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Mechanical properties of carbon-derived Si3N4+SiC micro/nano-composite

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Page 1: Mechanical properties of carbon-derived Si3N4+SiC micro/nano-composite

Int. Journal of Refractory Metals & Hard Materials 27 (2009) 438–442

Contents lists available at ScienceDirect

Int. Journal of Refractory Metals & Hard Materials

journal homepage: www.elsevier .com/locate / IJRMHM

Mechanical properties of carbon-derived Si3N4+SiC micro/nano-composite

Lucia Heged}usová a,*, Monika Kašiarová a, Ján Dusza a, Miroslav Hnatko b, Pavol Šajgalík b

a Institute of Materials Research, Slovak Academy of Sciences, Watsonova 47, 043 53 Košice, Slovak Republicb Institute of Inorganic Chemistry, Slovak Academy of Sciences, Dúbravská cesta 9, 845 36 Bratislava 45, Slovak Republic

a r t i c l e i n f o

Article history:Received 12 September 2008Accepted 12 September 2008

Keywords:Si3N4+SiCNano-compositeStrengthCreepFracture

0263-4368/$ - see front matter � 2008 Elsevier Ltd. Adoi:10.1016/j.ijrmhm.2008.09.012

* Corresponding author. Tel.: +421 557922111; faxE-mail addresses: [email protected], l

Heged}usová).

a b s t r a c t

In this study the mechanical properties of a recently developed carbon-derived Si3N4+SiC micro/nano-composite have been investigated with the aim to study the influence of the SiC addition on the room-temperature strength and high-temperature creep behavior. The bending strength values are in the inter-val from r0 = 675 MPa to r0 = 832 MPa with the Weibull modulus of 6.4 and 8.6 The contact strength ischanging from r0R = 1997 MPa to r0R = 1167 with mR = 17 and mR = 15. Fracture origins in the specimenstested in bending are clusters of pores and large SiC grains and in the specimens tested in contact modecone cracks. The composite exhibits high creep resistance thanks to the changed chemical compositionand viscosity of the intergranular phase comparing to the monolithic Si3N4 and due to the interlocked sil-icon nitride grains by the intergranularly located SiC nanoparticles.

� 2008 Elsevier Ltd. All rights reserved.

1. Introduction

Different approaches have been used during the last decadeswith the aim to improve the room- and high-temperature proper-ties, reliability and lifetime of silicon-nitride-based structuralceramics [1,2]. Using the ‘‘nanoparticle strengthening” approachSi3N4-SiC nano-composites have been developed in which nano-sized SiC particles are dispersed in the Si3N4 matrix which usuallyare located either intragranularly, with a size of approximately 30–40 nm, or intergranularly, with a size of approximately 150–170 nm [3].

The mechanical properties of Si3N4-SiC nano-composites havebeen studied by a number of authors during the last 15 years[4,5]. It was shown that these ceramics exhibit greater high-tem-perature strength and better creep resistance than the monolithicSi3N4. According to Rendtel et al. [6], the strength of the Si3N4-SiCnano-composite at 1400 �C lies between 80% and 100% of the roomtemperature strength, with creep rates as low as 1 � 10�9 s�1 at300 MPa and 1400 �C. This behavior can be explained by a thinnerand more refractory intergranular phase of the composites com-pared to the monolithic materials. This is caused by trapping of asignificant amount of intergranular glass in stress-free SiC grainpockets and by the presence of carbon in the intergranular phase[7]. Excellent oxidation resistance for Si3N4-SiC nano-compositescompared to monolithic silicon nitride was reported in [8]. Thisphenomenon was mainly connected with the change in the oxida-tion and damage mechanisms during operation at high-tempera-

ll rights reserved.

: +421 [email protected] (L.

tures. Cheong et al. [9] found a very high room temperaturestrength of a Si3N4-20 vol.% SiC nano-composite with Y2O3 + Al2O3

as sintering additives, however, a strength degradation occurredat temperatures higher as 1000 �C. Nano-composite with 4 wt%Y2O3 of sintering additive had the room temperature strength lower(approximately 1 GPa). However, this value of strength remainedup to 1400 �C due to the direct bond of the intergranularly locatedSiC particles to the Si3N4 matrix and due to the inhibition of grainboundary sliding and cavity formation. Dusza and Šajgalík reportedan increase in strength and Weibull modulus from 987 MPa to1203 MPa and from 6.7 to 19, respectively, by addition of 10% SiCto silicon nitride in a SiCN-derived nano-composite [10]. The goodstrength characteristics were associated with a fine and defect-freemicrostructure.

During the last years an inexpensive in-situ method utilizingcarbothermal reduction has been used for a preparation of thestudied Si3N4-SiC micro/nano-composite, which was characterizedmainly as regards its high-temperature properties [11].

The aim of the present contribution is to characterize thestrength and creep behavior of a carbon-derived Si3N4-SiC micro/nano-composite.

2. Experimental procedure

The studied C-derived Si3N4-SiC micro/nano-composite, withthe composition listed in Table 1, was prepared at Institute of Inor-ganic Chemistry, Slovak Academy of Sciences in Bratislava. Thesample contained SiO2 and C with the aim to achieve 5 wt% ofSiC by carbothermal reduction of SiO2 after densification. The start-ing mixtures were homogenized in polyethylene bottle with Si3N4

Page 2: Mechanical properties of carbon-derived Si3N4+SiC micro/nano-composite

Table 1Composition of the starting mixture

Compound Si3N4 Y2O3 C SiO2

Content (wt%) 84.1 4.4 4.1 7.4

L. Heged}usová et al. / Int. Journal of Refractory Metals & Hard Materials 27 (2009) 438–442 439

spheres in isopropanol for 24 h. The dried mixture was sievedthrough 25 lm sieve in order to eliminate large hard agglomerates.Green discs with a diameter of 48 mm and 5 mm thick were diepressed under the pressure of 30 MPa. Green discs were thenembedded into a BN powder bed and positioned into a graphiteuniaxial die. Samples were hot-pressed under a specific atmo-sphere, mechanical pressure, and heating regime at 1750 �C for 2 h.

The microstructure of the nano-composite was characterized byX-ray diffractometry (XRD), scanning electron microscopy (SEM),and TEM/HREM. XRD analysis was carried out using a Philips X-part diffractometer equipped with a CuK2 radiation source. Both,as-sintered and crept samples were analyzed. Polished and plas-ma–etched sections of the bulk materials were examined in SEM.To prevent surface charging during examination, the samples werecoated with a thin layer of gold. The overall structural and chemi-cal characterization of each specimen was carried out by conven-tional and analytical TEM investigation with a JEOL 2010microscope equipped with an ultra-thin window for energy–dis-persive X-ray spectrometer. For the structural analysis of grainboundaries and phase boundaries, as well as of intergranularphase, the HREM lattice imaging technique was applied using aJEOL JEM-12,000 microscope.

Thirty specimens with an effective volume of 18.4 mm3 and1.55 mm3 were tested in four-point bend on fixture with spans20 mm � 40 mm and 9 mm � 18 mm, respectively. Before testing,the specimens were ground to a 15 lm by a diamond grindingwheel. Both edges on the tensile surface were rounded with a ra-dius of about 0.15 mm in order to eliminate failure from the edgesof the specimens. The specimens were broken at a cross-headspeed of 0.5 mm/min, and the test environment was ambient air.Contact strength tests were realized, by using of the part of thespecimens after the bending strength test between spheres and be-tween rollers at the loading rate of 0.5 mm/min and calculatedaccording to Eq. (1) [12,13]. The characteristic strength and Wei-bull modulus were computed using the linear regression methodof two-parameter Weibull theory [14], SEM and EDX analyses wereused to localize and to characterize the fracture origin [15,16].

rmax ¼1� 2t2

3p6PE02

R2

!1=3

;1E0¼ 1� t2

1

E1þ 1� t2

2

E2ð1Þ

where E1; t1 are the elastic constants of the spheres and E2; t2 arethe related parameters of the ceramic.

Creep tests were performed in four-point bending using a fix-ture made of silicon carbide with inner and outer span length of20 mm and 40 mm, respectively. The measurements were carriedout in a creep machine with dead-weight loading system in airatmosphere at temperatures between 1200 �C and 1450 �C withouter fibre stresses in the range from 50 to 150 MPa. The sampledeflection was recorded continuously during the creep test. Fromthe deflection data, the outer fibre strain was calculated as a func-tion of time, t, and taken as the creep strain, e. The creep rate wascalculated from the slope of the e vs. t curve. The steady-state creeprate is usually described by the Norton equation

_e ¼ Arn 1dm exp �Q c

RT

� �ð2Þ

where A is a constant, depending on the respective material proper-ties and microstructure, r is the stress, n is the stress exponent, d is

the grain size, m is the grain size exponent, QC is the activation en-ergy of creep, and T and R have their usual meaning.

3. Results and discussion

The characteristic microstructure of the experimental materialis illustrated in Fig. 1a, and b. The microstructure analysis usingceramography and SEM revealed that technological defects arepresent in the form of clusters of SiC and porosity in the mate-rial only randomly. The Si3N4-SiC nano-composite consists of avery fine, homogeneously distributed Si3N4 grains with a low as-pect ratio. The composite additionally contains globular nano-and submicron-sized SiC particles located intragranularly in theSi3N4 grains or intergranularly between the Si3N4 grains. It wasdifficult to distinguish between the SiC particles located integra-nularly and the grain boundary phase because they are similarlyaffected by plasma etching. The mean diameter of Si3N4 grain is140 nm and grains with a diameter larger than 500 nm were ob-served in the microstructure only occasionally. The volume frac-tion of the SiC nanoparticles can be estimated approximately as5 vol%. X-ray analysis revealed that the main phase in material isb-Si3N4 with a small amount of b-SiC. Beside the Si3N4 and SiC,some additional crystalline phases were detected in the compos-ite, mainly Y2SiO7 and Si2N2O. HREM revealed different thicknessof the intergranular phase in the composite indicating that prob-ably the equilibrium state was not reached during the sinteringprocess. Analytical electron microscopy revealed different com-position of intergranular phase in composite and monolithic sil-icon nitride (prepared at a similar condition but without C andSiO2 additives), with higher carbon and lower oxygen contentin the composite.

Weibull distribution of the measured four-point flexurestrength values of the investigated nano-composite with the spec-imen effective volume of 18.4 mm3 is illustrated in Fig. 2a. Usingthe two-parameter Weibull statistics the characteristic strengthof material was r0 = 675 MPa and Weibull modulus was m = 6.4.Slightly higher value of characteristic strength r0 = 832 MPa andWeibull modulus m = 8.6 was obtained tested the specimens witheffective volume of 1.55 mm3.

Very high strength and Weibull modulus was obtained duringthe contact strength test of the material between spheres,r0 = 1997 MPa and m = 17.1 for the specimens with bigger volumeand r0 = 1167 MPa and m = 15 for the specimens with smaller vol-ume, respectively, Fig. 2b.

After the bending test the Weibull parameters, the characteris-tic strength and Weibull modulus of the carbon-derived compositeare lower in comparison with the parameters of the SiCN-derivedSi3N4-SiC micro/nano-composite tested in bending, [10]. Bendingtests and following fractographic analysis of the fracture surfaceof failure specimens revealed that a reason of the lower character-istic strength compared to the strength of the SiCN-derived mate-rials is a presence of the processing flaws in the investigatedmaterial. In most cases the location of the fracture origins/flawswere in the volume of the specimens (42%), but they were also lo-cated near the surface (23%), on the surface (20%), and at the edges(15%), Fig. 3a, and b. Fracture origins were in all cases processingflaws and machining induced flaws were not found. Study of thesize and shape of the fracture origins revealed that their size werein the range from 10 lm to 180 lm, with the mean value of half-minor axis length of about 35 lm. Their shape was mostly ellipti-cal, but circular-shaped fracture origins were also observed. In themajority of cases it was easy defined the shape of the defects atleast in two-dimension form. Fractographic analysis of the fracturesurfaces of failured specimens combined with EDX analysisrevealed two main types of processing flaws acting as a failure

Page 3: Mechanical properties of carbon-derived Si3N4+SiC micro/nano-composite

6.6 A 50 nm

Si3N4/Si3N4boundary

SiCnano

a b

Fig. 1. Characteristic microstructure of the material: (a) TEM and (b) HREM.

Fig. 2. Weibull distribution of strength values (a) in bending and (b) in contact mode.

400 μm 200 μm

Fig. 3. Characteristic fracture origin in bending specimen.

440 L. Heged}usová et al. / Int. Journal of Refractory Metals & Hard Materials 27 (2009) 438–442

initiating flaws in the studied material. The first type of fractureorigin was a porous region, often connected with another type offlaws. The second type of fracture origins was a cluster of largeSiC grains. An area of non-reacted carbon and SiO2 was found,too. However, this third type of fracture origin occurred rarelycompares to the previous ones.

The fracture origins after testing the specimens with an effec-tive volume of 1.55 mm3 were the same defects as above de-scribed, but they were located mainly at the tensile surface ofthe specimen.

According to the result of ceramographic/fractographic exami-nation the fracture during the contact strength test ‘‘sphere/sphere” has been initiated under the surface of the specimensand the fracture site is not connected with a technological defect.During the contact strength test ‘‘sphere/sphere” the fracture iscaused by creation and growth of the cone cracks during the testup to the critical size, Fig. 4. The higher Weibull moduli (the lowerscatter in the strength values) in contact strength ‘‘sphere/sphere”test is connected to the relatively similar sized cone cracks at thefailure.

Page 4: Mechanical properties of carbon-derived Si3N4+SiC micro/nano-composite

150 μm

Fig. 4. Cone crack formation during the load in contact strength.

L. Heged}usová et al. / Int. Journal of Refractory Metals & Hard Materials 27 (2009) 438–442 441

Comparing the results of ceramography of polished/etchedsection and fractography of fracture surface after the bendingtest we can say that bending test together with fractography re-veals the technological defects which cannot to be revealed byceramography. The presence of such a defects are clearly showthat for the developed material the processing is still not opti-mized and further work has to be done in this direction, mainlyas regards the homogenization of the mixture and the optimiza-tion of the carbothermal reduction. These defects are responsible

Fig. 5. Creep behaviour of the micro/nano-composite and the monolithic siliconnitride with the activation energies of the creep.

Fig. 6. Schematic illustration of creep in

for the low characteristic strength and low Weibull moduluscomparing to the similar systems developed during the lastdecade [10].

The creep deformation of the composite can be characterized bya primary creep range and a pronounced steady-state creep range.The composite ceramics exhibits only minimum creep deformationup to 1300 �C and significant creep deformation was only detectedat the temperatures from 1350 �C to 1450 �C. The creep resistanceof the nano-composite was found to be significantly higher com-pared to the creep resistance of the monolithic Si3N4, Fig. 5. Atthe temperature of 1350 �C the composite exhibits similar creepstrain at 150 MPa/150 h as monolithic material at 50 MPa/25 h.The stress exponents and activation energies together with the re-sults of TEM examination show that there exist different creepmechanisms in the monolithic silicon nitride and in the nano-com-posite. In monolithic silicon nitride the dominant mechanismswere probably cavitation together with grain boundary slidingand in the Si3N4-SiC micro/nano-composite probably solution/pre-cipitation mechanism, Fig. 6.

Besson et al. investigated the compressive creep behavior ofSi3N4 + 10 wt% SiC nano-composite and of reference monolithicSi3N4 in the temperature range from 1250 to 1450 �C under stres-ses from 45 to 180 MPa. Stress exponents of 0.8 and 1.0 and appar-ent activation energies of 514 and 590 kJ mol�1 were found formonolithic and composite systems, respectively [17]. Grain bound-ary sliding accommodated diffusion through the intergranularphase is considered the main creep mechanism. Rendtel et al.[18] investigated the creep behavior of Si3N4+SiC nano-compositewith different sintering additives and with different wt% of addi-tives from 2.5% to 30% in the temperature range from 1400 to1550 �C at stresses from 30 to 200 MPa. A minimum steady-statecreep rate was found at about 10–15 wt% SiC in the composite.The maximum detected rate reduction compared to the monolithicceramic is a factor of three. Recently Rendtel and Hubner [19]investigated the effect of the heat treatment on microstructureand creep behavior of Si3N4+SiC nano-composites. Chemical mod-ification of the intergranular phase, grain growth, and a pro-nounced interlocking of the grain facets in the vicinity of SiCparticles are suggested to explain the increased creep resistanceof the nano-composite compared to the creep resistance of themonolithic material. Recently Wan et al. [20] used amorphousSi–C–N powder, derived from pyrolysis of a liquid polymer precur-sor, with Y2O3 additive from 8 wt% down to 0 wt% for the process-ing of Si3N4+SiC nano-composites. They measured excellent creepresistance for such composites mainly for material with very lowvolume fraction of additives and clean silicon nitride boundarieswithout a glassy grain boundary phase. The results of the presentcontribution are in good agreement with the results of the

monolithic and composite ceramics.

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442 L. Heged}usová et al. / Int. Journal of Refractory Metals & Hard Materials 27 (2009) 438–442

mentioned authors obtained by investigating similar ceramiccomposites.

4. Conclusions

Strength and creep behavior of a carbon-derived in-situ rein-forced Si3N4-SiC micro/nano-composite have been investigated.The following main results were found:

� The Weibull parameters in bending were r0 = 675 MPa/r0 = 832 MPa and 6.4/8.6 and in contact sphere on sphere testr0 = 1997 MPa/r0 = 1167 MPa and m = 17.1/m = 15 for speci-mens with different effective volume, respectively.

� Processing related fracture origins located in bulk and near thetensile surface of the specimens were found with dimensions inthe range from 10 lm to 180 lm.

� The fracture during the contact strength test ‘‘sphere/sphere”was caused by initiation and growth of a cone cracks to criticalsize.

� The composite material exhibits significantly higher creepresistance compared to the monolithic material probably dueto the improved viscosity of the intergranular phase and theinterlocked silicon nitride grains by the intergranularly locatedSiC nanoparticles.

We can conclude that the influence of the SiC nanoparticles onthe strength is in two directions; due to the hindering of graingrowth of Si3N4 grains is positive, but due to the reason of the clus-ter formation is negative. The influence of the SiC nanoparticles onthe creep resistance is evidently positive. The processing steps forthe material preparation have to be improved with the aim to offerthis material for structural applications.

Acknowledgement

This paper was supported by Slovak Grant Agency via No. 2/7194/27, by Nanosmart, Centre of Excellence of SAS, APVV 0170-06 and by KMM-NoE project of the EU 6FP.

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