MECHANICAL RESPONSE AND RUPTURE MODES OF SiC/C CVD LAMELLAR COMPOSITES

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MECHANICAL RESPONSE AND RUPTURE MODES OF SiC/C CVD LAMELLAR COMPOSITES

M. Ignat, M. Nadal, C. Bernard, M. Ducarroir, F. Teyssandier

To cite this version:

M. Ignat, M. Nadal, C. Bernard, M. Ducarroir, F. Teyssandier. MECHANICAL RESPONSE AND

RUPTURE MODES OF SiC/C CVD LAMELLAR COMPOSITES. Journal de Physique Colloques,

1989, 50 (C5), pp.C5-259-C5-267. �10.1051/jphyscol:1989532�. �jpa-00229555�

MECHANICAL RESPONSE AND RUPTURE MODES O F S i C / C CVD LAMELLAR COMPOSITES

M. IGNAT, M. NADAL" , C. BERNARD, M. DUCARROIR*. and F. TEYSSANDIER*

L.T.P.C.M.-E.N.S.E.E.G., Domaine Universitaire, BP. 75, F-38400 St Martin d l H e r e s , France

*I.M.P.-CNRS, Universite Avenue de Villeaeuve, F-66025 Perpignan, France

Rbume

Des composites lamellaires SICK autoportants ont 6tt5 p&pa&s par CVD sous pression kduite.

AfindUvaluer leur comportement mecanique, des essais de flexions trois points ont ete kalises.

Cepaisseur des lamelles influe sur le comportement mhnique. Les

n5suttats

sont d i ien utilisant I'approche de Griffith ou I'effet des contraintes internes r6siduelles. Les grandeurs caracteristiques dktuites sont compa6es a quelques donn6es anterieures obterwes sur d'autres composites &rami- ques.

Abstract

SiCIC lamellar composites w r e prepared by CVD under reduced pressure. T h ~ e point bending tests were performed on samples with different lamellar thickness, in order to estimate their mechanical behavior. This parameter seems to influence the properties of such composites. The results of the deformatbn tests were analyzed in terms of Griffith's approach or residual internal stress effects. The results are compared to previous data obtained on similar composites.

The increase of the mechanical properties of materials has been obtained in the last ten years by the development of compites, and by improving their method of obtention. For metallic composite materials, unidirectional solidification methods have led to the obtention of oriented single crystalline superalloys.

Nowadays, the CVD techniques allow the preparation of fibrous or lamellar ceramickeramic materials. In the latter case, few studies have been carried out, atthough the interest for advanced materials has been emphasised by several authors /1,2/.

With metallic composites it has been clearly established that the role of interfaces is predominant. Detailed microstructural analysis showed that the controlling mechanisms fixing the transmission and the accomodation of plasticity are active at the interfaces /3,4/. In the case of similar ceramiclceramic composites, the role of interfaces is also expected to be predominant, but principally with respect to the crack opening process during deformation. In fact, the interfaces will act as barriers to transversal cracks : stopping or deviating them. In several previous works /5,6/, it has been pointed out that deviating the cracks at weak interfaces will increase the global toughness with respect to the toughness of the single ceramic phases constituting a biphased material. In this paper our interest is focused on a lamellar ceramic composite which is constituted by the akemance of pyrocarbon and silicon carbide layers. From bending tests, its mechanical response and rupture modes are analysed. This chemical pair was chosen because the differences in properties of the constituants, the eased of deposition, and the existence of previous mechanical data for structural SIC-C composites.

2

-

EXPERIMENTAL METHODS AND EXPERIMENTAL RESULTS

The deposits are cartled out in a classical cold wall reactor by tbw rate modulations of the gaseous vectors

-

Si(CH3)4 and C4HI0- under reduced pressure.

Figure 1 show the microstructural perfection of the samples otXained in our apparatus. The mean macroscopic

dimensim

of

the specimens

~ t35 lo3 e [m] long, 4.6 1 0 ~

[m]

wide and Wir thiimss was 0.34 lo3

[m]

for sample 1 and 0.42 10

-' ^[rn]

for sample 2. On both samples, a three point bending test was performed. The distance betwen the supports of the bending device was 25 [m], and the displacement of the loading zone was adapted to a comwter assisted testina machine.

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1989532

C5-260 JOURNAL DE PHYSIQUE

The load cell used during the bending runs allowed an accuracy inferior to 0.2N. The

two

tests were performed with the samt! constant central loading displacement rate of 5

lo4

[m.sec-ll. The data acquisition (force- displacement ) was carried out with a periodicity of four points per second.

The reconled curves comsponding to the tests be performed are shown in figure 2 a and b. They cornpond to samples, in which the lamellar t h m s were different For sample 1 it was 4

lo6

[m ]and for sample 2 it was 20 lo4[ m]; for both samples, the volume fraction of the C and Sic phases was 50°h.

Before the experiment, the testing device was layed down so that an instantaneous unloading was performed when the device detected a drop in stress of 0.25 N. Then the following unloading period was performed with the same displacement rate as the loading period 5.1 O~ [m.sec'l].

The sampfes w r e controlled by Scanning Ekctmn Microscopy (S.E.M) before the test and carefully oberved after. Figure 3 a and b presents the most striking features. For sample 1 (figure 3a), the only damage obsewed afterthe test was localized on the outer layer and in the central part of the sample, below the load application zone. A transversal crack was observed, and this unique crack intercepts

the

first interface producing a crack opening at the interface. These

mre

^{t k} only observations giving evidence that the sample was subm*Med to a test. For sample 2 (figure 3 b), it was surprising to observe no outer layer rupture, but on the contrary strong evidence of damage : cracks in the lamellae and at the interfaces. Debonding of lamellae and intense intralamellar transversal and longitudinal cracks, partiarlary in the Sic lamellae w r e observed. Near the outer layer of this sample, a large longitudinal crack remains after the test.Contrarily to sample 1 which showed no macroscopic permanent deformation, sample 2 presented a macroscopic curvature after the bending test.

Due to the lack of information on this type of material

,

some simple predicitve calculations were performed.

These calculations can give an idea about the potential composite behaviout; they were based on the nrle of mixtures formalism. The Young modulus and the maximum strengh were then calculated (relations [I] and [2] ) , assuming that :

-the elastic response of each phase followed Hooke's law.

-

the geometty and volume fraction of the constitutive phases are well defined.

-there is perfect bonding at the interface so that the strain in each phase ate equal when considering the parallel model or the stress in each phase are equal when considering the series model (these models correspond ^to the extreme cases when the applied force is parallel or perpendicular to the interfaces).

The equations for the Young modulus are

The equations for the maximum strength are :

-

parallel model : OR = V1 u1 R + V2 O R 2

1

-

v 1 v 2

-

series model : -

- - + -

OR

uk 4

In UES? equations

the

su#ndex i = 1 or 2 refers to

each

phase

(C

and Sic); Vi conesporrds

to the

volume frac-

TaMe 1 : Results of calculations with a rule of mbdure formalism : bulk values ate taken from references ti'/

and 181.

From

the

three point bending tests, \ h ~ deduced a flexural young modulus

(EF)

and an apparent yield stress (ay) these parameters ate presented in table II.

For the appafent yield stress (a ), we used the relation lbl :

9

FI is the maximum bading force ; D the distance k81nsen

the

supports in the bending test (25

lo4

[m]) ; b the sample width and e its thickness

.

Sic C Composite

(parallel) Composite

(series)

The Young modulus from the flexural test was given by the equation 161 : E [GPa]

250 180 21 5 209

oR .

10 3 [MPa]

2.5 18 10.25

4.39

D ~ A F

EF =

-

^Is]

4 b e 3 ~

volume fractions [ %I 100

100 50150 50/50

In which, the AFIAI term conesponds to the force- displacement gradient when the loading path is linear.

C5-262 JOURNAL DE PHYSIQUE

Table I1 : Results deduced fmm the bending tests with relations [5] and [6]. Md mam'mum loading displacement value, when the load drops instantaneously 0.25N.

4

-

DISCUSSION cornlasite

(m/=) Sample 1 ez4.10 6 [m]

Sample 2 e=20.10 [ml -6

Fmm these first ~sutts, several points may be emphasised and discussed.

Firstly, for this ceramic composite predicitve calculations based on the bulk properties of the carbon and silicon carbide constitutive phases (mixture rules formalism), do not represent the behaviour of the composite.

Results of predictive calculations are one or

two

orders of magnitude higher than the values calculated with experimental data (see ~ b l e I and 11).

Secondly, our experiments tend to suggest an increase in the flexural strength when

the

lamellae are thinner.

This behaviour has already been established for unidlrectionally solidified metallic composites 1'91 and is conelated with a Hall Petch type relation : a

d

-'I2, a is

the

strength and h the lamellar t h i i . But this is not a general behavior for ceramic composites : for example, in the TiC/TiB2 PVD lamellar composite, with a 50

% volume tractton of each phase when its lamellar thlcfm?ss decreases from 5 1 0 ~ [ m ] to 1 ~o~~II], no Hall Petch effect is detected 110 1. On the contrary, for the NirriC composite (82% volume fraction of Ni phase), when the lamellar thickness decreases from 8

lo4

[m] to 2.104[m], the strength of the material increases following a Hall Petch relation I 1 11.

In our case on the one hand, it appears that the l a m e l l a r t h i affects

the

stmnglh; but on the other hand, the lamellar thickness should also affect the deforming mechanisms. As a matter of fact even if deformed with exactly the same experimental parameters, both samples show totally different deformation curves and totally different deformation microstructures.

Fmm these remarks,stands out the third point of our discussion, which is an attempt to characterize the composite rupktre behaviour. This attempt is suppotted by Griffith's approach 11 2c when considering the results obtained with sample 1; while for sample 2 we based our analysis on the effects of residual internal stresses.

Taking into account Griffith's earlier

works /la,

and assuming that the loading and unloading paths are linear and correspond to full elastic behaviour, a strain energy release rate term can be calculated with the relation

EF [ GPa

1

65

9.9 21.2

A UeI is the elastic energy associated to the deformation cycle : a 3

~y

10 [MPa]

0.34

0.1

Md

lo1

[m]

0.25

0.27

AUeI=(FI

-

^{FJ. d}

¹⁸¹

Observations loading and unloading linear paths

Calculations of E performed w r e considering ^tw2 approximative linear paths when loading

in which d is a recoverable displacement, F1 and Fu are the maximum loading and unloading forces.

rectly deduced. They are 12 N, 7.7 N and 0.25 10-~[m] respectively.

The crack surface value is hard to evaluate directly from the SEM observations : the crack follows a transversal path on the outer layer (maximum tensile stress) then deviates at the first interface. Its total extension cornsponds approximatively to ten times t k lamellar thiiness. As the sample's span is 4.6 10-~[m], it results a crack surface of 1.84 We must insist that the value of the crack surface is extrapoled directly fmm surface observations by SEM; therefore they must be considered with care.

With the above mentioned parameters, the strain energy release rate term G is : 6.25 103 [Jouledd]. As the sample was instantaneously unloaded when a drop of 0.25N was detected, it seems reasonable to associate the strain energy release rate to a critical value for crack opening in the material. Moreover our value is of the same order of magnitude as the energy values of the work of fracture detected for CIC and SiCJSiC laminar composites : 5.3

lo3

[~ouledn?] and 3.4 l$[Jouledn?] respectively/l3/. Finally, taking into account these considerations and assuming that a plane swss condition prevails during the crack opening process, an apparent fracture toughness value can be deduced from the following relation I121 :

with E = 65 l o g [ ~ a ] (see table II) and G = 6.25 1 3 [~ouleslm~], the apparent fracture toughness of our composite (sample 1) is : 20 1 ~ ~ [ ~ a . m ~ / ~ ] .

For laminar composites reinforced by fibers with a sintered or CVD matrix the fracture toughness parameters given by S a w 1131for C/C and SiUSiC composites are respectively 11.5 [MPa mlR] and 8.46 [MPa m112]. For fibrous CISiC composites, the toughness parameter deduced by Gomina et a1 11 4/, from the analysis of three pint bending tests was 15 [MPa mlR] ; vvhich is of

the

same order of magnitude as our apparent value. On the contrary, the toughness value determined for other composites having an aluminum metal matrix and reinforced with Sic fibers and A1203 dispersed particles 119 is lower : 7 [MPa m112]-

At this point, the comparison among the different values of toughness seem to confirm the rate expected for regular distributed interfaces in this type of materials : deviate and stop crack opening.

For sample 2, this approach is unrealistic : the Fon=e/Displacement curves are not linear (loadinglunloading paths), and a permanent displacement remains at the end of the test. Furthermore an evaluation of the global rupture surface after the test is impossible. Meanwhile an attempt to explain the origin of the severe observed damage by considering the internal residual stress effects can be developed. This attempt is based on the following relation which has been often used for to estimate the order of magnitude of the thermal residual stress in a layer 116,'.

oth is then the Brmal residual stress developped in a layer having a young modulus E and a coefficient of Pois- son v, AT is the temperature variation of the sample and Aa is the linear thermal expansion mismatch hewn the considered layer and an adjacent one constituted by a different phase.

For o u r t h i i r samples the calculated residual thermal stress, in a SIC layer can reach 1.4 [GPa] (calculated with ESiC = 250 [GPa], v = 0.2, AT = 1473 [K] and

Aa

= 3 . 1 0 . ~ [K-'1).

Now considering these thermal stresses full elastjc the corresponding stored energy UeI in a lamellae of unit surface area and having a thickness A, can be calculated by the relation 1171 :

In the case of sample 2, with this relation , we

omin

for the SIC lamellae : '125 [~ou~!?s/m~].

JOURNAL DE PHYSIQUE

Compared to the above mentioned works of fracture energies for a transversal tupture in C/C and SiClSiC laminar cornposites, our value is one order of magnitude lower. Instead of that, we may point out that an interlaminar work of fracture close to our value was measured for the SiCJSiC laminar composite : 140 [~ouleslrn~] (13). Then if as observed in the case of in situ metallic lamellar composites, the lamellars faulted regions (bends, lamellar ends, kinks, ...) can act as sWss concentrators /3,4/; nearby such regions interlamellar crack opening could occur so to release the high level of residual thermal stresses.

Further deformation of the material presenting thick lamellae would then be controlled principally by the coalescence of cracks.

This last assumption is supported by our microsttuctural observations and also by the low level of work of deformation deduced from the curves corresponding to the test performed on sample 2 : approximatively 0.28

lom3

[Joules], remining that for sample 1 we obtained : 1.15 10-~[~oules].

5

-

CONCLUDING REMARKS

In spite of the restricted number of experiments, the results on this type of lamellar ceramic composites are somewhat promising.

With respect to the crack opening process, improvement is obtained when comparing the apparent fracture toughness value to values for other ceramic composites.

Our ~esuks seem to prove that thinner lamellae and numerous interfaces stop transverse crack opening. The influence of lamellar thickness and lamellar perfection appearto be crucial to the properties of the investigated composite. Even if it is still not clear whether or not the material follows a Hall Petch type relation, thicker lamellae increase the residual stress effects catastrophically.

Further investigations may clarify the mechanical behaviour in function of other parameters.

REFERENCES

I11 HOLLECK,H., J. Vac. Sci. Technol. A4(6) (1986) 266, DUBOIS 12~' IGNAT,M., METAIS,P., DUBOIS,J., BERNARD,C., DUCARROIR,M.

MRS Meeting 5-9 April (1988) Reno USA.

/3/ IGNAT,M., CAILLARD,D. MRS Proceedings

-

Vol 12 "In situ composites I V - November 1981- Boston

-

Elsevier Science Publishing (1982) 115.

141 IGNAT,M., BONNET,R. Acta Metallurgica- Vol31, 12 (1983) 1991.

IUTETELMAN, A.S., EVILY,A. J. "Fracture of Structural Materials" (1967) 16 I COOK,J., GORDON,J.E. Proc. Roy. Soc- London 282 (1964) 508.

i7l GEIR,H. "Guide pratique des maf5riaux compositesn- Ed Eyrolles (1982).

181 Mc LAREN,J.R., TAPPIN, G., DAVIDGE, R.W.

Proc.

Brit Ceram. Soc,

20

(1972) 259.

DESPANDEV, C., DOERR, J.H.

Thin solid films,

91

(1982) 215.

IIIISANS, C., DESPANDEV, C., DOERR, J.H., BUNSHAH, R.F., MOVCHAN, B.A., DEMCHISHIN, A.V.

Thin solid Films,

107

(1 983) 345.

1121 GRIFFITH, A.A

Phil. Trans. Roy. Soc. London- V.A 221 (1921) 163.

11 3/ SAKAI, M., TAKEVCHI, S., FISCHBACH,D.B, BRADT,R.C.

Proceedings fo the Conference of Ceramic Science- University of California, Berkeley- (July 1986) 1141 GOMINA, M., THEMINES,D., CHARMANT, J.L, OSTERSTOCK,F.

Internat. Journal of Fracture

34

(1987) 219.

11 51 SKINNER, A., KOCJAK, M.J, LAWLEY, A.

Met. Trans. 13A (1 982) 289.

116/ HOFFMAN, R.W.

Phys. Thin Films, 3 (1 %6) 21 1.

I171 PRESCOTT,J.J.

"Applied Elasticity"- Dever Publications- New-York (1 961) 187

Figure 1 Micrograph of the lamellar ClSiC CVD composite. The layers are perpendicular to the length of the sample.

JOURNAL DE PHYSIQUE

Force daN. 10-1

Displacement (mm) 10-2

a

Force daN. 10-2 75

Displacement (mm) 10-2

b

Figure 2 (a) and (b) : Force/Displacement curves. The dotted lines correspond to linear fitting. Curve (a) : sample 1. Curve (b) : sample 2.

micrograph (b), the small white arrow points to a wide longitudinal crack; the dotted lines schematize the sample contours.