High temperature behavior of periodic mesoporous ethanesilica glasses prepared from a bridged silsesquioxane and a non-ionic triblock copolymer

Journal of Non-Crystalline Solids 353 (2007) 1109–1119 www.elsevier.com/locate/jnoncrysol

High temperature behavior of periodic mesoporous ethanesilica glasses prepared from a bridged silsesquioxane and a non-ionic triblock copolymer S. Masse *, G. Laurent, F. Babonneau CNRS, UMR 7574 CMCP, Paris F-75005, France Universite´ Pierre et Marie Curie-Paris6, UMR 7574 CMCP, Paris F-75005, France Received 17 February 2006; received in revised form 25 October 2006 Available online 8 February 2007

Abstract Periodic mesoporous organosilicas (PMOs) with 2D-hexagonal structure have been prepared by the sol–gel route in acidic conditions from a bissilylated reactive as a silica-modified source, namely 1,2-bis(triethoxysilyl)ethane (BTEE), and Pluronic P123 non-ionic surfactant as a templating agent. Pyrolysis of the PMO material was performed under inert atmosphere up to 1000 C in order to assess the possibility to transform the hybrid network into a porous SiOC glass network. The pyrolysis process has been mainly followed by 29Si and 13C solid state NMR. Thermal degradation of the pore network was followed by XRD and TEM and measured by N2 adsorption– desorption experiments. This study shows that despite the Si–C bonds cleavage occurring after 600 C causes a phase separation into silica and free carbon, an ordered structure is retained up to 1000 C with a remaining specific surface area.  2006 Elsevier B.V. All rights reserved. PACS: 82.56.Ub; 61.10.Nz; 81.20.Fw Keywords: X-ray diffraction; Nuclear magnetic (and quadrupole) resonance; TEM/STEM; Porosity; Silicates; NMR, MAS-NMR and NQR; Sol–gel, aerogel and solution chemistry; Organic–inorganic hybrids

1. Introduction Despite the weak hydrothermal stability presented by the first generation of mesoporous silica and organosilica materials, it is now possible to overcome this property using high-molecular-weight block copolymers instead of conventional ionic surfactants [1]. This could have great potential for specific applications, such as filters, membranes or catalysts for severe operating conditions, that requires materials having large internal surface areas and narrow pore size distributions. The first synthesis of meso*

Corresponding author. Address: Chimie de la Matie`re Condense´e de Paris, UPMC-Paris6, Tour 54, Couloir 54/55, Etage 5, 4 place Jussieu, 75252 Paris cedex 05, France. Tel.: +33 1 44 27 23 98; fax: +33 1 44 27 47 69. E-mail address: [email protected] (S. Masse). 0022-3093/$ - see front matter  2006 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2006.12.020

porous silica using non-ionic surfactants was published in 1995 [2]. In the last 10 years, amphiphilic block copolymers like Pluronic-type triblock copolymers, namely [poly(ethylene oxide)]–[poly(propylene oxide)]–[poly(ethylene oxide)], were widely used instead of ionic surfactants to generate wider pores, thicker silica walls and finally a more stable material [3]. Actually, wall thickness – which is a key-parameter for the silica-based network stability – could also be improved using bridging organic functions provided by bissilylated precursors [4]. Now, a large variety of organic spacers was investigated: saturated ones, like methane or ethane bridges, or unsaturated ones, going from the simplest ethene bridge [5] to the more complex aromatic derived bridges. In particular, phenylene [6–9] bridges were used for the possibility of ‘chemistry of the channels’ they offer, or for their action as ‘ligand channels’ for metal complexes and organometallics suitable for

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catalysis applications. Some other investigated spacers are thiophene and ferrocene [6] for their applications in electrochemistry, or also biothiophene and acetylene. As related by Cho et al. [10,11], PMOs materials, that can be obtained by reacting an organosilane instead of pure silica, are unique in containing organic functional groups inside the channel wall. Actually, the organic modification of the mesoporous silica framework could provide diverse variations of optical, electrical or mechanical properties, which can give specific characteristics to the material. On top of that, the presence of Si–O and Si–C bonds inside these materials makes them perfect candidates as precursors for silicon oxycarbide glasses. Sol–gel derived SiOC glasses have already exhibited unique high temperature properties as mechanical strength [12,13] and chemical durability in previous studies. So, the choice of a basic spacer such as CH2–CH2 could be valuable if the simplicity of the spacer nature is counterbalanced by a further heating treatment which could generate SiOC powder ceramics from the sol–gel derived hybrid polymer. Such an approach is now well recognized to be a suitable way to obtain dense silicon oxycarbide glasses [14–17]. Another extension of sol–gel chemistry applied to this system was formerly undertaken for producing porous refractory silicon oxycarbide glasses [18,19] while retaining the excellent properties at high temperature. Indeed, the 2D-hexagonal and cubic Pm3n structures formed in that case from bis(trialkoxysilyl)ethane and ionic surfactant like cetyltrimethylammonium chloride (CTAC) have shown that the former collapses after pyrolysis at 800 C whereas the latter was retained up to 1000 C. In the present study, we describe first an approach for synthesizing 2D-hexagonal ethanesilica structures by the sol–gel route under acidic conditions using Pluronic P123, i.e. [PEO]20–[PPO]70–[PEO]20, as a structure-directing agent and bridged silsesquioxanes as silica sources. Then, the as-synthesized PMOs were submitted, after extraction of the templating agent, to high temperature under inert atmosphere (argon flow) in order to follow the thermal behavior from 400 C to 1000 C. The aim of this study is to assess the thermal stability of the Si–C and C–C bonds which are constitutive of the thick walls we have previously generated, in order to check whether or not these PMOs materials could be promising precursors to silicon oxycarbide glasses. Such an approach of the high temperature degradation of the PMOs is not usual. Despite the abundant literature dealing with PMOs, only a few authors [20] reported interesting studies about the temperature behavior of such materials. The use of neutral triblock copolymers as structuredirecting agents in PMOs synthesis began a few years ago. With Pluronic P123 and bistrimethylsilylethane (BTME) as organosilica source under acidic conditions, Muth et al. [21] synthesized for the first time a PMO material with hexagonal symmetry akin to SBA-15-type silica with large pores (6.5 nm), thick walls (5.9 nm) and high surface area (900 m2/g for the extracted material). The

same year, Burleigh et al. [22] synthesized a series of PMO whose pore size is ranging from 6 to 20 nm from bistriethoxysilylethane (BTEE) and P123 using trimethylbenzene as a micelle-swelling agent under acidic conditions. Some other authors [23] also used polypropylene glycol as a polymer-swelling agent in order to increase the pore size of an organo-modified SBA-15 material, as it can play a role in controlling the length of the hydrophobic block segment of the parent triblock copolymer. Zhu et al. [24,25] also reported a highly ordered large pore (in the range 4.0–7.7 nm) organosilica material working on the BTEE:P123:HCl:H2O system under the high block copolymer concentration conditions used for the Liquid Crystal Templating (LCT) approach. Matos et al. [26] also studied this system with a composition BTEE:P123:HCl:H2O equal to 1:0.02:5.81:200 to form a SBA-15-like PMO with 6.5 nm in diameter pore size. An alternative method for producing large pore PMOs (>4.5 nm), which could have a particular interest regarding to encapsulation of large molecules, laid on the use of inorganic salts, as it was investigated by Guo et al. through the addition of large amounts of NaCl [27] or K2SO4 [28] to a synthesis mixture containing BTME and Pluronic P123 (or F127, respectively) under acidic conditions to afford an hexagonal P6m (respectively, a cubic (Im3m)) ethane silica mesophase with large cavities of 6.5 nm (respectively, 9.8 nm) in diameter to be formed. 2. Experimental procedures 2.1. Synthesis In a typical synthesis, 1,2-bis(triethoxysilyl)ethane (BTEE) was used as silica source providing an ethane bridge between silicon atoms and Pluronic P123 ([EO]20[PO]70[EO]20) in acidic aqueous solution (HCl, pH  1) so that the molar composition of the final mixture was BTEE:P123:HCl:H2O = 1:0.0169:16:326. When using block copolymers, we must pay attention to choose the real molecular weight value of the P123 formula, which is represented by the MW value (also mentioned FW) because it takes into account the polydispersity of the formula, on the contrary to the average molecular weight, also mentioned MN (or MAV), which is often the value that is incorrectly reported by the authors in literature. In the case of P123, the value of 8400 (provided by BASF) was retained as MW for our syntheses. After the P123 acidic aqueous solution was stirred at 35 C during 0.5 h, BTEE was added. The mixture was covered for aging during 24 h at 35 C under stirring. Condensation stage was achieved by heating the mixture at 85 C in an oven 24 h more. The white precipitate was then filtered and washed several times with deionized water before being dried in an oven at 80 C during a few hours. The as-synthesized product was then weighted before extraction of the surfactant by reflux of ethanol at 80 C during 20 h. The extracted product was then washed several times with ethanol before drying at 85 C and weighting.

S. Masse et al. / Journal of Non-Crystalline Solids 353 (2007) 1109–1119

The extracted samples were then placed in an alumina crucible which was submitted to thermal treatment between 400 C and 1000 C under inert atmosphere through argon flow. In a first step, heating at 5 C min 1 was performed up to 100 C; this temperature was held during 2 h for adsorbed water elimination. In the next step, heating at 5 C min 1 went until the final temperature was reached; the dwelling time was 2 h at this temperature. The selected calcination temperatures are 400 C, 600 C, 800 C and 1000 C. Natural cooling under argon flow was then performed until room temperature was achieved. 2.2. Experimental techniques Thermogravimetric analyses (TGA-DTA) were performed under N2 atmosphere (12 cm3/min flow rate) with a TA Instruments SDT 2960 Simultaneous DSC-TGA equipment and were carried out loading approximately 10 mg of sample in an alumina crucible. The heating rate was 5 C min 1 until 1000 C. The infrared vibrational spectroscopy was used to follow the surfactant elimination after the extraction stage. For the X-ray diffraction experiments, the powdered specimens were pressed into the 20 mm diameter and 1 mm depth cavity of a polycarbonate sample-holder. The X-ray diffraction (XRD) patterns were collected with a powder Bruker AXS D8 Advance diffractometer in Bragg–Brentano configuration equipped with a graphite ˚) monochromator using Cu Ka radiation (k = 1.54178 A equipped with a rotary absorber and a Go¨bel focalization mirror. The running conditions for the X-ray tube were 40 kV and 40 mA. The data collection ranged from 2h = 0.6 to 5.0 by steps of 0.01 with a 10 s delay. N2 adsorption–desorption isotherms were recorded at 196 C using a Micromeritics ASAP 2010 surface analyser at relative pressures P/P0 ranging from 0.05 to 0.99. The specific surface area was determined by the Brunauer– Emmet–Teller method in the range 0.05 < P/P0 < 0.30. The pore distribution was determined using the BJH model. The pore diameter was taken off from the desorption curve. Microporosity was estimated through the a-plot method using a macroporous silica, namely Li-Chrospher Si-1000, as a reference. Solid state nuclear magnetic resonance, magic angle spinning (MAS NMR) as well as cross polarization coupled to magic angle spinning (CP-MAS NMR) experiments were performed on an Avance 300 Bruker spectrometer equipped with Bruker CP/MAS probes. Samples were finely ground and packed into ZrO2 rotors (7 mm diameter) closed with Vespel caps. For 29Si MAS NMR experiments, the powders were spun to 4 kHz and experiments were carried out at a frequency of 59.63 MHz with single pulse technique with high power decoupling during acquisition, using 90 pulse and recycle delays of 150 s. Such delays allow ensuring quite complete longitudinal relaxation for the silicium nuclei since the recycle delay will exceed 3 times the T1 value (estimated by a saturation-recovery experiment). FIDs were

1111

acquired with a sweep width of 30 kHz, 2048 data points and about 384 transients. All chemical shifts (d) were referenced to tetramethylsilane (TMS; d = 0 ppm). The FID data were processed using WinNMR-1D software and a line broadening of 50 Hz was applied before Fourier transformation. The experimental spectra were simulated using DmFit modelling software [29]. For 13C CP-MAS NMR experiments, the spinning rate was 5 kHz and experiments were carried out at a frequency of 75.48 MHz. The recycle delay was 3 s, the contact time 1 ms and the number of transients depends on the sample (from 160 for the H-rich samples to 1440 for the H-poor heat-treated specimens). The CP technique allows us to overcome the problem of very long 13 C longitudinal relaxation times while taking into account 1 H much shorter ones. It also represents a good way to enhance the detection of this low-sensitivity nucleus. On the other hand, it led us to a selective detection of C sites with protons in their close environment and to a qualitative measurement. Quantitative elemental analysis (EA) of the samples was performed by inductively coupled plasma spectrometry at the ‘Service Central d’Analyse du CNRS’ (Vernaison, France) for the determination of silicon, carbon and hydrogen amounts. The oxygen amount was determined by difference. Transmission electron microscopy (TEM) experiments were performed in bright field with a Philips CM12 field emission gun electron microscope operated at an accelerating voltage of 120 kV using a single-tilt sample holder and a single lanthanum hexaboride crystal as filament. Both imaging and diffraction modes were employed for specimen characterization. The specimens were dispersed on carboncoated copper grids by deposition of a drop of a diluted suspension of the particles in ethanol after sonification. Further experiments were then performed on microtomed samples using a JEOL 100 CX II electron microscope operating at 100 kV. The powdered samples were previously embedded in epoxy resin and cut in slices as thin as 70 nm using a diamond knife (Leica Ultracut UCT). The microtomic cuts were then laid on carbon-coated copper grids for analysis. 3. Results 3.1. DTA–TGA Preliminary experiments were undertaken first on the assynthesized hybrid gel in order to evaluate the thermal stability of the non-extracted material before heat treatment and to check the surfactant extraction process efficiency, and further on the extracted material to check the present of residual ethoxy groups resulting from an incomplete hydrolysis of the silica precursor. For the first experiment (as-synthesized material), a specimen containing 6.29 mg of the non-extracted sample powder was analyzed by DTA–TGA under N2 atmosphere. The classical experimental curves (not shown here)

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mainly show an intense and broad exothermic peak pointing out at 187 C associated with a shoulder at ca. 230 C, correlated to a total weight loss of ca. 40 wt%, that is consistent with the weight of P123 really incorporated in the system. These temperatures are in agreement with literature concerning P123 desorption from SBA-15-type materials [20,30]. Then, a continue weight loss of 10% from 415 C to 900 C is observed, which can be attributed to Si–CH2–CH2–Si decomposition [31]. For the second experiment (11.50 mg of the extracted material), only the second narrow component at 230 C is present on the thermogram. It may correspond to a fraction of P123 which cannot be removed by chemical extraction. In this study, it appeared realistic to us to follow the thermal behavior of these materials beginning at 400 C. 3.2. Elemental analysis The EA results obtained for the various samples are reported in Table 1. The conversion of the raw weight percentages into atomic ones allows calculating an average formula, with a silicon amount arbitrarily fixed to 1 to facilitate molar ratios comparison. We can note that for the extracted material, the C/Si ratio, determined as 1.2 by EA, is exceeding the theoretical value of 1.0 – available for a SiCH2O1.5 fully condensed siloxane network – because of the presence of a supplementary amount of carbon coming from residual P123 as mentioned above by DTA and as will be confirmed by a 13C MAS NMR experiment (spectrum not shown here). The decrease of C and H amounts with temperature rising is probably due to volatile hydrocarbons departure. It appears also that the O/Si ratio remains constant in the range 400–1000 C, because no loss of oxygen occurs during heating under inert atmosphere. 3.3. Solid state NMR In order to get a better insight into the structural changes that occur during pyrolysis, the samples were analysed by multinuclear solid state NMR. 3.3.1. 29Si MAS NMR experiments The characterization of the samples was performed by 29 Si MAS NMR as illustrated in Fig. 1. The spectrum of

the extracted PMO shows well-resolved signals due to T units at 46.5 ppm (T1:[CSiO(OH)2]), at 56.6 ppm (T2:[CSiO2(OH)]) and at 63.6 ppm (T3:[CSiO3]). The simulation of the spectrum gives relative percentages for T1/T2/ T3 = 3%/67%/30% (Table 2), that leads to a network condensation degreeP(D) = 76%, according to the formula P D = (nTn)/3 · Tn. We should notice that the corresponding fully condensed siloxane network would lead to a molecular formula of SiCH2O1.5 corresponding to a C/Si ratio equal to 1.0. Above 400 C, a Q-signal is appearing; it is principally composed of three components characteristic of the surface units (Q2:[SiO2(OH)2] at ca. 91 ppm and Q3:[SiO3OH] at ca. 101 ppm) and of the core units (Q4:[SiO4] at ca. 109 ppm). The simulation of the spectra allow calculating the formulation of an average composition taking into account the relative proportions between the different species and finally the molar ratio C/Si for each one (Table 2). The global condensation degree of the hybrid network couldP then be calculated, according P P to the general P formula D = [ (nTn) + (nQn)]/[(3 · Tn) + (4 · Qn)] that represents the ratio between the number of Si–O–Si linkages existing in the material and the number of Si–O– Si linkages which could be obtained if the network condensation would be at the maximum. During pyrolysis, it could be noticed that some of the Si–C bonds responsible for the T-signal are progressively transformed into Si–O–Si bonds. This global transformation of the T units into Q ones begins at 400 C and progressively converts the hybrid network into silica and free-C phases (above 600 C), that leads to a reduction of the C/Si ratio (0.91 at 400 C and 0.57 at 600 C). At the same time, new T-species appear at more negative values, corresponding to T units linked to Q ones, due to the change from T–T to T–Q entities during the progressive network condensation. Above 800 C, T-species are no longer present and the Q signal due to the siliceous matrix is now the unique signal of the spectrum, showing that the Si–C bond cleavage is then completely achieved, with Q3/Q4 relative proportions typical of hydroxylated silica. At this stage, the C/Si ratio is so equal to 0. 3.3.2. 13C CP-MAS NMR experiments As deduced above from 29Si MAS NMR, the Si–C linkages are progressively broken at high temperature.1H–13C cross-polarization MAS NMR experiments were then undertaken to study the C-species environment up to

Table 1 Composition of the PMOs at different temperatures as calculated from quantitative elemental analysis results PMO

Si (wt%) ± 0.1%

C (wt%) ± 0.1%

H (wt%) ± 0.1%

Oa (wt%) ± 0.1%

Average composition

Extracted 400 C-H.T. 600 C-H.T. 800 C-H.T. 1000 C-H.T.

31.5 36.0 38.2 39.7 41.3

16.0 15.0 12.4 9.1 7.1

4.3 2.9 1.8 0.7 0.4

48.2 46.1 47.6 50.5 51.2

SiC1.2H3.8O2.7 SiCH2.2O2.2 SiC0.8H1.3O2.2 SiC0.5H0.5O2.2 SiC0.4H0.3O2.2

H.T.: heat-treated. a The oxygen amount was determined by difference towards 100%.

S. Masse et al. / Journal of Non-Crystalline Solids 353 (2007) 1109–1119

Fig. 1.

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29

Si MAS NMR spectra of the PMO material at different calcination temperatures.

1000 C (Fig. 2). The signal of C atoms in the Si–CH2– CH2–Si linkage coming from the precursor has been detected through an intense signal at 5.0 ppm in the spectra of the extracted sample and of the 400 C heat-treated sample. After pyrolysis at 600 C under argon flow, a broad signal at 130 ppm due to amorphous carbon appears, as well as its related spinning side bands, which is attributed to an aromatic free-carbon phase. At the same time, the peak at 5.0 ppm coexists with some other new peaks. The first of them, the broadest one, is located at 15–30 ppm; the second one, located at 2.1 ppm, is the most intense; the third one appears as a shoulder at  6.2 ppm. These signals corresponds to the rearrangement products of Si–CH2–CH2–Si into Si–R, with R a

branched or not alkyl chain, as intermediates into the aromatic carbon phase. At 800 C, only the broad peak due to aromatic free-C is present in the spectrum, showing that all the Si–C bonds are now cleaved. Since the organic bridge cleavage is complete at 800 C, the hybrid character of the material is seriously affected above such a temperature. This result is in perfect agreement with 29Si NMR results which indicate that no more T-species were present on the spectra at the same stage. At 1000 C, the very low amount of protons remaining in the material (0.4 wt%) does not allow detection of the 13C signal from crosspolarization sequence so that the related spectrum is very noisy. But, in that case, a complementary 13C MAS single pulse NMR experiment (not shown here) has revealed that

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Fig. 2.

13

C CP-MAS NMR spectra of the PMO material at different calcination temperatures (SSB: spinning side bands).

Table 2 Composition of the PMOs at different temperatures as deduced from the simulation of the

29

Si MAS NMR experiments

PMO

T1 (mol%) ± 2%

T2 (mol%) ± 1%

T3 (mol%) ± 1%

T–Q (mol%) ± 2%

Q2 (mol%) ± 2%

Extracted 400 CH.T. 600 CH.T. 800 CH.T. 1000 CH.T.

3 2

67 32

30 57

– –

– 1

– 5

2

7

23

26

4

–

–

–

–

–

–

–

–

H.T.: heat treated.

Q3 (mol%) ± 1%

Q4 (mol%) ± 1%

Condensation degree (%) ± 1%

Molar ratio C/Si ± 1%

– 3

76 86

1.00 0.91

22

16

88

0.57

–

38

62

90

0

–

19

81

95

0

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the unique peak present at high temperature is still related to the aromatic free-C phase. 3.4. N2 adsorption–desorption measurements N2 adsorption–desorption isotherms have been recorded as a function of the pyrolysis temperature. The extracted sample shows a specific surface area of 1000 m2/g (±5%) and an hysteresis loop of type H1, with characteristic vertical and parallel branches, that is typical of cylindrical-shaped mesopores. The cumulative pore volume measured as a function of pore diameter (not shown here) shows a quite narrow pore size distribution ranging from 3.5 to 4.5 nm approximately. After pyrolysis at temperatures not exceeding 600 C, the shape of the curve is remained but the surface area decreases to 664 m2/g (Table 3). This phenomenon is mainly due to the cell contraction observed by XRD which is also correlated to a contraction of the pore volume. But after pyrolysis above 800 C, this progressive contraction of the structure is no more realistic and a breakdown is observed in the thermal behavior of the material, as measured and confirmed through repetitive measurements. First, the specific surface areas falls down to 214 m2/g at 800 C and 120 m2/g at 1000 C, and then, a curve shape transformation is observed for the isotherm curve shown in Fig. 3 which presents a novel hysteresis feature, extended to higher relative pressure than before. That should indicate that a secondary class of porosity is appearing between 600 C and 800 C probably related to microporosity generated during heat treatment. An a-plot method was used (see the curve Vads = f(a-plot) in inset) to estimate the proportion of micropores into the network. 3.5. X-ray diffraction The X-ray diffraction patterns were collected for the samples prepared from composition BTEE:P123:HCl:H2O 1:0.0169:16:326 after pyrolysis at 400 C, 600 C, 800 C and 1000 C (Fig. 4). For the extracted sample, an intense peak at 0.827 (2h) was obtained, as well as three secondary weaker peaks at 1.515, 1.690 and 2.271 (·30, see inset). These peaks were indexed as the (1 0 0), (1 1 0), (2 0 0) and (2 1 0) reflexions of a 2D-hexagonal structure [32]. As calcination temperature rises, a shift of the maximum of the peaks to higher values can be observed, as shown in

Fig. 3. N2 adsorption–desorption isotherm recorded for the 800 C treated PMO.

Table 4. This is in agreement with a cell contraction when pyrolysis temperature rise, with a cell parameter a ranging from 12.3 nm before any heat-treatment to 9.4 nm after pyrolysis at 1000 C. This decrease is gradual until 600 C. Then, between 600 C and 800 C, a total decrease of 24% seems to be the limit of structure contraction. 3.6. TEM investigation The channels of the 2D-hexagonal structure were first observed by TEM in bright field mode for the 400 C heat-treated PMO (Fig. 5(a)) and for the 1000 C heattreated PMO (Fig. 5(b)). These micrographs show an arrangement of parallel channels of mesopores, but it was impossible, by direct observation, to observe these channels in the transversal direction, neither to observe the core of the material due to the superposition of several layers of matter that prevent electron transmission. To overcome this problem, microtomes were cut in cone-shaped powder-embedded resin in order to observe the part of pore channels perpendicular to the surface. The pore diameter as well as the wall thickness could then been measured for the various PMOs, before heat treatment (Fig. 6a) and after pyrolysis at 600 C (Fig. 6b) and at 1000 C (Fig. 6(c) and (d)). The values of 4 nm and 6 nm respectively found for the pore diameter and the wall thickness of the 600 C heat-treated specimen are in very good agreement with the results previously provided by N2

Table 3 Specific surface areas, total pore volume and pore diameter as measured by N2 adsorption/desorption experiments Sample

SBET (m2/g) ± 2 m2/g

Sa-plot

Smeso (m2/g)

Vtotal (cm3/g)

Vmeso (cm3/g)

˚ ) ± 1% Umeso (A

Extracted 400 C-H.T. 600 C-H.T. 800 C-H.T. 1000 C-H.T.

999 938 664 – –

– – – 322 208

999 938 664 214 120

0.82 0.75 0.56 0.32 0.22

0.82 0.75 0.56 0.18 0.11

41 41 40 34 37

H.T.: heat-treated.

S. Masse et al. / Journal of Non-Crystalline Solids 353 (2007) 1109–1119 (110)

(100)

1116

(210)

(200)

x 30

1.3

1.5

1.7

1.9

2.1

2.3

2.5

1000˚C

Intensity(a.u.)

800˚C

600˚C

400˚C

Extracted

As-synthesized

0

1

2

3 2 θ(˚)

4

5

6

Fig. 4. XRD patterns of the as-synthesized and of the extracted sample calcined at various temperatures.

Table 4 Evolution of the (1 0 0) peak position as a function of calcination temperature as measured by XRD PMO

2h ()/d(1 0 0) ˚ ) ± 1% (A

Average a ˚ ) ± 1% (A

1 a/a0 (%)

As-synthesized Extracted 400 C-H.T. 600 C-H.T. 800 C-H.T. 1000 C-H.T.

0.859/102.7 0.827/106.7 0.888/99.4 0.945/93.4 1.102/80.1 1.086/81.3

119 123 115 108 93 94

– 0 6.5 14 24 24

H.T.: heat-treated.

adsorption/desorption experiments and XRD measurements. After calcination at 1000 C, numerous wellordered hexagonal arrays of 1D mesoporous channels surrounded by an amorphous worm-like matrix can still be observed, but the pore-center to pore-center distance is

now estimated to be 9 nm and the wall thickness to 5 nm, that is rather less than the values found previously for the 600 C heat-treated PMO. 4. Discussion We achieved in synthesizing a well-organized PMO material with a very good specific surface area (SBET  ˚ ), 1000 m2/g), a narrow pore size distribution (35–45 A and thick walls (e  8 nm). Despite porous ordered areas are still present in the material until 1000 C, the surface area and pore volume principally drop down in the range 600–800 C. At this stage, solid state 29Si and 13C NMR reveal that the C–C linkages of the organic bridges and Si–C bonds are progressively cleaved with the formation of an aromatic free-carbon phase. The poor stability of the Si–C bonds should probably be related to the rather

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Fig. 5. Bright field TEM micrographs of the 400 C heat-treated PMO (a) and 1000 C heat-treated PMO (b).

Fig. 6. TEM micrographs of microtomes of the extracted sample before calcination (a), after pyrolysis at 600 C (b) and 1000 C ((c), (d)).

large amount of silanol groups present in the as-prepared sample, as suggested by Hatton et al. [4] who proposed a reaction mechanism between Si–C bonds and silanol groups converting a bridging methylene groups into a terminal methyl group. By analogy, such a mechanism seems

us to be likely also for the ethylene group, finally leading to a terminal methyl group formation, in accordance with the intense broad peak we observed from d = 2 to 6 ppm by 13 C CP-MAS NMR. Similar results were indeed previously obtained in the same group [33] and this signal was

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Table 5 Estimation of the pore wall thickness calculated by difference between the cell parameter a as deduced from the angular deviation of the (1 0 0) reflexion measured by XRD and the average mesopore diameter Umeso as measured by N2 adsorption/desorption experiments ˚ ) ± 1% ˚ ) ± 1% ˚ ) ± 1% a (A e (A Sample Umeso (A Extracted 400 C-H.T. 600 C-H.T. 800 C-H.T. 1000 C-H.T.

41 41 40 34 37

123 115 108 93 94

82 74 68 59 57

pare porous SiOC glasses. Quantitative information concerning the structure evolution was collected with various techniques with the aim of probing: (i) the short range order and chemical composition (29Si MAS and 13C CP-MAS NMR, FT-IR and elemental analysis) and (ii) the long range order (X-ray diffraction) and microstructural features (TEM, porosity and surface area measurements). The main results concerning the thermal evolution of the 2D-hexagonal PMO material can be summarized as follows:

H.T.: heat-treated.

unambiguously assigned to Si–CH3 species with the aid of a 13C CP-MAS NQS NMR experiment. By 800 C, none of the Si–C bond survived the pyrolysis treatment leading to a sample that can be described as a two-phase system: amorphous SiO2 and free-C. The specific surface area decreased to 100–200 m2/g with appearance of microporosity. At 1000 C, it represents half of the total pore volume. XRD, N2 adsorption–desorption data and TEM observations are consistent with those previously reported for silica-based SBA-15 [32]. XRD experiments showed that the network is still organized after pyrolysis at 1000 C whereas N2 adsorption–desorption measurements seem to indicate a loss of open porosity above 800 C, likely due to the formation of free carbon inside the pores, that probably blocks their access to a certain level. All the methods used lead to the same estimation of the pore diameter and wall thickness for this material. The comparison of the pore size measurements with the cell parameter obtained from XRD powder patterns allows to estimate the wall thickness of the network and its evolution with pyrolysis temperature (Table 5). The structural changes induced by temperature have implications on the average wall thickness of the network, which drops down from 8.2 nm before heat-treatment to 5.7 nm at 1000 C. The excess carbon which appears as a free-carbon phase during pyrolysis process of polymeric precursors was precisely investigated by Trassl et al. [34,35] with a panel of spectroscopic techniques. According to these authors, this phase appears around 500–600 C after cracking and cyclization that occur during the pyrolysis of aliphatic hydrocarbons. Because paramagnetic centers exist in the material – as delocalized electrons coupled to sp2-C domains – ESR should be a powerful technique to undertake, as well as Raman spectroscopy, complementarily to 13 C NMR, to observe the aromatic C-nanodomains. In a complementary study based on the same approach but using Pluronic F127 block copolymer instead of Pluronic P123, we achieved in observing the so-typical turbostratic free-carbon regions by HR-TEM at high temperature. 5. Conclusion The pyrolysis under argon of SBA-15-like ethanesilicas has been studied in order to evaluate the possibility to pre-

• The C–C linkage of the Si–CH2–CH2–Si bridge is still existing in the network after pyrolysis at 600 C during 2 h. At the same time, a great part of the Si–C bonds are not yet broken. But evidence for the appearance of a free-C phase coexisting with the silica phase is observed. • A phase separation then occurs above 600 C between amorphous silica and free carbon. At 800 C, the two phase separation is complete and a large decrease of the specific surface area occurs with formation of microporosity. • From a microstructural point of view, increasing the pyrolysis temperature has logically led to a reduction of the pore size (from 4.1 nm before pyrolysis to 3.7 nm at 1000 C) as well as of the wall thickness (from 8.2 nm before pyrolysis to 5.7 nm at 1000 C) of the material as a consequence to the cell contraction. Pyrolysis under argon retains an ordered structure up to 1000 C and well-ordered areas are still present in the 1000 C pyrolyzed sample. Acknowledgments We are very grateful to Fabienne Warmont (CNRS/Service Commun de Microscopie de Paris 6) for preparing and examining the microtomed specimens by TEM. We would also like to thank Jocelyne Maquet for her precious contribution to NMR investigations. References [1] E.-B. Cho, K. Char, Chem. Mater. 16 (2) (2004) 270. [2] S.A. Bagshaw, E. Prouzet, T.J. Pinnavaia, Science 269 (1995) 1242. [3] P. Kipkemboi, A. Fogden, V. Alfredsson, K. Flodstro¨m, Langmuir 17 (2001) 5398. [4] B. Hatton, K. Landskron, W. Whitnall, D. Perovic, G. Ozin, Acc. Chem. Res. 38 (4) (2005) 305. [5] W. Wang, S. Xie, W. Zhou, A. Sayari, Chem. Mater. 16 (9) (2004) 1756. [6] C.-Y. Ishii, T. Asefa, N. Combs, M.J. MacLachlan, G.A. Ozin, Chem. Commun. (1999) 2539. [7] W. Wang, W. Zhou, A. Sayari, Chem. Mater. 15 (26) (2003) 4886. [8] S. Inagaki, S. Guan, T. Ohsuna, O. Terasaki, Nature 406 (2002) 304. [9] Y. Goto, S. Inagaki, Chem. Commun. (2002) 2410. [10] E.B. Cho, K. Char, Chem. Mater. 16 (2) (2004) 270. [11] E.B. Cho, K.W. Kwon, K. Char, Chem. Mater. 13 (2001) 3837. [12] T. Rouxel, G. Massouras, G.D. Soraru`, J. Sol–gel Sci. Technol., Special Issue on Silicon Oxycarbide Glasses by Chemical Route 14 (1) (1999) 87.

S. Masse et al. / Journal of Non-Crystalline Solids 353 (2007) 1109–1119 [13] G.D. Soraru`, S. Modena, E. Guadagnino, P. Colombo, J. Egan, C. Pantano, J. Am. Ceram. Soc. 85 (2002) 1529. [14] F. Babonneau, K. Thorne, J.D. Mackenzie, Chem. Mater. 1 (1989) 54. [15] F. Babonneau, G.D. Soraru`, G. D’Andrea, S. Dire, L. Bois, J. Mater. Soc. Symp. Proc. 271 (1992) 789. [16] See papers in: J. Sol–gel Sci. Technol., Special Issue on Silicon Oxycarbide Glasses by Chemical Route 14 (1) (1999). [17] H. Bre´quel, J. Parmentier, S. Walter, R. Badheka, G. Trimmel, S. Masse, J. Latournerie, P. Dempsey, C. Turquat, A. DesmartinChomel, L. Le Neindre-Prum, U.A. Jayasooriya, D. Hourlier, H.-J. Kleebe, G.D. Soraru`, S. Enzo, F. Babonneau, Chem. Mater. 16 (13) (2004) 2585. [18] B. Toury, R. Blum, V. Goletto, F. Babonneau, J. Sol–gel Sci. Technol. 33 (2005) 99. [19] B. Toury, F. Babonneau, J. Eur. Ceram. Soc. 25 (2005) 265. [20] F. Kleitz, W. Schmidt, F. Schu¨th, Micropor. Mesopor. Mater. 65 (2003) 1. [21] O. Muth, C. Schellbach, M. Fro¨ba, Chem. Commun. (2001) 2032. [22] M.C. Burleigh, M.A. Markovitz, E.M. Wong, J.-S. Lin, B.P. Gaber, Chem. Mater. 13 (2001) 4411. [23] B.-G. Park, W. Guo, X. Cui, J. Park, C.-S. Ha, Micropor. Mesopor. Mater. 66 (2003) 229. [24] H. Zhu, D.J. Jones, J. Zajac, J. Rozie`re, R. Dutartre, Chem. Commun. (2001) 2568.

1119

[25] H. Zhu, D.J. Jones, J. Zajac, R. Dutartre, M. Rhomari, J. Rozie`re, Chem. Mater. (2002) 4886. [26] J.R. Matos, M. Kruk, L.P. Mercuri, M. Jaroniec, T. Asefa, N. Coombs, G.A. Ozin, T. Kamiyama, O. Terasaki, Chem. Mater. 14 (2002) 1903. [27] W. Guo, J.-Y. Park, M.-O. Oh, H.-W. Jeong, W.-J. Cho, I. Kim, C.-S. Ha, Chem. Mater. 15 (2003) 2295. [28] W. Guo, I. Kim, C.-S. Ha, Chem. Commun. 21 (2003) 2692. [29] D. Massiot, F. Fayon, M. Capron, I. King, S. Le Calve´, B. Alonso, J.O. Durand, B. Bujoli, Z. Gan, G. Hoatson, Magn. Reson. Chem. 40 (2002) 70, . [30] S. Fiorilli, B. Onida, B. Bonelli, E. Garrone, J. Phys. Chem. B 109 (2005) 16725. [31] J. Liu, Q. Yang, M.P. Kapoor, N. Setoyama, S. Inagaki, J. Yang, L. Zhang, J. Phys. Chem. B 109 (2005) 12250. [32] D. Zhao, Q. Huo, J. Feng, B.F. Chmelka, G.D. Stucky, J. Am. Chem. Soc. 120 (1998) 6024. [33] T. Asefa, M.J. MacLachlan, H. Grondey, N. Coombs, G.A. Ozin, Angew. Chem. Int. Ed. 39 (10) (2000) 1808. [34] S. Trassl, G. Motz, E. Ro¨ssler, G. Ziegler, J. Am. Ceram. Soc. 85 (1) (2002) 239. [35] S. Trassl, H.-J. Kleebe, H. Sto¨rmer, G. Motz, E. Ro¨ssler, G. Ziegler, J. Am. Ceram. Soc. 85 (5) (2002) 1268.