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Impurities in silicon carbide ceramics and their role during 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 role during high temperature creep. Journal de Physique III, EDP Sciences, 1993, 3 (12), pp.2189-2210.
�10.1051/jp3:1993269�. �jpa-00249076�
Classification Physic-s Abstracts
20.790 61.16D 61.70W
Impurities in silicon carbide ceramics and their role during high 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 pr6cipit6s
de 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 respectivement 364 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 investigated
materials 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 conductivity
and 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 plane together 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 creep behaviour 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-point bending 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 grain
sizes [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,
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 Auger microprobe and electron energy loss spectroscopy. They identify carbon precipitates, micrometer-large B4C grains and, after annealing, small additional coherent B4C precipitates inside the grains. No enrichment of boron in the grain boundaries or formation of intergranular
films 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 stacking
faults 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 (typical
strain 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 by lattice diffusion. The authors explain that dislocation glide is the dominant mechanism at high
strain 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 part
be 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-negligible discrepancies 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 pure
form 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
2192 JOURNAL DE PHYSIQUE III N° 12
(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
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 particularly performant 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 grains
consist 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 high
resolution image, figure 2d, which shows the well aligned basal planes of carbon within a ramified structure.
ioo
80 G
r a
60
n u m b e r
20
O
o-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
2194 JOURNAL DE PHYSIQUE III N° 12
@' OO02
~~~ ~
' ~
,/"~
~ ~
~ C ~
' .
'
j
~~~~~
~-w~
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b)
coun~« lx10'j
~°~~~* ~~'°'~
4
graphite carbon grain of SiC
6 3
s, 5
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z i
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~' 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 (about 3.5 ~Lm see grain size distribution in Fig. 3) is larger than that of material A. The material
consists of a 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 m b e r
O
1,2 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 slightly
different 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.
