in good agreement with those of Blaha and Schwarz (ref. 12 and refs. cited therein) who gave a full account of the electronic structures of TiX (X=C, N, O) compounds by use of a linearized APW (augmented plane waves) method. The feature of the n(EF) minima at the Fermi level is related to the refractory nature of TiX and their stability.13 From -6 eV to EF, C 2p states predominate, whereas from EF to 8 eV, Ti 3d states with their t2g (4 eV) and eg (>4 eV) components show the major contribution to the DOS.In the valence band and from 4 eV to higher energies, C and Ti states have similar shapes, which is indicative of a covalent interaction between them. The DOS of ES follow the same evolution as the Ti and C ones, indicating that charge transfer into them is from both species.In the purely Oh point symmetry, Ti 3d orbitals split into two types of manifold: t2g and eg. The projection of the DOS along them is shown in Fig. 6. In the valence band, eg orbitals Fig. 6 Oh crystal-field decomposition of the Ti d-orbital DOS: (a) Ti d(t2g); (b) Ti d(eg) have a larger contribution with respect to t2g ones; they are involved with pds-type bonding with carbon whereas pdp bonding should be less involved.Since metal Ti–Ti interactions are of the dds type, one expects little bonding of this type in the valence band. This is explained more quantitatively by examining the COOP shown in the same energy window as the DOS in Fig. 5(b). They are resolved for three kinds of interactions in the lattice, namely Ti–C, Ti–Ti and C–C. The latter two types of interactions show fewer bonding features than the former; Ti–Ti interactions are clearly less predominant (for the reasons argued above) than C–C interactions, which exhibit bonding and antibonding states in the valence band whereas Ti–Ti bonding features can only be seen in the conduction band.Thus, in TiC 3d–2p bonding is the driving bonding force. Two types of 3d–2p bonding are found with increasing energy, i.e.pds predominates over pdp in the valence band, whence the directional character of the bonding in this compound. Since the former are stronger the antibonding counterparts are reversed, following energetical order: pdp* resembled by the antibonding peak at 4 eV and pds* at higher energies. In our two carbide systems, the directionality of the bond is reduced by the presence of Al, which acts through its p and d states in its bonding to Ti.From this there is an increase in the amount of d character in the valence band. This is indicated by the larger occupation of Ti 3d with 0.52 and 0.47 electrons in Ti2AlC and Ti3AlC, respectively, in comparison to the Ti d-band occupation in TiC. As a matter of fact, it explains the larger n(EF) of the former and its vanishing value in TiC.Discussion and Conclusion Experimentally, it was found that Ti–C bonding was the driving force which controlled the C–TiAl and C–Ti3Al interactions. 3 These experimental features agree rather well with Fig. 5 TiC: (a) site projected densities of states (solid line Ti; dashed the results of our calculations, indicating strong Ti–C inter- line C; dash-dotted line ES); (b) partial COOP for pair interactions: Ti–C (solid line), Ti–Ti (dashed line), C–C (dotted line) actions which are reminiscent of the formation of TiC in the 102 J.Mater. Chem., 1997, 7(1), 99–103immediate vicinity of the carbon fibre. This should destabilize lations was performed within the MNI pole of intensive computations.the actual alloy lattice, leading to the formation of precipitates. Our investigation has shown that carbon plays different roles in Ti2AlC and in Ti3AlC. While it bonds mainly to Ti in the former, both Ti–C and Al–C bonds are present in the References valence band for the latter. This is supported experimentally 1 D. Vujic, Z. Li and S. H. Wang, Metall.T rans. A, 1988, 19, 2445. because the solubility of carbon is higher in Ti3AlC when it 2 M. Morinaga, J. Saito, N. Yukawa and H. Adach, Acta Metall. enters the alloy lattice. Valence and conduction bands are Mater., 1990, 38, 25. largely separated by non-bonding states at EF in Ti3AlC, 3 L. Clochefert, PhD Thesis, Universite� Bordeaux 1, 1995. whereas bonding metallic Ti–Al interactions are present at EF 4 S.R. Chubb, D. A. Papaconstantopoulos and B. M. Klein, Phys. in Ti2AlC. This could point to a higher ‘ionic’ character in Rev. B, 1988, 38, 12120. 5 H. Erschbaumer, R. Podloucky, P. Ro�gl, G. Temnitschka and Ti3AlC and a higher ‘metallic’ character in Ti2AlC. R. Wagner, Intermetallics, 1993, 1, 99. Our comparison of the bonding within Ti2AlC and Ti3AlC 6 S. F. Matar and J. Etourneau, J. Alloys Compd., 1996, 233, 112. to its characteristics in TiC shows that Al reduces the direc- 7 A. R. Williams, J. Ku�bler and C. D. Gelatt Jr., Phys. Rev. B, 1979, tionality of the bonding mainly through its p states, as well as 19, 6094. its d ones which interact with Ti d states. From this there is 8 U. von Barth and L. Hedin, J. Phys. C, 1972, 5, 1629; J. F. Janak, an increase in the amount of d character in the valence band. Solid State Commun., 1978, 25, 53. 9 R. Hoffmann, Angew. Chem., Int. Ed. Engl., 1987, 26, 846. To conclude, the present study has brought a new insight 10 V. Eyert and S. F. Matar, 1994, unpublished results. into the bonding features in the C–Ti–Al ternary system, 11 W. B. Pearson, Acta Crystallogr., Sect. A, 1980, 36, 724. allowing for a more quantitative description on the chemical 12 P. Blaha and K. Schwarz, Int. J. Quantum Chem., 1983, 23, 1535. bonds. 13 J. Ha�glund, G. Grimvall, T. Jarlborg and A. Fernandez Guillermet, Phys. Rev. B, 1991, 43, 14400. Facilities provided by the computer centre of the Universite� Bordeaux 1 (CRIBx1) are acknowledged. Part of the calcu- Paper 6/05113H; Received 23rd July 1996 J. Mater. Chem., 1997, 7(1), 99&nda

ISSN:0959-9428    

DOI:10.1039/a605113h

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Metal sulfide preparation from a sol–gel product and sulfur Vesna Stanic�,a Thomas H. Etsell,*a Alain C. Pierrea† and Randy J. Mikulab aDepartment of Chemical and Materials Engineering, University of Alberta, Edmonton, Alberta T 6G 2G6, Canada bCANMET Western Research Centre, Devon, Alberta T 0C 1E0, Canada Synthesis of metal sulfides has been demonstrated by the example of monoclinic germanium disulfide produced by reaction of the sol–gel product with sulfur.The elemental sulfur, in turn, was obtained as the result of oxidation of H2S in the presence of concentrated sulfuric acid. This sulfur was transported into a toluene solution of germanium ethoxide and found to be homogeneously distributed within the gel. Heat treatment of the sol–gel product yielded single-phase GeS2.Products before and after heat treatment were characterized by IR, XRD, SEM and EDXA measurements. The electronic and optical properties of semiconducting metal medium for H2S oxidation. The sol–gel reaction product (before and after heat treatment) was characterized by IR sulfides are strongly affected by impurities and inhomogenities in the material, which may be caused by the preparation spectroscopy measured in the range 400–4000 cm-1 using an FTIR spectrometer (Bruker IFS 113V) equipped with a process.The sol–gel process is expected to provide a homogeneous product with a low level of impurities, and lead to a photoacoustic cell (Princeton Applied Research, model 6003, EG&G). The samples were analysed by X-ray diffraction product with a controlled structure and stoichiometry.However, when alkoxides and H2S gas were used as precur- (XRD) on a Rigaku RU-200B automated powder diffractometer with Cu-Ka radiation. The morphologies of the sors, the reaction products were mixtures of metal oxides and sulfides,1–5 since the alkoxides readily react with water. In dried gels and the heat-treated products were examined on a Hitachi S-2700 scanning electron microscope (SEM).order to avoid oxide formation, special precautions have to be applied to protect the reaction mixture from contamination Microanalysis of the Ge and S content in the samples was performed by energy dispersive X-ray analysis (EDXA) with a by water. Otherwise, an additional reaction has to be undertaken to eliminate the oxide impurity and purify the metal Link analytical eXL incorporated into the electron microscope.sulfide product. Kumta and Risbud,4 for instance, used H2S gas during heat Sol–gel processing treatment to obtain single-phase lanthanum sulfide. Sriram Germanium ethoxide and toluene were mixed under a dried and Kumta5 applied a similar method to produce crystalline nitrogen atmosphere in a glove box.The toluene5ethoxide TiS2, i.e. heat treating the powder, obtained by sol–gel pro- ratio was 9555 (vol%). Before introduction into the solution, cessing from titanium alkoxide and H2S gas, in the presence H2S was passed through concentrated sulfuric acid that half- of an H2S gas flow. filled an Erlenmeyer flask in an air atmosphere. An excess of Here, we present an original approach to metal sulfide H2S carrying elemental sulfur liberated in sulfuric acid was preparation as demonstrated through the synthesis of mono- bubbled through the toluene–ethoxide solution until complete clinic germanium disulfide.The method is a combination of gelation had occurred. The obtained gel enclosed in the reactor the sol–gel processing of germanium disulfide previously was aged for 24 h.Samples for analysis were dried in a vacuum reported1,2 and an H2S oxidation reaction to elemental sulfur oven at room temperature. and water in the presence of concentrated sulfuric acid. Sulfur forms a colloidal suspension in sulfuric acid, and it can be Heat treatment extracted and homogeneously distributed in a toluene solution of germanium ethoxide by an H2S gas flow.Since the sol–gel The gel was placed into a quartz ampoule evacuated to reaction product was a mixture of GeS2 and GeO2,1–3 sulfur 10-4 Torr pressure, heat-treated isothermally in a furnace at converted the oxide to sulfide during heat treatment yielding 630°C for 24 h to homogenize the solid mixture and then single-phase monoclinic GeS2. The reaction product was quenched in water to room temperature.The cooling rate was characterized by various methods before and after heat >17°C s-1. treatment. Results Experimental The obtained gel was yellow and of smooth appearance. The Materials and analysis XRD pattern of the gel [Fig. 1(a)] shows the presence of two For the sol–gel processing of GeS2, germanium ethoxide crystalline phases. One was identified as hexagonal GeO2,6 manufactured by CHEMAT [99.99% Ge(OEt)4] and hydro- while the other was orthorhombic sulfur.7 The XRD pattern gen sulfide, produced by Linde Canada (99.6% H2S), were of the gel after heat treatment [Fig. 1(b)] is identical to that used as precursors and toluene was used as solvent (Fisher of monoclinic GeS2.8 Scientific, HPGC grade). Even though the toluene had a low The peak positions from this XRD pattern were matched water content, <0.02%, it was dried by refluxing over Na with a standard8 using software PDF-2 Database sets 1–42.metal, for 24 h and then distilled. Concentrated sulfuric acid The 2h and d-spacings for the heat-treated product and the (98%), supplied by Fisher Scientific, was used as the acidic standard are shown in Table 1 along with their absolute deviations.The IR spectrum of the gel in the range 400–1000 cm-1 is † Present address: Universite� Claude Bernard – Lyon I, LACE, Ba�t. 303, 43 Bd. du 11 Novembre 1918, 69622 Villeurbanne Cedex, France. shown in Fig. 2(a). In the range 400–450 cm-1 shoulders at J. Mater. Chem., 1997, 7(1), 105–107 105Table 1 XRD peak positions (2h) and d-spacings of the heat-treated GeMOMGe,10 which confirms the presence of hexagonal product and of standard GeS28 with absolute deviations GeO2.A shoulder at 780 cm-1 suggests the existence of a GeMOH vibration.10 An absorption band at 473 cm-1 is 2h/degrees d/A° characteristic of an SMS bond stretch of elemental sulfur;11 absorption peaks due to toluene are also observed at ca. 694 standard heat-treated standard heat-treated hkl GeS2 product D(2h) GeS2 product Dd and 730 cm-1.11 The IR spectrum of the heat-treated gel is shown in Fig. 2(b) 200 15.479 15.477 -0.002 5.7200 5.7208 0.0008 and has absorption peaks close to those of monoclinic GeS2 111 16.372 16.377 0.005 5.4100 5.4082 -0.0018 at ca. 405, 430 and 450 cm-1.9 Furthermore, an absorption 211 21.239 21.245 -0.006 4.1800 4.1788 -0.0012 band at 483 cm-1 can be assigned to SMS vibrational 002 26.555 26.551 -0.004 3.3540 3.3544 0.0004 absorption. 151 31.832 31.839 0.007 2.8090 2.8084 -0.0006 260 36.978 36.975 -0.003 2.4290 2.4292 0.0002 A scanning electron micrograph of the dried gel [Fig. 3(a)] -332 39.223 39.228 0.005 2.2950 2.2947 -0.0003 reveals that the gel is colloidal, i.e.formed by linking of spherical colloidal particles into a porous network. From the micrographs, it is evident that particles of ca. 0.4 mm are linked by necks. As well as the gel network, the micrograph of the gel obtained at lower magnification [Fig. 3(b)] reveals large crystals of dimension ca. 60 mm. The morphology of the heattreated gel is shown in Fig. 3(c) and indicates a fused structure with large grains and the absence of macropores.Energy dispersive X-ray analysis (EDXA) shows that the particles have a Ge5S atomic ratio of 153.7. However, the large crystals with a ratio 1546.5 are almost pure sulfur. The microstructure changes after heat treatment, with the sulfur crystals completely disappearing and a homogeneous structure being formed. EDXA revealed that the atomic ratio Ge:S was constant when measured at different spots on the sample, not deviating significantly from the average value of 152.9.Discussion The IR spectrum of the gel has well defined absorption peaks, typical of hexagonal GeO2 the presence of which is also confirmed by the sharp peaks in the XRD pattern of the gel. Although the XRD pattern of the gelnot indicate amorphous GeS2, its presence was confirmed by IR spec- Fig. 1 XRD pattern of (a) the gel and (b) monoclinic GeS2 obtained troscopy. However, both the IR spectrum and the XRD pattern after heat treatment of the gel confirmed the presence of crystalline orthorhombic sulfur. The overall reaction for H2S oxidation can be proposed as: H2S+cO2 ,b) H+ S0+H2O (1) The sulfur formed by reaction (1) was carried into the reaction mixture by the H2S gas flow and crystallized into the stable orthorhombic structure within the gel.The influence of the acid on the H2S oxidation can be explained by the following electrochemical reactions: H2S=2H++S0+2e (1a) 2H++cO2+2e=H2O (1b) If reaction (1b) is the slow step then the presence of the concentrated acid aids the oxidation of H2S.The oxygen involved in reaction (1b) probably comes from air trapped above sulfuric acid in the Erlenmeyer flask. Since the gel contains both GeO2 and GeS2,1–3 the sol–gel processing of these compounds occurred simultaneously. Reactions (2)–(4) describe the formation of GeO2: OGeMOR+H2O=OGeMOH+ROH (2) OGeMOR+HOMGeO=GeMOMGeO+ROH (3) Fig. 2 IR spectrum of (a) the gel and (b) monoclinic GeS2 obtained after heat treatment OGeMOH+HOMGeO=OGeMOMGeO+H2O (4) The GeMOH bond was identified in the IR spectrum of the ca. 405, 415 and 430 cm-1 can be assigned to GeMS vibrational absorptions by vitreous GeS2.9 This spectrum shows a gel, providing evidence for this reaction mechanism during the sol–gel processing of GeO2. Water involved in the reaction vibrational absorption triplet of hexagonal GeO2 at 515, 555 and 587 cm-1.10 The gel also has a strong absorption peak at was formed according to reaction (1) and was introduced into the mixture by the H2S gas flow.3 ca. 885 cm-1, assigned to the asymmetric stretching of 106 J.Mater. Chem., 1997, 7(1), 105–107Fig. 3 Scanning electron micrographs of (a) the gel, (b) sulfur crystals deposited in the gel and (c) the sintered structure of monoclinic GeS2 The reaction mechanism for the sol–gel processing of GeS2, Conclusions proposed by Melling,1 consists of the following reaction steps: The efficient preparation of GeS2 from a gel mixture of GeOx and GeSx is possible by conversion of GeO2 with sulfur OGeMOR+H2S=OGeMSH+ROH (5) simultaneously introduced into the toluene–ethoxide solution by the H2S flow during sol–gel processing. Oxidation of H2S OGeMSH+ROMGeO=OGeMSMGeO+ROH (6) by oxygen from the air in the presence of concentrated sulfuric acid (pH-2) yields elemental sulfur and water which were OGeMSH+HSMGeO=OGeMSMGeO+H2S( (7) transported into a toluene solution of germanium ethoxide by the H2S gas flow and homogeneously distributed. The heat- In reaction (7), formation ofOGeMSMGeO is favoured since treated product shows two types of bonds, GeMSMGe and H2S gas is easily removed from the reaction system.GeMSMSMGe. Therefore, the complete consolidation of GeSx gel and reduction of GeOx gel with sulfur occurs upon heat Characterization results for the gel after heat treatment inditreatment at 630 °C. cate that the product is monoclinic GeS2.From Table 1 it is evident that the agreement between d-spacings and 2h values This research was supported by CANMET, Fuel Processing of the heat-treated product and the standard is excellent, show- Laboratory, Western Research Centre, Devon, Alberta, ing that the product is primarily GeS2. Its structure consists of Canada. The authors are grateful for their financial and large crystals with no evident porosity at the grain boundaries.technical support. The IR spectrum of the heat-treated product indicates that, besides consolidation in the gel during heating, reduction of References GeO2 by elemental sulfur occurred as proposed by reaction (8): 1 P. J. Melling, Am. Ceram. Soc. Bull., 1984, 63, 1427. GeO2+3S=GeS2+SO2 (8) 2 A. B. Seddon, S.N. B. Hodgson and M. G. Scott, J. Am. Ceram. Soc., 1991, 26, 2599. 3 V. Stanic�, R. Mikula, T. H. Etsell and A. C. Pierre, J. Mater. Res., An absorption peak at 473 cm-1 in the IR spectrum of the 1996, 11, 364. gel was assigned to the SMS bond in elemental sulfur. This 4 P. N. Kumta and S. H. Risbud, in Ultrastruct Processing of peak is shifted to higher wavenumber (483 cm-1) in the heat- Advanced Materials, ed.D. R. Uhlmann and D. R. Ulrich, Wiley, New York, 1992, p. 555. treated product, indicating the presence of SMS bridging in 5 M. A. Sriram and P. N. Kumta, J. Am. Ceram. Soc., 1994, 77, 1381. the latter. Therefore, it can be concluded that, besides conver- 6 JCPDS powder diffraction file, card 36–1463. sion, during heat treatment, curing of GeS2 molecules due to 7 JCPDS powder diffraction file, card 8–247. the excess of sulfur occurred producing GeMSMSMGe 8 JCPDS powder diffraction file, card 27–238. connectivities. 9 Y. Kawamotoand C. Kawashima,Mater. Res. Bull., 1982, 17, 1511. 10 S. P. Mukerjee and S. K. Sharma, J. Am. Ceram. Soc., 1986, 69, 806. Semiquantitative analysis of the heat-treated gel shows that 11 J. Poucher, T he Aldrich L ibrary of IR Spectra, Aldrich Chemical the obtained product is homogeneous, with a Ge:S atom% Co., 3rd edn., 1981. ratio of 152.9. This result also indicates that an excess of sulfur was introduced into the gel during synthesis. Paper 6/04938I; Received 15th July, 1996 J. Mater. Chem., 1997, 7(1), 105–107 1

ISSN:0959-9428    

DOI:10.1039/a604938i

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Pyrolysis chemistry of polysilazane precursors to silicon carbonitride Part 1.—Thermal degradation of the polymers Djamila Bahloul,*a Mario Pereiraa and Corine Gerardinb† aL aboratoire de Mate�riaux Ce�ramiques et T raitements de Surface, URA CNRS 320, Faculte� des Sciences, 87060 L imoges, France bL aboratoire de Chimie de la Matie`re Condense�e, Universite� P. et M. Curie, 75252 Paris, France The cross-linking reactions and the decomposition processes of three different polysilazane precursors to silicon carbonitride have been investigated systematically in the course of pyrolysis to 1400°C.The polymer to ceramic conversion chemistry was studied for the following precursors: (ViSiHNH)n (VS), (ViSiHNMe)n (VNMS), [(ViSiHNH)0.5–(MeSiHNH)0.5]n (VS/MS), by means of thermogravimetry, mass spectrometry, solid-state NMR and IR spectroscopies, elemental analysis and X-ray diffraction.The influence of the structure and the polymer backbone (the nature and distribution of the functional groups) on the conversion chemistry was investigated. The occurrence of the cross-linking reactions, the ceramic yields and the abundance of the free carbon phase are related to the polymeric precursor structures.The development of organometallic polymers as preceramic functional groups (vinyl, SiH, NH) have been recognized to be effective in the construction of the polymer backbone. The materials for the synthesis of silicon carbide (SiC), silicon nitride (Si3N4 ) and silicon carbonitride (SiCN) ceramics has possible reactions resulting from the high reactivity of these precursors are hydrosilylation, transamination, vinyl polymer- recently received considerable attention.1–11 The main advantage of the conversion from polymers to ceramics is that it is ization, dehydrocoupling and redistribution reactions.These reactions, which are very common in organosilicon chemistry, possible to prepare readily ceramic fibrous shapes which are difficult to synthesize by traditional inorganic chemical pro- will be detailed later. Note that the initial C/Si ratio varied from 3 to 1.5 (see Table 5, later) in the oligosilazane compounds cesses.These polymers may also be used as binders for sintering powders and for infiltration of porous ceramic bodies.12–15 studied and the N/Si ratio was the same for all polymers, i.e.N/Si=1. One of the limitations of the process is the loss of volatile organometallic species, frequently observed by pyrolysis, which may drastically lower the ceramic yield. Cross-linking of the Polymer pyrolysis preceramic polymers before pyrolysis is necessary to improve The samples were heated in an Al2O3 tube fitted with a the ceramic yields.16 It may be performed either chemically, programmable temperature controller and a control flowmeter in the presence of a catalyst, or thermally if the polymer in nitrogen.To obtain more details on the mechanisms of is reactive enough.17–21 The conversion of an organomet- polymer pyrolysis, thermogravimetry (TG) studies (B60 allic precursor into a ceramic depends on different parameters Setaram, Ugine-Eyraud System, Caluire, France) were carried such as the molecular structure of the polymer and the pyro- out under the same conditions.The gases produced during lysis conditions (temperature, duration, atmosphere).22–25 polymer pyrolysis under flowing helium were identified in a Numerous reactions occur during the thermal treatment. It is continuous process using a quadrupole mass spectrometer of great value to determine the nature, rate of evolution and (MS; Quadrex 200, Leybold Heraeus, Les Ulis, France; 70 eV, total amounts of the major volatile products at various stages electron impact).For quantitative chemical analysis (CNRS, of the thermal conversion process. This type of information is Service Central d’Analyse) of the precursors and the pyrolysis needed to understand the thermal conversion process in order products, the silicon content was determined by ICP (inductive to control the chemical composition and microstructure of the coupled plasma) from an aqueous solution of sodium and/or final ceramic.Mass spectrometry (MS) in conjunction with potassium silicate resulting from the chemical attack of the thermogravimetry (TG) helps to provide some suggestions sample.Carbon and hydrogen contents were established by about the reaction mechanisms responsible for the mass losses. high-temperature combustion using IR spectroscopic analysis. This paper presents some correlations between the chemical Similarly, the oxygen content was determined using IR spectro- compositions and structures of polysilazanes and their derived scopic analysis of the carbon monoxide formed after pyrolysis ceramic products, based on the results obtained using a of the sample at ca. 3000 °C, whereas that of nitrogen was combination of TG, MS, IR and NMR spectroscopies, elemen- derived according to a thermal conductivity detection method. tal analysis and XRD techniques. IR spectroscopic analyses of the cross-linked products were performed according to the conventional KBr pellet technique Experimental using an FTIR spectrometer (Nicolet 5DX, Madison, WI).The XRD patterns (Cu-Ka; PW 1130 Philips, Eindhoven) were Polymer preparation obtained from the pyrolysis residues ground into fine powders. The synthesis and the structural analysis of the three organometallic precursors chosen for this investigation have been Results described previously.11,25,26 A summary of the synthesis and cross-linking conditions is given in Tables 1 and 2. Three Thermogravimetry and ceramic yields TG studies, at 10°C min-1 under a nitrogen flow, of the three † Present address: Laboratoire de RMN et Chimie du Solide, UMR CNRS 50, Universite� Louis Pasteur, Strasbourg, 67000 France.cross-linked polymers (Fig. 1) reveal that the VS precursor J. Mater. Chem., 1997, 7(1), 109–116 109Table 1 Synthesis of the different oligosilazanes sample theoretical formula synthesisa VS (SiViH-NH)n SiViHCl2+NH3, toluene VNMS (SiViH-NMe)n SiViHCl2+NMeH2, toluene VS/MS (SiViH-NH)x–(SiMeH-NH)y 0.5 SiViHCl2+0.5 SiMeHCl2+NH3 , toluene aAll at 0 °C. Table 2 Cross-linking and pyrolysis of the different polymers be of the same order of magnitude as those of the relative silicon losses (DSi/Sip) determined from elemental analy- cross-linking ceramic sis.This temperature domain corresponds to the distillation of sample Fa conditions yield (%) volatile silicon-containing compounds, as mentioned previously.21 VS 3 toluene, 110°C 87.6 VNMS 2 hexane, 70 °C with Pt4+ 63.6 catalyst Elemental analyses VS/MS 2.5 heating at 120 °C 79.8 The evolutions of chemical compositions of the pyrolysed without solvent or catalyst products issued from the VS, VS/MS and VNMS precursors aAverage number of reactive functions (vinyl, SiH, NH) per silazane at various temperatures are displayed in Fig. 2, 3 and 4, unit. respectively, and suggest the following remarks: the main changes occur between 400 and 800°C with an important loss of hydrogen and carbon; hydrogen is still present in the material even at high temperature.The amounts of silicon present before and after the pyrolysis of 100 g of precursor are given in Table 4. Aside from possible errors in the elemental analysis, the silicon loss of 21% during the pyrolysis of VNMS is not solely due to the distillation of oligomers, in contrast to VS and VS/MS for which the silicon loss is equivalent to the first mass loss determined by TG within the temperature range 25–200 °C.For the VNMS polymer, this additional silicon loss probably occurred in the second temperature range, beyond 600 °C. Indeed, a significant mass loss is recorded by TG, showing an important modification within the polymeric backbone.As suggested previously,17,26,27 the reactions responsible for the loss of Fig. 1 TG profiles of the different precursors pyrolysed under N2, ', VS; #, VS/MS; %, VNMS having the highest functionality (F=3) leads to the best ceramic yield (87%) at 1400 °C. Only a small mass fraction (Dm/mo=0.5%) is lost below 500 °C, revealing the high deg-linking and the thermal reactivity of this precursor.The decomposition behaviours of the VNMS and VS/MS polymers are fairly similar. Three main regions are present in the TG curves: (1) from 100 to 300 °C, an initial 6% mass loss is observed, followed by a plateau; (2) from 300 to 450 °C, an additional 3% mass loss occurs; (3) above 450 °C, an increasing slope of the curve is observed.The mass loss is considerably greater for the VNMS polymer (29%) compared to only 11% Fig. 2 Elemental analysis (atomic composition) of VS precursor and for VS/MS and 13% for the VS polymer. those of intermediate products isolated at different temperatures. %, It appears that the number, the nature and the distribution H/Si, #, N/Si; 2, C/Si. of the functional groups in the polymeric structures are important factors that determine the ceramic yield.From the data in Table 3, it can also be seen that the mass loss is influenced by the pyrolysis heating rate. The lowest ceramic yield is obtained for the fastest heating rate (60°C min-1). The values indicate that the major mass loss variations occur in the low temperature range (25–250°C).These first mass loss values appear to Table 3 Mass loss of VS/MS pyrolysed under N2 at various heating rates heating rate/ ceramic yield, first mass loss °C min-1 Y (%) (25–250°C) by TG DSi/Sipa (%) 1 84.1 3.6 3 10 79.8 6.1 6 Fig. 3 Elemental analysis (atomic composition) of VS/MS precursor 60 71.5 13.5 17 and those of intermediate products isolated at different temperatures. %, H/Si, #, N/Si; 2, C/Si.aSee Table 4. 110 J. Mater. Chem., 1997, 7(1), 109–116The compositions of the residues at 1400 °C, assuming that only equilibrium phases, SiO2 , SiC, Si3N4 and free carbon, are present in the final ceramics are reported in Table 5. These data show that the pyrolysis of polysilazane precursors leads to ceramic residues containing high free carbon levels.It can already be noticed that this calculation leads to a larger amount of free carbon in VNMS than in VS than in VS/MS. These results will be confirmed and detailed by quantitative NMR measurements in Part 2.30 Note that elemental analysis of ceramic products is difficult and poses problems attributed on the one hand to a possible incomplete combustion of the compounds and, on the other hand, the pyrolysis products are porous, absorb moisture and volatiles very readily, leading sometimes to misleading results.Fig. 4 Elemental analysis (atomic composition) of VNMS precursor and those of intermediate products isolated at different temperatures. IR spectroscopy %, H/Si, #, N/Si; 2, C/Si. The IR spectra of unpyrolysed polymers and those of solid Table 4 Evolution of the silicon content of the different precursors samples isolated at intermediate temperatures during the pyrolysed under N2 up to 1400°C decomposition of polymers VS, VS/MS and VNMS are displayed in Fig. 5, 6 and 7, respectively. VS VS/MS VNMS ceramic yield, Y (%) 87.6 79.8 63.6 VS. Fig. 5 displays the IR spectra of the VS precursor and Sip 39.47 43.11 32.97 of samples obtained at different temperatures.When the pyrol- YSic 39.32 40.53 26.01 ysis temperature is increased, a decrease in the intensity is DSi/Sip (%) 0.4 6.0 21.1 observed as well as a broadening of the absorption bands. The first mass loss (%) 0.6 6.1 5.2 spectrum of the sample pyrolysed at 250 °C indicates a decrease by TG (25–200°C) in the band intensities of vinyl groups at 3047, 1592 and DSi/Sip (%)=(Sip-Y Sic)/Sip ×100, where Sip is the silicon content in 1406 cm-1.The SiMH stretching band at 2135 cm-1 is less mass% (theoretical) of the starting precursor, Sic is the silicon content modified. At 500 °C, the absorption bands arising from NMH in mass% of the final ceramic, Y is the ceramic yield. (3400, 1170 cm-1), SiMH (2130 cm-1) and vinyl groups (3050, organosilicon species during pyrolysis are redistribution reactions involving the exchange of SiMN and SiMHbonds and/or two SiMN bonds.The atomic compositions of the samples heated at 1400 °C under nitrogen (determined by chemical analyses) are compared with the theoretical elemental analyses of the precursors in Table 5. It is shown that the N/Si ratios of the VS and VNMS samples did not decrease from the precursors (N/Si#1) to the final ceramics, indicating that the transamination reactions are negligible.11,28,29 The distillation of organosilicon species and exchange reactions should not modify the N/Si ratio.In contrast, a decrease in the nitrogen content of VS/MS is consistent with the release of a large amount of NH3 issued from transamination reactions, as will be confirmed later in this paper by the mass spectrometry results.A comparison between the C/Si ratios in the precursors and in the final ceramics shows that the presence in VNMS of a methyl group on the nitrogen atom (NMe) instead of a hydrogen atom (NH) in the oligosilazane does not lead to a significant increase in the final carbon content: C/Si=1.75 compared to C/Si=1.59 for pyrolysed VS.This result reflects the high lability of methyl groups bonded to nitrogen atoms. The partial replacement of vinyl groups by non-reactive methyl groups bonded to silicon atoms leads to a small decrease in the total carbon content: C/Si=1.59 for pyrolysed VS compared to C/Si=1.08 for pyrolysed VS/MS. This is equivalent to a total C loss of 20% for VS compared to 27% for VS/MS and suggests that C atoms from vinyl groups are less readily Fig. 5 Evolution of the FTIR spectra of VS under N2 at various temperatures. (a) Precursor; (b) 250°C; (c) 500 °C; (d) 1400 °C. evolved as gaseous products than Me groups. Table 5 Atomic compositions of the oligosilazanes and ceramic compositions after pyrolysis under N2 at 1400°C N/Si N/Si C/Si C/Si sample (initial) (final) (initial) (final) theoretical formula VS 1 0.98 2 1.59 55.6 Si3N4 , 15.2 SiC, 2.9 SiO2, 26.3 C VS/MS 1 0.86 1.5 1.08 54.3 Si3N4, 20.2 SiC, 8.5 SiO2 , 17.0 C VNMS 1 1.19 3 1.75 61.4 Si3N4, 1.1 SiC, 7.1 SiO2 , 30.4 C J.Mater. Chem., 1997, 7(1), 109–116 111Fig. 6 Evolution of the FTIR spectra of VS/MS under N2 at various temperatures. (a) Precursor; (b) 250°C; (c) 500 °C; (d ) 650 °C; (e) 850 °C; ( f ) 1400 °C. Fig. 7 Evolution of the FTIR spectra of VNMS under N2 at various temperatures. (a) Precursor; (b) 250 °C; (c) 500°C; (d) 650°C; (e) 850 °C; (f ) 1400°C. 1594, 1404 cm-1) are reduced considerably. As the temperature increases, the residual SiMH, NMH and CMH bonds are eliminated. The spectra show only broad absorption bands arising from SiMN and SiMC bonds in the 1200–600 cm-1 decrease significantly.At the same time, the SiMH stretch range. (2100 cm-1) and NMCH3 (1460 cm-1) deformation decrease. The spectrum of the sample at 500°C shows a weak and VS/MS. On heating the cross-linked precursor from 25 to broadened SiMH stretching band (2160 cm-1). It is interesting 250 °C (Fig. 6), there is a decrease in the intensity of the asym- to note that a broadened peak centred near 1600 cm-1 appears; metric CHNCH2 stretch at 3048 cm-1 and those of Si–vinyl this suggests that a new structure with unsaturated carbon is deformation (1408 cm-1) and CNC stretch (1595 cm-1).A formed. Above 500 °C, there remains only a very broad absorp- broadening and a slight decrease in the SiMH band tion band between 1200 and 600 cm-1, attributed to SiMC (2122 cm-1), as well as a broadening of the NMH deformation and SiMN bonds.band (1179 cm-1) are observed. At 500 °C, the SiMH band has decreased but remains intense, whereas the bands arising Mass spectrometry analysis from the vinylic groups have disappeared almost completely. Correlations between MS results and TG profiles help to The intensity of the band attributed to the SiMCH2MSi provide some suggestions about the mass loss mechanisms.deformation (1046 cm-1) increases with increasing tempera- The TG–MS analyses under helium of the three precursors ture, suggesting the insertion of methylene groups into the VS, VS/MS and VNMS are given in Fig. 8, 9 and 10(a)–(c) silicon network. The band at 1179 cm-1 of the NMH defor- respectively.mation decreases significantly in intensity. Note that the SiMCH3 deformation band at 1260 cm-1 is still present. This suggests that SiMCH3 groups are more stable than Si–vinyl VS. The TG–MS curve (Fig. 8) shows a very small mass loss (0.5%) up to 400 °C. This mass loss is probably due to groups up to 500 °C. Above 500 °C, the IR spectra indicate that most of the SiMH, NMH and CMH functionalities are the distillation of low molecular mass compounds which condense rapidly at low temperatures in the gas line.Thus, lost. By 1200°C, only bands between 1200 and 600 cm-1 characteristic of SiMC and SiMN bonds are present. they could not be observed by MS. From 200 to 300 °C some traces of NH3 (at m/z=17) were detected, indicating a low degree of transamination. Hydrogen starts to evolve only VNMS.As shown in Fig. 7, the IR absorption peaks at 3400 and 1200 cm-1 arising from NMH bonds indicate the presence above 300 °C. In the 400–1200 °C range, the mass loss increases considerably (13%). This domain is characterized by the main of this group in the structure of polymers. With increasing temperature to 250 °C, the intensities of the IR absorption evolution of methane (m/z=15), ethane, ethene (m/z=26, 27, 28) and hydrogen (m/z=2).bands assigned to vinyl groups (3047, 1591 and 1402 cm-1) 112 J. Mater. Chem., 1997, 7(1), 109–116Fig. 8 TG–MS analyses of VS pyrolysed under a helium flow. #, TG; m/z=2 (%), 15 ($), 17 ('), 26 (– –), 27 (—–), 28 (- - - -). Fig. 9 TG–MS analyses of VS/MS pyrolysed under a helium flow.#, TG; m/z=2 (1), 15 ($), 17 (%), 26 (– –), 27 (—–), 28 (- - - -), 41 ('). VS/MS. The TG–MS analysis of VS/MS is displayed in Fig. 9. On heating from room temperature to 200 °C, a mass loss of 6% is observed, corresponding to the distillation of low molecular mass oligomers (not detected by MS). Between 200 and 400 °C, a second mass loss (3%) is observed and at the same time a large release of ammonia at m/z=17 is detected Fig. 10 TG–MS analyses of VNMS pyrolysed under a helium flow. by MS, which indicates the occurrence of transamination (a) #, TG; m/z=15 ($), 30 (—–), 31 (- - - -), 85 (%), 86 (– –). reactions [eqn. (1), (2)]. (b) m/z=2 (#), 15 ($), 27 (—–), 28 (- - - -), 41 (%). (c) m/z=15 ($), 42 (—–), 43 (#), 44 (- - - -), 56 (%).fractions in this case, the TG–MS results are displayed on three separate graphs in order to show a clearer presentation of the evolution curves. From 80 to 400°C the loss observed by TG corresponds to the escape of (HSiViNHMe)+ at m/z= 85and (SiViHNHMe)+ at m/z=86, arising from SiViH2NHMe and probably other oligomers which condense in the cold parts of the apparatus. In parallel the N-methylamine at m/z= 30, 31 is detected, arising from transamination of the SiNMeH groups corresponding to the ends of the chains of linear oligomers present in the product as observed by IR analysis. A lesser release of hydrogen (m/z=2), which indicates dehydrogenation reactions, is also detected.In this temperature Small amounts of hydrogen (m/z=2), ethane and ethene at m/z=27, 28 are also detected.In the 400–800 °C range, the range, the plateau observed by TG can be explained by reactions which do not lead to the release of a gas, such as the mass loss increases sharply (25%). The main gases detected by MS are hydrocarbons such as methane (m/z=15), ethylene, hydrosilylation reaction via vinyl and SiMH groups and/or polymerization reactions via vinyl polyaddition.These reac- ethane (m/z=27, 28) and a significant amount of propene (m/z=41, 42), butane and butene (m/z=56, 41). Under the tions will be discussed later. Likewise, some volatile species with smaller fractions are also detected [methane (m/z=15), same experimental conditions and with the same initial mass of product, the intensities of the signals at m/z=41 and 56 are ethane, ethene (m/z=26, 27, 28)].Above 400 °C, the main gases produced are methane, hydrogen and, to a lesser extent, higher for VNMS than for the other polymers. A second gaseous evolution at m/z=30, 31 with a maximal release speed ethene and propene (m/z=41). At higher temperature, only hydrogen is detected. at 500 °C is observed. It is difficult to assign this peak in the same way as in the first temperature range, i.e.to the Nmethylamine arising from transamination of the ends of the VNMS. The TG–MS analysis of VNMS is shown in Fig. 10(a)–(c). Because of the escape of numerous different chains, owing to the small number of NMH bonds present in J. Mater. Chem., 1997, 7(1), 109–116 113first stage (up to 400 °C), further cross-linking proceeds by various reactions.Likewise, the loss of low molecular mass oligomers occurs; (2) in the second stage (400–800 °C) the SiMC, SiMHand CMHbonds are broken and small molecules, mostly hydrocarbons and hydrogen, are evolved. This step is commonly called a mineralization step and consists of the organic–inorganic transition; (3) at temperatures higher than 800°C, a mineral ceramic material is obtained consisting of free carbon and an amorphous SiCxNy silicon network. As the temperature reaches 1450 °C, crystallization is induced.Cross-linking The difference in ceramic yields between the three samples stems primarily from the difference in the degrees of crosslinking occurring during the first step up to 400 °C. On the basis of the IR spectroscopic data, the chemical analysis of the Fig. 11 XRD patterns of the residues resulting from the pyrolysis at intermediate product at 250 °C and the amount of silicon lost 1400 °C of the three precursors. (a) VS; (b) VNMS; (c) VS/MS. during pyrolysis which is equivalent to the first mass loss, there is a clear similarity between the composition of the volatile organosilicon species and the starting precursor.Many kinds of reactions can account for the departure of volatile products. At the beginning, the escape of organosilicon species can be explained on the one hand by the distillation of low molecular mass oligomers when cross-linking is inadequate and, on the other hand, by depolymerization reactions such as exchange of SiMN bonds or SiMN and SiMH bonds, or SiMN and NMH bonds.Exchange of SiMN bonds according to eqn. (5) would lead to the formation of silazane oligomers having the same structural units as the precursor, whereas the exchange between SiMN and SiMH bonds [eqn. (6)] would lead to volatile silanes as suggested in the case of the thermolysis of the VNMS precursor [eqn. (3), (4)]. Fig. 12 Comparison of XRD patterns of pyrolysed products under nitrogen at 1450°C, 48 h.(a) VS/MS; (b) VNMS. $, a-Si3N4; %, b-Si3N4 . the product. The most rational hypothesis for the formation of such fragments is redistribution involving exchange of SiMN and SiMH bonds which may lead to the volatile silane CH2NCHSiH3 [eqn. (3), (4)] with abundant ions SiH2+ at m/z=30 and SiH3+ at m/z=31. 2 N Si N Me H Me N SiH2 Me Si (N)3 Me (3) + The exchange of SiMN and NMH bonds, called the trans- N SiH2 Me N Si N Me H Me SiH3 Si (N)3 Me + (4) + amination–condensation reaction [eqn.(1), (2)] involves the formation of new SiMN bonds and the release of NH3. Note Note that this redistribution reaction leads to the formation that the escape of NH3, detected by MS analysis, indicates of trisilylated nitrogen atoms, NSi3. that the transamination reaction occurs between 200 and The departure of these silanes would explain the loss of 500°C whatever the precursor.The presence of methyl groups silicon observed by chemical analysis (Table 4). on the nitrogen atoms (NMCH3) instead of NH groups pre- Beyond 800 °C, the release of hydrogen continues. vents the transamination reaction, as in the case of the VNMS precursor.In this case, the N-methylamine observed by MS X-Ray diffraction analysis corresponds only to transamination of the ends of chains. As shown in Fig. 11, the diffraction patterns of the three As mentioned already, the transamination reaction allows polymers pyrolysed under nitrogen up to 1400 °C do not the formation of trisilylated nitrogen atoms (NSi3). Such NSi3 exhibit any significant diffraction peaks.Thus, these products sites can also result from dehydrogenation reactions between consist of amorphous structures. The development of crystal- SiMH and NMH bonds [eqn. (7)] and thus can explain the linity was observed in samples fired at higher temperature27 release of hydrogen observed by MS during the thermolysis of (Fig. 12). Note that no SiC phase was observed by XRD.the polymers below 400 °C. Discussion From the data obtained using different analysis techniques, the pyrolysis process of the organometallic precursors may be considered to take place in three consecutive stages: (1) in the 114 J. Mater. Chem., 1997, 7(1), 109–116Dehydrogenation of SiMH and NMH contiguous bonds Si(C sp3)2N2 sites can be formed from duplicate hydrosilylation as shown in reactions (14) and (15). can also explain the departure of hydrogen through an intermediate silylimine as suggested previously.17,28 Then, the silylimine reacts by insertion into NMH and/or SiMH bonds following reactions (8), (9) and (10), leading to the formation of NSi3 sites.–HN Si H –HN Si Si NH– NH– –HN –HN Si(C sp3)2N2 2 (14) –HN Si H –HN –HN Si –HN Si NH– H NH– Si NH– NH– Si(C sp3)2N2 3 (15) Thus the resulting preceramic network is more stable towards redistribution/exchange of bonds than the SiMN skeleton, and this seems to be an important inhibiting factor in the depolymerization reactions.The small mass loss, the results of elemental analysis and the few traces of ammonia observed by MS during the thermolysis of VS reinforce this Si N Si H Si Si N H Si Si + (10) Insertion into Si—H bond: suggestion.In the cases of VNMS and VS/MS, the departure of gaseous hydrocarbons (ethane, methane) detected by MS analysis in this temperature range (T<400 °C) could be assigned to 1–2 migration of hydrogen with formation of silylene species29 according to reaction (16) rather than via the formation of Hydrogen was also postulated to arise from reaction (11), radical groups R and H .producing silene species30,31 which could insert into SiMH bonds to form new �SiMCH2MSi� bridges. Thus, carbon atoms can be incorporated into the silicon network. This reaction is consistent with our 29Si NMR observations, detailed in the following paper.30 It is shown that the silicon carbonitride SiCnN4-n sites are more numerous in VS/MS These silylene species could then, in turn, insert into other than in VS and VNMS ceramic products.SiMH, CMH or NMH bonds producing SiMSi bridges (less In addition, the higher mass retention of VS relative to those probable), SiMCH2MSi or MSiM(N)3Nbridges respectively. of VS/MS and VNMS precursors during thermolysis is attri- The combined results for the samples heated to 400°C show buted to the greater degree of branching and cross-linking that a lower cross-linking degree corresponds to more numer- issued from the conversion of vinyl groups to saturated hydro- ous redistribution reactions.The various losses are consistent carbon functionalities according to eqn. (12) and (13). with the occurrence of exchange reactions (transamination, exchange of SiMN bonds with SiMH or SiMN ones) which are facilitated in VS/MS compared to VS.It is clear that the presence of SiMMe groups alternating with SiMVi ones in the precursor backbone limits the occurrence of cross-linking reactions involving Vi groups (hydrosilylation, Vi polyaddition) and thus gives less rigidity to the silicon network below 400°C.These conclusions are in agreement with the results previously reported by Choong Kwet Yive et al.26 Mineralization step A considerable change occurs within the temperature range 400–800 °C. The mass loss observed, whatever the starting precursor, is mainly assigned to an evolution of hydrocarbons and hydrogen resulting from broken SiMC and NMC bonds and also SiMH, CMH and NMH bonds.A detailed study of these decomposition reactions, which involve radical mechan- This is supported, on the one hand, by the IR study which indicates the formation of more saturated CH2 groups isms, has been published previously.25 A comparison between the evolution during pyrolysis of the three precursors indicates (2900 cm-1) and a decrease in the number of vinyl groups (3048, 1595 and 1408 cm-1) and SiMH bonds (2120 cm-1) a significant difference in where the carbon atoms bind to the Si skeleton atoms.Indeed, the high molecular mass hydro- and, on the other hand, by the 13C CPMAS NMR characterization which reveals that substantial amounts of vinyl groups carbons detected by MS analysis, such as propene, butane and butene, are only obtained in the case of the VNMS precursor.were consumed (resonance centred at d 135) in favour of essentially aliphatic carbon atoms (resonance centred at This can be explained by the presence of the methyl group on nitrogen atoms which limits cross-linking processes such as d 10–30) below 500°C in the three routes. The 29Si MAS NMR spectra detailed in Part 2 confirm these results;30 the the hydrosilation reaction.This reaction is limited because of the steric hindrance due to methyl groups. Consequently, the silicon environments are described as a distribution of Si(C sp3)(C sp2)N2, Si(C sp3)HN2 which can be formed degree of polymerization via polyaddition of vinyl groups [reaction (13)] increases, which produces bridged or ring-type by simple hydrosilylation or vinyl polymerization, whereas J.Mater. Chem., 1997, 7(1), 109–116 115hydrocarbon products as shown schematically in eqn. (17) between the structure of the polysilazane precursor and the ceramic yields and the global atomic compositions of the final and (18). ceramics. The results reveal phase separation during the pyrolysis process, the formation of a structure with an SiMNMC backbone and the separation of free carbon.The stability of SiMC bonds and the stable incorporation of carbon into the SiMNMC network during pyrolysis appear to depend strongly on the substituents on the Si and N atoms in the precursor backbone. The differences between the thermochemical evolutions of the three preceramic polymers will be clarified further in the following paper on the basis of quantitative 29Si MAS NMR studies which examine SiM(C,N) bonds directly.References 1 T. F. Cooke, J. Am. Ceram. Soc., 1991, 74, 2959. 2 S. Yajima, J. Hayashi and M. Omori, Chem. L ett., 1975, 9, 931. These six-membered rings formed with C atoms may be 3 (a) Y. Hasegawa, M. Iimura and S. Yajima, J.Mater. Sci., 1980, 15, precursors for the free carbon phase.Such structures lead to a 720; (b) Y. Hasegawa and K. Okamura, J. Mater. Sci., 1983, 18, low carbon content in the silicon network. This agrees with 3633. the low fraction of silicon carbonitride sites observed in the 4 R. West, in Ultrastructure Processing of Ceramics, Glasses and 29Si NMR spectra of VNMS heated at 600 °C and the appear- Composites, ed. L. L. Hench and D.R. Ulrich, Wiley Interscience, New York, 1984, p. 235. ance of C sp2 sites shown by 13C CPMAS NMR results. The 5 D. J. Carlson, D. J. Cooney, S. Gauthier and D. J.Worsford, J. Am. limitation of the hydrosilylation reaction due to the presence Ceram. Soc., 1990, 73, 237. of NMe groups would explain the higher free carbon content 6 M. Arai, S. Sakurada, T. Isoda and T. Tomizawa, Polym.Prepr. in VNMS compared to VS, if we accept that Me groups (Am. Chem. Soc., Div. Polym. Chem.), 1987, 28, 407. bonded to N atoms are evolved as a gaseous product. The 7 O. Funayama, T. Isoda, T. Suzuki and Y. Tashiro, Polym. Prepr. occurrence and abundance of free carbon in the three materials (Am. Chem. Soc., Div. Polym. Chem.), 1991, 32, 542. 8 G. E. Legrow, T. F. Lim, J. Lipowitz and R.S. Reaoch, Am. Ceram. will also be discussed in the following paper.30 Soc. Bull., 1987, 66, 363. 9 D. Seyferth and G. H. Wiseman, in Ultrastructure Processing of Crystallization Ceramics, Glasses and Composites, ed. L. L. Hench and D. R. Ulr Wiley Interscience, New York, 1984, p. 265. From 800°C, after the organic-to-inorganic transformation, 10 D. Bahloul, M. Pereira and P.Goursat, Ceram. Int., 1992, 18, 1. the three different materials are constituted of several structures 11 A. Lavedrine, D. Bahloul, P. Goursat, N. S. Choong Kwet Yive, (free carbon, silicon carbonitride, silicon nitride) as suggested R. Corriu, D. Leclercq, H. Mutin and A. Vioux, J. Eur. Ceram. by the nature of local environments which are observed by Soc., 1991, 8, 221.NMR (detailed in Part 2).30 The only change in chemical 12 R. R. Wills, R. A. Markle and S. P. Mukherjee, Am. Ceram. Soc. composition is the decrease in H/Si ratio with increasing Bull., 1983, 62, 904. 13 E. Bouillon, D. Mocaer, J. F. Villeneuve, R. Pailler, R. Naslain, temperature, which is confirmed by the release of hydrogen. M. Monthioux, A. Oberlin, C. Guimon and G. Pfister, J.Mater. XRD studies of the residues issued from pyrolysis under Sci., 1991, 26, 1517. nitrogen at 1400 °C for 1 h indicated the completely amorphous 14 K. Sato, T. Suzuki, O. Funayama and T. Isoda, J. Ceram. Soc. Jpn., structure of the silicon carbonitride ceramics (Fig. 11). With 1992, 100, 444. increased temperature and holding time (pyrolysis at 1450 °C 15 K. S. Mazdiyasni, R.West and L. D. David, J. Am. Ceram. Soc., for 48 h), the development of a crystalline phase corresponding 1978, 61, 504. 16 K. J. Wynne and R. W. Rice, Annu. Rev. Mater. Sci., 1984, 14, 297. to a-Si3N4 was observed (Fig. 12). Compared to other routes, 17 P. Peuckert, T. Vaahs and M. Bru�ck, Adv. Mater., 1990, 9, 398. the VNMS-derived product seems to be more crystallized in 18 D.Seyferth and G. H. Wiseman, J. Am. Ceram. Soc., 1984, 67, a- and b-Si3N4 phases. The high free carbon content, lost by C132. the silicon network, does not prevent the crystallization of 19 G. T. Burns, T. P. Angelotti, L. F. Hanneman, G. Chandra and Si3N4 phases. J. A. Moore, J.Mater. Sci., 1987, 22, 2609. Note that, beyond 1400 °C, the stability and the crystallis- 20 B. Kanner and R.E. King III, Adv. Chem. Ser., 1990, 224, 607. 21 Y. D. Blum, G. A. Mc Dermott and A. S. Hirschon, in Inorganic ation of the SiCN phase depends on the nature of the pyrolysis and organometallic oligomers and polymers, Proc. 33rd IUPAC atmosphere. As shown in previous work,27 the decomposition Symp. Macromolecules, ed. J. F. Harrod and R. M. Laine, Kluwer under argon of the amorphous SiMNMC material yields a Academic Publishers, Netherlands, 1991, p. 161. pyrolytic residue almost totally free of nitrogen but enriched 22 W. H. Atwell, Adv. Chem. Ser., 1987, 224, 593. by an SiC phase, involving the formation of gaseous species 23 G. T. Burns and G. Chandra, J. Am. Ceram. Soc., 1989, 72, 333. (N2, Si) and also oxygen-based species31 (SiO, CO). Oxygen 24 R. J. P. Corriu, D. Leclercq, P. H. Mutin and A. Vioux, Chem. Mater., 1992, 4, 711. contamination mainly comes from exposure to the air during 25 D. Bahloul, M. Pereira, P. Goursat, N. S. Choong Kwet Yive handling. In contrast, during the heat treatment under nitrogen and R. J. P. Corriu, J. Am. Ceram. Soc., 1993, 76, 1156. the decomposition of the SiMNMC structure leads to Si3N4 26 N. S. Choong Kwet Yive, R. J. P. Corriu, D. Leclercq, P. H. Mutin and free carbon. Si3N4 formation can also result from a and A. Vioux, Chem.Mater., 1992, 4, 141. reaction involving SiO and nitrogen according to the following 27 D. Bahloul, M. Pereira and P. Goursat, J. Am. Ceram. Soc., 1993, equation:31 3SiO+2N2�Si3N4+3/2 O2. 76, 1163. 28 Y. D. Blum, K. B. Schwartz and R. M. Laine, J. Mater. Sci., 1989, It seems that the decomposition of the SiMNMC(O) phase 24, 1707. is temporarily impeded under a nitrogen atmosphere in com- 29 B. J. Aylett, Organomet. Chem. Rev., 1968, 3, 151. parison with argon. This nitriding process will be discussed in 30 C. Ge�rardin, F. Taulelle and D. Bahloul, J. Mater. Chem., follow- detail in a future paper.32 ing paper. 31 D. Mocaer, PhD Thesis, University of Bordeaux, 1991. 32 D. Bahloul, M. Pereira and P. Goursat, J. Am. Ceram. Soc., 1996, Conclusion to be submitted. The combination of different analysis techniques has led to the following conclusions: there exist strong relationships Paper 6/03165J; Received 7thMay, 1996 116 J. Mater. Chem., 1997, 7(1), 109&nda

ISSN:0959-9428    

DOI:10.1039/a603165j

出版商:RSC    

年代:1997

数据来源: RSC

摘要:

Pyrolysis chemistry of polysilazane precursors to silicon carbonitride Part 2.†—Solid-state NMR of the pyrolytic residues Corine Ge�rardin,*a‡ Francis Taulellea‡ and Djamila Bahloulb aL aboratoire de Chimie de la Matie`re Condense�e, Universite� P. et M. Curie, 75252 Paris, France bL aboratoire de Mate�riaux Ce�ramiques et T raitements de Surface, URA CNRS 320, Faculte� des Sciences, 87060 L imoges, France The chemistry of pyrolytic conversion has been studied for three polysilazanes, (ViSiHNH)n, (ViSiHNMe)n and [(ViSiHNH)0.5–(MeSiHNH)0.5]n, precursors to silicon carbonitride ceramics. 13C and 29Si MAS and CPMAS NMR spectroscopies were used to clarify the processes leading to the formation of the silicon-based mineral network as well as the segregation of a free carbon phase.The assignment of 29Si NMR signals corresponding to SiCnN4-n sites was essential to follow the number of SiMC and SiMN bonds that are cleaved or formed. It was shown that at the organic–mineral transition temperature (ca. 900 °C) the final amount of free carbon as well as the final composition of the silicon-based network were already reached. Above this temperature, redistribution reactions around silicon atoms inside the amorphous silicon-based matrix take place in order to favour nitrogen-rich environments, i.e.crystallized Si3N4 regions. Above 1400 °C, all ceramics contain a similar amorphous silicon carbonitride structure, whose composition is close to SiN0.85C0.35 and which coexists with a crystallizing Si3N4 phase. Finally, the relative amounts of the three possible final carbon states, at 1400 °C, i.e.gas products, carbon incorporated in the silicon-based network and free carbon, could be related to the nature of the carbon-containing substituents in the precursor backbones and to the occurrence of the cross-linking reactions below the mineral transition temperature. Polysilazane polymers whose backbones consist of alternating 29Si NMR spectroscopy can lead to quantitative studies of local Si sites.The fractions of SiCnN4-n environments in the SiMN bonds with carbon-containing substituent groups are widely used as precursors of silicon carbonitride ceramics.1–8 pyrolytic residues can be obtained and those values can be used to estimate a chemical composition of the mineral silicon- The conversion chemistry of these polymeric precursors is still not completely understood.The derived ceramics obtained at based network. The purpose of this paper is to show how NMRcharacteriza- 1400 °C are often formed with three phases: a silicon carbide or nitride phase, an amorphous silicon carbonitride structure tions of the solid residues associated with data obtained previously from chemical analyses, IR spectroscopy, TG, MS and a free carbon phase.The thermomechanical properties of the ceramics depend largely on the proportions of these phases analysis and XRD21 can provide a better description of the major chemical transformations involved in the pyrolytic con- and a goal is to be able to predict the amount and the nature of the different combined structures knowing the backbone version.It will show how local structural details can give an insight into the chemical bonding and also into the chemical structure of the precursor. To achieve this objective, one must elucidate the various chemical transformations of the precer- compositions of the structures present in the multiphased materials at all intermediate temperatures. Special emphasis amic polymer in the course of pyrolysis. The difficulties arise from the characteristics of the intermediate solid residues which will be given to the influence of the nature of the precursor on the final ceramic structure.The polymer to ceramic conversion cannot be completely characterized by conventional solid-state analysis methods such as X-ray diffraction techniques.The chemistry was thus studied for three precursors containing various silazane units, with the same N/Si ratio but different lack of long-range order in the ceramics obtained below 1400 °C severely limits the investigation by XRD. The most C/Si ratios: (ViSiHNH)n (VS), (ViSiHNMe)n (VNMS) and [(ViSiHNH)0.5–(MeSiHNH)0.5]n (VS/MS). common investigation techniques used to characterize the pyrolytic conversion are thermogravimetry (TG), mass spectrometry (MS), IR spectroscopy and chemical analyses.4–12 Experimental These techniques give insights into the global changes in the materials; they answer the questions: what are the gas-phase The synthesis and curing step of the starting oligosilazanes products? When are the gaseous products released and in what have been reported elsewhere.8,10,21,22 The three main func- amounts? But very little information is generally obtained on tionalities present in the polymers are the NMH, SiMH, and the chemical bondings present in the solid inorganic intermedi- Vi (vinyl) groups which can give rise to the typical thermal ates and in the final amorphous ceramics.cross-linking reactions: transamination, hydrosilylation, Few studies centred on the NMR characterizations of the dehydrogenation and vinyl polyaddition.These reactions were solid-phase intermediates have been reported.13–20 Solid-state detailed in other papers.8,21 Note that the VS polymer contains NMR studies of the pyrolytic residues provide a description the three reactive functions per silazane unit; in VNMS, the of local arrangements around atoms. 13C NMR spectroscopy non-reactive NMe group replaces the NH function, and in is very useful for examining the evolutions of C sp3 and C sp2 VS/MS, half of the Vi functions are replaced by Me groups. environments. It can indicate when carbon atoms leave the The precursors were heated under a nitrogen flow at a silicon-based network and segregate as a free carbon phase.pyrolysis temperature in the range 250 to 1450 °C and were kept at that temperature for 1 h.21 The pyrolysis products were analysed by various techniques such as IR, XRD and chemical † Part 1, preceding paper. analyses and the results were detailed in the preceding paper ‡ Present address: Laboratoire de RMN et Chimie du Solide, UMR CNRS 50, Universite� Louis Pasteur, Strasbourg, 67000 France.(Part 1).21 The atomic ratios C/Si and N/Si of the pyrolytic J. Mater. Chem., 1997, 7(1), 117–126 117residues at 1400°C, which were obtained by chemical analyses, are indicated in Table 1. The NMR spectra were recorded on a Bruker MSL 400 spectrometer, operating at 100.62 MHz and 79.5 MHz, for 13C and 29Si nuclei respectively. Magic angle spinning (MAS) NMR spectra were acquired at a rotation frequency of 4 kHz.The 29Si one-pulse experiments required a recycle delay varying from 20 to 60 s with a 25° pulse angle. Some spectra were recorded with a recycle delay of up to 1 h and showed no difference in relative intensities from those of spectra acquired with a recyle delay of 1 min. The 29Si MAS NMR spectra could be quantitatively analysed, they were deconvoluted using the Winfit Bruker software program.Lorentz–Gaussian peaks were used to fit the experimental spectra; the three parameters, position, linewidth and amplitude, could be varied or fixed. The 13C MAS NMR spectra were acquired with a recycle delay of 60 s and a pulse angle of 20°. The 13C repetition times Fig. 1 NMR spectra of the VS/MS residues at different pyrolysis were not optimized in order to gain quantitative 13C MAS temperatures. A, 13C MAS (without cross polarization): (a) precursor, NMR spectra.A contact time of 1.5 ms and a repetition time (b) 250°C, (c) 500 °C, (d) 650 °C; B, 13C CP MAS (with cross polarization): (a) precursor, (b) 250°C, (c) 500°C, (d) 650 °C, (e) 850 °C. of 10 s were used for both 13C and 29Si cross-polarization (CP) MAS NMR spectra.In 29Si and 13C NMR, the compound used as chemical shift reference was tetramethylsilane. For all spectra, between 600 and 1000 scans were accumulated. a decrease of the peak may come from the loss of the methyl group through the evolution of methane; secondly, the loss of resolution may reflect the linking leading to Results progressively less protonated C sites such as CH2Si2 bridges or CHSi3 environments as the temperature increases.In the 13C NMR spectroscopy latter case, carbon insertion into the silicon-based network can The chemical shifts of the signals present in the 13C NMR occur. The 13C NMR resonances corresponding to C sp3 sites spectra lie between d 0 and 200. Signals from d 0 to 50 are are still observed up to 650°C under CPMAS conditions and due to carbon atoms in sp3 configuration, and signals from up to 500 °C under MAS conditions.At temperatures higher d 100 to 200 to carbon atoms in sp2 configuration. More than 850 °C, the number of protons close to carbon atoms is precisely, between d 0 and 25–30 appear signals from too low to obtain a sufficient proton magnetization transfer SiMC(sp3) sites; at d ca. 30 signals from NMC(sp3) sites; and towards 13C nuclei and cross-polarization is inefficient. The at d 135 signals from CNC sites in vinyl groups. The very signal-to-noise ratio of 13C CPMAS NMR spectra at 650 and broad signal centred at d 110 observed in 13C MAS NMR 850°C is rather low, but no significant signal corresponding spectra without cross-polarization (indicated with an asterisk to a still-protonated graphitic-type carbon structure can be on the figures), is due to carbon atoms present in the probe, detected in these samples. and so it does not represent the sample.VS. The 13C, 29Si, 1H, 15N and 14N NMR spectra of the VS VS/MS. Fig. 1A and B show the 13C MAS and CPMAS precursor in solution in C6D6 have been detailed elsewhere.13 NMR spectra of VS/MS pyrolysates.The spectra of the It was shown that 15% of the vinyl groups already transformed precursor show the presence of vinyl groups (signal at d 135) into C sp3 sites during the curing step. Fig. 2A shows the 13C and of SiMCH3 groups (at d ca. 0). The 13C CPMAS spectrum CPMAS NMR spectra of VS residues pyrolysed from 500 to enhances the resonances between d 10 and 30 characteristic of 900°C.At 500°C, no signal from vinyl groups can be observed SiMCHn (n=1 and/or 2) sites; the 1H�13C polarization trans- but a broad signal at d 9 is present, it is due to the C sp3 sites fer is indeed more efficient in these groups compared to the formed from the transformation of vinyl sites. This resonance case of mobile methyl groups.These sites are the result of the does not change much up to 600 °C. Broad and weak signals transformation of vinyl groups by polyaddition or hydrosilyl- appear at low field; they are characteristic of protonated CNC ation leading to SiMCMCMSi or SiMCMSi bridging groups. bonds and are certainly due to polyaromatic compounds which The transformation of vinyl groups is shown to take place are precursors of a free carbon phase.The occurrence of such mainly below 250 °C and it continues up to 500 °C; at this carbonaceous units was discussed in a previous paper.21 Vinyl temperature, all vinyl sites have disappeared. Between 250 and polymerization can lead to ring-type hydrocarbon species; 500 °C, a major rearrangement in the C sp3 environments is then, SiMC cleavage can give rise to six-membered rings observed, the carbon sites newly formed from vinyl groups containing C sp2 atoms.At 900 °C, it can be seen that C sp3 (signal centred at d 30) change considerably. The loss of sites are present in a minor amount compared to the abundant resolution of the peak representing the methyl function C sp2 atoms.The very broad signal centred at d 120 is due to (d ca. 0) may be the result of two different phenomena: first, still protonated C atoms such as CMCHNC. VNMS. Fig. 3A and B show the 13C MAS and CPMAS Table 1 N/Si and C/Si atomic ratios in the precursors (theoretical) and in the ceramics at 1400 °C (determined by chemical analyses) NMR spectra of VNMS pyrolysates. The weak signal at d 135 in the 13C CPMAS NMR spectrum shows that most vinyl global material elements were consumed at 250 °C, they transformed into starting silazane at 1400 °C aliphatic carbon sites whose signal appears at d ca. 10. A large (theoretical) chemical analyses) number of carbon atoms in NMMe groups (at d 30) are still sample N/Si C/Si N/Si C/Si present at this temperature, but most are transformed between 250 and 500 °C; the consumption of NMMe bonds still con- VNMS 1 3 1.19 1.75 tinues up to 650 °C.At 650 °C, a broad signal centred at d 130 VS 1 2 0.98 1.59 appears. It presents a large chemical shift anisotropy; several VS/MS 1 1.5 0.86 1.08 spinning sidebands separated by the spinning speed frequency 118 J. Mater. Chem., 1997, 7(1), 117–126Fig. 2 NMR spectra of the VS residues at different pyrolysis temperatures.A, 13C CP MAS: (a) 500°C, (b) 600 °C, (c) 900°C; B, 29Si MAS: (a) 500 °C, (b) 600 °C, (c) 900 °C, (d) 1200°C, (e) 1300°C, (f ) 1400 °C, (g) 1450°C; C, 29Si CP MAS: (a) 500 °C, (b) 600°C. are observed. These resonances do not appear in spectra without cross-polarization and are due to protonated C sp2 atoms. This observation shows the presence of abundant C atoms in an aromatic carbon structure.These structures may be formed by breaking of SiMC bonds, as explained earlier. From 650 to 850°C, the signals due to C sp3 sites disappear and the resonances due to C sp2 sites become weaker; the main reason for this is the progressive deprotonation of all carbon atoms leading to a less efficient {1H}-13C cross-polarization.By 850 °C, a complete loss of sp3 carbon occurs in the 13C CPMAS NMR spectrum of VNMS, in contrast to the case of VS pyrolytic residues for which a C sp3 signal is still observed at 900 °C. Fig. 3 NMR spectra of the VNMS residues at different pyrolysis temperatures. A, 13C MAS (without cross polarization): (a) 250 °C, 29Si NMR spectroscopy (b) 500°C, (c) 650 °C; B, 13C CP MAS (with cross polarization): (a) 250 °C, (b) 500 °C, (c) 650 °C, (d) 850°C. 29Si NMR spectra were registered under MAS conditions with and without cross-polarization. The comparison between both types of acquisitions was helpful for the assignment of peaks; The method to determine 29Si NMR chemical shifts corresponding to SiCnN4-n first coordination spheres was already it was essentially used to indicate strong 1H–29Si dipolar couplings which arise from direct SiMH bonds. 29Si CPMAS developed elsewhere,23,24 it will not be detailed here. The chemical shifts of SiCnN4-n sites were calculated using the NMR spectra were not acquired for samples pyrolysed above 800 or 900 °C because the hydrogen contents were too low at partial charge model24 and the classical theory of nuclear shielding.It was shown that replacing a carbon atom by a these temperatures. In the spectra analysis, only SiHxCyNz sites will be con- nitrogen atom in the Si environment does not lead to a regular chemical shift variation. The shape of the curve of d(29Si) sidered. The oxygen amount in the materials is indeed small. As was shown in Part 1, the chemical analyses indicated less (SiCnN4-n) vs.n is mainly governed by the nature of the first coordination sphere. The nature of the second coordination than 5 atom% of oxygen. As was also discussed in Part 1, the oxygen contamination occurs during handling and the oxygen sphere mainly shifts that curve to lower or higher field depending on the substituents on the C and N atoms. This is atoms are then present in the amorphous SiMCMN(O) structure.The low oxygen content leads to a distribution of different clearly observed in Fig. 4 which represents the chemical shifts of two series:the first one considers SiMen(NMe2)4-n molecules SiCxNyOz sites (with z=1 or 2) which exist in too small an amount to be identified in the NMR spectra as peaks dis- and the second one presents SiCnN4-n environments with C atoms being CSi4 sites (as in SiC) and N atoms NSi3 sites (as tinguishable from the main SiCnN4-n ones.As a consequence, the oxygenated sites cannot be taken into account in our in Si3N4). In a non-protonated silicon carbonitride structure with only Si, C (as CSi4 sites) and N (as NSi3 ) atoms (which spectr analysis and it would be too hazardous to discuss the oxygen evolution from the NMR results.is expected at high pyrolysis temperatures such as 1400 °C), J. Mater. Chem., 1997, 7(1), 117–126 119Fig. 4 29Si NMR chemical shifts corresponding to SiCnN4-n sites as a function of n for two series: SiMen(NMe2)4-n molecules ($) and SiCnN4-n sites hacing CSi4 and NSi3 environments (%) the chemical shifts of SiCnN4-n sites were calculated to be: d -15 (n=4), -10 (n=3), -19 (n=2), -34 (n=1) and -49 (n=0) with a precision of ±2 ppm.The exact positions of the peaks corresponding to these environments are then experimentally adjusted while fitting the spectra. As will be seen later, the main Si environments present in the samples studied here, from 500 to 1400 °C, are SiC2N2, SiCN3 and SiN4. Fitting the spectra using these three components leads to the exact positions of the individual signals as the temperature increases.It is shown that SiCnN4-n chemical shifts decrease when the temperature increases and this corresponds to the change in the Si second coordination sphere, i.e. the decrease in the number of protons. At temperatures higher than 500 °C, a very small amount of N atoms exist as NHSi2 sites, most of them were already transformed into NSi3 environments.The protonation state of nitrogen atoms changes only slightly from Fig. 6 29Si CP MAS (with cross polarization) NMR spectra of the 500 to 1400 °C, which is why the chemical shift corresponding VS/MS residues at different pyrolysis temperatures: (a) precursor, to SiN4 sites does not change much with temperature (Fig. 5). (b) 250 °C, (c) 500°C, (d) 650°C, (e) 850 °C In contrast, C atoms are still highly protonated at 500 °C. As T increases, carbon deprotonation occurs. In our samples, the non-cross-linked oligomers. All the sites formed by hydro- it is clear that SiCnN4-n 29Si chemical shifts increase with a T silylation and/or Vi polyaddition give the NMR signals shown increase in proportion with n, the number of C atoms in the in Fig. 6. It is emphasized that both types of cross-linking Si sphere (Fig. 5). reactions can lead to the same types of Si first coordination spheres, and so it is not easy to distinguish here which cross- VS/MS. Figs. 6 and 7 show the 29Si CPMAS and MAS linking reaction is predominant. The comparison between 29Si NMR spectra of VS/MS pyrolysates.The assignment of the CPMAS and 29Si MAS NMR spectra shows whether Si first main signals present in the CPMAS spectra is given in Fig. 6. coordination spheres are protonated. It appears that Numerous environments are present in the precursor: SiH(C sp2)N2 and SiH(C sp3)N2 environments are still present SiH2(C sp2)N sites are due to ends of chains, the other Si at 250 °C, with resonances at d ca.-35 and -20 respectively. environments can be explained as shown in Fig. 8. The Si sites Si(C sp3)N3 sites may be present at 500 °C and appear in the indicated with an asterisk represent the initial sites present in same chemical shift range as SiH(C sp3)N2 sites. SiN4 environments appear at 500 °C, at d ca.-45. With increasing temperature to 850°C, Si(C sp3)N3 and Si(C sp3)2N2 become the major Si first coordination spheres, while SiMH bonds are consumed.Between 850°C and 1400 °C, the intensity of the signal due to SiN4 sites increases but never becomes predominant. VS. The NMR study of the VS precursor in solution13 showed the presence of SiH(C sp2)N2, SiH(C sp3)N2 and Si(C sp3)(C sp2)N2 sites in the precursor. Let us add that the nitrogen environments present in the precursor were determined by 15N and 14N NMR spectroscopy; the following distribution was found: 30% NH2Si (at d -356.2), 46% NHSi2 sites (at d -346.8) and 24% NSi3 (at d -330).The chemical shifts were relative to CH3NO2. This result reflected the occurrence of the transamination reaction during the curing step of the starting oligosilazanes.Fig. 2B and C show the 29Si Fig. 5 29Si NMR chemical shifts corresponding to SiCnN4-n sites MAS and CPMAS NMR spectra of VS pyrolysates. At 500°C, (n=0, 1 and 2) in the pyrolytic residues as a function of temperatures. $, SiN4; &, SiN3C; +, SiN2C2 . the main environments are Si(C sp3)2N2, Si(C sp3)N3 and 120 J. Mater. Chem., 1997, 7(1), 117–126SiN4, (signals at d -2, -24 and -46 respectively); there rich sites form.Between 500 and 1200 °C, the intensity of the signal due to SiN4 sites increases and becomes predominant. certainly remain some protonated sites of the type SiH(C sp3)N2 but no SiH(C sp2)N2 since all vinyl groups were Note that among all SiCnN4-n environments, only SiC2N2 sites can be explained by solely cross-linking reactions such as consumed before 500 °C.Between 600 and 900 °C, a marked consumption of Si(C sp3)2N2 sites is observed while nitrogen- hydrosilylation and vinyl polyaddition. One possibility is that SiCN3 environments form from SiHCN2 sites by dehydrocoupling between SiMH and NMH bonds. Another explanation is that SiMC cleavage around SiC2N2 environments can lead to C-deficient Si sites and the formation of new SiMN bonds, thus explaining the occurrence of SiCN3 sites and particularly the increase of SiN4 environments.VNMS. Fig. 9A and B present the 29Si MAS and CPMAS NMR spectra of VNMS pyrolysates. The main component of the spectra at 250 °C, present at d ca. 0, is due to Si(C sp3)2N2 sites, the component at d ca. -15 to -20 corresponds to SiH(C sp3)N2 environments and there remain some initial sites of the type SiH(C sp2)N2 (d ca.-30). The increase of the component at d ca. -20 may correspond to the formation of Si(C sp3)N3 sites, which can be connected to the consumption ofNMMe bonds leading to new NMSi bonds. This proposition is in agreement with 13C NMR spectroscopic observations. The change is abrupt between 500 and 650 °C when a large amount of SiN4 sites appears as a broad signal at d ca.-45. This observation suggests the formation of abundant SiMN Fig. 7 29Si MAS (without cross polarization) NMR spectra of the VS/MS residues at different pyrolysis temperatures: (a) precursor, (b) 250 °C, (c) 500 °C, (d) 650°C, (e) 850°C, (f ) 1400°C Fig. 9 NMR spectra of the VNMS residues at different pyrolysis Fig. 8 Possible firstcoordination spheres of silicon atoms afterhydrosi- temperatures.A, 29Si MAS (without cross polarization): (a) 250 °C, (b) 500 °C, (c) 650°C, (d) 850 °C, (e) 1400 °C, B, 29Si CP MAS (with lylation and/or vinyl polyaddition starting from the initial silicon sites (indicated with a star) cross polarization): (a) 250°C. (b) 500°C, (c) 650°C. J. Mater. Chem., 1997, 7(1), 117–126 121bonds at the expense of SiMH or, more probably SiMC bonds.this temperature range.21 The decrease in C/Si is mainly due to the breaking of NMMe bonds leading to the release of CH4 The change observed in the 13C NMR spectra was also very important in this temperature range. With increasing tempera- and this probably results in the formation of new NMSi bonds at the expense of SiMC and SiMH bonds.This is the first ture, SiN4 sites become more abundant; they are largely predominant at 1400°C. Again, it is observed that SiMN main origin of the formation of nitrogen-rich silicon sites. The second one is the expulsion of C groups from the silicon-based bonds form at the expense of SiMC bonds in the silicon structure observed by 29Si NMR spectroscopy.network to form a carbon phase, which is clearly evidenced through 13C NMR studies from 650 to 850 °C. Vinyl functions first transform into C sp3 groups essentially by polyaddition Discussion reactions. These carbon chains partly lead to the evolution of high molecular mass hydrocarbons as observed by mass spec- This discussion is divided in three parts, the first concerns the formation of the mineral network taking place below the trometry.21 Another large number of SiMC bonds are cleaved leading to free C sp2-type carbonaceous species and more organic–inorganic transition temperature (ca. 850 °C), the second one will be devoted to the evolution of the silicon-based numerous SiMN bonds in the silicon matrix. Carbon segregation thus leads to nitrogen enrichment of the silicon network.matrix at temperatures higher than 900 °C (silicon carbonitride structure and silicon nitride phase) and the third part to the In the case of the VS route, there is a first increase in the amount of SiN4 sites up to 500–650 °C (ca. 30%). This phenom- free carbon phase, its formation and its amount which is determined from the combination of quantitative NMR results enon is associated with a small mass loss and a small C/Si decrease.21 Cross-linking reactions (hydrosilylation, Vi poly- and chemical analyses.merization, dehydrocoupling of SiMH and NMH bonds) have led to SiC2N2 and SiCN3 sites. These sites can transform to Formation of the mineral network N-rich sites (SiN4) through expulsion of C groups as a solid During the first pyrolysis step, the mineral network forms; this free carbon structure.This is what is observed by 13C NMR stage was characterized in detail by MS, TG and IR spec- studies, showing the formation of new C sp2 sites at 600 °C troscopy in the preceding paper.21 It was shown that in the which become very abundant at 900 °C. In this case, the first stage, up to 400 °C, further cross-linking proceeds while a expulsion of carbon from the silicon-based network to give loss of low molecular mass oligomers occurs.In the second free carbon mainly explains the formation of SiN4 sites. stage, from 400 to 800 °C, the mineralization step is charac- In the VS/MS route, the increase in the fraction of SiN4 is terized by a mass loss corresponding to the evolution of much lower up to 900°C compared to the other two precursors.hydrocarbons and hydrogen; it is associated with the breaking This is in accordance with the higher stability with temperature of SiMH, CMH, NMH, NMC and SiMC bonds. Using 29Si of the SiMC bond for Me groups compared to the SiMC bond NMR spectroscopy, it is possible to observe the consequences obtained by hydrosilylation or polyaddition of vinyl species of the cross-linking reactions and the gas evolution on the (stability shown by IR and 13C MAS NMR studies).The local transformations taking place in the solid residues. Silicon formation of SiN4 sites can be mainly related to the departure environments are very numerous below 900 °C (SiHpNnCm of carbon as a gaseous product; the evolution of methane is sites) and it would be hazardous to try to obtain precise values indeed observed up to 900 °C.The formation of free carbon of proportions of Si sites present in the materials because of plays a lesser role in this route, the signal corresponding to the low resolution of the resulting NMR signals. The 29Si MAS new C sp2 environments is indeed very weak in the 13C NMR NMR spectra were completely simulated but only the compo- spectra, in agreement with the results obtained by Raman nent present at high field (d -45) was interpreted quantita- spectroscopy.25 tively.The fraction of SiN4 sites was thus estimated from 500 From 13C and 29Si NMR results, the structural state of the to 900 °C for the three routes, it is the only signal appearing preceramic network can be deduced.The presence of in the chemical shift range d -45 to -50. Fig. 10 shows that SiMCMCMSi bridges formed at low temperature (below at 900 °C, the fraction of SiN4 sites reaches 54% in the VNMS 600°C) renders the matrix more stable towards redistribution route, 33% in the VS route and 23% in the VS/MS route. around Si atoms. Note that redistribution mainly involves In the VNMS route, it is observed that SiN4 sites increase SiMN bonds with SiMN or SiMH ones below 900 °C.Cross- rapidly from 500 to 900°C and become predominant. This can linking involving vinyl groups reduces the segment mobility be related first to the large mass loss observed in TG (ca. 25%) and hinders exchange reactions. It was shown21 that transamin- together with the large decrease in C/Si (from 3 to 1.75) in ation is hindered in the VS and VNMS routes (where vinyl groups lead to a high cross-linking degree) compared to the VS/MS route; the departure of NH3 was reported to be considerable only in the case of VS/MS polymers.At ca. 500°C, considering that all vinyl groups were transformed to C sp3 units, it is observed that VS and VS/MS intermediates contain about the same number of SiMC sp3 bonds per silicon.Note that the 29Si CPMAS NMR spectra at 500 and 600°C of VS and VS/MS are indeed very similar, showing that the proportions of SiMC sp3 and SiMN bonds in the silicon network are about the same in the two routes. But SiMC sp3 bonds, appearing at low temperature by vinyl transformation are not very stable and partly disappear when the temperature is increased to 900 °C. The SiMC cleavage leads to SiN4 sites and enrichment in N atoms of the silicon network.It was observed that the 29Si MAS NMR spectra at 900 °C of VS and VS/MS are different, SiN4 sites being more abundant in VS than in VS/MS. At 900 °C, the silicon network is mainly built of SiMN bonds but it is richer in SiMC bonds in VS/MS than in VS than in VNMS.The combination of 13C and 29Si NMR results was helpful Fig. 10 Fractions of SiN4 sites as a function of the pyrolysis temperature in the three routes. #, VS/MS; ×, VS; &, VNMS to clarify the next point: it is clear that SiMN bonds are 122 J. Mater. Chem., 1997, 7(1), 117–126formed at the expense of SiMC bonds, but do SiMC bonds samples; this is supported by the fact that the fraction of SiC3N sites (which would give a signal at even lower field) is found break to form a solid carbonaceous residue or to lead to the evolution of carbon-containing gaseous products? In each to be zero from 900 to 1400 °C.Moreover, X-ray diffraction peaks corresponding to SiC phases were never observed in route, we showed that it was possible to evaluate which phenomena predominate.those samples pyrolysed under nitrogen, even at high temperature (1450 °C).21 The low oxygen content is neglected in these calculations, and so only SiMC and SiMN bonds are taken Evolution of the silicon-based network from 900 to 1400 °C into account in the silicon matrix. The silicon-based network is now examined in the temperature Fig. 12 shows the evolution of the fractions of SiCnN4-n range 900–1400 °C.It is constituted of all Si atoms and N and sites from 900 to 1400°C. It is observed that a redistribution C atoms directly bonded to Si atoms, and so it excludes the between SiMC and SiMN bonds takes place from 900 to free carbon phase formed by the breaking of SiMC bonds. Up 1400°C. SiN4 and SiC2N2 sites are favoured at the expense of to the organic–inorganic transition temperature, deprotonation SiCN3 sites.A tendency to segregate a silicon nitride phase is in the silicon-based network was considerable.21 From 900 °C, thus observed. The linewidth of the peak due to SiN4 environ- it is assumed that the number of SiMH bonds is negligible. Assuming that at temperatures higher than 900°C, CMH, NMH and NMMe bonds are also negligible, atomic compositions of the silicon matrix can be estimated.C atoms are taken into account as CSi4 sites and N atoms as NSi3 sites. The proportions of Si environments are obtained from deconvolution of the 29Si NMR spectra considering SiCnN4-n firstcoordination spheres; Fig. 11 shows the results obtained at 1400 °C. It is assumed that SiC4 sites are negligible in these Fig. 12 Fractions of silicon sites (SiN4, SiN3C and SiN2C2) in the Fig. 11 29Si MAS NMR spectra of the three samples at 1400 °C: (a) VNMS, (b) VS, (c) VS/MS. Deconvolution into three individual silicon-based networks of (a) VS/MS, (b) VS and (c) VNMS samples pyrolysed from 850 to 1400°C. #, SiN2C2; ×, SiN3C; &, SiN4. peaks corresponding to SiN4, SiN3C and SiN2C2 sites.J. Mater. Chem., 1997, 7(1), 117–126 123Table 3 Estimated proportions of silicon nitride and silicon carbo- ments shows that, at temperatures lower than 1400 °C, the nitride structures in the silicon-based networks at 1400°C silicon nitride structure is still mainly amorphous, which agrees with XRD results.21 With increasing temperature in the VS sample silicon nitride+silicon carbonitride route, the formation of more numerous SiN4 sites is clearly observed in the 29Si NMR spectra at 1400 and 1450 °C VNMS 0.64 SiN4/3+0.36 SiN0.85C0.36 VS 0.44 SiN4/3+0.56 SiN0.88C0.34 (Fig. 2B). This agrees with the XRD results22 which show a VS/MS 0.32 SiN4/3+0.68 SiN0.84C0.37 better crystallization of silicon nitride at 1450 °C. Atomic compositions of the Si networks are calculated using SiCnN4-n fractions.The calculations are performed as follows: if p=SiC2N2 (%), q=SiCN3 (%) and r=SiN4 (%), It is also interesting to compare the composition of the we have: Si(%)=(p+q+r)/S; C(%)=(2p+q)/4S and silicon carbonitride structure of these materials to the composi- N(%)=(2p+3q+4r)/3S with S=(p+q+r)+(2p+q)/4 tion of another silicon carbonitride structure obtained in a +(2p+3q+4r)/3. The deduced C/Si and N/Si atomic ratios material which leads to the major crystallization of SiC and are indicated in Table 2.The compositions are found to be not Si3N4 at high temperature. From carbosilazane precursors quite stable with temperature from 900 to 1400°C, which prepared from thermolysis of (SiMe2)n–(NHSiMeHNH)m means that carbon is not lost from the Si network in this copolymers,27,28 silicon carbonitride ceramics with SiC crystals temperature range.This observation proves that segregation were obtained at 1400 °C.16 Following a similar strategy to of the free carbon phase takes place mainly before the organic– characterize these ceramics,16,23 it was shown that the main Si inorganic transition.After this transition, the main reactions sites present in that silicon network were SiC4, SiC3N and are local rearrangements in the amorphous silicon matrix SiC2N2 from 900 to 1400°C. Again, if we assumed that all which permit the creation of N-rich environments in order to SiC4 sites are only part of the SiC crystalline phase at 1400°C, allow the further silicon nitride crystallization.Exchange it was possible to estimate an atomic composition characteriz- involving SiMC and SiMN bonds is the main reaction. ing the amorphous silicon carbonitride structure that coexists The atomic compositions characterizing the silicon networks with the SiC crystals. The whole silicon phase was found to in the three samples present a similar N/Si ratio, close to unity.have the composition SiC0.71N0.40, and the silicon carbonitride It is seen that C/Si ratios decrease from VS/MS to VS to structure (excluding in this case SiC4 sites) to have the composi- VNMS, suggesting a much higher ability to retain C atoms in tion SiC0.61N0.52, which can be formally written as the silicon phase in VS/MS than in VS than in VNMS. This 0.6SiC+0.4SiN1.33. Note that the atomic composition of the trend is opposite to the initial C/Si values in the precursors.amorphous silicon carbonitride structure is rich in carbon Compared to the atomic compositions in the precursors, atoms in this case and thus seems to be related directly to the ca. 17% of total carbon atoms were retained in the VS/MS nature of the coexisting crystallizing phase, which is thermo- silicon phase compared to 10% in VS and 4% in VNMS.It dynamically expected. Note that carbon atoms initially present is evident that an increase in the total carbon content in the as SiMMe groups in this route lead to a stable incorporation precursor, whatever the polymeric backbone structure, does of carbon in the silicon carbonitride structure and to <5% of not lead to a proportional increase in the carbon content of free carbon.the silicon-based network in the ceramic. Let us now consider the materials pyrolysed at 1400 °C The free carbon phase which present the formation of minor amounts of crystalline silicon nitride, and let us assume that all SiN4 sites segregate It was reported, in Part 1,21 that the global N/Si and C/Si atomic ratios in the pyrolysed materials were almost constant in order to favour Si3N4 crystallization. In this case, it is possible to characterize the silicon carbonitride structure and from 900 to 1400 °C; the corresponding values obtained from chemical analyses at 1400°C are indicated in Table 1.We have calculate an estimated atomic composition of that amorphous structure. The compositions of the silicon carbonitride just shown that these ratios calculated now for the siliconbased network from 29Si NMR results were also constant from structure are then calculated as explained earlier, but only the strictly mixed Si sites (SiC2N2 and SiCN3) are taken 900 to 1400 °C (Table 2).Both results suggest that carbon atoms have left the Si network mainly before the mineral into account.The results are (SiN0.85C0.36) for VNMS, (SiN0.88C0.34) for VS and (SiN0.84C0.37) for VS/MS. Note that transition at 900°C to form a free carbon phase. This combination of quantitative analyses obtained from 29Si MAS NMR the compositions are very similar, close to SiN0.85C0.35 which can be also formally expressed as 0.65SiN1.33+0.35SiC. This studies and chemical analyses is in accordance with the qualitative 13C MAS and CPMAS NMR spectra.It was shown composition is rich in nitrogen atoms, which seems to agree with the fact that the amorphous silicon carbonitride structure that changes in C chemical environments were considerable before 900 °C: a large increase in the fraction of new C sp2 actually coexists with an Si3N4 phase in the process of crystallization.If we compare the relative abundances of the silicon sites was observed from 600 to 900°C at the expense of C sp3 atoms bonded to Si atoms. The new C sp2 sites were attributed nitride phase and the silicon carbonitride structure (Table 3), we observe that the amorphous silicon carbonitride structure to protonated cyclic precursors of a graphite-type structure formed by SiMC bond cleavages; the protons play the role of in the three materials is more abundant in VS/MS than in VS and than in VNMS, which remains in agreement with the poisons in the growth of graphitic carbon cages.These results were supported by TEM data from Delverdier.29 When free order of ability to retain C atoms in the Si matrix. carbon appears, it is present as small aromatic carbon units.The first primary aromatic entities, called basic structural units Table 2 N/Si and C/Si atomic ratios in the silicon-based matrices at (BSU), are embedded in the amorphous silicon carbonitride 900 and 1400 °C, determined by NMR phase. When the temperature increases, the BSUs rearrange into carbon stacks giving rise to more or less complete cages silicon matrix silicon matrix formed from graphite planes.at 900 °C at 1400 °C The abundances of the free carbon phases in the different sample N/Si C/Si N/Si C/Si routes can be estimated from the fractions of C atoms which are no longer bonded to Si atoms. These values can be obtained VNMS 1.11 0.13 1.15 0.13 from the comparison of the C/Si ratios in the global materials VS 1.09 0.19 1.09 0.19 obtained by chemical analyses (CA) and in the Si networks VS/MS 0.97 0.27 1.00 0.25 obtained by NMR studies (Tables 1 and 2).The percentages 124 J. Mater. Chem., 1997, 7(1), 117–126of atoms in the free carbon phase are calculated as follows: mate proportions of phases in the ceramics desired for specific properties. free C(%)=[(C/Si)CA–(C/Si)NMR]/[1+(C/Si)CA+(N/Si)CA]. The results are as follows: 41% (VNMS), 39% (VS) and 28% (VS/MS), which correspond to the following fractions of total Conclusions C atoms: 93% for VNMS, 87% for VS and 76% for VS/MS.It is shown that the abundance of free carbon is inversely The present study shows that solid-state NMR permits the related to the abundance of the amorphous silicon carbonitride elucidation of different aspects of the pyrolytic conversion of structure.This observation is true for pyrolytic residues at polysilazane precursors into silicon carbonitride ceramics. 850 °C and holds up to 1400 °C. The relative abundances of The assignment of the signals corresponding to SiCnN4-n the three different structures present in the materials at 1400 °C (n=0, 1, 2, 3, 4) sites allows us to follow the different reactions are summarized in Table 4.of cleavage or formation of bonds in the silicon-based matrix. The abundance of free carbon can be related directly to the The quantification of these sites leads to a more precise nature and amount of cross-linking reactions occurring below description of the structural evolution with temperature of the 900 °C. The combined results suggest that Vi polyaddition different phases forming the material, namely the silicon-based mainly leads to carbon structures with very little carbon network and the free carbon phase, which at high temperature insertion into the Si network.Polyaddition creates new CMC separate in three distinct structures: the crystallizing silicon bonds and the formation of polymeric carbon chains (CH2)n nitride phase, the amorphous silicon carbonitride structure with n3 gives SiMC sp3 bonds with very low stability.and free graphitic carbon. Hydrosilylation is the main reaction that creates new SiMC From the combination of quantitative NMR results and bonds at low temperature leading to stable carbosilane bridges chemical analyses, the following points are clearly shown.from vinyl groups. Hydrosilylation leads to SiMCMCMSi Below the organic–inorganic transition temperature the carbon chains or SiMCMSi bridges. As the temperature increases, groups leave the silicon network to segregate as a free carbon SiMCMCMSi chains can rearrange into SiMCMSi carbosilane structure. The carbon-bearing functions, i.e. the SiMMe, SiMVi units.Stable carbon incorporation into the silicon network and NMMe groups, are compared by evaluating in the three can thus occur. In the VS route, Vi polyaddition in preference routes the different relative amounts of the three final carbon to hydrosilylation consumes Vi groups, and it is even more states, which are gaseous products, carbon incorporated in the true in VNMS, for which hydrosilylation is certainly hindered silicon matrix and free carbon.Another major point is that by the presence of NMMe groups;21 this leads to high free the chemical composition of the amorphous silicon carbo- carbon contents. Vi groups are not readily evolved as gaseous nitride structure can be estimated; in the present cases, it is products, they lead to high final carbon contents in the close to SiN0.85C0.35.It is shown that that composition is ceramics and that carbon is very slightly incorporated in the essentially related to the nature of the coexisting phase which silicon matrix when polyaddition is preferred to hydrosilyl- is crystallizing (Si3N4 or SiC). Also, the proportions of all the ation. In the VS/MS route, the methyl groups give rise coexisting phases in the ceramics can be deduced.preferentially to either gaseous products such as methane or to stable SiMCH2MSi carbosilane bridges which progressively References deprotonate into CHSi3 and finally transform into CSi4 environments. These processes lead to a low free carbon 1 R. M. Laine, Y. D. Blum, D. Tse and R. Glaser, Inorganic and content but to a reasonable amount of silicon carbonitride Organometallic Polymers; ACS Symp. Ser. 360, ed. M. Zeldin, structure. It is possible to summarize schematically the evol- K. J. Wynne and H. R. Allcock, ACS, Washington, DC, 1988, p. 124. utions of carbon atoms from the precursors to the ceramic 2 D. Seyferth, in Silicon-based Polymer Science. A Comprehensive materials obtained by pyrolysis at 1400 °C for the three routes Resource; Adv.Chem. Ser. 224, ed. J. M. Zeigler and F. W. Fearon, (Table 5). Three possible final carbon states are considered: ACS,Washington, DC, 1990, p. 565. gaseous products, carbon in a free carbon phase and carbon 3 K. J.Wynne and R. W. Rice, Annu. Rev. Mater. Sci., 1980, 14, 297. as CSi4 sites in a silicon-based structure. This Table is very 4 M. Peuckert, T.Vaahs and M. Bu�ck, Adv.Mater., 1990, 2, 398. informative about the transformations of the carbon-contain- 5 D. Seyferth and G. H. Wiseman, J. Am. Ceram. Soc., 1984, 67, C132. ing substituents in the precursors. It appears that NMe groups 6 Y. D. Blum, K. B. Schwartz and R. M. Laine, J. Mater. Sci., 1989, lead to a completely inefficient carbon insertion in the solid 24, 1707. material.Tables 3 and 5 may be very helpful in the design of 7 K. B. Schwartz and Y. D. Blum,Mater. Res. Soc. Symp. Proc., 1988, the backbone of polymeric precursors if one knows the approxi- 121, 483. 8 N. S. Choong Kwet Yive, R. J. P. Corriu, D. Leclercq, P. H. Mutin and A. Vioux, Chem.Mater., 1992, 4, 141. Table 4 Atomic fractions of the three structures constituting the cer- 9 R.J. P. Corriu, D. Leclercq, P. H. Mutin and A. Vioux, Chem. amics at 1400 °C Mater., 1992, 4, 711. 10 A. Lavedrine, D. Bahloul, P. Goursat, N. S. Choong Kwet Yive, sample Si3N4 (atom%) SiNxCy (atom%) C (atom%) R. Corriu, D. Leclercq, H. Mutin and A. Vioux, J. Eur. Ceram. Soc., 1991, 8, 221. VNMS 38 21 41 11 G. T. Burns, T. P. Angelotti, L. F. Hanneman, G. Chandra and VS 28 33 39 J. A. Moore, J.Mater. Sci., 1987, 22, 2609. VS/MS 25 47 28 12 H. N. Han, D. A. Lindquist, J. S. Haggerty and D. Seyferth, Chem. Mater., 1992, 4, 705. 13 C. Ge�rardin, F. Taulelle and J. Livage,J. Chim. Phys., 1992, 89, 461. Table 5 Estimated distributions of carbon atoms in the three possible 14 C. Ge�rardin, M. Henry and F. Taulelle, Mater. Res. Soc. Symp. final states: carbon incorporated in the silicon carbonitride structure, Proc., 1992, 271, 777. carbon in free carbon and carbon in gaseous products 15 C. Ge�rardin, F. Taulelle and J. Livage, Mater. Res. Soc. Symp. Proc., 1993, 287, 233. no. of C atoms per Si 16 D. Mocaer, R. Pailler, R. Naslain, C. Richard, J. P. Pillot, J. Dunogue`s, C. Ge�rardin and F. Taulelle, J. Mater. Sci., 1993, Si free gaseous 28, 2615. sample initial carbonitride carbon products 17 W. R. Schmidt, P. S. Marchetti, L. V. Interrante, W. J. Hurley, R. H. Lewis, R. H. Doremus and G. E. Maciel, Chem.Mater., 1992, VNMS 3 0.13 1.62 1.25 4, 937. VS 2 0.19 1.40 0.41 18 R. H. Lewis and J. E. Maciel, J.Mater. Sci., 1995, 30, 5020. VS/MS 1.5 0.25 0.81 0.44 19 E. Bacque, C. Richard, J. P. Pillot, M. Birot, J. Dunogue`s, M. J. Mater. Chem., 1997, 7(1), 117–126 125Pe�traud, C. Ge�rardin and F. Taulelle, J. Inorg. Organomet. Polym., 25 D. Bahloul, M. Pereira, T. Merle, P. Goursat, C. Gerardin and F. Taulelle, 3rd Int. Conf. Ceramic–Ceramic Composites, October 1995, 5, 169. 20 G. E. Legrow, T. F. Lim, J. Lipowitz and R. S. Reaoch, Am. Ceram. 1994, Mons, Belgium. 26 A. Lavedrine, PhD Thesis, University of Limoges, 1992. Soc. Bull., 1987, 66, 363. 21 D. Bahloul, M. Pereira and C. Ge�rardin, J. Mater. Chem., preced- 27 E. Bacque, J. P. Pillot, J. Dunogues and P. Olry, Eur. Pat., 296028, 1988. ing paper. 22 D. Bahloul, M. Pereira and P. Goursat, J. Am. Ceram. Soc., 1993, 28 E. Bacque, J. Dunogues, C. Biran, P. Olry and J. P. Pillot, Fr. Pat., 2589037, 1986. 76, 1156. 23 C. Ge�rardin, PhD dissertation, Universite� P. et M. Curie, Paris, 29 O. Delverdier, PhD Thesis, University of Pau, 1991. 1991. 24 M. Henry, C. Ge�rardin and F. Taulelle, Mater. Res. Soc. Symp. Paper 6/03181A; Received 7thMay, 1996 Proc., 1992, 271, 243. 126 J. Mater. Chem., 199

ISSN:0959-9428    

DOI:10.1039/a603181a

出版商:RSC    

年代:1997

数据来源: RSC