sic preparation

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Applied Catalysis A: General 187 (1999) 255268

Preparation and characterization of SiC microtubes

Nicolas Keller a, Cuong Pham-Huu a, Marc J. Ledoux a,, Claude Estournes b, Gabrielle Ehret ca Laboratoire de Chimie des Matriaux Catalytiques, ECPM, Universit, Louis Pasteur, 25, rue Becquerel 67087 Strasbourg Cedex 2, France

b Groupe des Matriaux Inorganiques, IPCMS, UMR 7504 du CNRS, 23 rue de Loess, 67037 Strasbourg Cedex, Francec Groupe Surface & Interface, IPCMS, UMR 7504 du CNRS, 23 rue de Loess, 67037 Strasbourg Cedex, France

Received 24 March 1999; received in revised form 2 June 1999; accepted 2 June 1999

Abstract

Silicon carbide microtubeswith medium surface area,3070 m2 g1, were successfully prepared by shapememory synthesis

involving the reaction between SiO vapors and low surface area, 12 m2 g1, carbon microfibers. The gross morphology of

the carbon microfibers was conserved during the carburization process. After calcination at 600C in order to eliminate the

unreacted carbon, hollow SiC microtubes were obtained. The pore size distribution of the material was centered around 10 nm,

allowing a high accessibility of potential reactants to an active phase during catalytic reactions. The surface was covered by

an amorphous layer 3 nm thick. XPS measurements revealed that this amorphous phase was composed of a mixture of SiO2and SiOxCy. Soda treatment at 80

C allowed complete removal of this surface phase without any change in the material

morphology. A similar result was also obtained after treatment with an aqua regia medium. 1999 Elsevier Science B.V. All

rights reserved.

Keywords: Silicon carbide; Microtubes; High surface area; Catalyst support

1. Introduction

Activated carbon has been widely used as a cata-

lyst support for active phases in liquid-phase processes

because of its high resistance towards aggressive en-

vironments, its possible separation from the reaction

media and also because of the relatively simple recov-

ery of the metals in the active phase at the end of the

catalysts life. However, attrition problems under vig-

orously agitated liquid can lead to the loss of active

components and to a decrease in the catalyst particle

size, which renders more difficult the separation of

the catalyst from the liquid phase. For these reasons,

Corresponding author. Tel.: +33-3-88-13-68-81; fax: +33-3-88-

13-68-80

E-mail address: [email protected] (M.J. Ledoux)

a growing interest has appeared for the use of a new

type of catalyst support based on carbon fibers [1,2].However, the high affinity of carbon for oxygen at rel-

atively low temperatures also renders the use of car-

bon as a catalyst support less attractive, i.e. support

loss during oxidative regeneration. It was then of in-

terest to search for substitutes for carbon having the

same positive properties but without the drawbacks.

Silicon carbide (SiC) exhibits a high thermal con-

ductivity and is chemically inert until 800C. The high

thermal conductivity avoids the formation of hot spots

during regeneration and the chemical inertness allows

easy recovery of the active phase by acidic or ba-

sic washing. In addition, when used as catalyst sup-

port, silicon carbide offers a unique advantage overtraditional materials such as alumina or silica since

it is also an excellent electrical conductor. Therefore,

0926-860X/99/$ see front matter 1999 Elsevier Science B.V. All rights reserved.

PII: S 0 9 2 6 - 8 6 0 X ( 9 9 ) 0 0 2 2 3 - 9


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256 N. Keller et al. / Applied Catalysis A: General 187 (1999) 255268

there is the possibility that metal crystallites intro-

duced onto silicon carbide will form an interaction

with the surface, and as a consequence, modifications

in the morphological characteristics of the catalyst

particles could give rise to unexpected activity and

selectivity patterns. Similar results were reported by

Kaneko [3] for carbon nanofiber catalyst support.

However, for silicon carbide to be useful as cata-

lyst support, it must be prepared in a medium surface

area form (20100 m2 g1) with an appropriate pore

size distribution and the inability to do this has been

the limiting factor in its application as support for het-

erogeneous catalysis. For this reason, considerable at-

tention has been focused on developing methods for

preparation, either directly via thermal treatment of

the precursors or assisted by a gasification catalyst,

that will yield high surface area materials [410]. The

high surface area SiC synthesized from these tech-

niques was successfully used as catalyst support for

many reactions such as the hydrodesulfurization reac-tion [6], CH4 reforming [9], automotive exhaust pipe

reactions [10,11], isomerization of linear saturated hy-

drocarbons [12] and selective oxidation of hydrogen

sulfide for Claus tail-gas treatment [13]. More de-

tailed studies concerning the synthesis, characteriza-

tion and catalytic use of silicon carbide as catalyst

support have been published by Lednor and Chorley

[14], by Ledoux and Pham-Huu [15] and by Ledoux,

Pham-Huu and Chianelli [16]. Finally, due to the low

sinterability of silicon carbide, it must be convenient

to find a method for preparation which can directly

yield the material in its final form (grain, cylinder or

honeycomb) without an additional need for shaping.The aim of the present article is to report the prepa-

ration of high surface area silicon carbide microtubes

from low surface area carbon fibers (12 m2 g1)

according to the shape memory synthesis method

developed by Ledoux et al. [5,6] for use as a hetero-

geneous catalyst support material. The silicon carbide

microtubes (after synthesis and after air calcination

in order to burn off the remaining carbon) were

characterized by different techniques such as powder

X-ray diffraction (XRD), thermal gravimetry anal-

ysis (TGA), and surface area and pore distribution

measurements by nitrogen adsorption. The morphol-

ogy and microstructure of the different solid materi-als were observed by scanning electron microscopy

(SEM) and high-resolution transmission electron mi-

Table 1

Characteristics of the carbon microfibers

Carbon 9497 wt.%

Ash 0.10.3 wt.%

Volatile compounds 1.54.0 wt.%

Sulfur 0.30.5 wt.%

Average diameter 1015mBET surface area 12 m2 g1

croscopy (HRTEM). The nature of the phases which

composed the uppermost layers of the material was

investigated using the XPS technique. The resistance

of the sample in an aggressive environment (strongly

acidic) was also investigated. The morphology of the

sample after such a treatment was observed by SEM.

2. Experimental section

2.1. Raw materials

The carbon microfibers were supplied by Carbone

Lorraine Com. (purity 97.5%) and their characteristics

are summarized in Table 1. Silicon (Janssen, 99.9%)

and silica (Merck, 99.5%) were used in a powder form.

2.2. SiC synthesis apparatus

A sketch of the apparatus is given in Fig. 1(a). The

carbon microfibers and the Si and SiO2 mixture were

located in two volumes separated by 3 cm in a den-

sified alumina crucible (inner diameter 60 mm, length

100 mm), which was not active towards the reactants

and the material formed (Fig. 1(b)). The SiO vapor

was generated by heating the mixture of silicon and

silica powder, located in the lower part of the alumina

crucible. The SiO vapor was then pumped through the

carbon fiber bed maintained at ca. 1250C where sil-

icon carbide was formed according to the following

reaction:

SiO(gas) + 2C(solid) SiC(solid) + CO(gas) (1)

The alumina crucible was introduced into an imper-

meable densified silicon carbide tube (inner diameter

130 mm, length 800 mm) in which a dynamic vacuum(p = 0.05 Torr) could be maintained. The dynamic vac-

uum was achieved via a turbo pump, the CO formed


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N. Keller et al. / Applied Catalysis A: General 187 (1999) 255268 257

Fig. 1. (a) Schematic diagram of the apparatus used in the synthesis of SiC microtubes. (b) Schematic diagram of the crucible used in the

synthesis of SiC microtubes.

was pumped out of the reactor and allowed the dis-placement of the reaction in favor of SiC formation.

Before the reaction, the carbon microfibers were

evacuated at 1000C (heating rate = 10Cmin1) for

2 h in order to desorb the impurities from its surface.

The temperature was increased from 1000C to the

synthesis temperature and kept at this value for 15 h.

After synthesis, the sample was cooled to room tem-

perature under vacuum and then stored in a closed

vessel before characterization.

It has been reported in our previous publications

that, depending on the reaction conditions, a part of

the unreacted carbon remained in the core of the ma-

terial. In order to obtain clean pure SiC microtubes, itwas necessary to burn off this unreacted carbon. The

sample was separated into two batches: one was char-

acterized without any treatment and the second wascalcined in air at 600C for 2 h.

2.3. Characterization techniques

Structural characterization of the samples was done

by powder XRD measurements, carried out with a

Siemens Diffractometer Model D-5000, using a Cu K

radiation. The measurements were made with long du-

ration scan (10 s) and a small step scan (0.02 2). The

mean crystallite sizes were determined from the Scher-

rer equation with the normal assumption of spheri-

cal crystallites. The nature of the crystalline phases

present in the different samples was checked using thedata base of the Joint Committe on Powder Diffrac-

tion Standards (JCPDS).


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TGA experiments were performed using a Balzer

thermo-analyser to determine the total amount of

carbon remaining in the material. The sample was

placed in a platinum crucible and heated from room

temperature to 1000C (heating rate= 10Cmin1)

using a 20% (v/v) O2/N2 mixture at a flow rate of

5 cm3 min1. During the heating process, the weight

change of the catalyst was monitored continuously.

The pore size and surface area measurements were

performed on a Coulter SA-3100 porosimeter using N2as adsorbant. The sample, after treatment, was trans-

ferred to the BET cell via a glove-box under dry nitro-

gen. Before each measurement, the sample was evacu-

ated at 300C for 3 h in order to desorb the impurities

adsorbed on its surface. The cell was equipped with a

greaseless valve in order to avoid air exposure of the

sample during the transfer to the porosimeter. SBET is

the surface area of the sample calculated from the ni-

trogen isotherm using the BET method. SBJH is the sur-

face area of all the pores except micropores calculatedfrom the N2 desorption isotherm. The micropore sur-

face area and volume were calculated using the t-plot

method developed by de Boer and co-workers [17]. A

more detailed study has been published by Mikhail et

al. [18] concerning the correctness of the different pa-

rameters used in the method. The t-plot consists of the

analysis of the vlt plot curve where vl is the volume

of nitrogen adsorbed as liquid at a given pressure P/P0by the BET surface and t is the statistical thickness

obtained by dividing the volume of nitrogen adsorbed

as liquid at a given pressure P/P0 by the BET surface.

The combination of the t-plot and the BJH method for

narrow and larger pores allows the analysis of a nearlycomplete pore volume and pore surface distributions

in the sample studied.

The morphology of the material was observed by

SEM using a Jeol microscope Model JMS-840 oper-

ated with an accelerating voltage of 20 kV. The sam-

ples were pre-covered by a thin layer of gold in order

to avoid the charging effect problem during analysis.

Surface characterization was carried out by XPS.

XPS spectra were recorded on a Cameca Nanoscan

50 using the Al K source at 1486.6 eV. The en-

ergy scale of the instrument was calibrated using Ag

3d5/2 = 368.0 eV and Au 4f7/2 = 84.0 eV. The C 1s at

283.3 eV and the Si 2p at 103.0 eV, corresponding, re-spectively, to the carbon engaged in the carbidic phase

and the silicon in the silica phase, were both used as

reference, by comparison with the published literature.

The resolution in energy was 0.7 eV and the vacuum

was maintained at below 5 109 atm. The sample

was pressed on an indium foil and then transferred

under air into the XPS preparation chamber. Only C

1s and Si 2p were recorded, and for a given emission,

the number of accumulations was selected in order to

achieve a good signal-to-noise ratio. The samples were

analyzed without any pretreatment, i.e. ion etching, in

order to avoid any surface modification such as selec-

tive reduction or composition modification. Indeed, it

has been reported by several authors in the literature

that Ar+ ion bombardment induced preferential ele-

mental sputtering thus leading to an inaccurate analy-

sis.

TEM and EDS were carried out in a Topcon Model

EM200B operating at 200 kV with a point to point res-

olution of 0.17 nm. To prevent artifacts due to contam-

ination, no solvents were used at any stage and sam-

ples were prepared by grinding the catalysts betweenglass plates and bringing the powder into contact with

a holey carbon-coated copper grid. Great care was

taken during the TEM experiments in order to avoid

heating effects from the incident beam. EDS analysis

was carried out with a 5 nm and 100 s life time elec-

tron beam. The detector was equipped with beryllium

windows, and thus, it allowed the detection of light

elements such as carbon or oxygen.

3. Results and discussion

3.1. Synthesis of cubic -SiC microtubes

Typically, 4.0 0.1 g of carbon microfibers was

used for each experiment. The reaction temperature

was set between 1200 and 1300C and the duration

fixed at 15 h. The reaction conditions and the corre-

sponding CSiC conversion, given as examples, are

summarized in Table 2. The total amount of carbon

was never transformed into SiC even with a long du-

ration of reaction, meaning that SiC conversion could

be influenced by modifying the reaction conditions.

This was explained by several factors: (i) as a function

of the transformation, the SiC layers formed on thecarbon surface acted as a barrier for the diffusion of

SiO and CO vapors, consequently decreasing the rate


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Table 2

Influence of the synthesis parameters on the C SiC conversion and the surface area of the SiC microtubes synthesized by the gassolid

reaction between SiO vapor and carbon fibers

Temperature, Si/SiO2 molar ratio,

(Si+SiO2)/C mass ratio

Conversion

(wt.%)aSBET after synthesis

(m2 g1 /sample) (m2 g1/SiC)

SBET after calcination

(m2 g1)

1300C/1/12 85 18 (21) 30

1300C/1/8 60 18 (30) 45

1200

C/1/8 18 25 (138) 66a The conversion was calculated based on the starting weight of the carbon fibers.

of carburization; and (ii) the carbon fibers used in this

preparation have a very low surface area (few square

meters which justified the corrected value of the sur-

face area in Table 2) compared to the high surface area

activated charcoal used in other works [6]. Such an

ordered and non-porous structure could mitigate SiC

nucleation and also limit the vapor diffusion in the

material. Under similar reaction conditions, the total

conversion CSiC was lower when starting from

the low surface area carbon fibers than when startingfrom high surface area activated charcoal [19]. Simi-

lar results have already been reported by Moene et al.

[9] during his investigation on SiC synthesis. He has

observed that the number of nucleation sites on acti-

vated charcoal was higher compared to that observed

on the ordered graphite material, resulting in a more

grained structure of the final SiC. In addition, it has

been reported by Benaissa et al. [20] that SiC forma-

tion was initiated in several domains throughout the

activated charcoal matrix and especially on the defect

points and micropores of the activated carbon surface

which had a high adsorption energy. The higher rate

of SiC formation observed on the activated charcoal

was also attributed to the high porosity which could

favor SiO vapor diffusion.

3.2. Crystalline structure and oxidation behavior

The XRD patterns of the SiC microtubes prepared

at 1300C for 15 h are presented in Fig. 2. Whatever

the reaction temperature and duration, XRD showed

the presence of only silicon carbide as the sole crys-

talline reaction product. The synthesized SiC exhib-

ited diffraction lines corresponding to -SiC crystal-

lized in a face centered cubic structure. In addition tothe -SiC diffraction lines, a broad peak located at a

low two-theta angle was observed for the sample syn-

Fig. 2. XRD patterns of the SiC microtubes after synthesis (a) and

after calcination in air at 600C for 2 h (b). The diagrams were

recorded with a step of 0.02 (2 ) and a scan time of 10 sec per

step.

thesized at 1300C, which could be attributed to the

unreacted remaining carbon. The reaction temperature

used (12001400C) allowed the formation of small

amounts of stable -SiC along with -SiC, evidencedby a peak located at around diffraction angles of 33.7

corresponding to -SiC in the 4H and 15R structures.


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This phenomenon was attributed to the formation of

stacking faults along the [111] direction in the SiC

particles, as will be shown by TEM below, and accord-

ing to the SiC microstructure published previously us-

ing HRTEM [20]. No traces of SiO2 or silicon were

detected along with the carbide XRD pattern, which

means that such species, if present, were only in an

amorphous form or in very small amounts.

The XRD patterns of silicon carbide synthesized at

1300C for different durations (not reported) showed

the increase in the silicon carbide diffraction lines at

the expense of those corresponding to the starting car-

bon, meaning that a relatively long time was needed

to transform more of carbon microfibers into SiC mi-

crotubes. It was significant to note that carburization

was realized by direct exchange between carbon and

oxygen atoms and no traces of silicon were detected

on the intermediate samples. The relatively long time

needed in order to obtain a high CSiC conversion

was attributed to the nature of the transformation, i.e.a Shrinking Core Model [21,22], where SiC formation

depended on the diffusion of the SiO vapor through

the SiC mantle and also to the relatively low reaction

temperature used in this study. The Shrinking Core

Model states that the reactional interface is maximum

at the beginning of the transformation and then de-

creases regularly as the reaction takes place. At the

beginning of carburization, all of the carbon surface

was accessible and available to the SiO vapor; as soon

as SiC formation started, the transformation rate de-

creased due to the diffusion limitation of the SiO and

CO vapors through the first SiC layers. The final ma-

terial structure could be described as silicon carbidemicrotubes containing a carbon core.

The behavior of a support under an oxidizing atmo-

sphere is important in order to check its resistance dur-

ing oxidative regeneration. The TGA diagrams of the

carbon fibers and of the as-synthesized material under

a 20% (v/v) O2/N2 flow are shown in Fig. 3. The car-

bon microfibers remained stable up to around 500C

and then rapidly burned completely (Fig. 3(a)). The

SiC microtubes (containing carbon) were stable up to

600C (Fig. 3(b)). At around 600C, a weight loss was

observed due to complete combustion of the remain-

ing carbon located in the core of the tubes. For higher

temperatures, the SiC microtubes remained stable andno weight loss was observed up to 800C where the

oxidation of SiC by air could begin to result in a weight

Fig. 3. Thermogravimetry analysis of the carbon microfibers (a)

and SiC microtubes after synthesis (b).

gain. Similar observations have already been reported

by Stegenga et al. [23] during their investigation of the

thermal resistance of a support for automotive exhaust

pipe reactions. In conclusion, it seemed that carboncovered by a layer of SiC (as shown by the SEM ob-

servation below) hindered oxygen penetration, mean-

ing that the SiC coating acted as a diffusion barrier

for carbon oxidation. The small increase in weight ob-

served at high temperatures was attributed to the oxi-

dation of the SiC surface according to Eq. (2) and as

per the results reported by Moene [41].

SiC(s) +32 O2(g) SiO2(s) + CO(g) (2)

The SiC after calcination in air at 600C for 2 h ex-

hibited an XRD pattern (Fig. 2(b)) similar to that ob-

served before calcination. The broad peak located at

the low two-theta angle disappeared, meaning that allof the remaining low surface area carbon was elimi-

nated by the oxidative treatment. In addition, no trace


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Fig. 4. Nitrogen adsorptiondesorption isotherms of carbon mi-

crofibers (1), SiC microtubes after synthesis (2) and SiC micro-

tubes after calcination in air at 600C for 2 h (3). Filled points

correspond to adsorption curves and empty points to desorption

curves.

of silica or other oxides were observed. Such a phe-nomenon could be explained by the fact that silicon

dioxide, if formed, was located mainly on the mate-

rial surface (see XPS results below) or in an amor-

phous form, neither of which is clearly shown by the

XRD technique. The relatively low oxidation temper-

ature used in the present work also explained the lack

of bulk oxidation of the sample.

3.3. Surface area and porous measurements

To be used as catalyst supports, materials must

be prepared with a relatively high surface area, i.e.10 m2 g1, which allows the exposure of the active

phase to the reactant mixture. The material must also

contain as few micropores as possible, which could

hinder the access to the reactants. In this section,

investigation of the evolution of the surface area of

the SiC microtubes as a function of the calcination

treatment along with the pore size distribution modi-

fication is described.

The nitrogen isotherm adsorptiondesorption

curves measured for the initial carbon fibers, the

as-synthesized SiC microtubes and the material

after calcination in air are shown in Fig. 4. The

adsorptiondesorption isotherms of the starting car-bon fibers were found to be of Type III according

to the well-accepted isotherm classification [24,25],

Fig. 5. Pore size distribution of SiC microtubes after synthesis and

after calcination in air at 600C for 2h.

and this shape indicated the existence of macropores.

The initial specific surface area of the carbon fibers,

calculated using the standard BET method, was verylow (12 m2 g1).

The nitrogen isotherm of the SiC tubes was drasti-

cally modified compared to that obtained on the start-

ing carbon fiber, with an increase in the porous volume

from 0.010 to 0.058 cm3 g1. The Type IV isotherms

observed were typical of a solid containing mesopores

ranging from 2.5 to 50 nm. The t-plot method per-

formed on the material showed a straight line cross-

ing the origin, meaning that almost no micropores

(2.5 nm) were present in the material.

The pore size distributions of the SiC tubes before

and after calcination in airare presented in Fig. 5. After

synthesis, an important fraction of the macropores ofthe carbon fibers precursor had disappeared, leading

to the formation of a large number of mesopores lo-

cated around 1011 nm and to an increase in the sur-

face area from 12 m2 g1 to around 12 m2 g1. The

increase in the material surface area when going from

C to SiC was probably due to the formation of sev-

eral nuclei at the beginning of the synthesis, followed

by the formation of small domains of SiC, and thus,

creation of pores which increased the surface area of

the material. Carbon consumption during the course

of the transformation also created voids inside the ma-

terial and contributed to the increase in the sample

surface area. Such a phenomenon has already been re-ported by Benaissa et al. [20] during their investigation

on the mechanism of SiC growth by HRTEM. This


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explanation will be confirmed by the SEM micro-

graphs reported below. However, it is significant to

note that the value of the surface area given here was

underestimated due to the presence of unreacted car-

bon with low surface area inside the material.

The remaining carbon was eliminated from the SiC

material by calcination of the sample in air at 600C

(according to the results obtained by the TGA tech-

nique) for 2 h. The surface area of the sample after

calcination was significantly increased from 12 to

45 m2 g1, and thus, it highlighted the advantage of

the gassolid synthesis method employed in this work

to produce high surface area SiC tubes from the start-

ing low surface area carbon fibers. The average distri-

bution of the pores was left unchanged and centered

around 1011 nm which could allow a good disper-

sion of an active phase and its accessibility without

microdiffusion phenomena. The absence of modifica-

tion of the pore size distribution after carbon removal

strongly supported the fact that such a treatment onlyremoved carbon inside the fiber, without sintering or

additional formation of other compounds with a dif-

ferent pore size distribution. The higher final surface

area of the calcined sample was also due to the disap-

pearance of carbon, leading to the formation of empty

fibers or tubes, as will be observed by SEM later. Such

a transformation is sketched in Fig. 6. However, one

could also suggest that the increase in surface area

after air treatment could also be attributed to the for-

mation, on the SiC surface, of a layer of amorphous

silica with high surface area. This hypothesis was re-

jected as the calcined material surface area, 25 m2 g1

in one sample, remained unchanged after soda treat-ment at 80C which removed the silica or oxycarbide

layers present on the surface, 24 m2 g1 (see XPS

results below). These results were in line with those

reported by Moene [9] who observed that almost no

SiO2 appeared on the SiC surface at a temperature

around 800C even after more than 2 h of oxidation.

It is significant to note that the material surface

area was strongly dependent on the reaction tempera-

ture, 66 m2 g1 at 1200C compared to 45 m2 g1 at

1300C. Such a phenomenon was attributed to the sur-

face diffusion process, as already reported by different

authors in the literature [26,27]. Elder and Kristic [26]

and Hase et al. [27] have shown that SiC started tosinter by surface diffusion at a temperature lower than

the theoretical sintering temperature. The corrected

values of the BET surface in Table 2 clearly illustrate

the sintering phenomenon. At low conversion (short

reaction time), SiC was not sintered and the surface

area was as high as 139 m2 g1, while for a longer

reaction time, this surface fell to 30 m2 g1. The rate

of the loss of surface area by a surface diffusion pro-

cess was proportional to the initial surface area of the

treated sample. This explained the relatively small de-

crease in the SiC surface area with increasing reaction

temperature. This behavior was completely different

from that observed when the starting carbon material

was a high surface area activated charcoal, for which

one could notice an important drop in the SiC surface

area with increasing reaction temperature [19].

3.4. Morphology and microstructure

Figs. 79 show the SEM micrographs of the carbon

precursor and the resulting SiC (after synthesis andafter air calcination at 600C). The SEM micrographs

of the carbon fibers showed an average diameter of

1015m with a smooth surface (Fig. 7(a) and (b)) in

agreement with the low surface area obtained by the

BET technique. Low magnification SEM micrograph

of the corresponding SiC microtubes synthesized at

1300C is presented in Fig. 8(a) and it shows that the

gross morphology of the carbon microfibers was re-

tained. Similar results have already been reported in a

previous article [6] during the silicon carbide synthesis

from other carbon shapes. Such a result was explained

by the fact that the transformation reaction between the

SiO vapors and carbon proceeded via the step-by-stepreplacement of C atoms by Si atoms, with the release

of CO, leading to the formation of SiC. Such a trans-

formation mechanism provides a final material with

the same gross morphology as that of the precursor.

High magnification SEM images of the resulting

SiC material are presented in Fig. 8(b) and (c). The mi-

crostructure of the SiC microtubes was strongly mod-

ified when compared to that observed on the starting

carbon microfibers: the smooth surface carbon fibers

were transformed into an homogeneous porous SiC,

with an average network structure of around 50 nm.

The morphology, in different magnifications, of

the SiC material after calcination in air at 600C ispresented in Fig. 9 and it shows that the gross mor-

phology of the sample was retained. Cut calcined SiC


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Fig. 6. Formation of SiC microtubes during calcination in air, starting from the SiCcarbon composite.

material revealed the presence of empty SiC micro-

tubes, formed by combustion of the unreacted carbon

inside the core of the material (Fig. 9(c)). Similar

results have already been observed during the syn-thesis of vanadium carbide, where the carbide fibers

obtained were empty [28]. The thickness of the SiC

wall and the tube internal diameter could be modified

by acting on the CSiC conversion, i.e. synthesis

parameters (Fig. 10).

The TEM image of the SiC material is presented in

Fig. 11. The lattice image of the as-synthesized SiC

tube is reported in Fig. 11(a), showing a high density of

stacking faults along the 111 direction. The stacking

faults can be viewed in the TEM micrograph as black

stripes in the grains. The high resolution TEM image

showed that the interplanar distance along this 111

direction was constant and equal to 0.25 nm. No longrange periodicity was observed, which led to the con-

clusion that the SiC formed was a one-dimensionally

disordered polytype, which grew preferentially along

the 111 direction, according to the literature and pre-

vious work [20,29]. Stacking faults are often classified

as growth faults during the layer-by-layer growth of a

crystal with close-packed structure when a new layer

is stacked incorrectly. This confirmed the step-by-step

mechanism of synthesis highlighted above. Similar re-

sults have already been reported by Wang et al. [30]

and by Koumoto et al. [31] during their investigations

of SiC formation and structure.

The high magnification TEM image also showedthat the SiC surface was covered by an amorphous

phase, 3 nm thick (Fig. 11(b)), which contained Si,

Fig. 7. SEM micrographs of the carbon microfibers.

O and C, as identified by EDS. The presence of an

amorphous phase, attributed to a silicon oxidic phase

on the SiC surface, has already been reported by

many authors in the literature [3236]. Zhu et al.

[37] have shown, by HRTEM, the presence of an

amorphous layer which seems to be an inherent partof the SiC whisker. The nature of the amorphous

phase observed on the SiC microtubes will be inves-


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Fig. 8. SEM micrographs of the SiC microtubes after synthesis.

tigated and discussed in the next section using XPS

analysis.

3.5. Surface investigation

In order to obtain further insight about the na-

ture of the different components present on the SiC

microtubes, surface XPS was performed. Dissym-

metries and line widths of the main C 1s and Si 2p

peaks showed that these signals do not correspond

to a unique chemical form. In order to resolve thedifferent chemical components, theoretical C 1s and

Si 2p spectra were constructed by addition of the

GaussianLorentzian spectra, each component being

representative of a chemical entity. The XPS spec-

tra (Figs. 12 and 13) showed that the SiC surface

was covered by an amorphous phase composed of at

least two silicon based components and not simply

by an SiO2 layer. The binding energy of these dif-

ferent phases which composed the amorphous layer

are reported in Table 3, in close agreement with those

reported in the literature. The existence of silicon sub-

oxide along with SiO2 was highly improbable since,

according to thermodynamic considerations, silicon

suboxide (Si


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Table 3

XPS of the C 1s and Si 2p binding energy and attribution of the C 1s and Si 2p components recorded from SiC microtubes (comparison

with the published literature)

C 1s (eV) Si 2p (eV)

I II III I II III

SiC SiOxCy CO SiC SiOxCy SiO2

SiC after synthesis 283.3 284.7 287.0 100.5 101.6 103.0SiC after synthesis followed by soda treatment 283.3 100.5

Literature

Portre et al. [33] 283.3 100.5 101.5 103.2

Ramahan et al. [32] 100.2 102.8

Bouillon et al. [34,35] 283.3 284.2 101.1 102.1 103.2

Fig. 10. SEM micrograph of the SiC microtubes synthesized at

1200C for 15 h, and calcined in air at 600C for 2 h, corresponding

to a CSiC conversion of 18%.

[33]. The second peak could be assigned to a silicon

based compound containing both carbon and oxygen,i.e. a silicon oxycarbide phase, SiOxCy: the presence

of oxygen provides an electron charge transfer from

silicon to oxygen, which increases the Si 2p binding

energy (Fig. 12(a)). In Fig. 13(a), the XPS spectrum

of C 1s is reported. Along with the carbon engaged

in the carbidic phase (BE = 283.3 eV), two other C 1s

peaks can be resolved unambiguously and the second

one attributed to the oxycarbide species.

Similar results have been reported by several au-

thors working on silicon based materials such as SiC

and Si3N4. Benaissa et al. [36], using a combination of

HRTEM and EDS analysis, have shown the presence

of a silicon containing amorphous layer with a thick-ness of 13 nm. Mozdrierz et al. [38] have reported the

stoichiometry of SiO1.52C0.61.05 for the amorphous

phase present on the SiC surface using a combination

of several techniques, i.e. EELS, HRTEM and EDS.

This resulthas also been confirmed by Moene [42] dur-

ing his investigations on Mo and Ni supported on SiC,

using the temperature-programmed technique. He ob-

served that oxides supported on SiC display a lower re-duction temperature compared to those obtained with

the same catalysts supported on bulk SiO2 prepared

and treated under similar conditions. Rahaman et al.

[32], using XPS analyses of the amorphous layer cov-

ering a silicon nitride or silicon carbide surface, have

suggested that such a layer was probably made up of a

silicon oxynitride or oxycarbide rather than pure SiO2.

XPS measurements have allowed Porte and Sartre [33]

to state the existence of a silicon oxycarbide phase

in silicon carbide fibers. Similar results were also re-

ported by Bouillon et al. [34] using the same tech-

nique, concerning silicon carbide fibers prepared by

pyrolysis of a polycarbosilane precursor, or by Pam-puch [39] during the oxidation of SiC.

The amorphous mixture of SiO2 and SiOxCy present

on the SiC surface could be removed by soda (20%

in an aqueous solution) treatment at 80C for 1 h. The

material, after treatment and water washing, was ex-

amined by XPS (Fig. 12(b) and Fig. 13(b)). After soda

treatment, the Si 2p peaks in SiO2 and SiOxCy disap-

peared totally and Si 2p engaged in SiC was the only

form observed (Fig. 12(b)). It must be pointed out that

the material after soda treatment at 80C was exposed

to air for a relatively long period, and that XPS anal-

ysis revealed almost no additional superficial oxida-

tion. Such a result could be explained by the fact thata large part of the oxidic amorphous phases present

on the SiC surface was formed during the course of


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Fig. 11. TEM micrographs of the SiC microtubes after synthesis,

showing the presence of a high density of stacking faults along

the [111] axis and an amorphous layer on the surface.

material synthesis, probably by the reaction between

C, SiO and oxygen in the reaction zone. The C 1s

XPS spectrum of the soda treated sample presented

in Fig. 13(b) only showed the carbon engaged in the

SiC form. The vanishing of the third peak observed

in the C 1s spectrum could be explained by the fact

that this peak was due to carbonate species dissolved

during the soda treatment, as already reported in the

literature [33].In summary, the SiC surface synthesized by the

gassolid reaction was not pure but covered by an

Fig. 12. XPS spectra of Si 2p of SiC microtubes after synthesis

(a) and followed by soda treatment (b).

amorphous layer with a thickness around 3 nm. Such

a layer was not only composed of SiO2 but proba-

bly also a mixture of SiO2 and SiOxCy phases which

strongly affected the surface properties of the mate-

rial, as compared to those observed on a bulk silica

material resulting in a higher ability to disperse metal

species deposited on it.

3.6. Stability of the SiC fibers in an aggressive

environment

As already pointed out in Section 1, the best advan-

tage offered by the usage of SiC as a catalyst support

is its total chemical inertness which allows its use foroperations in aggressive media. The stability of SiC in

such an aggressive environment was investigated by


13/14


Fig. 13. XPS spectra of C 1s of SiC microtubes after synthesis

(a) and followed by soda treatment (b).

treating SiC in a boiling aqua regia solution for 2 h.

The sample was then washed with demineralized wa-

ter until the pH equaled 7, filtered and dried in an oven

at 100C for 2 h. SEM observations of the morphol-

ogy of the sample, before and after treatment, show

that this morphology was kept unchanged. Similar re-

sults have been reported by Rubio et al. [40] and by

Moene [41] with boiling HNO3 solution.

4. Conclusion

Low surface area carbon fibers (12 m2 g1) werereadily converted to high specific surface area SiC mi-

crotubes (3070 m2 g1) after reaction with SiO va-

por at 12001300C under dynamic pumping accord-

ing to the shape memory synthesis. The results ob-

served underline the high potential of such a method

for preparing high surface area support material from

low surface area precursors. The material displayed a

high surface area without the presence of micropos-

ity, a good resistance towards attrition and a suitable

pore distribution (10 nm) for catalytic applications, es-

pecially in liquid-phase reactions. Furthermore, the

sample exhibited a high resistance towards oxidation

compared to the corresponding carbon fibers and an

excellent stability in an aggressive environment. The

catalytic applications of materials based on SiC micro-

tubes, especially for the liquid slurry dewaxing process

and selective oxidation, will be presented elsewhere.

Acknowledgements

The TGA and SEM experiments were performedat the Groupe des Matriaux Inorganiques of the

IPCMS (UMR 7504 of the CNRS). Helen Lamprell

is gratefully acknowledged for having performed part

of the experiments. J. Ohmet (Groupe Surfaces et

Interfaces) and Prof. J. Guille (Groupe des Matriaux

Inorganiques) of the IPCMS-UMR 7504 of the CNRS

are gratefully acknowledged for XPS experiments

and helpful discussions.

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