Nanostructured Silicon-based Powders and Composites

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1 Objectives and state-of-the-art ofnanocomposites

Michel Cauchetier and Andr Pierre Legrand

1.1 Introduction

Processes to make new ceramic materials are many and depend on the final desirable structureand macroscopic shape to be obtained. So silicon carbide can be used as an abrasive, refrac-tory, structural material or semiconductor with a negative temperature coefficient. Processesdeveloped to produce such materials are extremely diversified. A process is selected basedon the required properties and structural shape of the parts to be obtained. As an example, thepreparation of carbon fibres, which needs specific polymer precursors for spinning followedby heat and oxidation treatments, had been adapted to the preparation of silicon carbidefibres. Other processes of transformation, descended from physics or chemistry domainsare chemical vapour deposition, chemical liquid deposition, etc. (Figure 1.1). They are ableto present improved mechanical, thermal or electrical properties, not solely depending ona peculiar chemical composition but on the arrangement of the crystalline phases or the sizeof the grains. Nanocomposites constitute a new class of extremely diversified materials whichhave appeared in the last decade.

Monomer unit polymer precursor

Solid state thermolysis

Mechanical, electrical and thermal properties

Characterisation by XPS X-ray and neutron diffraction TEM, SEM, optic solid state NMR

Laser pyrolysis

Chemical vapour or liquid deposition

Liquid or gaseous precursor molecules

Synthesis

Analysis

Sinte

ring

Figure 1.1 Schematic representation of ceramic preparation showing the feedback process lead-ing to a new ceramic material involving three main steps: synthesis of a precursor;analytical methods; and sintering and validation of the physical and chemicalproperties.

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Vitreous grain boundary (YSiAION)

Si3N4 (SiAION-) SiC

80 nm

Micronano composite Nanonano composite

1 m

Figure 1.2 Ceramic/ceramic micronano and nanonano composites aspects (from Figures 7.8and 7.20(b)).

Sajgalik et al. (1996) proposed a classification into two main systems:

micronano composites in which the matrix is constituted by micrometric crystalsembedded into a nanometric phase. Moreover, Niihara (1991) distinguishes threedifferent configurations: intragranular, intergranular, intra/intergranular (Figure 1.2);

nanonano composites in which the nanometric grains of the matrix and the secondphase are spread uniformly (Figure 1.2).

What makes a nanocomposite especially interesting is that at least one of its phases hasdimensions in the nanometer range (10100 nm). In this range, chemical and physical inter-actions have critical length scales and, if the nanoscale building block is made smaller thanthis one, the corresponding fundamental properties can be to changed. An example of this islight scattering, where by controlling the size of pores and nanocrystals in the range of 8 nm,Nanophase Technologies Corporation was able to produce transparent ceramics in the visibledomain. Other enhanced properties have been developed (superplasticity, magnetoresistance,low temperature densification, enhanced and finer homogeneity, etc.).

1.2 Nanostructured ceramic elaboration

Two main routes are generally used, depending on the densification rate and the propertiesto be obtained.

1.2.1 Amorphous ceramics obtained by solid state thermolysis ofpolymer precursors

Pyrolysis, under a controlled atmosphere, of polymers, such as polysilazane, polycarbosilane,etc. is an efficient process for preparation of ceramics. According to Bill et al. (1999) Thein-situ crystallisation of these materials permits the preparation of nanocrystalline materialsby a completely powder-free process. The structure and composition of the grain boundariesthat originate from these crystallisation processes strongly depend on the molecular structure


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of the initially used precursors. Such processes are particularly interesting in the preparationof fibres.

1.2.2 The powder-making processes

To obtain nanocomposites it is necessary to prepare nanopowders so as to increase thereactivity of the different phases during the sintering process. This is due to a high sur-face/volume ratio (generally diameter < 100 nm). To obtain such powders, which haveto present ideal characteristics (nanometric in size, spherical in shape, monodisperse, lowagglomerated, high or controlled purity), it is necessary to develop new processes. Thesenew methods must be innovative and need to be adapted to industrial production.

Such recently developed processes deal with reactions such as

Solid phase: Evaporationcondensation, laser ablation, ball milling (or mechanicalalloying), self-propagating high-temperature synthesis (SHS);

Liquid phase: Solgel, spray-drying of solutions, aerosol pyrolysis; Gaseous phase: Chemical vapour condensation at low pressure, plasma and laser

synthesis.

Some nanosized powders are now commercially available in the USA and in Japan:

Nanophase Technologies Corp. has developed a gas condensation process to producedifferent oxide powders (aluminium, titanium, zirconium oxides).

Nanodyne is focused on the production of cobalttungsten carbide, a ceramicmetalnanocomposite used to make tools for cutting and wear-resistant devices.

Ultram International produces ceramic and composite powders by a super-high-frequency plasma chemical process.

MarkeTech International proposes nanosized silicon carbide, silicon nitride, siliconcarbonitride, boron carbide and silica powders obtained by a CO2-laser pyrolysis process.

Mitsubishi Mining & Cement Co. Ltd has prepared silicon/carbon/nitrogen nanocom-posite powders from pyrolysis in a furnace of an organosilicon compound.

A second problem which needs to be solved concerns the sintering for the preparation ofcomplex shaped components so as to keep the initial size of the crystals.

Similarly, the identification of the most efficient methods of characterisation, all along thedifferent steps of preparation, is of significant importance.

1.3 Nanostructured Si3N4/SiC materials

Different methods have been developed for the preparation of Si3N4/SiC nanocomposites byhot pressing sintering using:

Micrometric silicon nitride, previously recovered with carbon, so as to produce thecarboreduction of silicon oxide present at the surface of the initial material:

SiO2 + 3C SiC+ 2CO (Ishizaki and Yanai 1995; Watari 1989; Bahloul et al. 1993).


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Micrometric silicon nitride with nanometric Si/C/N or Si/C powders and sintering addi-tives (Al2O3, Y2O3) (Niihara 1990). Nevertheless, the mixture, in a liquid medium, ofsuch different sized powders containing so many different chemical groups on the sur-face, is difficult to obtain. This produces agglomerates and a heterogeneous distributionof the second phase in the matrix.

Nanometric Si/C/N or Si/C powders with sintering additives (Al2O3, Y2O3) (Sasaki1993; Kaiser 1996; Kennedy 1996).

1.4 Analytical methods

Different methods can be used for the characterisation of initial to heat-treated and/orpressurised precursors such as:

Chemical analysis (CA) Electron paramagnetic resonance (EPR) Extended X-ray absorption fine structure (EXAFS) Infrared spectroscopy (IR) Soft X-ray spectroscopy (SXS) Solid state nuclear magnetic resonance (NMR) Specific surface area determination (SA) Transmission electron microscopy and electron diffraction (TEM) X-ray (XD) and neutron (ND) diffraction X-ray photoelectron spectroscopy (XPS).

Such methods are able to provide information of varied nature, depending on the originsof the physical phenomena involved (Table 1.1).

Another important remark concerns the method of preparation of the sample. From theas-formed, heat treated and sintered sample, it is necessary to arrive at a condition adaptedto the method of measurement. Crushing the material and filling, dispersing, spreading itin/on the sample holder are able to influence the results. This is why cross measurements arenecessary.

Table 1.1 Available information depending on the method used

Method Surface Bulk Crystalline Amorphous Local or Average or Sample typeshort order long order

CA PowderEPR PowderEXAFS PowderIR PowderND PowderNMR PowderSXS AnyTEM NanopowderXD PowderXPS AnySA AnyNote Indicates the dominant domain of expertise of the method.


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Some examples shown in this book demonstrate such phenomena:

X-ray, TEM and NMR. Figure 2.3, section on Si/C-based powders in Chapter 4 andFigure 5.1.6: --SiC and amorphous detected in different proportions;

NMR and TEM. Figures 4.65, 4.66 and 5.1.27: Si3N4 detected in the bulk and not intograins;

XPS and NMR. Figures 5.3.10 and 5.1.10: SiO2 detected on the surface and not in thebulk;

where the different phases are not evaluated with the same efficiency.

1.5 Conclusion

Although till now many studies on that subject have been published, results are contradictoryand fragmentary. No interdisciplinary studies, till now, have been done on the silicon nitridesilicon carbide system. A better knowledge of the behaviour and potential of such materialsis needed.

Elaboration of high-temperature ductile nanocomposites is one of the objectives of thiswork. Consequently, systematic studies and analyses of the preparation proceed from stepby step. Chemical bonding, nanostructure, microstructure and monolith material propertiesare determined.

Such an approach allows us to obtain original results for each step of the preparation and toestablish relations between the synthesis of powders, their sintering aptitude, the developmentof the microstructure and the properties obtained. The spiral in Figure 1.1 represents thissymbolically.

References

Bahloul, D., Pereira, M. and Goursat, P. (1993) J. Am. Ceram. Soc., 76(5), 115668.Bill, J., Wakai, F. and Aldinger, F. (1999) Precursor-Derived Ceramics. Wiley-VCH.Ishizaki, K. and Yanai, T. (1995) Silic. Indus., 78, 215.Kaiser, A. (1996) Silic. Indus., 56, 215.Kennedy, T. (1996) Silic. Indus., 910, 201.Niihara, K. (1990) J. Mater. Sci. Lett., 10, 112.Niihara, K. (1991) The Centennial Memorial Issue of the Ceramic Society of Japan, 99(10), 974.Sajgalik, P., Dusza, J., Hofer, F., Warbichler, P., Reece, M., Boden, G. and Kozankova, J. (1996)

J. Mater. Sci. Lett., 15, 72.Sasaki, G. (1993) Mater. Res. Soc. Symp., 287, 335.Watari, K. (1989) Mater. Sci. Eng., A109, 89.


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2 Laser synthesis of nanosized powders

Michel Cauchetier, Emmanuel Musset, Michel Luce,Nathalie Herlin, Xavier Armand and Martine Mayne

2.1 Generalities on the synthesis of nanosized powders bylaser pyrolysis

2.1.1 The laser pyrolysis process

Infrared laser pyrolysis (IRLP) is a highly versatile method for the production of a wide rangeof nanopowders including SiC, SiCN, SiCO, BN and carbon (Haggerty and Cannon 1981;Cauchetier et al. 1988, 1994; Herlin et al. 1996; Boulanger et al. 1995; Ehbrecht et al. 1993;Voicu et al. 1996). A summary of the syntheses performed so far for materials applicationsis given in Table 2.1 and completes the review appearing in Knudsen (1997).

In the IRLP process, the reactants are heated by IR laser radiation and decompose, causingclusters to nucleate and grow rapidly. This process is inherently very clean because homo-geneous nucleation occurs in a well-defined reaction zone without interaction with chamberwalls. The small reaction volume and the ability to maintain steep temperature gradients(106 C/s) allow precise control of the nucleation rate, growth rate and residence times. Thephysical and chemical properties of the particles can be controlled by changing the molecularprecursors and the synthesis parameters (laser power, pressure, etc.). The resulting powdersare very fine, spherical, extremely pure, more or less agglomerated and nearly monodispersedin size. The mean particle size can be adjusted from 10 to 100 nm. All these character-istics (i.e. nanometric size (

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Table 2.1 CO2 laser synthesized nanometric powders

Powders Chemical systems References

Si SiH4 Haggerty and Cannon 1981; Cauchetieret al. 1987

SiC SiH4 + C2H4 Sumaya et al. 1985; Cauchetier et al.1987; Frster et al. 1991

SiH4 + C2H2 Cauchetier et al. 1987; Curcio et al.1989; Fantoni et al. 1990; Tougneet al. 1993

SiH4 + CxHy with x = 4 Cauchetier et al. 1988SiH2Cl2 + C2H4 Suzuki et al. 1992(CH3)2Si(C2H5O)2 vapour Li et al. 1994a

Si3N4 SiH4 + NH3 Haggerty and Cannon 1981; Kizakiet al. 1985; Symons and Danforth1987; Buerki et al. 1990

SiH2Cl2 + NH3 Bauer et al. 1989SiC+ Si3N4 orSi/C/N

[(CH3)3Si]2NH vapour Rice 1986; Li et al. 1994bSiH4 + CH3NH2 + NH3 Cauchetier et al. 1989; Cauchetier et al.

1991SiH4 + (CH3)2NH+ NH3 Alexandrescu et al. 1991; Borsella et al.

1992SiH4 + C2H4 + NH3 Suzuki et al. 1993(CH3SiHNH)x with x = 3 or 4

(aerosol)Gonsalves et al. 1992

[(CH3)3Si]2NH (aerosol) Cauchetier et al. 1994; Herlin et al.1994; Musset et al. 1997

SiO2 Si(C2H5O)4 (aerosol) Luce et al. 1994

SiO2 + SiC+ Cor Si/C/O

(RO)4xSi(R)x with0 = x = 3,R = C2H5O,R = CH3and [(CH3)3Si]2O (aerosol)

Armand et al. 1995; Martinengo et al.1996; Herlin et al. 1996; Fusil et al.1997

BN BCl3 + NH3 Luce et al. 1993; Baraton et al. 1994;Willaime et al. 1995

B4C BCl3 + CH4 + H2 Knudsen 1987aBCl3 + C2H2 + H2 Luce et al. 1993

TiB2 TiCl4 + B2H6 Casey and Haggerty 1987aTiCl4 + BCl3 + H2 Knudsen 1987b

ZrB2 Zr(BH4)4 Cauchetier et al. 1987; Rice andWoodin 1988

TiO2 Ti[OCH(CH3)2]4 Casey and Haggerty 1987b; Rice 1987;Curcio et al. 1991

TiC TiCl4 + C2H4 Alexandrescu et al. 1997Al2O3 Al(CH3)3 + N2O+ C2H4 Borsella et al. 1993WC WF6 + C2H2 + H2 + SF6 Bourgeois et al. 1995

electronic grade remains too expensive (Cauchetier et al. 1994) for use in large quantities(>1 ton/year) even if energy and laser equipment costs are low compared to the raw materialcosts. In the case of Si3N4 laser synthesis, a study of the economics shows that a price forSiH4 less than one-fourth of its current price is needed to compete with a classic method ofsynthesis (Schoenung 1991). Alternative methods have been proposed with the substitution

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of chlorosilanes (Bauer et al. 1989; Suzuki et al. 1992) or organosilicon compounds (Rice1986; Gonsalves et al. 1992; Cauchetier et al. 1994) for silane. Technical improvementsin two different fields have enhanced the versatility of the method. First, new introductionsystems for the silicon (Si) precursors have been tested. Liquid organosilicon compoundsare introduced into the laser beam in the aerosol form, fine liquid droplets being obtainedby ultrasonic nebulisation (Gonsalves et al. 1992; Cauchetier et al. 1994). Second, the useof high-power tunable continuous wave (c.w.) CO2 lasers, now available, has increased thearea of applications of the process (Luce et al. 1994).

In this chapter, the synthesis of Si-based powders by two methods, using gaseous or liquidprecursors, will be presented:

SiH4 as gaseous precursor for the synthesis of Si, SiC, Si3N4 and SiCSi3N4 mixturepowders and also methylsilane, SiH3CH3, for the synthesis of SiC powder,

hexamethyldisilazane [(CH3)3Si]2NH or HMDS as liquid precursor for the synthesisof Si/C/N (silicon carbonitride) powders; mixture of organosilicon precursors for thesynthesis of Si/C/N powders containing the elements of sintering additives.

2.2 Synthesis of Si, SiC, Si3N4 and SiC Si3N4 mixtureswith SiH4 and SiH3CH3 as silicon precursors

2.2.1 Experimental setup

The experimental device is presented in Figure 2.1 (Cauchetier et al. 1991). The unfo-cused Gaussian beam (diameter = 12 mm) of a high-power c.w. CO2 laser (CILAS CI1000,Marcoussis, France) enters the reaction cell through a KCl window and crosses the path ofthe gaseous flow of the reactants injected through an inlet capillary (inner diameter = 2 mm).

Pumping IR spectrometer powder collector

CO2 laserbeam

(600 W)

SiH4+ CH2NH2(+C2H2+ NH3)Ar

Figure 2.1 Schematic laser irradiation cell for gaseous precursors. (Reprinted from Cauchetier,M., Croix, O., Luce, M., Baraton, M. I., Merle, T. and Quintard, P., (1991) Jour-nal of the European Ceramic Society, Nanometric Si/C/N Composite Powder: LaserSynthesis and IR Characterisation, 8, 215, with permission from Elsevier Science.)


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A resonance effect between the emission line of the laser at 10.6 m and an infrared absorptionband of one of the reactant gases (silane, SiH4 or methylsilane, CH3SiH3) causes the reactionto occur. This laser-driven reaction leads to high temperatures (up to 1800C) and a brightflame. An argon flow prevents powder deposition on the windows and guides the productsinto the collection chamber. The powders are collected and then stored in a glove-box underargon atmosphere to avoid contamination by air or water vapour.

2.2.2 Synthesis results

Table 2.2 summarises the synthesis conditions of the different powder samples used in dif-ferent parts of this book (Si, SiC, Si3N4 and mixtures of SiC/Si3N4) and also presents firstcharacterisation results. For all the experiments reported here, the cell pressure was keptconstant at 1 atmosphere (105 Pa). In Table 2.2, most of the results concern SiC and Si/C/Nsynthesis, but some results shows the synthesis of Si from pure SiH4 (Si50), Si3N4 from a mix-ture of SiH4 and NH3 (SiN7). Structural investigations on both these powders are presentedin Chapter 5. SiC powders have been obtained from methylsilane (SiC151, SiC152) or frommixtures of SiH4 and C2H2. In the run, SiC171 diborane was added in the silane/acetylenemixture, as a boron source used for sintering aid in densification tests (Croix et al. 1991), andafter pressureless sintering at 2000C, densities up to 0.95 were achieved. The last samplespresented in Table 2.2 are Si/C/N samples, obtained from a mixture of SiH4,CH3NH2, and,in some cases, NH3 has been added to the reactive mixture. In two runs (Si50, SiC163), aninert gas (Ar, He) has been added to the reactive gases in order to increase the linear velocity.After comparing the chemical composition of the gas phase with the chemical compositionof the powder, the C/Si and C/N ratios were obtained (see Table 2.2). In some cases, theflame temperature was measured using an optical pyrometer. Specific surface area (S) wasdetermined from BrunauerEmmetTeller (BET) measurements and the equivalent diame-ter (D) was calculated from these measurements. The time of residence in the laser beamis calculated from the gas flow rates (2 mm nozzle diameter, 12 mm laser beam diameter)assuming room temperature. From the results presented in Table 2.2, some information canbe obtained on the effect of experimental parameters.

Laser power

Increasing the laser power, and keeping all the other parameters constant, leads to anincrease in temperature, which is expected to have an influence on the powders produced.In experiments SiC151 and SiC152, the laser power has been decreased by a factor of three(600/220 W) and the measured temperature decreases from 1530C to 1150C. The samecrystalline structure is identified by X-ray diffraction (XRD). Therefore, in this temperaturerange, the flame temperature does not influence the structure of the powder. In contrast, BETmeasurements indicate a decrease in the powder size as the temperature decreases. Thus, thetemperature has a direct influence on the kinetics of the powder growth.

Flow rates

In runs SiC163177, the flow rates have been changed keeping the laser power constant.Increasing the flow rates leads to a decrease in the residence time as shown in Table 2.2.Figure 2.2 presents the variation in the particle size of SiC powders with the residence time forsamples SiC163177. Figure 2.2 also presents results (broken line) obtained five years before


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Table 2.2 Gas phase synthesis conditions and characterisation results of the different samples

Samples Flow rates (cm3/min) Atomic ratio(gas phase)

Laserpower (W)

Temperature(C)

Residencetime (ms)

BET results XRD

SiH4 C2H2 He, Ar CH3NH2 NH3 C/N C/Si S (m2/g) D (nm)

Si50 120 0 480 (Ar) 0 0 600 930 4.3 29 89 SiSiC151 CH3SiH3 : 600 1530 13.0 34 55 -SiC

200SiC152 CH3SiH3 : 220 1150 13.0 52 36 -SiC

200SiC163 540 300 1000 (He) 0 0 1.1 600 1.2 117 16SiC171a 1080 600 0 0 0 1.1 600 1.3 135 14SiC173 200 110 0 0 0 1.1 600 7.3 61 31SiC174 350 192 0 0 0 1.1 600 4.2 74 25SiC177 120 66 0 0 0 1.1 600 12.1 36 52SiC212 596 306 0 0 0 1.0 640 2.5 70 27SiN7 120 0 0 0 480 600 1250 4.3 71 26SiCN12 340 0 0 200 0 600 1550 4.8 47 40 , -Si3N4,

-SiCSiCN27 280 0 0 70 250 600 1615 4.3 26 72 , -Si3N4SiCN29 360 0 0 90 320 0.22 0.25 400 2.9 40 49SiCN35 350 0 0 200 0 1 0.57 600 4.1 48 39

Notea 40 cm3/min of diborane B2H6 added as boron source which acts as sintering aid for the densification of SiC.


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50

25

05 10

Residence time (ms)

SiC163

SiC174

SiC173

SiC177

BE

T d

iam

eter

(nm

)

15

Figure 2.2 SiC particle size v. the residence time in the laser beam (laser power, 600 W; cellpressure, 1 atm). Similar results have already been observed previously ( ) witha laser power of 640 W. (Reprinted from Tougne et al. (1993) Diamond and RelatedMaterials, Evolution of the Structure of Ultrafine SiC-Laser-formed Powders withSynthesis Conditions, 2, 486490, with permission from Elsevier Science.)

the present results, with a higher laser power (640 W) (Tougne et al. 1993). It must be notedthat the value obtained for the SiC212 (640 W) experiment (not plotted) is in good agreementwith these early results. Figure 2.2 clearly shows a linear decrease in the size of the particlesas the residence time decreases. Figure 2.3 presents the XRD diagrams of the SiC163177samples. The lines become sharper and the full width at half maximum (FWHM) decreasesas residence time increases, which indicates an increasing size of crystallites, parallel to theincreasing particle size (BET measurements).

TEM observations also confirm the BET and XRD results. Sample SiC163 (Figure 2.4(a))presents fine spherical particles with a mean diameter of 20 nm (BET measurement: 16 nm)and with a unimodal distribution. They are not fully crystallised and appear as fine crystallitesin an amorphous matrix. It clearly shows that the size of the particles (BET) and the size ofthe crystallites (XRD) must be distinguished. Sample SiC177 (Figure 2.4(b)) also presentsfine spherical particles with an increased mean diameter of 60 nm (BET measurement:52 nm). They are well crystallised with some black streaks. In correlation with the chemicalcomposition C/Si = 1.1, some fine filaments of turbostratic carbon (C) can be noticed.

In conclusion, it must be noted that both laser power and flow rates have a noticeable effecton the size of the particles. A good control of these experimental parameters allows a goodcontrol of the size and the structure of powders. Some other points which are relevant formaterials applications must be noted:

The amount of synthesised powders is significant. For example, in the run SiC171, aproduction rate of about 100 g/h indicates that the process can be easily scalable. The


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20 40

SiC163

SiC174

SiC173

SiC1773 C

3 C

3 C

3 C

3 C

602 (CuK)

Inte

nsity

(a.

u.)

Pol

ytyp

es

80

Figure 2.3 XRD patterns of the as-synthesised SiC powders (arbitrary units (a.u.)). (Reprintedfrom Tougne et al. (1993) Diamond and Related Materials, Evolution of the Structureof Ultrafine SiC-Laser-formed Powders with Synthesis Conditions, 2, 486490, withpermission from Elsevier Science.)

duration of the experiment can be 3 h or more in order to collect powder quantities inthe 150200 g range for the elaboration of materials.

Also, a very good reproducibility of the experiments can be achieved; for example,Table 2.2 shows BET results for the runs SiCN12 and SiCN35 undertaken after a delayof 4 years.

Another interesting point is the good yield of the reaction. For example, in the SiC212run, and supposing 100% conversion of SiH4 and C2H2 to the solid phase, the maximumcalculated production rate is 64 g of SiC per hour and the experimental result is 59 g/h this reaction yield (almost 100%) is very good.

2.2.3 Chemical analyses of the as-formed powders

The chemical composition in stoichiometric compounds (Si3N4, SiC, SiO2, Si, C) has beencalculated from elemental analysis results, with the classical assumption that all O is in theform of SiO2, then all N forms Si3N4 and the remaining Si is in the form of SiC. Free C(or Si) is the difference between total C (or Si) and C (or Si) bonded to Si (or C) in SiC.Such a calculation is useful to evaluate the composition of future sintered materials, or tocompare several powders but it is obvious that it is only an approximation, especially whenthe powders are not fully crystallised.


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100 nm

100 nm

(a)

(b)

Figure 2.4 TEMs of two as-formed powders: (a) SiC163; (b) SiC177.

SiC powders

In order to avoid the presence of free Si in the powders which makes the powder very sensitiveto oxygen contamination, most of the SiC samples were synthesised with a gaseous mixturecontaining excess C (C/Si = 1.1 in the gaseous mixture). In Table 2.3, two examples show thatthis ratio is also found in the powders produced. In agreement with the chemical composition,XRD patterns (similar to Figure 2.3) correspond to the -SiC phase. The O and free Ccontents are low and correspond to the values usually encountered for commercial powders.Table 2.3 also presents the chemical composition calculated in stoichiometric compoundswhich indicates, in good agreement with the XRD pattern, that the powders are mostlycomposed of SiC.

Si/C/N powders

For the synthesis of Si/C/N composite powders, two gaseous mixtures were investigated:a binary mixture SiH4+CH3NH2 (SiCN35) and a ternary mixture: SiH4+CH3NH2+NH3


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Table 2.3 Chemical characterisations of gas phase synthesised powders

Powder Samples

Chemical analysis (wt%) Atomic ratio Chemical composition (wt%)

Si C N O C/Si C/N Si3N4 SiO2 SiC C Si

SiC173 68.3 30.2 0 0.9 1.03 0 1.5 96 2.0 0SiC212 68.3 30.3 0 1.4 1.03 0 2.7 95.7 1.6 0SiCN29 58.3 6.5 34.6 0.6 0.27 0.22 86.6 1.1 8.3 4.0 0SiCN35 66.9 13.6 18.2 2.3 0.47 0.87 45.1 4.3 45.1 0 5.5

Si3N4 Si3N4

SiCSi

SiCN35

SiCN29

10 20

Inte

nsity

(a.

u.)

40 50

Angle (2 degrees)

60 70 80 9030

Figure 2.5 XRD patterns of as-formed SiCN nanopowders used in material elaboration.(Reprinted from Mayne 1997.)

(SiCN29) (Table 2.2). In the case of SiCN29, the presence of NH3 is known to make thenitriding introduce an increase in the nitrogen (N) content. For SiCN35 a similar content inC and N is expected in the resulting powder due to the C/N initial atomic ratio of 1 in theformula of monomethylamine. Tables 2.2 and 2.3 show a very good correlation between C/Siand C/N ratios in the gas phase and in the powder for SiCN29, but the correlation is not sogood for SiCN35 probably due to evolution of carbonaceous gas of C in the gas phase duringthe synthesis.

Chemical composition calculations show that the powder SiCN35, containing free Si isthe most sensitive to oxygen contamination. XRD patterns (Mayne 1997) are reported inFigure 2.5 and confirm the calculated compositions. The SiCN29 sample contains mainly- and -Si3N4; the quantity of SiC is too small and is not detectable. The SiCN35 samplepresents the characteristic lines of Si and -SiC and a broad peak between 2 = 30 and40, which can be attributed to an amorphous phase of silicon carbonitride.

From the results presented in this section, it appears that the decomposition in stoichiomet-ric compounds is a good representation of the powder and also that the chemical compositionof the powders produced is most often well controlled by the chemical composition of thegaseous reactants. Together with the possibility of controlling the size and the structure of


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the powder, it makes this synthesis technique attractive. But, as explained before, in orderto increase the safety and decrease the cost, it has been applied to the liquid precursorscontaining Si, C and N elements.

2.3 Synthesis of silicon carbonitride (Si/C/N) powders usinghexamethyldisilazane as Si precursor

2.3.1 Experimental

Choice of the precursor

As noted before, investigations based on costs of various silicon precursors for the deliveryof large quantities have shown that silane remains too expensive for use in ceramic powdersynthesis (Cauchetier et al. 1994). In our case [(CH3)3Si]2NH or HMDS with a Si cost verylow compared to the Si cost of SiH4 (43 US$/kgSi in HMDS and 343 US$/kgSi in SiH4) waschosen as a model of the liquid precursor in order to demonstrate the possibility of using liquidwastes of the organosilicon compound chemistry for the synthesis of nanometric Si-basedpowders. Indeed, HMDS has a strong IR absorption band at 10.6 m due to the Si N bond,which was, therefore selected. It is a common liquid Si compound (boiling point of 125C,density of 0.76, and a viscosity of 0.69 cP at 25C) largely used for the silylation of a widerange of functional groups and it was the precursor used by Niihara and co-workers in thesynthesis of Si/C/N powders leading to non-oxide superplastic ceramics (Wakai et al. 1990).The first experiments using HMDS in the vapour phase as Si precursor for Si/C/N powdersynthesis by CO2-laser pyrolysis were performed by Rice (1986). A conversion efficiency(liquid powder) of about 60% has been obtained but with low flow rates of HMDS (7 g/h)and with a low-powered CO2 laser (135 W). Higher production rates have been attained(80120 g/h) when the laser power has been increased to the 5001000 W range (Li et al.1994b).

The experimental device

The apparatus is shown in Figure 2.6. Liquid HMDS is placed in a special glass jar containinga piezoelectric transducer (Pyrosol 7901 type from RBI, Meylan, France, apparatus developedin collaboration with the CEA (Spitz and Vigui 1970)). The focusing of the intense beamof ultrasonic energy delivered by the transducer near the surface of the liquid yields uniformdroplets whose diameter d is given by

d = 3

4f 2

where and are the surface tension and the density of the liquid, respectively, and f thefrequency of the transducer (here 850 kHz). In order to increase the mass of the displacedliquid, the glass jar is heated near 90100C with a heating ribbon. The aerosol dropletsand vapour are injected into an irradiation cell very similar to those presented in Figure 2.1.The reaction cell is maintained at a regulated pressure of 105 Pa using a flow of argon or anargonammonia mixture through a glass inlet tubing (inner diameter of 13 mm). The laserbeam diameter is increased from 12 to 24 mm with a beam expander in order to cover thereactant flow entirely. The initial laser power is in the range 480520 W.


chap02 2002/9/6 page 16 #11

Powder collector

Heating wires

Vacuum pump or IR spectrometer

KCl window

Argon

Chimney gas

Gas for the propulsion of the aerosol

PrecursorRF power supply

PYROSOL system

Piezoelectric ceramic

Argon

Laser CO2

Filter

Figure 2.6 Schematic of the aerosol generator with the irradiation cell and the powder collector.

2.3.2 Synthesis results

Table 2.4 compares the influence of two parameters, the nature of the carrier gas (argon andnitrogen) and heating of the liquid precursor (near 100C), on the synthesis results. It can beseen immediately that the yield of the reaction (expressed as the ratio between the mass ofpowders produced and the mass of displaced liquid), does not appear to be as good as in the gasphase. Both the amount of displaced liquid and production rate increase when the precursoris heated. The effect of the heating is to combine the enrichment of the gaseous phase inHMDS vapour and decrease the viscosity of the liquid. In order to work with significantamounts of powders, most of the experiments were carried out with liquid heating. The yieldseems independent of the heating of the precursor, but is strongly dependent on the nature ofthe gas. It is 4649% for Ar carrier gas flow and 3839% for N2 carrier gas flow.


chap02 2002/9/6 page 17 #12

2.3.3 Effect of experimental parameters

Effect of carrier gas (Ar or N2)

Table 2.4 shows that the chemical composition of the powder does not depend on the natureof the carrier gas flow (Ar or N2) and the C/N atomic ratio remains stable in the 2.52.9 range.The activity of atomic N (N2 2N) is low at the measured flame temperature: 126020C.Table 2.4 also shows that all powders are sensitive to pollution by oxygen, but the powdersobtained from the heated precursor seem the most sensitive to pollution. It must be notedthat the chemical analysis presented in this section is not complete due to the presence ofhydrogen in the powder and due to the uncertainty of the measurements.

Effect of chemical evolution of the reactive mixture

Table 2.5 presents the synthesis parameters of all the samples studied in the different partsof this book. Only samples HMDS4045 where the effect of chemical composition has been

Table 2.4 Effect of the carrier gas nature and precursor heating

Run Carrier gas flowrate (cm3/min)

Displacedliquid(cm3/h)

Powderproductionrate (g/h)

Yield(wt%)

SBET(m2/g)

Chemical analysis(wt%)

C/N(Atomicratio)Ar N2 Si C N O

HMDS28a 2430 108 38 46 88 49.8 30.3 12.3 5.6 2.9HMDS36b 2500 164 61 49 82 45.8 28.0 11.7 9.7 2.8HMDS38b 2500 166 48 38 125 47.2 28.4 12.8 11.6 2.6HMDS39a 2500 95 28 39 79 49.9 29.7 14.1 5.2 2.5

Notesa Liquid precursor at room temperature.b Liquid precursor heated near 100C.

Table 2.5 Synthesis conditions and chemical analysis of HMDS samples (chemical analyses were alwaysperformed after several weeks of exposure to air)

Run Carrier gas flowrate (cm3/mn)

Displacedliquid(cm3/h)


Yield(wt%)

SBET(m2/g)

Chemical analysis(wt%)

C/N(Atomicratio)

Ar N2 Total Si C N O

HMDS40 1935 205 2140 182 59 43 110 46.7 25.3 21.8 6.1 1.35HMDS41 1770 410 2180 160 61 50 93 46.1 20.5 25.2 6.7 0.95HMDS42 1570 600 2170 175 75 56 98 44.4 18.5 29.7 7.7 0.73HMDS43 1370 800 2170 179 74 53 95 48.3 15.5 26.8 8.0 0.67HMDS44 1150 1040 2190 210 81 51 95 47.1 13.1 26.7 10.9 0.57HMDS45 830 1346 2180 167 66 52 106 47.6 8.3 28.9 13.4 0.34HMDS35 1600 1000 2600 87 21 32 144 47.3 6.6 33.4 9.8 0.23HMDS66 1920 1920 114 32 36 66 51.2 31.9 15.1 1.8HMDS67a 1440 1920 115 31 35 66 52.0 30.8 14.8 2.4HMDS73a 862 207 1920 56 14 32 115 51.9 12.6 25.3 10.1

Notea H2 in the carrier gas.


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0.5

35 45

44

4342

41

40

22

HMDS28

1 2 3 4C/N atomic ratio (gas + aerosol)

5 6

1.0

C/N

ato

mic

rat

io (

pow

der)

1.5

2.0

2.5

3.0

Figure 2.7 C/N atomic ratio in as-formed HMDS powders v. C/N atomic ratio in the precursors(liquid + ammonia). (Reprinted from Musset et al. 1997 with the permission of theBelgian Ceramic Society, Mons, Belgium.)

closely studied will be commented on in this section; the main part of the results appear in thePhD thesis of Emmanuel Musset (Musset 1995). For this set of six experiments, the liquidtemperature is about 100C, the total flow rate of the carrier gas mixture remains constantand the corresponding residence time of the precursor in the laser beam is about 0.1 ms.

Table 2.5 shows the influence of the addition of ammonia in the argon used as carriergas, namely, on the chemical composition of the resulting powders (samples HMDS4045).Figure 2.7 shows a good correlation between the C/N ratio in the as-formed powder and in thereactive mixture; the C/N atomic ratio decreases from 1.35 to 0.36 when the ammonia volumeratio increases from 10% to 60%. This correlation confirms the possibility of controlling thechemical composition of the powders. Table 2.5 shows that the contamination by oxygenbecomes more and more important when the N content in the powder increases, because NHbonds are easily hydrolysed.

2.3.4 First characterisations

Specific surface area

The specific surface area is always high, about 100 m2/g and taking into account the meanvalue of the density i.e. 2.00 (pycnometry measurement), the corresponding mean value forthe diameter is 30 nm, in agreement with TEM observations.


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NH

SiNSi

SiH

SiC

SiN

KBr

NHOH

CH (CH3) SiCH3

HMDS45

HMDS44

HMDS43

HMDS42

Abs

orba

nce

(a.u

.)

HMDS41

HMDS40

3500 3000 2500 2000 1500

Wavenumber (cm1)

1000 500

Figure 2.8 IR spectra of the as-formed HMDS powders. (Reprinted from Musset 1995.)

Infrared spectroscopy

IR spectra of the samples HMDS4045 are shown in Figure 2.8. They present a broad absorp-tion band between 750 and 1250 cm1 with a maximum at 920940 cm1 due to a Si/C/Namorphous phase, corresponding to a flat XRD diagram. A weak band at 550 cm1 is due tothe Si N bond. Other bands are related to hydrogenous species: N H at 1180 cm1, Si Hat 2050 cm1 and Si CH3 at 1250 cm1, indicating an incomplete pyrolysis of the precursoror absorption of volatile species obtained from the pyrolysis. The presence of hydrogen is ingood agreement with the chemical analysis, which is often less than 100% (Table 2.4).

The results presented in this section show that it is possible to produce Si/C/N nanopowderswith controlled chemical composition from a liquid precursor. The powders are obtained insignificant quantities for materials applications. The main differences compared to powdersobtained from the gas phase are the amorphous character and high hydrogen content of thepowders obtained from the liquid phase.

2.4 Silicon carbonitride (Si/C/N) powders containing in situthe elements of the sintering aids (Al, Y): pre-mixedpowders

2.4.1 Introduction: elaboration of SiCSi3N4 nanocomposite materials

Materials in the SiC/Si3N4 system are elaborated by heat treatment of powders in the presenceof liquid-forming sintering aids such as Y2O3, MgO, Al2O3, etc. singly or in combination.Such processes need a previous mixing step of all powders having different natures andsometimes different sizes. Recently, some studies report the incorporation of metal elements


chap02 2002/9/6 page 20 #15

(Al, Y, etc.), necessary for the sintering step, into silicon nitride or carbonitride systems, mostof them being focused on chemical processing routes using organometallic precursors (Soraruet al. 1992; Iwamoto et al. 1998; Koyama et al. 1998). A few other publications concernphysico-chemical processes: laser irradiation of gaseous mixtures (SiH4,NH3, (CH3)3Al)allows the formation of nanosized powders with Al in a Si3N4 matrix (Borsella et al. 1993b).Aerosol of Si/C/N particles with yttrium and aluminium compounds were prepared under dualirradiation of CO2 and excimer laser beams of a SiH4,C2H4 and NH3 mixture, aluminiumbeing incorporated in the form of alumina powder in the reactant flow and yttrium being intro-duced as a low volatile organometallic compound in the reaction zone (Yamada et al. 1997).

2.4.2 Synthesis of pre-mixed powders

In order to prepare Si/C/N powders containing uniformly dispersed sintering aids inves-tigations were undertaken to set up a laser spray synthesis process of Si/C/N/O/Al(Y)nanopowders using solutions of a liquid organosilicon precursor HMDS with solidmetallic alcoxides, namely aluminium and yttrium isopropoxides [(CH3)2CHO]3Al and[(CH3)2CHO]3Y.

Table 2.6 summarises the synthesis conditions of 14 runs which can be divided into3 groups:

Group I includes experiments carried out with solutions having different Al (+Y)/Siatomic ratio.

Group II includes experiments carried out with solutions having the same Al (+Y)/Siatomic ratio. All runs are realised with a binary mixture of carrier gases (Ar + NH3).

The third is composed of experiments made with the same solutions as in group II, butthe carrier gas is composed of Ar+NH3+H2 in order to see if the chemical compositionof the powders is modified by the presence of H2.

Some changes occur compared to the experiments described above in Section 2.3. Theinitial laser power is in the range 300330 W, but, here, with an unfocused laser beam(diameter = 13 mm) which covers the aerosol flow sufficiently. For comparison, anexperiment with pure HMDS carried out under similar conditions is also reported (HMDS66).

2.4.3 Effect of experimental parameters

Effect of the solutions of precursors

In Table 2.6, most of the experiments were carried out with metallic isopropoxides in HMDSbut, in two cases (HSAl04 and HSAlY05), isopropanol was added to the reactive solution inorder to increase the solubility of the alcoxides and to decrease the viscosity of the solutions.In these experiments, more liquid is displaced than in other experiments but the productionrate remains at the same level (30 g/h). Also, the amount of oxygen incorporated in theas-formed powders is very high (>10%). In conclusion, isopropanol does not improve theproduction or quality of the powders and this solution was not developed.

Precursor solutions are characterised by their atomic ratios: [Al(+Y)]/Si. For the firstexperiments, the Al/Si ratio in the precursor solution was adjusted to 0.069 (e.g. HSAl03);this value corresponds to 6 wt% Al2O3 content in a final SiC/Si3N4 material according to thedifferent oxide contents reported in the literature from about 10 (Kennedy et al. 1996) to 1 wt%


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Table 2.6 Synthesis conditions and results of pre-mixed Si/C/N powders

Run Carrier gas flow rates(cm3/min)

Displacedliquid (g/h)


SBET(m2/g)

Chemical analysis (wt%) Atomic ratio

Ar NH3 H2 Si C N O Al Y Al + Y/Si Al/YHSAl03 1900 0 0 87 34 66 46.5 33.9 12.9 5.7 2.5 0HSAl04 1900 0 0 119 32 81 43.5 33.7 10.7 13.2 2.4 0HSAl07 1900 0 0 88 24 110 47.1 29.1 12.8 8.6 2.5 0 0.05HSAl08 0 0 1900 49 11 280 48.7 26.6 14.1 8.9 1.6 0HSAlY05 1900 0 0 115 30 79 43.7 33.4 11.2 11.5 1.7 1.4HSAl09 1140 760 0 71 23 112 49.5 6.8 36.0 6.1 1.5 0 0.03HSAl12 1520 380 0 87 28 96 47.2 14.8 29.0 7.5 1.5 0 0.03HSAl13 1710 190 0 87 30 83 49.8 23.0 19.2 5.9 2.0 0 0.04HSAl11 475 570 855 64 27 134 49.8 5.8 37.2 5.7 1.5 0HSAl14 855 190 855 74 26 89 48.8 17.3 24.5 7.7 1.6 0HSAl15 665 380 855 75 22 99 50.9 11.7 26.8 8.8 1.8 0HSAlY16 1140 760 0 81 26 100 53.0 7.5 30.3 8.2 0.4 0.4 0.01 3.0HSAlY17 1520 380 0 89 25 91 50.7 16.9 25.4 5.4 0.9 0.7 0.02 4.2HSAlY18 1710 190 0 79 28 77 49.3 21.9 22.4 5.1 0.5 0.7 0.01 2.5


chap02 2002/9/6 page 22 #17

alumina and yttria (Burger et al. 1997). The first result is that it is possible to incorporate Aland Y additives in the powders, but Table 2.6 clearly shows that the ratio Al/Si, in this rangeof experimental conditions, does not reach the 0.069 ratio. This is due to the preferentialpulverisation of the most volatile compound of the mixture, that is, HMDS.

So, in the following experiments, the initial Al/Si and (Al+Y)/Si ratios were decreased to0.04 (HSAl09, 12, 13) and 0.02 respectively (HSAlY16, 17, 18). The similarity of the Al/Siatomic ratio in the starting solution (0.04) and in the as-formed HSAl powders (0.030.04)suggests that the nebulisation of the mixture of precursors is homogeneous. In the case ofHSAlY powders, the Al/Y atomic ratios are different in as-formed powders (2.54.2) andin the starting solution (1.1), though (Al + Y)/Si atomic ratios remain similar (0.010.02)to the starting solution. This is due to the incomplete dissolution of yttrium isopropoxide inHMDS and probably to an inhomogeneous nebulisation of the mixture of precursors.

Effect of carrier gas

The nature and composition of the carrier gas mixture play a significant role in the chemicalcomposition of powders. Whatever the powder (HSAl or HSAlY), an increase of ammoniarelative flow rate induces a decrease in C content and C/N atomic ratio as obtained for powderswithout sintering aids (Musset 1995).

Table 2.6 also shows that when hydrogen is added partially in the place of argon inargonammonia mixtures, the C content decreases. This seems to indicate that more volatilehydrocarbon species are formed during the synthesis. For example, the C content varies from14.8 to 11.7 wt% in a 80% argon20% ammonia mixture (HSAl12 and 15 samples) and from23.0 to 17.3 wt% in a 90% argon10% ammonia mixture (HSAl13 and 14 samples). It mustbe noticed that a large quantity of free C (1014%) remains in powders when the ammoniarelative flow rate increases from 0 to 20.

2.4.4 First characterisations

Specific surface area

Powder production rates in the 2234 g/h range and specific surface area values in66134 m2/g range are very similar to the values obtained for powders without metallicadditives.

Infrared spectroscopy and X-ray diffraction

As for powders without metallic additives, XRD patterns show a large background withoutany distinct peak indicating an amorphous structure. This is confirmed by infrared spectrashowing a large band around 900 cm1. This one is essentially composed of Si C bondswhen powders are synthesised with argon-rich carrier gas while Si N bonds in an amorphousenvironment (large band at 400 cm1) appear with ammonia-rich carrier gas. Moreover,Si CH3 bonds are present (1250 cm1) in all powders.

2.5 Conclusion

In this chapter, it has been shown that laser pyrolysis is a versatile method well suited for theproduction of significant amounts of Si-based nanopowders. Due to their properties (purity


chap02 2002/9/6 page 23 #18

limited only by the purity of the reactants, low size dispersion, etc.), the nanopowders aregood candidates for applications in ceramics. Nanopowders can be obtained from gaseous orliquid precursors. The powders obtained from the gas phase are crystallised while the powdersobtained from the liquid phase are amorphous. In both cases, the chemical composition of thepowders is controlled by the chemical composition of the reactive mixture. Finally, an interest-ing result is the incorporation of the elements of sintering aids (Al, Y, O) during the synthesis.

References

Alexandrescu, R., Morjan, I., Borsella, E., Botti, S., Fantoni, R., Dikonimos-Makris, T., Giorgi, R. andEnzo, S. (1991) Journal of Materials Research, 6, 2442.

Alexandrescu, R., Borsella, E., Botti, S., Cesile, M. C., Martelli, S., Giorgi, R., Turtu, S. and Zappa, G.(1997) Journal of Materials Science, 32, 5629.

Armand, X., Herlin, N., Martinengo, H., Musset, E. and Cauchetier, M. (1995) Fourth Euro-Ceramics,vol. 1 (Faenza Editrice S.p.A., Faenza), pp. 3744.

Baraton, M. I., Boulanger, L., Cauchetier, M., Lorenzelli, V., Luce, M., Merle, T., Quintard, P. andZhou, Y. H. (1994) Journal of the European Ceramic Society, 13, 371.

Bauer, R. A., Smulders, R., Becht, J. M. B., Van der Put, P. J. and Schoonman, J. (1989) Journal of theAmerican Ceramic Society, 72, 1301.

Borsella, E., Botti, S., Fantoni, R., Alexandrescu, R., Morjan, I., Popescu, C., Dikonimos-Makris, T.,Giorgi, R. and Enzo, S. (1992) Journal of Materials Research, 7, 2257.

Borsella, E., Botti, S., Giorgi, R., Martelli, S., Turtu, S. and Zappa, G. (1993a) Applied Physics Letters,63, 1345.

Borsella, E., Botti, S., Alexandrescu, R., Morjan, I., Dikonimos-Makris, T., Giorgi, R. and Martelli, S.(1993b) Materials Science and Engineering A, 168, 177.

Bourgeois, L., Barbier, G., Vigui, J. C., Herlin, N. and Cauchetier, M. (1995) Fourth Euro-Ceramics,vol. 1 (Faenza Editrice S.p.A., Faenza), pp. 225232.

Buerki, P. R., Troxler, T. and Leutwyler, S. (1990) High Temperature Science, 27, 323.Burger, P., Duclos, R. and Crampon, J. (1997) Materials Science and Engineering, A222, 175.Casey, J. D. and Haggerty, J. S. (1987a) Journal of Materials Science, 22, 737.Casey, J. D. and Haggerty, J. S. (1987b) Journal of Materials Science, 22, 4307.Cauchetier, M., Croix, O., Luce, M., Michon, M., Paris, J. and Tistchenko, S. (1987) Ceramics

International, 13, 13.Cauchetier, M., Croix, O. and Luce, M. (1988) Advanced Ceramic Materials, 3, 548.Cauchetier, M., Croix, O., Robert, C., Lance, M. and Luce, M. (1989) Euro-Ceramics, vol. 1, Elsevier

Applied Science, London, pp. 130134.Cauchetier, M., Croix, O., Luce, M., Baraton, M. I., Merle, T. and Quintard, P. (1991) Journal of the

European Ceramic Society, 8, 215.Cauchetier, M., Croix, O., Herlin, N. and Luce, M. (1994) Journal of the American Ceramic Society,

77, 993.Cauchetier, M., Armand, X., Herlin, N., Mayne, M., Fusil, S. and Lefevre, E. (1999) Journal of

Materials Science, 34, 1.Croix, O., Gounot, M., Bergez, P., Luce, M. and Cauchetier, M. (1991) Ceramics Today Tomorrows

Ceramics, Elsevier Science Publishers, Amsterdam, pp. 14471455.Curcio, F., Ghiglione, G., Musci, M. and Nanetti, C. (1989) Applied Surface Science, 36, 52.Curcio, F., Musci, M., Notaro, M. and Quattroni, G. (1991) Ceramics Today Tomorrows Ceramics,

Elsevier Science Publishers, Amsterdam, pp. 25692578.Fantoni, R., Borsella, E., Piccirillo, S., Ceccato, R. and Enzo, S. (1990) Journal of Materials Research,

5, 143.Frster, J., von Hoesslin, M., Schfer, J. H., Uhlenbussch, J. and Vil, W. (1991) Proceedings of the

10th International Symposium on Plasma Chemistry (FRG) vol. 1 (eds U. Ehlemann, H. G. Lergonand K. Wiesemann), pp. 16.


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Fusil, S., Armand, X., Herlin, N. and Cauchetier, M. (1997) Key Engineering Materials, 132136, 141.Garifo, L. (1995) Laser Focus World, 30.Gonsalves, K. E., Strutt, P. R., Xiao, T. D. and Klemens, P. G. (1992) Journal of Material Science, 27,

3231.Haggerty, J. S. and Cannon, R. W. (1981) Laser-induced Chemical Processing, Plenum Press, New York,

pp. 165241.Herlin, N., Musset, E., Luce, M. and Cauchetier, M. (1994) Journal of the European Ceramic Society,

13, 285.Herlin, N., Armand, X., Musset, E., Martinengo, H., Luce, M. and Cauchetier, M. (1996) Journal of

the European Ceramic Society, 16, 1063.Iwamoto, Y., Kikuta, K. and Hirano, S. (1998) Journal of Materials Research, 13, 353.Kennedy, T., ONeil, J. P., Hampshire, S., Poorteman, M. and Cambier, F. (1996) Silicates Industriels,

910, 201.Kizaki, Y., Kandori, T. and Fujitani, Y. (1985) Japan Journal of Applied Physics, 24, 800.Knudsen, A. K. (1987a) Advances in Ceramics, 21, 237.Knudsen, A. K. (1987b) US Patent 4, 689, 129.Knudsen, A. K. (1997) Carbide, Nitride and Boride Materials: Synthesis and Processing, Chapman &

Hall, London, pp. 343358.Koyama, S., Nakashima, H., Sugahara, Y. and Kuroda, K. (1998) Chemistry Letters, 191.Li, Y., Liang, Y., Zheng, F. and Hu, Z. (1994a) Journal of the American Ceramic Society, 77, 1662.Li, Y., Liang, Y., Zheng, F. and Hu, Z. (1994b) Materials Science and Engineering, A174, L23.Luce, M., Croix, O., Zhou, Y. H., Cauchetier, M., Sapin, M. and Boulanger L. (1993) Euro-ceramics

II, vol. 1, Deutsche Keramische Gesellschaft, Kln, pp. 233238.Luce, M., Herlin, N., Musset, E. and Cauchetier, M. (1994) Nanostructured Materials, 4, 403.Martinengo, H., Musset, E., Herlin, N., Armand, X., Luce, M. and Cauchetier, M. (1996) Silicates

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Ceramics: Getting in the 2000s, Part B, Techna, Faenza, pp. 211218.Musset, Emmanuel (Nov. 1995) PhD Thesis, Paris-Sud Orsay University, 95PA112537;

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7, 1266.Spitz, J. and Vigui, J. C. (1970) French Patent 2, 110, 622.Suzuki, M., Nakata, Y., Okutani, T. and Kato, A. (1992) Journal of Materials Science, 27, 6091.Suzuki, M., Maniette, Y., Nakata, Y. and Okutani, T. (1993) Journal of the American Ceramic Society,

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1356.Symons, W. and Danforth, S. C. (1987) Advances in Ceramics, 21, 249.Tougne, P., Hommel, H., Legrand, A. P., Herlin, N., Luce, M. and Cauchetier, M. (1993) Diamond and

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3 Thermal behaviour of as-formedsilicon-based nanopowders

General introduction

This chapter presents the effects of thermal treatments on Si-based nanopowders obtained asdescribed in the previous chapter from the gas phase or from the liquid phase. It is mainlyfocused on the behaviour of Si/C/N and Si/C/N/Al/Y/O powders (0.2 < C/N(at) < 2.6)because such nanopowders may be used to obtain Si3N4/SiC materials after thermal treatment.

This chapter is divided into two parts. The first part is devoted to thermal treatment in ahigh-temperature graphite furnace in order to know the evolution of the physico-chemicalcharacteristics. The second part reports the study of the thermal behaviour using TGA/MSin order to determine the chemical reactions and mechanisms involved during annealing.


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3.1 Thermal behaviour in ahigh-temperature graphite furnace

Emmanuel Musset, Martine Mayne, Michel Cauchetier,Xavier Armand, Michel Luce and Nathalie Herlin-Boime

3.1.1 Introduction

The aim of the first part of this study is to determine the influence of the annealing treatmentsperformed in a high-temperature graphite furnace both on the weight of samples and ondifferent physico-chemical properties of nanopowders such as the chemical composition,structure, reactivity and size. It is mainly focused on Si/C/N samples obtained from the liquidphase using the aerosol method. The results of annealing treatments and the consequencesof these treatments on the chemical composition, size, organisation, etc. of silicon-basedpowders are presented with the intent of showing the similarities and the differences betweenthe different powders, to correlate these properties to the synthesis conditions, etc., and tobegin to compare the properties of these powders with the necessary properties for ceramicapplications.

3.1.2 Weight evolution of SiC, Si/C/N and Si/C/N/Al/Y/Onanopowders

The as-formed powders are annealed at atmospheric pressure in a high-temperature graphitefurnace under argon or nitrogen (N) atmospheres. The temperature generally varies in therange 10001700C, and the dwell time is 1 or 4 h.

3.1.2.1 Weight evolution under inert atmosphere

Thermal treatments were performed on SiC or Si/C/N nanopowders in an argon atmospherefor temperatures up to 1600C. SiC powders exhibit a weight loss whatever the temperature.The maximum weight loss, observed near 1500C, remains low (usually

p1chap3 2002/9/6 page 27 #3

loss is high (>20%). This also suggests the decomposition of the Si/C/N chemical system,which appears to be more important than in the case of Si/C/N obtained from the gas phase.Detailed mechanisms will be proposed and discussed in the TGMS section.

The decomposition of the Si/C/N chemical system under an inert atmosphere could bepartially related to the degradation of the nitrogen-containing species inducing an evolutionof nitrogen during the thermal treatment. Usually, the elaboration of silicon nitride basedmaterials is achieved by sintering powders under a nitrogen atmosphere in order to limit theweight loss and the degradation of the silicon nitride based system. That is why a nitrogenatmosphere is used for heat treatment in the following sections.

3.1.2.2 Weight evolution under nitrogen atmosphere

The aim of the use of nitrogen is to limit the weight loss by preserving the N content in theas-formed powders in order to further obtain dense silicon nitride based ceramic materials.

Heat treatment was performed at 1000C and then in steps of 100C between 1300Cand 1700C with a heating rate of 10C/min. Most often, the dwell time is 4 h, unlessotherwise indicated in the text. Previously, the as-formed Si/C/N powders (about 2 g) areslightly pressed in a Teflon mould to form pellets (diameter = 20 mm; height 10 mm)which are then placed in graphite crucibles located inside the graphite furnace.

For Si/C/N nanopowders synthesised from the gas phase (SiCN29 and SiCN35), a signifi-cant mass change (+8 wt%) is obtained for the SiCN35 sample between 1200C and 1400Csuggesting the nitridation of the free Si present in the powder. Contrary to the SiCN35 sam-ple, the SiCN29 sample exhibits a relative mass stability up to 1600C (Mayne 1997). Thiscan be related to the high N content and to the absence of Si that cannot enable a nitridationprocess (Mayne 1997; Mayne et al. 1998). X-ray diffraction (XRD) patterns of the annealedpowders show the presence of crystalline phases (- and -Si3N4 for SiCN29 sample and amixture of - and -Si3N4 with -SiC for SiCN35 sample).

In contrast to Si/C/N nanopowders obtained from the gas phase, Si/C/N and Si/C/N/Al/Y/Opowders obtained from the liquid phase (HMDS or mixtures of HMDS with organic com-pounds) exhibit weight losses that are of the same order of magnitude (2030% at 1600C).Therefore, the weight evolution of Si/C/N/Al/Y/O powders (Cauchetier et al. 1999; Mayneet al. 1999), will not be specifically presented in this chapter. Such weight losses suggest adecomposition of the Si/C/N chemical system even under nitrogen atmosphere. In Table 3.1.1,the weight losses and chemical compositions of six HMDS samples are given. It shows thatthe weight loss is dependent on the chemical composition of the as-formed powder. Afterannealing at 1600C, the weight loss is the highest for the samples with the highest C/Nratio. This difference could be related to different chemical mechanisms involved in thedecomposition process and interpretations will be suggested in the TGMS section.

This study on the weight evolution of Si-based samples during heat treatment enables usto show the following:

a strong difference in the behaviour of powders obtained from the gas phase or from theliquid phase using the aerosol method especially under a nitrogen atmosphere. Nanopow-ders obtained from the gas phase are more stable than those obtained from the liquidphase, which suggests that they are more suitable for ceramic conversion than thoseobtained from the liquid phase;

that the thermal behaviour of powders obtained from the liquid changes as a function ofthe chemical composition of the powder, suggesting different mechanisms.


p1chap3 2002/9/6 page 28 #4

Table 3.1.1 Results of chemical analyses for HMDS40 to HMDS45

Heating Mass Si C N Ob N/Si C/Si C/N

treatment (C) loss (%)

HMDS40AFa 46.7 25.0 21.8 6.1 99.6 0.93 1.24 1.331000 5.9 48.4 23.6 20.9 5.0 97.9 0.86 1.14 1.311300 5.7 51.3 23.7 23.7 2.2 100.9 0.92 1.07 1.161400 6.0 51.5 23.7 22.5 1.5 99.2 0.87 1.07 1.231500 19.4 59.6 26.8 12.6 0.7 99.6 0.42 1.05 2.471550 27.0 65.7 28.2 4.7 0.8 99.4 0.14 1.00 7.021600 27.9 68.6 29.1 2.9 0.5 101.1 0.08 1.00 11.15

HMDS41AFa 46.1 20.6 25.2 6.7 98.6 1.09 1.04 0.951000 4.4 50.4 19.3 23.5 4.6 97.8 0.93 0.89 0.951300 5.8 51.4 19.2 27.3 2.4 100.3 1.06 0.87 0.821400 7.4 52.5 18.7 27.0 1.7 99.9 1.03 0.83 0.811500 17.2 59.2 21.0 19.6 1.1 100.9 0.66 0.83 1.251550 28.6 61.6 23.1 11.8 1.1 97.6 0.38 0.87 2.281600 28.2 67.4 25.2 8.5 1.4 102.4 0.25 0.87 3.50

HMDS42AFa 44.4 18.7 29.7 7.7 100.5 1.34 0.98 0.731000 4.0 48.9 17.9 27.1 4.4 98.3 1.11 0.85 0.771300 5.9 51.6 18.3 29.5 1.5 100.9 1.14 0.83 0.721400 8.0 52.6 18.0 29.8 1.6 102.0 1.13 0.80 0.701500 14.6 55.8 18.4 22.1 0.8 97.1 0.79 0.77 0.971550 23.4 65.2 20.9 13.1 1.0 100.2 0.40 0.75 1.851600 27.0 67.2 22.5 11.5 1.0 102.2 0.34 0.78 2.271700 28.8 66.6 23.3 9.7 0.3 99.9 0.29 0.81 2.80

HMDS43AFa 48.3 15.5 26.8 8.0 98.6 1.11 0.75 0.671000 4.7 50.2 13.8 30.3 5.9 100.2 1.21 0.64 0.531300 8.0 53.2 14.0 31.5 3.8 102.5 1.18 0.61 0.521400 10.8 54.3 14.4 28.2 2.9 99.8 1.04 0.62 0.601500 18.0 58.3 13.0 26.9 1.4 99.6 0.92 0.52 0.561550 25.3 58.8 13.2 27.3 0.7 100.0 0.93 0.52 0.561600 25.2 65.4 16.7 19.3 1.0 102.4 0.59 0.60 1.011700 28.8 64.2 19.4 17.1 0.4 101.0 0.53 0.71 1.33

HMDS44AFa 47.1 13.0 26.7 10.9 97.7 1.13 0.64 0.571000 6.2 49.8 11.1 32.5 7.4 100.8 1.30 0.52 0.401300 9.7 53.3 11.1 32.4 4.5 101.3 1.21 0.48 0.401400 12.7 52.6 10.5 32.4 3.4 98.9 1.23 0.46 0.381500 19.1 59.0 9.8 29.7 2.0 100.5 1.01 0.39 0.381550 27.7 56.9 10.4 31.2 0.5 99.0 1.10 0.42 0.391600 24.4 62.4 12.1 23.3 1.1 99.0 0.75 0.45 0.611700 25.7 62.6 13.7 25.6 0.5 102.4 0.82 0.51 0.62

HMDS45AFa 47.6 8.3 28.9 13.4 98.2 1.21 0.41 0.331000 7.4 51.8 6.1 34.2 9.7 101.8 1.32 0.27 0.211300 11.6 55.4 5.7 32.4 5.9 99.4 1.17 0.24 0.201400 14.1 52.9 4.7 35.8 4.9 98.3 1.35 0.21 0.151500 17.0 58.5 3.6 35.9 0.9 99.0 1.23 0.14 0.121550 23.4 54.7 1.6 34.4 0.6 92.3 1.26 0.07 0.051600 20.4 60.3 4.2 33.8 1.3 99.6 1.12 0.16 0.14

Notesa As-formed powder.b The sum of the different element masses from 100% due to the uncertainty of the chemical

analyses and the hydrogen is not taken into account, namely for the as-formed powders.


p1chap3 2002/9/6 page 29 #5

The as-formed powders obtained from the gas phase are at least partially organised, andthe weight changes occurring during thermal treatments are much less noticeable than forthe powders obtained by the aerosol method from the liquid phase that are amorphous.

3.1.3 Chemical evolution during annealing in anitrogen atmosphere

As in the previous chapter, the chemical composition in stoichiometric compounds of thedifferent samples annealed under a N2 atmosphere has been calculated from chemical analy-sis obtained from CNRS analysis laboratory (Vernaison, France). The knowledge of thecomposition in stoichiometric compounds (Si3N4, SiC, SiO2, Y2O3, Al2O3), especially forannealed powders is important, because it enables us to evaluate roughly the composition offuture sintered materials. Moreover, this allows a more precise choice of the adapted powdersto obtain the desired material with specific properties.

The determination of the chemical composition in stoichiometric compounds for the dif-ferent samples is as follows: for Si/C/N/Al/Y samples, all Al and Y are assumed to be inthe form of Al2O3 and Y2O3, the remaining O being in the form of SiO2, then all N formsSi3N4 and the remaining Si is in the form of SiC. Free C is the difference between totalC and C bonded to Si in SiC. The rules are the same for Si/C/N samples apart from Aland Y that are not present. It must be noted that such a calculation is approximate both foras-formed and annealed powders. Effectively, in the case of amorphous as-formed powders,compounds such as Si3N4, SiC, Y2O3, Al2O3 do not really exist as crystalline compounds. Inannealed Si/C/N powders, crystalline compounds Si3N4 and SiC exist but the crystallisationis not always complete, depending on the annealing temperature (as shown by XRD mea-surements). In Si/C/N/Al/Y/O powders, Al and Y are not only bonded to O (Y2O3, Al2O3)but may also be combined in SiAlON or YSiAlON amorphous phases (see Chapter 5), asfor classical sintering of silicon nitride compounds in the presence of sintering aids (Y2O3,Al2O3) (Ekstrm and Persson 1990; Hoffmann and Petzow 1993). Note that the calculatedSi3N4 and SiC contents should be very similar to that which exists in materials because ofthe low quantity of Y and Al bonded to N. Furthermore, H present in as-formed powders hasnot been taken into account for these chemical measurements. The chemical composition instoichiometric compounds is presented for Si/C/N and Si/C/N/Al/Y powders, both obtainedfrom the liquid phase in Tables 3.1.2 and 3.1.3. Tables 3.1.2 and 3.1.3 also give the C/Nratios in the as-formed powders, because they are often used in the following sections tocharacterise a sample.

3.1.3.1 Si/C/N samples

The samples studied in this section have an initial C/N ratio in the range 0.22.6 (HMDS4045and HMDS36, 66 and 67). The chemical composition in atomic ratio (C/Si, N/Si and C/N) andits evolution as a function of annealing treatment between 1000C and 1600C is presentedin Table 3.1.1 for samples HMDS4045. The results for stoichiometric compounds afterannealing at 1600C, are reported in Table 3.1.2 for HMDS4045 and HMDS36, 66 and 67.Part of these results appear in the PhD thesis of Emmanuel Musset (Musset 1995).

Up to 1400C, the chemical composition of Si/C/N samples presented in Table 3.1.1 showsno significant changes in the different atomic ratio values. Meanwhile, at higher temperatures


p1chap3 2002/9/6 page 30 #6

Table 3.1.2 Chemical composition (wt%) of various HMDS samples afterannealing 1600C assuming stoichiometric composition

Samples Si3N4 SiC SiO2 C C/N (at)initial ratio

HMDS36a 5.3 91.5 1.9 1.3 2.8HMDS40 7.7 87.6 2.1 2.6 1.3HMDS41 19.1 75.1 3.3 2.6 0.9HMDS42 28.6 67.4 3.0 1.0 0.7HMDS43 56.4 37.7 3.0 2.9 0.7HMDS44 65.9 29.9 2.8 1.4 0.6HMDS45 88.9 8.4 2.5 0.3 0.3HMDS66 6.2 83.0 1.5 9.9 1.8HMDS67 6.3 82.4 1.5 9.8 2.4

Notea Without ammonia in the carrier gas during synthesis.

Table 3.1.3 Chemical composition of Si/C/N/Al/Y/O pre-mixed powders after annealing under N(1 h, 1600C) assuming stoichiometric composition

Powdersample

Chemical composition (wt%) (Al + Y )/Si Al/Y C/N initialratio

Y2O3 Al2O3 SiO2 Si3N4 SiC C

HSAl03 2.8 0 17.8 70.2 10.4 3.1HSAl04 7.1 0 18.3 66.7 8.1 3.7HSAl07 3.1 0 17.6 72.2 5.9 0.04 2.7HSAl09 3.3 1.7 92.3 0.2 2.5 0.03 0.2HSAl12 2.6 0.4 52.7 44.0 0.2 0.02 0.6HSAl13 3.3 0.8 17.0 78.5 0.4 0.03 1.4HSAl14 3.5 0.5 46.5 48.1 1.4 0.03 0.8HSAl15 3.9 0 77.5 16.1 2.4 0.04 0.5HSAlY05 1.3 3.8 0 19.3 64.1 9.0 3.5HSAlY16 0.1 0.3 1.9 92.4 3.1 2.1 0.01 14.8 0.3HSAlY17 0.6 0.9 1.4 51.9 43.7 1.5 0.01 3.2 0.8HSAlY18 1.1 1.1 1.5 14.5 80.9 0.8 0.01 2.3 1.1

different evolutions occur depending on the initial C/N ratio (Tables 3.1.1 and 3.1.2):

For the C-rich samples (HMDS36, 66, 67 and HMDS40), the C/N atomic ratios comparedto the as-formed powders increase and the powders are converted essentially towards SiC(between 80% and 90%) during annealing (C/Si atomic ratio near 1, see Table 3.1.1).Free C is always present in these samples after annealing at 1600C (Table 3.1.2).

For the N-rich samples (HMDS44 and 45), the C/N atomic ratios decrease with a sta-bilisation of the N content and the powders are converted essentially towards Si3N4(between 66% and 89%).

For the two samples with intermediate composition (HMDS42 and 43, initial ratio C/N =0.7), the C/N atomic ratios remain relatively stable with a stabilisation or a slight decreaseof the C/Si atomic ratios and a conversion of the powder towards a composite powdercontaining both SiC and Si3N4.


p1chap3 2002/9/6 page 31 #7

Whatever the composition of the initial powders, annealing induces a decrease in silicaand in free C content because of the carbonitridation of this compound which involves COevolution.

For C-rich powders, the change in composition (Table 3.1.2) to SiC-rich compounds, duringannealing at 1600C under a N2 atmosphere, can be essentially linked to the decomposition ofnitrogenous species (Si/C/N, Si3N4) occurring from 1450C. Such decomposition is impor-tant because of the high C content in as-formed powders (C/N > 0.75) (Musset 1995; Mussetet al. 1997). Effectively, extended X-ray absorption fine structure (EXAFS) (Chapter 5.4)and nuclear magnetic resonance (NMR) (Chapter 5.1) analysis of C-rich powders synthesisedfrom HMDS have shown that the atomic local arrangement of Si is essentially composed ofmixed SiC3N tetrahedra and of SiC4 tetrahedra which are already close to the SiC structure.This suggests the decomposition of slightly nitrogenous compounds (SiC3N) during anneal-ing in order to reach the SiC structure. Moreover, annealing induces a decrease in silicaand in free C content because of the carbonitridation of this compound which involves COevolution.

For N-rich powders, EXAFS and NMR analysis have shown that the atomic local arrange-ment is essentially composed of SiN4 tetrahedra which are the local environment of crystallineSi3N4. Therefore, the evolution towards a major phase of Si3N4 after annealing is notsurprising.

The drastic change between the evolution of Si/C/N towards SiC or Si3N4 compounds isobserved for the initial C/N atomic ratio in the as-formed powders near the 0.75 value andcorresponding to a 3SiCSi3N4 mixture (Figure 3.1.1).

3.1.3.2 Si/C/N/Al/Y/O samples

For samples obtained from a mixture of organometallic compounds as precursors, the com-ments are mainly focused on the aspects specific to the presence of Al and/or Y in the powders.Table 3.1.2 presents the results in stoichiometric compounds after annealing at 1600C for1 h under a N2 atmosphere for Si/C/N/Al/Y/O samples with an initial C/N ratio in the samerange as HMDS samples (0.22.6).

The most important difference between Si/C/N samples obtained from only HMDS andSi/C/N/Al/Y samples obtained from a mixture of organometallic precursors is the incorpora-tion of Al and Y metallic elements into the Si/C/N system, which could modify the evolutionof the chemical composition of the Si/C/N system during annealing. That is why a comparisonbetween Si/C/N and Si/C/N/Al/Y samples with a similar C/N ratio is presented below.

Tables 2.6 and 3.1.3 clearly show that, in the Si/C/N/Al/(Y) systems, the Al+(Y)/Si ratiosremain similar to those of the corresponding as-formed powders. This attests that the thermaltreatment under a nitrogen atmosphere (1600C) does not induce the degradation of all Aland Y containing species. Thus, Al and/or Y additives remain in the powder after annealing,which is important for the future sintering process, in which sintering aids (Y2O3, Al2O3)are necessary in order to obtain dense nanomaterials.

By comparing, after heat treatment, HSAl or HSAlY samples (Table 3.1.3) with HMDSsamples (Table 3.1.2), one can see that the evolution of the chemical composition inSi/C/N/Al(Y) system as a function of the C/N ratio is comparable to that in HMDS samplespreviously discussed. In this way, C-rich Si/C/N/Al/Y samples (initial C/N ratio between1.1 and 2.7) are converted essentially to SiC (content between 64% and 81%, see samplesHSAl03, 04, 07 and HSAlY05) while N-rich samples (C/N ratio between 0.2 and 0.3) are


p1chap3 2002/9/6 page 32 #8

45

44

43

Si3N4

SiC

SiO2C

42

41

40

0.0 0.5 1.0 1.5 2.0 2.5 3.0

0

C/N atomic ratio (as-formed powder)

Mas

s co

nten

t (%

)

20

40

60

80

100

36

Figure 3.1.1 Change in phase content (SiC, Si3N4, SiO2 and C) in HMDS powders after 4 hannealing under nitrogen at 1600C. (Reprinted from Musset 1995.)

converted essentially to Si3N4 (content of 92%, see samples HSAl12 and HSAlY17). Sam-ples exhibiting an intermediate composition (C/N ratio between 0.6 and 0.8) are convertedto a mixture containing both SiC and Si3N4 (SiC around 44 wt%, Si3N4 between 52% and53%, see HSAl12 and HSAlY17 samples). Note that C-rich samples also contain a noticeablequantity of free C (between 6% and 10%). Meanwhile, even if the evolution of the chemicalcomposition is similar between Si/C/N and Si/C/N/Al(Y)O samples, noticeable differencesappear where the C/N values and also the Si3N4 and SiC contents are concerned, especiallyfor Si/C/N and Si/C/N/Al/Y C-rich samples (C/N = 2.6). Figure 3.1.2 shows the evolu-tion of the C/N ratio after a 1 h annealing treatment at 1600C under a N2 atmosphere forsome C-rich samples (HMDS66, HSAl03, 04 and HSAlY05). In this figure, the C/N ratiois around 15 for HMDS powders while it is only of 4.35.1 for HSAl(Y) powders. Theseresults indicate that the loss of N is lower for powders containing Al and Y additives andexhibiting high C/N ratio. The SiC content is then lower and the Si3N4 content is higher inthese Si/C/N/Al/(Y)O samples than those in Si/C/N samples (see Tables 3.1.2 and 3.1.3). Thissuggests that the decomposition of N containing species (Si/C/N, Si3N4) is limited duringheat treatment of Si/C/N/Al/Y powders, probably due to the formation of a liquid phase in the


p1chap3 2002/9/6 page 33 #9

10

00 500

HMDS66HSAl03HSAl04HSAlY05

1000Temperature (C)

C/N

(at

. rat

io)

1500

Figure 3.1.2 Evolution of the C/N atomic ratio with annealing temperature. (Reprinted fromCauchetier et al. 1999 with permission from Kluwer Academic/Plenum Publishers.)

quaternary SiAlON or in the quinary YSiAlON diagram preventing the degradationof nitrogenous compounds.

3.1.4 Structural changes

XRD and infrared (IR) measurements were carried out in order to obtain information aboutthe nature of the different phases present and the evolution of structure and of the differentchemical bonding after heat treatment. They were carried out on both Si/C/N and Si/C/N/Al/Ysamples obtained from the liquid phase.


For annealing temperatures below 1500C, the XRD diagrams of Si/C/N nanopowders(HMDS samples) are flat indicating that the powders are still amorphous. The IR spectraare very similar to the spectra of as-formed powders (see for example Figure 2.8), apartfrom the absorption bands at 1170, 2950 and 1255 cm1 involving hydrogenated bonds thatdisappeared.

At 1500C, peaks are present on the XRD diagram (Figure 3.1.3), indicating the pres-ence of crystallised -SiC in HMDS40, 41 and 42 samples and - and -Si3N4 phases inHMDS45. For HMDS43 and 44, the peaks are small and broad indicating a lower degree ofcrystallisation. Figure 3.1.4 presents IR spectra for three samples (C/N = 0.73, 0.67 and 0.57,respectively for HMDS42, 43 and 44) at the same temperature (1500C). It shows two broadpeaks centred near 900 and 500 cm1 indicating the presence of an amorphous material andcorresponding to Si C N and Si N bonds, respectively. These results, together with theXRD result, indicate that crystallisation is not complete at this temperature, especially forsamples with an intermediate C/N ratio (HMDS43 and 44). This late crystallisation will bediscussed in more detail in Chapter 5.5.

Between 1500C and 1600C, noticeable changes occur both in the XRD diagram and inthe IR spectra (Figures 3.1.4 and 3.1.5). Figure 3.1.5 presents the XRD diagrams of samplesannealed at 1600C. C-rich powders (HMDS41 sample) lead preferentially to the -SiCphase, as was already noticed at 1500C. N-rich powders (HMDS45 sample) mainly show


p1chap3 2002/9/6 page 34 #10

-Si3N4

-Si3N4

Al Support

Al Al Al

Al Al Al Al

Al

-SiC

-SiC -SiC

-SiC

-SiC

-SiC

-SiC

-SiC -SiC

-SiC-SiC

-SiC

-SiC

HMDS45

HMDS44

HMDS43

HMDS42

HMDS41

HMDS40

20 30 40 50

Angle (2)

60 70 80

Figure 3.1.3 XRD patterns of HMDS powders after annealing at 1500C. (Reprinted from Musset1995.)

the presence of - and -Si3N4 phases and very small -SiC peaks. Samples exhibitingintermediate composition (HMDS43 and 44) lead to a mixture of SiC and Si3N4. In theIR spectra (Figure 3.1.4), the broad peak centred near 900 cm1 becomes sharper between1500C and 1600C. The shape becomes similar to the signature of crystallised SiC nanopow-ders obtained from the gas phase. The fine absorption bands present in the 800400 cm1range indicate the crystallisation of N-rich powders only present in HMDS43 and 44 samples(- and -Si3N4 phases, Luongo 1983). These results (XRD and IR) are in good agreementwith the calculated chemical compositions given in Table 3.1.2.


p1chap3 2002/9/6 page 35 #11

1700C

1600C

1550C

1500C

1500 1000

Wavenumber (cm1)

HMDS42

500

1700C

1600C

1550C

1500C

Abs

orba

nce

(a.u

.)

Abs

orba

nce

(a.u

.)

1700C

1600C

1550C

1500C

1500 1000

Wavenumber (cm1)

HMDS44

5001500 1000

Wavenumber (cm1)

HMDS43

500

Abs

orba

nce

(a.u

.)

Figure 3.1.4 Change in IR spectra during annealing under nitrogen between 1500C and 1700C.(Reprinted from Musset 1995.)


As mentioned above for Si/C/N samples, annealing treatment under a nitrogen atmospherealso induces a drastic change in the structure of Si/C/N/Al/Y powders. From amorphous,in their as-formed state, they become crystallised or at least partially crystallised after a 1 htreatment at 1600C under nitrogen atmosphere. Figure 3.1.6 shows XRD diagrams of Si/C/Nsamples (0.2 < C/N < 1.7) annealed under nitrogen at 1600C for 1 h. The nature and thecrystallographic variety of the phases that crystallise during heat treatment are very similar tothose obtained from Si/C/N powders (HMDS samples). Thus, for a C-rich Si/C/N/Al powder(HSAl13), the XRD diagram shows an evolution towards- and-SiC compounds while for aN-rich Si/C/N/Al powder (HSAl09), - and -Si3N4 are formed with a higher quantity of the phase compared to the phase. For samples having intermediate composition (HSAl12), -and -Si3N4 and -SiC phases are formed. The IR spectra of the same samples (Figure 3.1.7)are in good agreement with the XRD diagram, that is to say that SiC bonds (around 900 cm1)are detectable in a C-rich system, while Si N bonds in a Si3N4 structure are detectable ina N-rich system. The sample with intermediate composition contains the Si N bonds inthe Si3N4 structure and also Si C bonds. IR and XRD results are very consistent with thechemical composition mentioned above.

For C-rich samples (C/N = 2.6, HSAlY05 and HMDS66 samples), a comparison betweenthe XRD diagrams of Si/C/N and Si/C/N/Al/Y powders annealed at 1600C is presented(Figure 3.1.8). Moreover, the evolution of the IR spectra as a function of the temperature ispresented for the same samples (Figure 3.1.9).

XRD patterns of HSAlY05 and HMDS66 annealed at 1600C are shown in Figure 3.1.8.The Si/C/N/Al/Y powder annealed for 1 h (HSAlY05 sample) presents broad diffraction


p1chap3 2002/9/6 page 36 #12

Al

Al

Al

AlAlAl

Al

Al

-SiC

-SiC

-SiC

20 30 40 50

Angle (2)

60 70 80

-SiC

-SiC

-SiC

-SiC

-SiC

-SiC

-SiC

-SiC

-S

iC

-SiC

-SiC-SiC

-SiC

-SiC

-S

iC

-S

iC

-SiC-SiC

HMDS45

HMDS44

HMDS43

HMDS42

HMDS41

+

-Si3N4

-Si3N4

Al Support

Figure 3.1.5 XRD patterns of HMDS powders after annealing at 1600C. (Reprinted from Musset1995.)

lines of -SiC that are indicative of a low degree of crystallisation (Figure 3.1.8(a)). Thereis also evidence for the presence of the -SiC phase (2H polytype). When annealing timeis increased, well-crystallised -SiC (2H polytype) and -SiC are present (Figure 3.1.8(b)).In contrast, for the Si/C/N powder (HMDS66 sample) there is evidence of - and -SiCeven after only 1 h of annealing and no difference can be noticed in the diffractogram for anannealing duration of 1 or 4 h (Figure 3.1.8(c) and (d)). These results suggest that for C-richpowders, the presence of Al and Y either limits the crystallisation or enables the formationof an amorphous phase initiated in the YSiAlON quinary system.


p1chap3 2002/9/6 page 37 #13

Inte

nsity

(a.

u.)

10 20 30 40 50 60 70 80 90Angle (2 degrees)

-SiC +-SiC

(+) Si3N4+-SiC

-Si3N4>>-Si3N4

HSAI09

HSAI12

HSAI13

Figure 3.1.6 XRD analysis of annealed HSAl powder. (Reprinted from Mayne et al. 1999.)

Abs

orba

nce

(a.u

.)

2000 1800 1600

HSAI13

HSAI12

HSAI09KBr

Si-C

Si-C

Si-N in -Si3N4

Si-N in -Si3N4

1400 1200

Wavenumber (cm1)

1000 800 600 400

Figure 3.1.7 IR spectra of annealed HSAl powder. (Reprinted from Mayne et al. 1999.)

Figure 3.1.9 presents examples of IR spectra of the as-formed and annealed Si/C/N andSi/C/N/Al/Y/O powders (high C/N ratio). Up to 1400C, the spectra of Si/C/N/Al/Y/Opowder are very similar to the spectra obtained for Si/C/N powder. The only differenceseems to be a flat base line for the powders containing Al(+Y) additives whereas a pro-nounced dip appears near 1000 cm1 in the IR spectrum of powders without additives


p1chap3 2002/9/6 page 38 #14

20 40 602 (CuK) 2 (CuK)

80 20

Cou

nts

(a.u

.)

40 60 80

(a) HSAIY05/1 h (b) HSAIY05/4 h

(c) HSDS66/1 h (d) HSDS66/4 h

Figure 3.1.8 XRD patterns of samples (a) and (b) HSAlY05 and (c) and (d) HMDS66 after 1and 4 h annealing under N at 1600C. (Reprinted from Cauchetier et al. 1999 withpermission Kluwer Academic/Plenum Publishers.)

1600C/1 hrAbs

orba

nce

(a.u

.)

1600C/4 hr

1500 1000

Wavenumber (cm1)

500

1400C/1 hr

Si-CH3

Si-C

HMDS66

Si-N

As-formed

1000C/1 hr

1600C/1 hr

Abs

orba

nce

(a.u

.)

1600C/4 hr

1500 1000

Wavenumber (cm1)

500

1400C/1 hr

Si-CH3

Si-C

HSAI03

Si-N

As-formed

1000C/1 hr

1600C/1 hr

Abs

orba

nce

(a.u

.)

1600C/4 hr

1500 1000

Wavenumber (cm1)

500

1400C/1 hr

Si-CH3

Si-C

HSAIY05

Si-N

As-formed

1000C/1 hr

Figure 3.1.9 Changes in IR spectra with annealing temperature. (Reprinted from Cauchetier et al.1999 with permission Kluwer Academic/Plenum Publishers.)

indicating crystallisation or grain growth as observed previously in the annealing of SiCpowders obtained by a laser gas-phase driven reaction (Tougne et al. 1993) or from pyrolysisof a polycarbosilane (Sasaki et al. 1989). No significant change is observed when the dwelltime increases from 1 to 4 h and no signature characteristic of Al or Y is observed.

This study of chemical and structural changes during annealing of Si/C/N andSi/C/N/Al/Y/O samples under nitrogen atmosphere has shown a comparable evolution ofSi/C/N and Si/C/N/Al/Y/O samples with comparable C/N ratio. A more detailed comparisonseems to indicate that the presence of elements of sintering aids limits the degradation ofthe nitride phase during heat treatment of Si/C/N/Al/Y/O samples. Another point is that theelements Al, Y and O aids remain present in the powder after thermal treatment, which isencouraging for future elaboration of materials.


p1chap3 2002/9/6 page 39 #15

3.1.5 Information about changes in the grain sizes

BET measurements were carried out in order to measure the specific surface area, whichenables determination of the changes in reactivity and grain size after heat treatment. Thesemeasurements should be related to transmission electron microscopy (TEM) analysis (seeChapter 4).


Figure 3.1.10 shows the variation of only two Si/C/N samples (HMDS41 and HMDS44).Below 1000C, the SBET decrease can be explained by the smoothing of the particles dueto the reticulation reactions, which occur during the ceramisation of the powders. Then theSBET value remains constant up to 14001500C and the decrease at 1600C is due to thecrystallisation and growth of the particles, as confirmed by TEM (Chapter 4).


Annealing (nitrogen atmosphere, 1600C, 1 h) of Si/C/N/Al/Y/O powders induces a decreaseof the specific surface area compared with that of initial powders attesting to a structural andmorphological change during heat treatment such as crystallisation, structural modification offree C and particle growth. The difference between the specific surface areas of as-formed andannealed powders increase when the C/N atomic ratio decreases. Thus for C-rich Si/C/N/Al/Opowders (HSAl07 sample), the specific surface area varies from 110 m2/g (as-formed) to54 m2/g (annealed) while for N-rich Si/C/N/Al/O powders (HSAl09 sample), the decreaseof the specific surface area is more important (11212 m2/g). A similar evolution is alsotrue for the series of Si/C/N/Al/Y/O samples. This phenomenon could suggest a differentgrain growth process in C-rich powders compared to N-rich powders, that is to say that thegrain growth could be more pronounced in N-rich samples. Apart from a low grain growth,

100

80

60

40

20

01600140012001000800

Temperature (C)6004002000

BE

T s

urfa

ce a

rea

(m2

g1 )

HMDS41

HMDS44

Figure 3.1.10 Change in specific surface area during annealing under N for a C-rich HMDS powder(HMDS41 sample) and a N-rich HMDS powder (HMDS44 sample). (Reprintedfrom Musset 1995.)


p1chap3 2002/9/6 page 40 #16

this could also be related to the low crystallisation degree (see XRD patterns above) and toa structural modification of free C giving rise to some C ribbons as was already observed(Musset 1995) and as is usually formed from the heat treatment of amorphous C between1500C and 2000C (Dresselhauss et al. 1988).

These interpretations must be checked by TEM observations in

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Nanostructured Silicon-based Powders and Composites

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