Impurities in silicon carbide ceramics and their role

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Impurities in silicon carbide ceramics and their roleduring high temperature creep

M. Backhaus-Ricoult, N. Mozdzierz, P. Eveno

To cite this version:M. Backhaus-Ricoult, N. Mozdzierz, P. Eveno. Impurities in silicon carbide ceramics and their roleduring high temperature creep. Journal de Physique III, EDP Sciences, 1993, 3 (12), pp.2189-2210.�10.1051/jp3:1993269�. �jpa-00249076�

https://hal.archives-ouvertes.fr/jpa-00249076

https://hal.archives-ouvertes.fr

J. Phys. III France 3 (1993) 2189-2210 DECEMBER 1993, PAGE 2189

Classification

Physic-s Abstracts

20.790 61.16D 61.70W

Impurities in silicon carbide ceramics and their role duringhigh temperature creep

M. Backhaus-Ricoult, N. Mozdzierz and P. Eveno

Laboratoire de Physique des Mat6riaux, CNRS, 92195 Meudon, France

(Received ii May J992. rei~ised 5 July J993, accepted14 September J993)

Rksumk. La microstructure de deux cdramiques, SiC sans ajouts et SiC avec bore et carbone,

est 6tud16e par microscopie dlectronique h transmission dans le but d'6valuer l'influence des

additifs sur les propr16t6s m6caniques h haute temp6rature. Dans tous les mat6riaux, des pr6cipit6sde graphite de diff6rentes tailles sont observds. Le carbure de silicium fabriqud avec du bore

contient des grands pr6cipit6s de B25C et des petites poches de silice amorphe. A partir de nos

observations de la microstructure, une pr6vision des propr16tds m6caniques des matdriaux est

possible. Ces pr6visions sont compar6es aux rdsultats de fluage h haute tempdrature. Les mat6riaux

sans ajouts et ceux avec carbone et bore sont ddformds entre 1773K et 1973K, sous des

contraintes de loo h I loo MPa. Le comportement des deux matdriaux suit une loi de puissance

avec un exposant de contrainte de 1,5 pour les faibles contraintes et de 3,5-4 pour )es fortes

contraintes. Les valeurs d'dnergie d'activation des deux types de matdriaux sont respectivement364 et 453 kJ/mole dans le domaine de faibles contraintes et 629 kJ/mole aux forte~ contraintes.

L'observation de la microstructure des mat6riaux ddformds montre comme mdcanisme principal de

fluage le glissement aux joints de grains accommod6 par la diffusion et accompagnd par une faible

cavitation due h la non-perm6abilit6 des joints de grains pour )es dislocations. Aux plus fortes

contraintes, )es joints de grains deviennent plus perm6ables, et la deformation par m6canismes de

dislocations devient alors prdpond6rante.

Abstract, The high-temperature compressive creep behaviour of hot-pressed silicon carbide

ceramics with different additive packages (boron and carbon or no additive) is investigated as a

function of several parameters : the microstructure, the nature of the additives and that of the

impurities. Additional carbon is present in all the materials investigated, as graphite precipitates of

various size and amount. In materials densified with addition of boron, large precipitates of B25C

and small amorphous silica pockets are identified. In the case of materials containing impurities,small precipitates of Fesi, Fe or Ti~si~ are detected. Creep experiments are conducted on

materials with no additives and on others containing boron and carbon additives, at temperatures

ranging from 773 K to 973 K and under stresses from loo to I loo MPa. A comparison of the

creep behaviour of the various materials points out to the destructive effect of carbon precipitates

on the creep rate : the stationary creep rate of the material containing carbon (and boronj additives

is by a factor 2.5-5 faster, eventhough its grain size is much larger The creep of both investigatedmaterials is described by a power law with a stress exponent of 1.5 in a low stress range and 3.5-4

2190 JOURNAL DE PHYSIQUE III N° 12

in a high stress range. The corresponding activation energies are 364 kJ/mole and 453 kJ/mole in

the low stress range and about 629 kJ/mole in the high stress range. At low stresses the materials

deform by grain boundary sliding compensated mainly by diffusion along the grain boundaries and

to a lesser extent by limited cavitation, as a result of the barrier role played by grain boundaries for

dislocations. At high stresses the grain boundaries are no longer an obstacle to dislocation motion,

which becomes the dominant deformation mechanism.

1. Introduction.

Due to its high performance in terms of hardness, chemical resistance, thermal conductivityand thermal shock resistance, silicon carbide has become one of the most often used ceramics

for structural applications at high temperatures. Its domain of application extends from heat

exchangers, seals or burners, to a wide range of coupling pieces, whose main purpose is the

prevention of friction and wear. In many instances, silicon carbide pieces must bear

mechanical or thermal stresses in temperature gradients. In this context, it is valuable to

experimentally measure fracture toughness and creep behaviour of the material at high

temperatures.The literature data available for fracture toughness and creep of different types of silicon

carbide are widely scattered. For example, reported fracture stress data range from 20 MPa to

414 MPa, at temperatures between 500 K and 2 500 K [1, 2]. A similar conclusion can be

drawn from creep data obtained in compressive mode [3-16].

After literature results [5-9, 26-29], single crystals, chemical vapour deposited (CVD) and

under certain conditions even polycrystalline silicon carbide deform at high temperature by a

dislocation mechanism. The basal slip system of silicon carbide is activated at temperatures as

low as 1300K, while the prismatic slip system becomes activated only at much higher

temperatures (about 2300K). Glide of perfect or partial dislocations in the basal planetogether with cross-slip, dislocations climb or short range diffusional transport are normally

considered at possible «dislocation» mechanisms responsible for a macroscopic creep

deformation of these silicon carbide materials.

It has been shown in the past that small quantities of second phases can effect the creepbehaviour of polycrystalline silicon carbide : the presence of B4C leads to embrittlement of the

material ; aluminium, as an additive, can result in the formation of continuous ductile silicate

films, which enhance grain boundary sliding [1, 17].

In the following, we want to mention the major results reported in the literature on the

deformation of polycrystalline silicon carbide, without attempting to give a complete literature

review on the subject.Farnsworth and Coble [10] investigate the creep behaviour of a hot-pressed, high density

SiC, containing 1.6 9b Al and 9b Fe and having an average grain size of 3 ~Lm by four-pointbending experiments between 2 173 and 2 473 K, under stresses ranging from 20 to 200 MPa.

Based on the measured strain rates (10-6 s-' at 2 300 K and loo MPa), the authors conclude

that diffusion of carbon along the grain boundaries is the controlling mechanism for creep and

confirm their hypothesis by further experiments with materials having different grainsizes [11]. No details on the microstructure of the materials art given.

Compressive creep experiments on sintered and hot-pressed silicon carbide containing boron

and carbon as sinter additives and having an average grain size of 3.5 ~Lm are reported by

Djemel et al, for temperatures between 1500 and 1773 K and stresses between 500 and

700MPa[12, 13]. These authors identify two deformation ranges a low stress range,

N° 12 IMPURITIES IN SiC 2191

characterized by a stress exponent close to one and an activation energy of 292 kJ/mole,correlated to Coble creep, and a high stress range with large strain rates and stress exponents

larger than lo, explained by a cavity formation mechanism in the tensile grain boundaries and

to microcrack propagation. In this source as well, the microstructure of the materials before

and after creep is not analyzed at a fine scale and no information is provided about further

contributing creep mechanisms or the role played by impurities.

Hamminger et al. [14] characterize the microstructure of four different sintered silicon

carbides doped with tither aluminium and carbon, or boron and carbon and creep some of

those materials at temperatures between 1743 and 1933 K and under stresses between loo

and 190 MPa Ii 5]. The microstructure after creep is not described. From the measured stress

exponent, close to I, and the high activation energy, 796 kJ/mole, the authors conclude that

creep is controlled by volume diffusion.

Moore et al. [16] analyze a boron-containing sintered silicon carbide by scanning Augermicroprobe and electron energy loss spectroscopy. They identify carbon precipitates,micrometer-large B4C grains and, after annealing, small additional coherent B4C precipitatesinside the grains. No enrichment of boron in the grain boundaries or formation of intergranularfilms is reported. Creep experiments (1 670-2 073 K, 138-414 MPa) performed on the same

materials are also reported by Davis et al. [17]. After creep, an increased density of stackingfaults is observed together with dislocation glide bands, which interact with the precipitates.Two temperature regions are distinguished by the authors : I) a low temperature region (typicalstrain rates of 3.7 x

10-1° s-' at 670 K and 179 MPa) with a stress exponent between 1.44

and 1.7 and an activation energy ranging from 338 to 434 kJ/mole, where, according to the

authors, grain boundary sliding is accommodated by grain boundary diffusion iii a high-

temperature range (typical strain rate of 6.3 x10-? s-' at 2073 K and 179 MPa) with a

higher activation energy, 802 to 914 kJ/mole, where grain boundary sliding is compensated bylattice diffusion. The authors explain that dislocation glide is the dominant mechanism at highstrain and high temperature, whereas dislocation climb only plays a minor role.

The scatter of the reported strain rates, stress exponents and activation energies can in partbe explained by the difference in density of the materials investigated, corresponding to

differences in the fabrication process used (natural sintering, uniaxial pressing or hot-isostatic

pressing). However, even when similarly dense materials are compared, non-negligiblediscrepancies remain among the mechanical properties reported by various authors [10-1?, 14,

16]. These differences are generally ascribed to the type of additives used for processing and to

the nature of the impurities introduced during the fabrication of the material. Since in its pureform silicon carbide powder does not easily sinter to a fully dense state, elemental carbon and

boron or aluminiumliron are normally used as sintering aids. Even though these additives or

impurities are normally present at very small concentration levels (typically 9b or less), they

can play a critical role for the mechanical properties of the material. Impurities and additives

can be present as solid solutions with silicon carbide or as second phase particles.The objective of the present study is to identify the creep mechanisms of different silicon

carbides, and their relation to the microstructure of the material, especially to the nature of

additives and impurity phases.

The materials selected for this investigation are hot-pressed or hipped silicon carbide

polycrystals fabricated either without additives or with carbon and boron as additives. They are

deformed at temperatures between 1773-1973K and for stresses ranging from 100 to

100 MPa, in compressive mode. The microstructure of the various silicon carbide ceramics

is analyzed before and after creep, to more precisely locate the additives and the impurities, to

identify their chemical composition and crystalline structure and, finally, to evaluate their role

during high temperature creep. Given the possibly very small size of the impurity precipitates


(down to loo nm), transmission electron microscopy (TEM) combined with energy dispersive

X-ray analysis (specially adapted for light element detection) is selected, which enables to

resolve both structure and chemical composition (light elements like boron, carbon and

oxygen) of these precipitates.

2. Experimental procedures.

2. I MATERIALS.

Material A is prepared by hot isostatic pressing of very pure silicon carbide powders

(Starck A lo) in tantalum containers, at 2 273 K and 2 000 bar, without any additive [18]. It is

densified to about 96 9b of the theoretical density. Traces of free carbon are found in the

material, which are related to the fabrication process itself. During the hipping process the

sealed tantalum containers are in direct contact with the silicon carbide powder compacts.

Tantalum is known as a good getter material for oxygen, but at high temperature it reacts as

well with SiC and forms tantalum silicides. A thin layer of silicide forms at the metal container

wall during the short duration (I h) of HIP at 2 273 K, thereby leaving excess carbon. Silicon

carbide being a very stoichiometric compound, the precipitation of carbon occurs at low

energy locations, like grain boundaries, where it results in the formation of intergr,anular

graphite films.

Material B, a commercial silicon carbide, consists of more than 98.5 vol9bSiC, about

vo19b boron and impurities of silicon, iron, oxygen and carbon at a total concentration below

2 000 ppm.

2.2 HIGH TEMPERATURE ANNEALING. Materials A and B are annealed in the creep

apparatus without applied stress, at 873 K for 25 h, under argon flow. The microstructure of

these samples is investigated by transmission and scanning (SEM) electron microscopy.

2.3 CREEP EXPERIMENTS. Compressive creep tests of materials A and B are performed at

temperatures between 1773 and 1973 K, stresses between loo and I loo MPa and under

argon flow. The 1.5 mm x1.5 mm x 4 mm samples with planparallel sides are deformed

under constant load in compressive mode. Most samples are subsequently deformed under

different loads or at different temperatures, to strain levels up to about 3 fl for each condition.

Samples are then cooled in the furnace under applied load. For microscopy investigations the

samples are crept only under one condition.

The constant-load creep apparatus used for this investigation is described in detail in [19]. It

comprises a furnace with lanthanum chromite heating elements. The load is applied onto the

sample i~ia silicon carbide push rods. Two silicon carbide platelets (10 mm x lo mm x 4 mm)

are placed between the push rod and the sample ends. The entire set-up is enclosed in a gas-tight alumina tube located inside the furnace, so that the atmosphere around the sample can be

controlled. All tests are run under continuous argon flow. A thermocouple located near the

sample is used to control the furnace temperature. The shortening of the sample is recorded by

a direct current deformation transducer. At the end of each experiment, the total recorded

deformation is compared to the final length of the sample.Thin slices are cut parallel and perpendicular to the deformation axis in the crept samples.

Some of them are observed by SEM and others, after further thinning, by TEM. For SEM and

TEM observations, special care is taken to keep track of the deformation axis of the sample.

2.4 CHARACTERIZATION OF THE MICROSTRUCTURE-A scanning electron microscope

(Philips 100B) equipped with an energy dispersive X-ray analyzer (EDX) is first used to

analyze the bulk samples. Their more detailed investigation is subsequently performed by

N° 12 IMPURITIES IN SIC 2193

transmission electron microscopy (TEM). For this purpose, thin slices of the materials are cut

with a diamond saw, mechanically grinded and polished to a thickness of 40 ~Lm, before final

ion milling by argon bombardment. The so-prepared thin samples are investigated in a

transmission electron microscope (JEOL 2000FX) equipped with an EDX analyzer particularlyperformant for the detection of light elements, such as boron, carbon and oxygen. Electron

diffraction, microdiffraction with a beam size of lo nm, chemical identification by EDX, and

high resolution lattice imaging, specially of carbon, are performed on as-received, annealed

and deformed samples.For most of the materials the grain size distribution is determined by measurements made on

SEM overview micrographs and on a large number of TEM micrographs taken at low

magnification, a technique which provides a good precision for such fine grained materials.

3. Experimental results.

3.I MICROSTRUCTURE OF THE NON-DEFORMED MATERIALS.

3. I. I Material A. The average grain size of this material is 0.5 ~Lm, with only a very small

fraction of grains larger than I ~Lm, as can be seen from the grain size distribution of this

material given in figure I. Most grains are equiaxial. The major polytypes present in the

material are determined by X-ray diffraction to be 51R, 15 R, 6 H and 4 H. All the grainsconsist of SiC. In most cases, no grain boundary film can be observed, leaving the SiC lattices

of adjacent grains in direct contact. However, in rare instances (evaluated to about 0.2 9b of the

grain boundaries), an intergranular film of thickness close to 20-50 nm is detected, see

figures 2a and 2b, which consists of pure carbon. This thin intergranular layer is crystalline and

has a well-ordered structure, as attested by its diffraction pattem (Fig. 2c) and its highresolution image, figure 2d, which shows the well aligned basal planes of carbon within a

ramified structure.

ioo

80G

r

a

60

n

u

m

b

e

r20

Oo-o,2 o,4-o,8 o,8-1 1,2-1,4 1.8-1,8 .2

Grain 8ize (vm)

I7oo reoieved -annealed fl%&defo,med

Fig. I. Grain size distribution of material A as processed, after annealing for 25 h at 873 K and after

deformation (25 h, 870 MPa at 873 K).

JOURNAL DE PHYSIQUE Iii T 3, N' 12, DECEMBER 1993 79


@'OO02

~~~~

' ~

,/"~~ ~

~C ~

' .'

j

~~~~~

~-w~

a)

b)

coun~« lx10'j~°~~~* ~~'°'~

4

graphite carbon grain of SiC

6 3

s,

5

4 z

3

z i

I

«e

~'ae

Z 4 6 10 Z 3 4 5

RanBe 'keVI RanR© lkevi

c)~~

Fig. 2. a) TEM bright field image of material A, showing graphite grain boundary films, b) EDX

spectra of the graphite film and of adjacent grains. c) Diffraction pattern of the grain boundary film.d) HREM image of the graphite film, showing the alignment of the graphite basal plane parallel to the

grain boundary and parallel to the SiC basal plane (parallel to the visible stacking faults in SiC).


3.1.2 Material B. The density of this commercial material approximates 97 9b of the

theoretical value, resulting in only rare pores. The average grain size of material B (about3.5 ~Lm see grain size distribution in Fig. 3) is larger than that of material A. The material

consists ofa

SiC. Its major polytypes are again detected by X-ray diffraction : 51 R and 6 H.

EDX analysis in the scanning electron microscope does not indicate any additional elements

besides silicon and carbon. However, the impurity concentrations below the detection limit of

the SEM-EDX system give rise to specific precipitates, visible in TEM.

~20

u

mb

e

r

O1,2 2-3 3,4 4-5 ,5

Grain size (vml

Fig. 3. Grain size distribution of material B as processed.

A representative low magnification image of the material is shown in figure 4. Even at this

scale, several impurity phases can be distinguished from the typical SiC grains :

Carbon precipitates. -As in the case of materialA, carbon is identified as one of the

impurity phases and three types of occurrence are found, see figure 5.

. Large (up to I ~Lm) bundles of well-aligned carbon fibres. Each fibre has a slightlydifferent orientation, corresponding to various amounts of departure from the carbon basal

plane (Fig. 5b). This small misorientation is not only demonstrated by the fluctuations in image

contrast, but even more by the diffraction pattern, where the 0002 reflections are distributed in

two concentrated half moons. This feature is typical of turbostratic carbon. High resolution

imaging confirms the graphite character of the precipitates. The carbon basal planes are well

aligned inside each fibre of the bundle, but distorted between the different fibres, see figure 5c.

. Long graphite fibres. -They are similar to those observed in material A, but more

common and of larger dimensions (d=

400 nm,f

~2 ~Lm), see figure 5d. These fibres consist

of turbostratic carbon without any misorientation inside the entire precipitate, contrarily to the

observation reported for the previous type of precipitates. These graphite fibres are often

located along silicon carbide grain boundaries.

. Globular precipitates of amorphous carbon. As seen in figure 5a, these precipitates

are characterized by typical homogeneous rings in the diffraction pattern. Their size averages

200 ~Lm x 300 ~Lm. This type of precipitate is rare.

2196 jOURNAL DE PHYSIQUE III N° 12

*

%

.~'~b

".~

~/

~ '-

'~ ~%,/+%/,~~t~~ ~

p2 pm

Fig. 4. TEM overview micrograph of material B, showing the frequency and size of impurity phases.

In summary, the fibrous carbon bundles are the most common type of carbon precipitatesand the very small amorphous carbon precipitates are rare. The overall frequency of occurrence

of carbon precipitates is about one per loo grains of SiC.

Metal compounds. Although they are not indicated by EDX/SEM analysis, metal

compounds containing impurities exist in the materials and can be occasionally detected byTEM as precipitates, always located in the neighbourhood of carbon precipitates.

Triangular shaped iron- and silicon-containing grains with a typical size of loo nm are

frequent. They are always located at triple junctions, next to carbon precipitates, see figure 5d.

Noticeable compositional variations are measured from one such precipitate to another ; some

of them contain several percents of either chromium and nickel, or zirconium. Furthermore,the total metal to silicon ratio varies one precipitate out of two contains only small amounts of

silicon. For a number of precipitates a non-negligible amount of oxygen is indicated in the

EDX spectra, but this oxygen signal comes in all cases from amorphous silica present in the

neighbourhood of the precipitate. Due to the variation in thin foil thickness from one

precipitate to another (absorption from the heavier metals affecting the silicon signal) and to

the small size of the precipitates (superposition of precipitate and surrounding matrix or

amorphous silica film), the EDX data cannot always be quantified. Therefore, the precipitate


OO04

( )/-/

OO02

_

~

'i.

_

loo nm

ioo nm b)~

'

-

ysnm

c) loo Dm ~~

Fig. 5. Typical TEM bright field images and electron diffraction patterns of the different types of

carbon precipitates occurring in material B : a) globular precipitate of amorphous carbon, b) fibrous

bundle of turbostratic carbon, c) HREM image of the fibrous boundle~ showing the carbon basal planes,d) graphite fiber with a small Fesi precipitate in its interface.


phases are identified by microdiffraction, using a small spot size of 10nm and several

orientation conditions. The following phases are detected.

. Cubic iron silicide (see Fig. 6a). The phase is identified as Fesi, with a measured

lattice parameter of 0.42 nm. The ASTM 38-1397 file indicates 0.449 nm. An explanation for

coo-"<u'o')

'V"~~'

?

_~)f~~~~ ~i~i~

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~

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,

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i

~

~

_~~~.

100nm

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lii~ll~

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~~

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_~

.~$

-~

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~

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F*

<ycu

cr

~~'

&--~~~ -'"' S' "'

0 2 4 6 8 IO Z 4 6 8 lo

Rang- (k*Vi R»n8. ik*V)

b)

Fig. 6. TEM bright field image, electron diffraction pattem and EDX spectrum of typical precipitates

in material B : a) Fesi, b) iron, cl titanium silicide and dj B~~C.


count*ioio'i cowl* iXlo'i

i~

~[[ ~ I

'~ 100 nm

~ z 5 3 IO z IO

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metal-containing impurity precipitates (see Fig. 6d). The EDX spectra of such boron-

containing precipitates indicate that boron is neither present as borate or boron oxide, nor as

B4C (since the ratio of boron to carbon signal is too high). Microdiffraction for different crystal

orientations allows to identify the boron-containing phase as tetragonal B~SC with

a =

0.8722 nm and c =

0,508 nm (ASTM 25-96).

The B~SC precipitates are rare and far less numerous than the carbon precipitates. Since

boron in soluble in SiC, part of the boron present in the material (1 9fi after the manufacturer)

may have dissolved in the silicon carbide phase during hot-pressing, and, therefore, be below

our detection limit.

3.2 MICROSTRUCTURE OF THE ANNEALED MATERIALS. A comparison by TEM of

materials A and B before and after a high temperature anneal in the creep apparatus without

applied stress reveals absolutely no evolution of their microstructure, besides a small grain size

increase, as can be seen by the grain size distribution after annealing, figure1.

3.3 CREEP RESULTS. The creep curves of both materials comprise the typical stages shown

in figure 7. After a primary creep phase of about four to six hours, where the creep rate

decreases with time, a steady-state with constant creep rate is reached for all experimentalconditions. For very high levels of deformation and high stresses, a ternary stage can be

reached, as demonstrated in figure 7, which finally ends in sample failure.

~

tl,00EZ5~~ A B670MPa

~~D A700MPa I

i MEW( 4~ ~~ fi

. BM2MPaI,~

(I,OOEW %~~ A B670MPa

~

~ ~ A700MPa

~o 1,ooE48

o m 1000 1W0 2000 0 2 4 6

time(mn) WaIn(%)

Fig. 7. -Typical creep curves of materials A and B.

The stationary creep rates measured for materials A and B are presented in figures 8a and 8b

for stresses and temperatures ranging from 100 to 1100MPa and 1773 to 1973K,

respectively. For a given applied stress (between 100-500 MPa), the creep rate is of the same

order of magnitude, 10-9 to 10-7 s-I for both materials. However, the creep rate of

material B is larger by a factor 2.5-5 than that of material A under the same creep conditions.

The logarithmic representation of the creep rate as a function of stress yields a set of parallel

lines for both materials and all temperature conditions, for stresses up to approximately

500 MPa. It can be concluded that the creep rate follows a power law- : =

Au ~ e~ ~~~~~ with I

being the strain rate, «the applied stress, n the stress exponent, AG the activation energy for

the creep process, R the gas constant, T the temperature and A a constant. The stress exponent

n in this power law expression is about 1.5 for both materials, exactly 1.5 ± 0.1 for A and

1.5 ± 0.3 for B. As can be seen in figures 8a and 8b, the stress exponent increases at higher

stresses to n =

3.5 ± 0. I for A and n =

4.0 ± 0.5 for B, which indicates the presence of other

creep mechanisms.

A presentation of the creep rate for each applied stress as a function of the creep temperature

allows to compute the activation energy for creep. For stresses up to 500 MPa an activation


energy of 364±40kJ/mol and 453 ± 30kJ/mol has been found for materials A and B,

respectively. In the higher stress range the activation energies increase to 629 ± 30 kJ/mol.

3.4 MICROSTRUCTURE OF THE DEFORMED MATERIALS.

3.4.1 Material A. A small increase in grain size is recorded as a result of creep annealing(see Fig. I). After 25 h at 873 K, the average grain size has increased from 0.5 to 0.8 ~Lm.

Deformation in the low stress range already yields visible changes of the microstructure of

material A. TEM observations of a sample deformed at 1873 K under 300 MPa reveal an

increase of the dislocation density and formation of cavities (Fig. 9a). In sample sections

parallel to the axis of the applied stress most of these cavities are elongated along a grainboundary and therefore, yield preferential debonding of the grain boundaries oriented

perpendicularly to the applied stress. Sample sections cut perpendicular to the direction of the

applied stress show rarely elongated cavities, most of them being located at triple junctions.

D A 1773K A A 1873K ° A 1933K * A 1973K

5

5.5

6

O

~_

6.5

~l ~

~

U$~ 7

O5f

~©~

~f~

7.5

n

8°

~

n

n

8.5

9

1.6 1,8 2 2.2 2.4 2.8 2.8 3 3.2

log stress ( MPa)

a)

Fig. 8. Creep rate i>eisus applied stress fir materials A (Fig. 8a) and B (Fig. 8b) between 773 K and

1973 K. (Literature data of an industrial material with B~C precipitates [17] and [20] are added in

pointed lines.)


D B 1773K ~ B 1873K ° B 1973K ° [1711820K

© jl 7] 1923 K . [17] 2020 K x j20] 1820K x j20] 2020K

5

5.5

~/

~~/

~, ~/Q/ -6.5

/~~ .~IX /

Q /$ / o

/.I 7 /~

"

a / /) /

~f / /

~

'?'~"

/~'

~

/°

/

~8 /° ~l'

~$' ,~~~'/

x ~/ / p~

85 ,x/ /(

//°

/

9

1.6 8 2 2,2 2,4 2,6 2.8 3 3.2

log stress MPa)

b)

Fig. 8 (continued).

The dislocations are always limited to one grain, they do not cross the grain boundaries and do

not form extended networks.

TEM observations of samples deformed under 700 MPa and 871 MPa at 873 K, indicate

that, at higher stress, the dislocation density increases to a large extent in the material

(Fig. 9b). Dislocations interact with the stacking faults of SiC and with the grain boundaries. In

these conditions, the grain boundaries are partially crossed by dislocations. Again, cavitation

can be observed, but in the high stress range the cavities are smaller and are limited to triplejunctions or, more rarely, to small lense-like flaws in the grain boundary.

Thin intergranular film~ of carbon, present in rare instances as second phase in material A,

are still present after deformation, but cracks are occasionally observed to propagate inside

these films (Fig. 9c).

3.4.2 Material B. The distribution of the different phases present in the material remains

the same, except for boron. Upon annealing under stress, new spherical boron-containingprecipitates form inside the silicon carbide grains. These preciptates are identified by EDX

analysis and by microdiffraction as hexagonal B4C. The average size of these precipitates is


~K~'4

ii

/

l~

<

'~

~f'"

~

'

' j~_

+.> fi

~

d_/_' -

~ i~ 500nm ib- 500nm

-

a)b~

~

~zJ

# ~

:.L '.

'jj'$

_4w

W

~~jr /

~6~

~i.

I

j%f~~~ ~li

~ ?l500f~_j

tm-~ ~~NW

C)

Fig. 9. -Typical TEM view~ of the deformed material A, a) at low stress (1873 K, 300MPa,

F0.5 §b), b) at high stress (1873 K, 870 MPa,

F1.6 i~), c) showing the failure of the carbon

intergranular film under stress.


200nm. A strong interaction of these precipitates with the dislocations is observed, see

figure lo.

012~

. .

. . _. .

. o

B C

so0nm

Fig. 10. TEM view of B4C precipitates in material B after stress annealing.

TEM observations of samples deformed in the low stress range at 873 K (Fig. I la) show,

as for material A, a slight increase of the dislocation density compared to the annealed

material, but the dislocation distribution is very irregular, yielding small grains without anydislocations and larger ones with dislocation agglomeration. Cavitation is rarely observed

under these creep conditions.

On the contrary, samples deformed at the same temperature in the high stress show again a

very high dislocation density with dislocations spreading across multiple grains, figure I16.

The dislocation densities observed are even higher than in material A.

4. Discussion.

As a first part of this section, the above detailed observations are shortly summarized. In the

two silicon carbide ceramics processed with different levels and types of additives (« no

additives ») (material A) and«

carbon and boron»

(material B), carbon constitutes the major

impurity phase. It is mainly present as thin grain boundary films or as elongated intergranularprecipitates and has a graphitic character.

Other secondary phases result from the presence of the other sintering additives in

material B :

. boron forms micrometer-large second-phase grains of B~SC in the as-processed

material B, which transform by a dissolution and reprecipitation process ~into thermodynami-


'~°>

~' ~' ~

$~- fl/ jO

I

~ fi~'" ~,~~

Q- ~,~ '~~ '[ ~~j'~'

j

~ j ~~ i $f# V..~ .q~ .'

It~'~

~ ~

~~~,_

~.~~

I' I,

,

zjl~ ~j~ 500nm

--~--.

r..-

~

$

aJ b)

Fig. ll, -Typical TEM views of the deformed material B, al at low stress (1873 K, 200MPa,

F0.5 §b), blat high strew ii 873 K. 700 MPa,

e =2.2 ill,

cally stable intergranular B4C precipitates of smaller size during stress annealing (as observed

by Lane et al. in sintered boron-containing material [17])

. amorphous silica pockets and small impurity particles containing metallic iron, iron

silicides or titanium silicides are observed in material B. The occurrence of these secondaryphases is due to the fact that material B is prepared from standard grade silicon carbide and

additives. On the contrary, a high purity starting material is used to prepare material A, which

makes it free from such secondary phases.In the following, possible effects of the second phase particles on mechanical properties will

be discussed.

During low- and high temperature fracture, small metal or metal silicide particles can

« attract »and deflect propagating cracks and play a toughening role. On the contrary, these

small precipitates are not expected to play a significant role during high temperature creep,

since their frequency of occurrence is low and their possible interaction with dislocations, in

the case of dislocation~controlled creep mechanism, is not very efficient due to their location at

triple junctions. The small triple point precipitates cannot contribute to other deformation

mechanisms based on diffusion or grain boundary sliding and should therefore not alter the

creep behaviour of a material.

The boron carbide precipitates which form inside the silicon carbide grains upon stress

annealing of material B can play a more important role than the smaller metal alloyprecipitates, because they are located inside the grains and can therefore more easily interact

with propagating dislocations, in the case of dislocation creep. As a matter of fact, a strong


interaction of the dislocations with the B4C precipitates is observed in material B, which in

some instances even yields dislocation pinning. In the case of high precipitate concentrations,

this effect is expected to lead to a decrease of the creep rate of the material, if creep is

dislocation-controlled.

Aluminium which is currently used as a possible additive for silicon carbide should have a

much larger impact on the mechanical behaviour, since it occurs as an aluminosilicate film in

the grain boundaries, which it wets over large distances. Due to the low viscosity of the

aluminosilicate, grain boundary sliding and fast diffusion along the grain boundary films

(compared to non-wetted grain boundaries) are favoured. Furthermore, the intergranular phase

promotes void formation and crack propagation, due to preferential stress concentration. As a

consequence. in both the low- and high~stress regimes, an increase of the creep rate should be

expected. Preliminary creep studies of such materials confirm this thesis.

The influence of graphite films on mechanical propertie~ is a well-known parameter for

composite engineering. Toughening of silicon carbide fibre composites is often enhanced bytailoring the interfaces in such a way that a thin intergranular carbon film is present. Cracks

propagate preferentially along this interphase, yielding debonding of the fibres and load

transfer on the stronger fibres. In our materials, fracture cracks can similarly propagate alongelongated graphite grains or graphite grain boundary films. However, no load transfer is

possible and the effect is different fracture occurs more easily, corresponding to a lower

toughness. The role of cjrbon is not as ambiguous for high temperature creep. The well alignedintergranular graphite sheets make creep easier, if the controlling mechanism is grain boundary

sliding. The same conclusion applies if grain boundary diffusion is the controlling mechanism

for creep, since the diffusion of carbon and silicon is likely faster through the graphite phasethan along a perfect silicon carbide grain boundary. At high stresses, high strain or high

temperature, voids and cracks may form as a result of stress concentrations in the graphitephase. In summary, the presence of graphitic precipitates can only be expected to lead to an

increase of the creep rate of the material.

Based on the hypothesis presented in the previous paragraph, the presence of graphiticcarbon always results in a degradation of the creep properties. On the contrary, for a

dislocation-controlled mechanism, in the a high stress range, the boron carbide precipitates

may yield a decrease of the creep rate by interaction with the propagating dislocations. Since

carbon and B4C constitute the major impurity phases in material B, these two opposite effects

must be taken into account. The deteriorating effect of the carbon precipitates can be seen bycomparing material B to material A, the latter containing only rare graphite grain boundaryfilms. Even though the average grain size of material A is 7 times smaller than that of

material B, its creep rate is lower by a factor 2.5-4.5.

Describing creep by a power law and considering the microstructural observations, the

following can be concluded on the controlling mechanisms for the creep of materials A and B.

Low-stress regime. In the low stress range, up to about 500 MPa, both materials have a

stress exponent of 1.5 with activation energies of 353 kJ/mole and 453 kJ/mole, respectively.

A stress exponent of I would be characteristic of a diffusion mechanism, either by volume

diffusion (Nabarro-Herring creep), grain boundary diffusion (Coble creep) or grain boundarysliding compensated by diffusion. The measured activation energies in the low stress range,

either comparable or smaller than values found in the literature [lo, 17, 20], will now be

compared to activation energies reported for carbon and silicon diffusion (volume and grainboundary diffusion) in silicon carbide [21-251. Hong et al. [21, 221, Hon and Davis [231 and

Goshtagore et al. [26] measure diffusion of silicon and carbon in silicon carbide single crystalsand polycrystals at temperatures ranging from 2100 to 2 600K. They compute volume

diffusion coefficients and, in one case, grain boundary diffusion coefficients [231. The


activation energies are rather high, between 550 and 200 kJ/mol. In all these experiments,the tracer layer deposited at the surface was discontinuous and its interface with the silicon

carbide material is not necessarily perfect. These two characteristics may indeed have stronglyaffected the experimental results. Measurements by Eveno et al. [24, 25] are derived from

implanted carbon or silicon tracer profiles in the temperature range 1800-2 loo K. Theyindicate diffusion of silicon and carbon along the grain boundaries at a much slower rate than

the values reported in [21-23, 26], with as well smaller activation energies on the order of

300 kJ/mol. These activation energies are in agreement with the values measured in our creepexperiments. Since obstacles are present in the grain boundaries, which yield decohesion

during grain boundary sliding and therefore limited cavitation, we conclude that creep is

controlled by grain boundary sliding compensated by diffusion along these grain boundaries.

The creep rate is controlled by the slower diffusing species, which may be either carbon or

silicon. Literature data [23~261 are widely scattered and do not make it possible to determine

whether carbon or silicon is the fastest diffusing species.

The difference in creep rate and activation energy observed between materials A and B can

be explained by microstructural differences, namely grain boundary chemistry and structure.

Material A has mainly clean grain boundariei, while material B, which contains many carbon

precipitates, has besides the clean boundaries a large part of boundaries containing carbon or

being silicon carbide-carbon interfaces. In those carbon-containing boundaries, diffusion of

silicon and carbon should be more rapid, therefore yielding a globally higher creep rate (even

though the grain size is larger- corresponding to a smaller number of grain boundarydiffusion paths per material volume). The presence of the silicon carbide-carbon interfaces in

material B makes diffusion of carbon and silicon through these chemically and structurallydifferent boundaries possible, a process which may have a different activation energy than

simple diffusion in grain boundaries of material A.

Our creep result~ and conclusions are in agreement with those made in II 7, 201 for a material

similar to material B, though having a slightly different grain size. The creep rate values in [17]

and [20] are smaller than ours, due to the larger grain size of the material used for those studies.

The creep rate values of [171 and [201 are drawn together with our values in figure 9b.

A less good agreement is obtained with measurements by Farnsworth and Coble for

materials similar to material B an activation energy on the order of 600 kJ/mole in a similar

stress region is experimentally obtained and explained by the authors by a grain boundarydiffusion mechanism.

Within the uncertainty of the diffusion data available for silicon carbide, it can be concluded

that grain boundary sliding accommodated by grain boundary diffusion of carbon and silicon is

the dominant step for creep in the low stress range. As can be seen from the microstructural

investigations, other mechanisms also contribute to creep. An increase of the dislocation

density was observed in deformed samples, but the dislocation networks were limited to a

single grain. Piling-up of dislocations at the grain boundaries as perturbation of grain boundarysliding by grain boundary obstacles may eventually yield grain boundary decohesion, as

observed in material A (without additives) deformed in the low stress regime, or to cavitation

at weak triple junctions containing carbon precipitates (material B). Therefore, we conclude

that the stress exponent in excess of one is related to limited contribution of a cavitational

mechanism in both materials.

High-stress regime. Our results in the high stress range are quite unique and cannot be

compared to literature results since all references described above are limited to lower stresses

(maximum 400 MPa). The only work made at higher stresses is done by Djemel [12, 13], but

the investigated temperature range is much lower.


We obtained in the high stress range an activation energy of 629 kJ/mole l'or both materials

and an increased stress exponent of 3.5 or 4, which must be related to the presence of extended

dense dislocation networks and, to a smaller extent, to cavitation (both observed by TEM). We

infer from these results and our TEM observations that, in this stress range, the probability for

dislocations to cross grain boundaries increases, and that, as a result, debonding of grainboundaries and cavitation become limited. Therefore, at high stresses a mechanism must

become activated or dominant that allows this penetration of dislocations through the grainboundaries. We propose as one possibility that the partial dislocations in the basal plane can

interact with the grain boundary, propagate by climb in the grain boundary and, by their

stresses give rise to the creation of a new dislocation further away from the initial one. The

activation energy of such a mechanism would correspond to the climb energy of dislocations in

the grain boundaries. Even though the observed stress exponents of about 4 are consistent with

this proposal, it cannot be concluded whether or not the measured activation energy on the

order of 600 kJ/mole is consistent with the proposed mechanism. Other mechanisms could be

considered only detailed TEM analysis of the dislocations present after deformation would

give further insight.Nixon and Davis [201 investigated a material similar to material B and found at temperatures

higher than ours and stresses between 40 and 200MPa an important contribution from

dislocation glide in the basal plane, giving rise to a stress exponent of 1.4 and an activation

energy on the order of 876 kJ/mol. They interpret this activation energy either as beingrepresentative of volume diffusion or as the energy necessary for a dislocation to overcome the

obstacles, considering the interaction of the dislocations with the small B4C particles presentin their material. Since in our material B4C particles are much less frequent, the interpretationproposed by Nixon and Davis does not apply. We then suggest that the grain boundaries

themselves act as obstacles to dislocation glide.Literature results on compressive creep of commercial silicon carbide containing B4C

precipitates [17] and Si~SiC [201 are presented as dotted lines in figure 9b, at 1820 and

2 020 K, respectively. The former material has a grain size only slightly larger than that of

material B used in the present report. Graphite precipitates up to I ~Lm in size are also reportedin this material. Figure 9 shows that the creep rate of material A is lower and the stress

exponent very similar to that measured for material B. The difference in average grain size can

only in part account for the difference in creep rate, especially since our material was deformed

at lower temperatures. A possible explanation to this difference may be a lower concentration

of graphite precipitates or a different graphite precipitate size distribution, as compared to

material B.

The role played by grain boundary films is also obvious from figure 9. Si-SiC materials can

be considered as a silicon carbide containing a large amount of pure silicon in the grainboundaries. Compared to materials A and B, where such a continuous intergranular film was

not observed, this Si-SiC material has a higher creep rate (see Fig. 91. An intermediate

behaviour should be expected for materials with aluminium additives which show thinner but

continuous grain boundary films.

5. Conclusions.

The microstructure of various silicon carbides prepared either without additives or with boron

and carbon was investigated by TEM/EDX. The role of the second phases like graphite in

varying size and quantity of B~SC and B4C precipitates (material with boron as additive) on

the mechanical properties was evaluated. High-temperature creep experiments were conducted

on a material without additives and on material with boron and carbon as additives. Both

materials showed a similar behaviour, but the material with additives (and with a grain size 7


times larger) had a creep rate larger by a factor 2. For both materials, a stress exponent 1.5 with

activation energies of 353 and 453 kJ/mole was observed ina low stress range up to 500 MPa.

After creep, grain boundary decohesion was observed in materials without additives as well as

triple point cavitation in materiah with additives. At the same time, an increased dislocation

density was found in the grains. These observations are interpreted as being indicative of Coble

creep with contribution of a cavitation mechanism.

At higher stresses, higher dislocation densities were observed in the samples and grainboundary decohesion was no longer visible. Under high stress, the grain boundaries become

« transparent » to dislocations, and creep is now controlled by a dislocation mechanism,characterized by the increase of the stress exponent to 3.5-4.

Acknowledgments.

The authors want to thank B. Pellissier from Laboratoire de Physique des Matdriaux for

performing the creep experiments and for the development of the creep rigs, J. Deschamps(Lab. Phys. Matdriaux) for her assistance in the microscopy observations, and Cdramique et

Composites, Tarbes (France), for providing some material.

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Pre~s, 19851 pp. 255-3 lo.

[2] Cannon W. R. and Langdon T. G., J. Mat Sri. IS l1983j 1-50.

[31 Hasselmann D. P. H. and Batha H. D.. Appt. Phys. Lett. 2 j1963) it1.

[41 Frantsevich I. N., Kravets V. A. and Nazarenko K. V., Sov. Powder Metall. Met. Ceram. 14 (1975)

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Documents

Impurities in silicon carbide ceramics and their role