NATIONAL INSTITUTE OF STANDARDS &

TECHNOLOGY

Research Information Center

Gaithersbiirg, MD 20699

NBSIR 88-3798

Mechanical Property Enhancement

in Ceramic Matrix Composites

S.W. Freiman, T.W. Coyle, E.R. Fuller, Jr., P.L. Swanson, D.C. Cranmer, and W. Haller

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

Institute for Materials Science and Engineering

Electrical/Electronics Ceramics Group

Ceramics Division

Gaithersburg, MD 20899

April 1988

Interim Report

7$ Yssrs Stimulating America's Progress 1913-1988

U.S. DEPARTMENT OF COMMERCE

NATIONAL BUREAU OF STANDARDS

NBSIR 88-3798

MECHANICAL PROPERTY ENHANCEMENT

IN CERAMIC MATRIX COMPOSITES

S.W. Freiman, T.W. Coyle, E.R. Fuller, Jr., P.L. Swanson, D.C. Cranmer, and W. Haller

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

Institute for Materials Science and Engineering

Electrical/Electronics Ceramics Group

Ceramics Division

Gaithersburg, MD 20899

April 1988

Interim Report

U.S. DEPARTMENT OF COMMERCE, C. William Verity, Secretary

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director

MECHANICAL PROPERTY ENHANCEMENT IN CERAMIC MATRIX COMPOSITES

S. W. Freiman, T. W. Coyle, E. R. Fuller Jr., P. L. Swanson,

D. C. Cranmer, and W. Haller

Ceramics Division

National Bureau of Standards

Gaithersburg, MD 20899

M. J. Koczak, M. Barsoum, U. V. Deshmukh, and T. R. Palamides

Department of Materials Engineering

Drexel University

Philadelphia, PA 19104

INTERIM REPORT

This program supported by the Strategic Defense Initiative

Office/Innovative Science and Technology under ONR Contract

CONTRACT NUMBER: N00014-86-F-0096

Mechanical Property Enhancement

In Ceramic Matrix Composites

Table of Contents

Page No

.

ABSTRACT 2

BACKGROUND 3

RESULTS 4

CRACK- FIBER INTERACTIONS 6

COMPOSITE PROCESSING 12

SUMMARY 17

FUTURE PLANS 18

REFERENCES 18

TABLE 1 21

FIGURES 23-27

Reprint Entitled: Fracture Mechanics Characterization of Crack/

Fiber Interactions in Ceramic Matrix Composites 39-44

Proceedings of the 1987 Northeast Regional ASM Meeting 46-55

1
ABSTRACT

The interfacial strength between AVCO SCS-6 filaments and a borosilicate glass

matrix was measured using an indentation push-in technique. Crack- filament

interactions were investigated using a fracture mechanics technique (DCDC)

.

An ultraviolet curing technique, coupled with hot pressing was developed to

prepare SiC filament-borosilicate glass composites.

2
MECHANICAL PROPERTY ENHANCEMENT IN CERAMIC MATRIX COMPOSITES

BACKGROUND

The key parameters responsible for the stress-strain behavior and damage

tolerance of fiber-reinforced ceramic composites are the fracture toughness of

the matrix, the strength of the fiber/matrix interface, and the properties and

distribution of the fibers themselves. Increases in the fracture resistance

of the matrix and the strength of the interface are predicted to lead to an

increase in the stress at which the matrix begins to microcrack. Balanced

against this increase in microcracking stress is the necessity for maintaining

some degree of fiber pullout in order to achieve the desired stress-strain

behavior in the composite. Quantitative experimental verification of the

models developed to explain composite fracture (1-5) has not yet been

obtained.

The primary objective of this program is to experimentally determine the

relationships between the fracture behavior of fiber-reinforced ceramic

composites, the properties of the individual components, and the interface

between the fiber and the matrix. A second objective is to establish

mechanical test procedures appropriate for these composite systems. The

increased understanding of the fracture process will be used to develop

strong/damage tolerant materials.

The initial composite system chosen for study in this program was a

borosilicate glass matrix reinforced with continuous SiC filaments (SCS-6,

AVCO Corp.). This system was chosen for a number of reasons including ease

of fabrication, thermal expansion match between fiber and matrix, and the

AVCO Corporation, Waltham, Massachusetts

3
handlability and reproducibility of fibers. The work to date is divided into

three parts: 1) measurement of the interfacial fiber/matrix strength, 2)

determination of the effect of fibers in retarding crack extension, and 3)

development of composite fabrication capability. Results of each of these

activities are described in the following sections.

RESULTS

Interfacial Strength Measurement

The method used to determine the strength of the fiber/matrix interface

is based on the indentation technique proposed by Marshall (6)

.

A Vickers

hardness indenter is used to apply a force, P, to a fiber parallel to the

fiber axis in a polished cross section of the composite. The method is an

extension of Marshall's technique in that the indenter is instrumented to

measure both the load applied to the fiber as well as the displacement of the

diamond tip. The load-displacement curves can be used to calculate

interfacial strength.

The experimental load-displacement curves must be corrected to account

for penetration of the indenter diamond into the matrix. During the initial

part of the loading, the diamond tip penetrates the surface of the fiber,

while at the same time, the fiber begins to slip in the matrix. To determine

the displacement, u, of the fiber surface relative to the matrix, the

penetration of the indenter, a, must be subtracted from the total

displacement. This penetration is given through Equation 1.

a = (F/2Hf)Vlan $ (1)

where F is the force applied to the fiber, Hf is the fiber hardness, and $ is

the half-angle (74°) of the Vickers indenter. Initially, all of the applied

load from the indenter acts on the fiber, F = P. As the fiber slips into the

4
matrix, the edges of the indenter penetrate the surrounding matrix.

Penetration begins at d = 2R, where d is the length of the diagonal made by

the indenter in the fiber surface, and R is the fiber radius. The portion of

the load now carried by the matrix must be subtracted from the measured load.

The calculation of the interfacial strength, r, when fiber slippage

occurs through the entire thickness of the sample, i.e., for thin samples, is

relatively straightforward if it is assumed that the applied shear stress is

constant along the interface. When r is constant, a simple force balance

yields

:

r = F/2?rRt (2)

where t is the specimen thickness. In this case, the applied load should be

constant for displacements less than that which causes the indenter to contact

the matrix.

The calculation becomes more complicated when the sample thickness is

larger than the fiber slippage length, 2 . Marshall (6) has presented an

approximate analysis for the case of t » i. Stress concentrations and

Poisson's ratio effects are ignored, and r is assumed to be independent of

position and slippage distance. If axial compressive stresses in the fiber

are neglected, then:

47T 2R

3 Ef F

[1 - 7rB(R/J!) 2

] 2 rU (3)

where Ef is the Young's modulus of the fiber and B is a constant ~ 1 (6).

O

A plot of versus u following Equation 3 is shown in Figure 1 for a

Nicalon SiC fiber in a glass ceramic (LAS-III) matrix. The slope of this

curve is proportional to r. The curve is essentially linear over the range

United Technologies Research Center, East Hartford, CT

5
0.25 /zm < u < 1.25 /zm, with a slope of 3.41 x 10^ N^/m. The fiber radius was

8.5 /im and Ef = 200 GPa. r was calculated to be 7.0 MPa, in agreement with

other measurements (6)

.

The interfacial strength between AVCO SCS-6 monofilaments and a

borosilicate glass matrix was also measured. Thin (0.16 mm) specimens were

cut from a larger billet so that the simpler analysis shown in Equation 2

could be employed as well as to limit the force required to displace the

filament. Values of the debonding (i.e., initial values) shear strength of

0.8 to 1.0 MPa were measured in this way.

CRACK- FIBER INTERACTIONS

Fracture Mechanics Procedure

This investigation made use of a test known as the Double Cleavage

Drilled Compression (DCDC) procedure. The DCDC specimen has been used to

study crack growth in brittle materials for several years (7)

.

The DCDC

specimen (Figure 2) is a rectangular column having a center hole drilled

through the thickness such that the bore axis is perpendicular to the width of

the specimen. The DCDC specimen is unique in that when the column is loaded

in axial compression, cleavage cracks initiate simultaneously from the top and

bottom of the drilled hole and grow in a stable manner as the load is

increased.

Some characteristic features which make the DCDC specimen desirable are:

A. stable centerline propagation of cracks without the aid of guide

grooves

,

B. the decaying Kj field present under remote load conditions,

C. well defined crack plane perpendicular to the direction of

propagation in a cleavage opening, and

6
D. specimen loading which eliminates detrimental torsion or bending

often present in other fracture mechanics specimens.

In all cases the stress intensity factor Kj has the general form:

KX = f(c/r)y3C 1/2 O)

where: f(c/r) = a dimensionless geometric coefficient,

c = the crack length.

r = the hole radius , and

P = the nominal applied stress.

The dimensionless geometric coefficient was experimentally determined by

plotting normalized stress versus normalized crack length. A linear

regression was performed to determine the stress intensity factor as a

function of crack length. The equation for the dimensionless geometric

coefficient used in this experiment was:

The DCDC specimen was used to investigate the effects of SiC filaments on

the stress intensity necessary to drive a crack. SiC filaments (SCS-6) were

placed at two distances from the center hole (Figure 3) . It was hypothesized

that the presence of the filaments would cause Kj to vary as a function of

crack length. If a propagating crack bows around the isolated fiber and

passes it without fiber fracture occurring, the fiber then acts as a bridge

between the two planar surfaces, inhibiting their separation and causing the

stress intensity factor to increase. In this study the bridging force

produced by the filaments was determined as a point force, and not a surface

shear stress, due to the unmeasured length of fiber pull-out between the crack

surfaces. This bridging, or traction force, is a function of the frictional

bonding between the fiber surface and the glass matrix.

f (c/r) = 0 . 138 (c/r) + 1.852 (5)

7
Assuming that the critical stress intensity, Kjc, is equal to the

fracture toughness of the glass at the tip of the crack (this is valid in the

absence of environmentally assisted slow crack growth) , a simple mathematical

relationship can be stated:

K+ "applied

+

^traction

where K

K

tip ^IC ^applied ' "'traction (6)

is the stress intensity for a given applied stress, and applied

is the stress intensity imposed by the fiber traction

By obtaining ^ from Equation (4) and knowing Kjc for borosilicate

glass (0.77 MPa^m3*5 ) from previous experiments, ^

tract ^ on

was determined.

K . can be related to a point force acting at a distance S behind the traction °

crack tip, by the following equation (8):

fraction

= - (j2)F/ ()r5)3/2 (7)

where F is a point force on the fiber.

Stress -wave fractography (9) was used to characterize the propagation of

the crack front and to make crack velocity measurements. By introducing a

periodic stress oscillation perpendicular to the direction of the maximum

principal tensile stress, optically observable undulations are produced on the

fracture surface. The spacing between the undulations gives a measure of the

crack growth rate

.

Experimental Procedure

The DCDC specimen was fabricated by the diffusional bonding of two

borosilicate*** glass plates. The specimen preparation route is illustrated

in Figure 4. Four 140-/zm-diameter SiC filaments were placed between the

plates prior to bonding. A U-shaped stainless steel bracket with notched

***CGW 7740, Corning Glass Works, Corning, NY

8
grooves cut cross-wise was used to ensure proper fiber alignment within the

matrix as well as to provide specimen reproducibility.

The SiC fibers were given an oxidation treatment (950 °C for 4 hrs .

)

prior to their incorporation in the matrix in order to minimize the presence

of bubbles occurring along the fiber/matrix interface during sintering. This

treatment is believed to promote the formation of a protective layer of Si02

on the fiber surface, which prevents deleterious reactions from occurring

during consolidation, but can also influence the interfacial strength between

the fiber and the glass.

The diffusional bonding operation was conducted in a tube furnace in air

at * 850 °C for 25-30 minutes. This fabrication route caused a thin

crystalline layer to form on all free surfaces. (No effect of this thin

crystalline layer on crack growth was noted.) After bonding, the specimens

were annealed and polished. The final preparation step was to bore a hole

through the thickness of the specimen.

In the stress -wave fractography study, a transducer was mounted to the

lower end of the DCDC specimen with vacuum grease. Resonance frequencies were

observed by monitoring the electrical output of the transducer. The signal

decided upon for the transducer was an 80-90 KHz sine wave having a maximum of

30 V peak- to-peak. That signal was gated with a burst frequency of 5 KHz

which consisted of 24% of the low frequency period. This signal gave clear

undulations in most regions along the crack plane. An illustration detailing

the arrangement of the equipment set-up and schematic of the wiring is found

in Figure 5

.

The DCDC specimens were tested in compression with a screw-driven

mechanical testing machine (Figure 6) . The crosshead was operated at a

9
constant displacement rate of 0.005 cm/min for all specimens. A double-cup

ball joint was placed between the specimen and the crosshead to minimize the

possibility of misalignment during testing. Polarized light was used to

follow the crack tip. Interference fringes in the polarized field were

enhanced with quarter-wave plates and complete documentation of the crack

route was made using 35 mm photography.

Results and Discussion

Force versus crosshead displacement curves were plotted having

incrementally numbered marks corresponding to photographs taken of the

propagating crack through the polarized field of view. The applied stress

intensity was determined from measurements of the applied load and the crack

length through the use of Equations 4 and 5

.

In Figure 7 the applied stress intensity is plotted versus the normalized

crack length. Once the crack passes the filament, indicated by the dark

vertical line, the applied stress intensity on the propagating crack increased

monotonically to a value ~ 20% higher than that for the base glass. This

increase is hypothesized to be due to the bridging effect of the filament. As

noted earlier, the bridging force due to the filament can be estimated via

Equation 7. The bridging force is produced by the frictional adhesion of the

filament to the matrix. The maximum force is reached prior to fiber pull-out

or fracture. A bar graph illustrating the variation in point force between

specimens is shown in Figure 8. Specimens C and E are not indicated in this

figure because they were misaligned during testing. The calculated force

ranges from 29 N to 63 N with specimens A and B identical at 29 N and specimen

D showing the highest point force of 63 N.

10
Figure 9 shows a photomicrograph of the fracture plane in the vicinity of

a SiC fiber. It should be noted that most fibers did not fracture at the

crack plane, indicating that some pull-out had occurred. Figure 10 gives a

closeup of the formation of a thin alpha-cristobalite crystalline layer which

formed on all free surfaces during the diffusional bonding process. This thin

crystalline layer did not appear to affect fracture in the vicinity of the

fiber

.

The interaction of the crack front with the fiber can be seen in

Figures 11 and 12. The leading edge of the crack plane appears to bow around

the fiber, leaving striations in its wake. Once past the fiber, the

propagating crack hesitates, leaving thin demarcations on the fracture

surface

.

Photomicrographs taken during the stress -wave fractography experiment are

shown in Figures 13 and 14. Figure 13 shows the onset of fracture, which can

be traced to the edge of the black region toward the center of the

photomicrograph. Figure 14 is a higher magnification of the undulations

occurring along the fracture surface from which crack velocities can be

determined. In this experiment, the velocity of the crack decreased from 6.0-

8.0 x 10 ^ m/s in the bulk of the specimen to 1. 5-3.0 x 10 m/s near the

fiber, indicating the inhibiting effect of the fiber on crack growth.

Conclusions

* The DCDC specimen was determined to be a viable specimen for the

characterization of crack/fiber interactions in ceramic matrix composites.

* The incorporation of SiC fibers in the DCDC specimen caused the

applied stress intensity to be increased by approximately 20%.

11
* Stress-wave fractography provided insight into the crack front

behavior and crack velocity. The velocities were 6 . 0-8.0 x 10 ^ m/s in the

borosilicate glass region of the specimen and 1. 5-3.0 x 10 ^ m/s in the

immediate vicinity of the fiber.

COMPOSITE PROCESSING

Most composite processing technology is proprietary; thus, some effort

was directed towards establishing a processing route to make composites in a

reproducible way. Consolidation temperature and pressure were identified as

the two most important parameters in producing fully dense composites. The

raw materials for this research were 7740 glass powder supplied by Corning

Glass Works and AVCO SCS -6 SiC monofilaments. The glass powder has a nominal

composition reported by Corning to be 81% Si02 , 2% AI 2 O3 , 13% B2 O3 , and 4%

Na20 . The SiC monofilaments are produced by a chemical vapor deposition (CVD)

process. Each filament consists of a 33 /xm diameter graphite core surrounded

by fine grains of /3 -SiC, which in turn is coated with carbon (AVCO)

(Fig. 15). The smooth graphite coating increases filament strength by

eliminating surface flaws and closing the open grain boundaries on the

surface

.

Because of the tendency of the glass powder to crystallize in a

temperature range 800-900 °C, processing temperatures were kept below 800 °C.

Above 900 °C the glass flows easily, and as a result can get squeezed into the

die clearances on application of pressure, causing die wear and making

composite removal difficult.

The composites were hot pressed in a carbon-carbon die in an argon

atmosphere. The punches were made of high strength graphite. Molybdenum

12
sheets were used in regions that came in direct contact with the glass, since

borosilicate glass does not stick to molybdenum.

The sequence of composite fabrication was as follows. The die with the

green composite in the cavity was assembled outside the hot press; then the

die assembly was placed in the hot press, aligned with the pressing rams and

held in place by the application of sufficient pressure. The hot press

chamber was then evacuated using a mechanical pump and flushed with helium at

least three times. Helium was kept flowing for the entire period of operation

to minimize die assembly oxidation. The die was heated to the required

temperature at a rate of 5 °C/min. After the die was at the required

temperature for about 15 to 30 minutes, pressure was applied. The reason for

this temperature hold is to allow the die interior to attain the surface

temperature. In the assembly that was used, only the surface temperature of

the die could be monitored. The pressure was applied in steps with

consolidation being monitored by LVDT movement which recorded the ram

displacement. Usually when the LVDT signal leveled off, it indicated maximum

possible consolidation at that temperature and pressure. Caution must be

exercised though because at high temperature the glass is less viscous, so

that under the applied pressure it may continue to flow in the die clearances.

As a result the LVDT could continue to register ram displacement. A typical

processing parameter chart is shown in Figure 16.

Initial composites were fabricated from dry powders, i.e., no wet binder

was used. Monofilaments ~ .5-m-long were cut from the spool and hand layed up

to form a 13-mm-wide tape. This tape was cut into ~ 40-mm-long sections, the

filaments being held by tape at the ends. Glass powder was poured into the

bottom of the die cavity and spread evenly; a section of filament tape was

13
placed on it. This process was repeated to build up the composite. The top

layer was glass. As this experiment was designed to optimize the processing

parameters, viz., temperature and pressure, only a few layers of filaments

were used.

Green composites fabricated as outlined above were hot pressed at 800 °C,

700 °C and 675 °C under a final pressure of 6.8 MPa. To monitor the

consolidation, the bulk density of the composite was measured using

Archimedean principles (Table 1). As can be seen, maximum density was

obtained at around 700 °C. Subsequent composites were hot pressed at a

slightly higher temperature.

The next step in the process was to increase the filament volume fraction

in the composites. The composites were manufactured exactly as above, except

that the amount of glass that was poured between the successive filament

section was considerably reduced. After assembling the die, the composites

were hot pressed at 750 °C under 6.8 MPa pressure.

There are drawbacks to this technique of composite manufacture. First,

the resulting composite tends to be inhomogeneous in terms of filament spacing

both within a row as well as between rows. The filament spacing within a row

is controlled by the accuracy of the lay-up. The spacing between rows is

controlled to an extent by the glass flow during hot pressing. Another factor

contributing to inhomogeneity comes from unequal quantities of glass between

filament layers. Secondly, the apparent volume fraction of the filaments in

the composites is reduced, since there is excess glass at the top and bottom.

In order to test these composites the excess glass at the surface must be

removed. Thirdly, making composites by this technique is fairly labor

intensive

.

14
In order to overcome these problems, it was necessary to adopt a

processing route in which the fibers are continuously coated by the matrix

powders and can be wound at a desired spacing. A typical route is shown

schematically in Figure 17. In this technique the fiber/yarn is unwound from

the spool and passed through a flame to burn off the sizing. It is then

passed through a slurry consisting of a fugitive organic binder, matrix

powder, and a few additives such as wetting, dispersing and anti- foaming

agents. The purpose of these additives is to facilitate good powder

dispersion in the slurry. The coated fiber/yarn is then wound on a mandrel to

form a tape. The tape is removed from the mandrel, cut to the size of the die

cavity, and stacked according to requirement, viz., 0 degrees or 0/90 degrees.

Next, this stack is transferred to a low temperature oven where the binder is

burned out. It is essential to burn out the binder in the presence of air to

remove all the carbonaceous matter from the matrix. It is also beneficial to

increase the temperature after the burnout so that the matrix powder sinters

slightly, which aids in handling. This "green" stack is put in the carbon

mold and hot pressed in an inert atmosphere at the required temperature and

pressure

.

It may be noted that the crucial aspect of this technique is to evenly

coat the fiber with the slurry. In the case of yarn, even though the wetting

may not be good, capillary action helps in the dispersion of matrix powder

among the individual fibers in the tow. The volume fraction of the fibers

obtained is mainly dependent on the slurry "pick-up" by the yarn. The volume

fraction can also be manipulated by changing the binder to matrix powder

ratio, or by changing the drawing speed. It should be noted that the "pickup" cannot be controlled precisely. This technique has been fairly well

15
established for the processing of composites with continuous fibers. Another

technique very similar to the one noted above was also used to manufacture

composites in this program. In this process the fiber is first wound on a

mandrel at the desired spacing to make a tape. The tape is then removed from

the mandrel and the slurry is hand painted. Further processing, viz., binder

burnout and hot pressing, is the same as above. Several composites were made

by this technique to demonstrate the process applicability.

The SCS-6 monofilaments used in this research have a carbon- rich surface

making them non-wetting to glasses. Thus, when a monofilament is drawn

through the slurry, the slurry film tends to break-up into a string of

globules along the fiber length. This surface instability results from the

effect of surface tension. The viscosity of the slurry dictates the amount of

time it takes for the slurry film to bead. However, too thick a slurry cannot

be used since it leaves a nonuniform film on the surface. Therefore the key

in processing appears to be the "hardening time", i.e, the time to dry/harden

the coating before it breaks up can be controlled by using a rapidly

evaporating solvent for the binder. Experiments were also conducted with

melt-binders such as agar-agar which were applied hot and solidified on

cooling. This technique was successful but process optimization was rather

tedious. Therefore, a set-up involving ultraviolet light curing was developed

(Figure 18). The set-up consists of a cup with a small hole at the bottom,

with a guide for centering the filament in the hole fixed on the top. The

slurry, consisting of a dispersion of the matrix powder in an organic monomer,

is poured into the cup. The slurry must be viscous to prevent it from running

down through the bottom hole. The cup can be pressurized to make the slurry

flow evenly. A filament is introduced through the top of the cup, passed

16
through the slurry, and the coated filament emerging from the bottom is

exposed to ultraviolet radiation. This instantly polymerizes the monomer,

thereby hardening the coating. The process can be operated continuously, and

the coated filament can be wound on the mandrel at a desired spacing. The

process is capable of reproducibly coating the monofilament. The amount of

matrix that can be applied to the fiber surface can be controlled in two ways.

First the ratio of glass to the monomer can be changed; second, the drawing

speed can be varied to obtain different coating thicknesses. Thus, this

process is more flexible and can be operated to obtain a range of monofilament

volume fractions.

In order to prepare composites, the coated filament tape is layed up,

sectioned to appropriate size, and stacked to obtain the composite thickness.

The binder is burned off in air, and the green composite is hot pressed.

Composites with SiC monofilament and borosilicate glass matrix have been

successfully manufactured by this route. It is important to note that the

same monomer can be used for any matrix. Thus, the technique can be used to

manufacture almost any monofilament reinforced composite as long as the matrix

is available as particulates.

SUMMARY

The three primary areas of research in this program are measurement of

interfacial strengths, determination of crack-fiber interactions, and

fabrication development for ceramic matrix composites. Significant progress

has been achieved in all three areas.

1. An instrumented hardness machine has been used to measure the

interfacial strengths between AVCO filaments and borosilicate glass. Values

of the debonding strength of 0.8 to 1.0 MPa were determined.

17
2. The double cleavage drilled compression (DCDC) specimen was used to

study crack-fiber interactions in this same SiC-borosilicate glass system. A

fracture mechanics analysis demonstrated the role of the bridging filament in

retarding crack propagation. The analysis also showed the quantitative

relationship of the bridging force to the interfacial strength between

filament and matrix.

3. An ultraviolet, polymer curing technique coupled with a hot pressing

procedure was developed to prepare SiC monofilament-borosilicate glass

composites

.

FUTURE PLANS

A major effort during the coming year will be to determine effects of

fiber surface modifications on both the interfacial strength as well as the

crack fiber interactions in the SiC-borosilicate glass system. Various

technique for surface modification will be employed, including coatings,

oxidation, etc. Effects of thermal expansion differences between fiber and

matrix will also be determined by varying the glass matrix composition. A

portion of our effort during the coming year will be to develop tensile test

procedures for these composites so that testing of the actual composites can

be performed at both ambient and elevated temperatures. Nondestructive test

techniques for detecting the onset of microcracking in these composites will

be investigated.

REFERENCES

1. J. Aveston, G. A. Cooper, and A. Kelly, The Properties of Fibre

Composites , pp. 15-26, National Physical Laboratory, Science and

Technology Press Ltd.

,

London, (1971)

2. J. Aveston and A. Kelly, J. Mat. Sci., 8 , 352, (1973)

18
3. D. B. Marshall and A. G. Evans, J. Am. Ceram. Soc., 68., 225 (1985)

4. D. B. Marshall, B. N. Cox, and A. G. Evans, Acta Metall.; 33., 2013 (1985)

5. B. Budiansky, J. W. Hutchinson, and A. G. Evans, J. Mech. Phys . Solids,

34, 167 (1986)

6. D. B. Marshall, Comm. Am. Ceram. Soc., 67, C-259 (1985)

7. C. Janssen, Tenth International Congress on Glass, 23, Kyoto, Japan (1974)

8. H. Tada, P. C. Paris, and G. R. Irwin, The Stress Analysis of Cracks

Handbook . Del Research Corp . , Hellertown, PA (1973)

9. H. Richter, Proc. of Eleventh International Congress on Glass, Vol. 2,

447, Prague, Czechoslovakia (1977)

19
Table 1

-

.

20
Table 1. Bulk Density and Open Porosity in SCS-6/Borosilicate

Glass Composites

HOT PRESSING

TEMPERATURE<°c) BULK DENSITY

(g/cc)

OPEN POROSITY

(%)

800 2.151 0.118

700 2.219 0.266

675 2.092 1.677

21
Figures
Force

Squared

(N‘2)

Figure 1. Plot of the square of the force applied to a fiber versus the

displacement of the end of the fiber relative to the matrix surface

during an instrumented indenter test. The fiber diameter is

17.0 /im.

23
H/W a 10

2 H/r - 75.0

2 W/r - 7.5

Figure 2. Geometry of a Double Cleavage Drilled Compression (DCDC)

Specimen.

24
°-A

t I I

!

1

i ;

4-9

L

o

O 11 1 -

4

m £. YY

I I t

0\

Figure 3. Loaded DCDC specimen with a crack emanating from

the top and bottom of the hole.

25
SPECIMEN PREPARATION

7740 Borosilicate Glass

Diffusional Bonding Polished

6 mm

/

Final Dimensions

Figure 4. Illustration of the fabrication sequence for the DCDC

specimen.

26
Figure 5. Schematic of Stress-Wave Fractorgraphy Set-up.

27
Figure 6. Mechanical testing operation for the DCDC specimen.

28
Applied

K

T

(MPaVm)

Applied K versus Normalized Crack Length

fiber

Sample A

Sample B

Sample D

Sample F

Figure 7. The change in the applied stress intensity with respect to the

position of the filament in the DCDC specimen.

29
Maximum Fiber Traction for Each Specimen

80 n

Specimen

Figure 8. The point force (N) variation from specimen to specimen that i

imposed by the fiber at maximum traction as calculated from

equation 7

.

30
Figure 10. SEM photomicrograph of the alpha-cristobalite layer which

forms during fabrication and surrounds the fiber.

31
Figure 11. Crack front bowing upon impingement with fiber.

(Nomarski polarizing filter in reflected light)

Figure 12. Crack front bowing upon impingement with fiber,

(transmitted light)

32
Figure 13. Stress -Wave Photomicrograph, Black region is the center hole.

Dark-grey region is the Wallner phenomenon and the vertical

undulations follow. Product of a 5 Hz low frequency 80 KHz high

frequency gated signal.

Figure 14. High magnification of the undulation produced with a 5 Hz low

frequency, 80 KHz high frequency gated signal along the free edge

of the fractured surface.

33
CL

Figure 15 Schematic of a section of AVCO SCS-6 fiber showing various layers and the surface composition.

34
TEMPERATURE

* >

E

35

Figure

16

Typical

processing

parameter

chart.

Outputs

from

the

various

sensing

devices

such

as

thermocouple

and

LVDT

are

in

millivolts.

Ram

displacement

is

measured

by

the

LVDT.
FIBER BURNER

SPOOL

SLURRY

WINDING MANDREL

0 DEGREES

(UNIDIRECTIONAL)

:r t - 1

HOT PRESS

Figure 17 Schematic of a composite making process,

are stacked and hot pressed.

After winding the tapes

36
Figure 18 Schematic of polymer coating/UV curing system to coat SCS-6

filaments with glass slurry.

37
Reprint Entitled

Fracture Mechanics Characterization

Crack/Fiber Interactions in Ceramic Matrix Composites

38
Reprinted from (he Ceramic Engineering and Science Proceedings, Vol. 8, No 7 -8. July-Augusi 1987 Copvrighi 1987 by The American Ceramic Socieiy, Inc Fracture Mechanics Characterization of Crack/Fiber

Interactions in Ceramic Matrix Composites

T. W. Coyle, E. R. Fuller, Jr., and P. Swanson

Ceramics Div., Inst. Materials Sci. and Eng.

National Bureau of Standards, Gaithersburg. MD 20899

T. Palamides

Drexel Univ.

Philadelphia, PA

A crucial factor in the structural performance of ceramic matrix composites is the

influence of the fiber/matrix bond on the interaction of a matrix crack with the rein- forcing fibers. To elucidate the character of this interaction under controlled fracture conditions, glass fracture mechanics specimens were fabricated in the doublecleavage, drilled compression (DCDC) configuration with simple arrays offibers.

Propagating cracks were observed in cross-polarized illumination to characterize

delamination of the fiber ahead of the crack and bridging interactions behind the

crack tip. Stress wavefractography was employed to analyze the shape and relative

velocity of the crack front.

Introduction

A critical issue in the fracture behavior and mechanical performance of ceramic

matrix composites is the influence of the strength of the fiber/matrix interface on the interaction of a crack with individual fibers. It is generally accepted that

if a strong bond is present between fiber and matrix, cracks run directly through

the reinforcing fibers which then have little effect on retarding the propagation

of the crack. Although the strength of such a composite may be enhanced through

load transfer, the toughness is not appreciably improved and failure is completely

brittle. A weak interface leads to crack deflection, debonding along the fiber/matrix

interface, and bridging of the crack by the fibers. If the fibers do not break behind

the crack tip failure is non-catastrophic. However, if the interface is too weak the

fibers slide freely through the matrix and no improvement in strength or toughness

is obtained. Thus, a better understanding of the details of this behavior is required

before the optimum levels of interfacial strength can be determined for a given

fiber-matrix system.

This issue can be investigated experimentally by direct observation of the in- teraction of a crack with a single fiber in an appropriate fracture mechanics

specimen. The double cleavage drilled compression (DCDC) configuration 1 was

employed because of the simple shape and loading geometry. No guiding side

grooves are required to keep the plane of the crack in the center of this compressively

loaded specimen.

Stable crack propagation is obtained in the DCDC specimen since the applied

stress intensity (Kapp) experienced at the crack tip decreases with crack extension

at a constant load. From the measured crack length (c) at a given remotely applied

compressive stress (aapp), Kapp can be calculated from the following expression: 1

K(ipp = aapp'TT /[(c/r)/(c/r)] (1)

39
where the specimen dimensions are as defined in Fig. 1. The function /(c/V) can

be determined experimentally by testing specimens of known Kjc

Experimental Procedure

Transparent specimens were produced by sandwiching SiC fibers* between

borosilicate glass plates 1 which were then fused together by heating in air at ap- proximately 850 °C for 10-15 min to produce transparent specimens. The samples were loaded in compression with a screw-driven mechanical testing machine* at a crosshead displacement rate of 0.005 cm/min. The strain contours visible in cross- polarized illumination were used to monitor the length of the growing crack and

to characterize the crack-fiber interactions. Stress wave fractography 2 was utilized to monitor the shape of the crack front

and the relative velocity of the growing crack along the specimen. Acoustic waves

were propagated through the specimen during crack extension by use of a piezoceramic transducer coupled to the specimen. The transducer was driven with

bursts of an 80-90 kHz sine wave at a maximum of 30 V peak-to-peak. The burst

frequency was 5 Hz; the burst duration was 25% of the low frequency period.

This technique produce ripples on the fracture surface at a spacing determined by

the frequency of the acoustic excitation. These ripples can then be observed by

optical microscopy in reflected light using Nomarsky contrast. With a constant

excitation frequency, widely spaced bands indicate a region of rapid crack extension while closely spaced bands indicate relatively slow crack growth.

Results and Discussion

Direct, qualitative observations of several aspects of the fracture behavior in the specimen could be made in transmitted, cross-polarized illumination at various

magnifications. Typically the fiber could be seen to bridge the crack for some

distance behind the crack tip before breaking. Evidence of localized strain at the

fiber-crack surface intersection (Fig. 2) indicated that the bridging fiber was ap- plying tractions across the crack faces. The shape of the strain contours ahead of

the crack tip or behind the crack tip away from the crack surfaces gave only occasional faint indications of localized strain which would show load transfer to the

fiber. Apparent delamination of the fiber ahead of the crack (Fig. 3) was observed

when a higher magnification was used for one specimen.

An optical micrograph of a region of the fracture surface which includes a

fiber, taken with Nomarsky contrast, is shown in Fig. 4. The light and dark bands

visible under these conditions represent snapshots of the position and shape of the

crack front at periodic intervals. It can be seen that the crack front bowed around

the fiber which was bridging the crack before the center portion broke away and

accelerated to reform a nearly straight front. A scanning electron micrograph of

this area (Fig. 5) shows that after bowing around the fiber the two sides of the

crack did not meet on the same plane. The resulting ligament would have to have

fractured before the center portion of the crack front could extend. From the direct

observations discussed above, the fiber appears to have broken at some later time,

after the ligament had fractured and the crack front had moved well beyond the

fiber. The chipping apparent around the fiber in Fig. 5 would have occurred as

the broken fiber was pulled out of the glass matrix.

SCS-6 Fibers. Avco Corp . Wilmington. MA.

Corning 7740. Corning Glass Works. Corning. NY.

Instron Corp.. Canton. MA.

40
To quantitatively assess the influence of the fibers on the crack propagation

the crack length was measured from photographs, such as shown in Fig. 2. and

used to calculate the apparent Kapp from the applied load using Eq. ( 1 ). The results for several specimens are shown in Fig. 6. Assuming that the crack is growing

at a constant stress intensity, K,ip = Kic for the glass matrix, the increase in the

calculated Kapp for crack lengths beyond the fiber position reflects the closure trac- tions exerted by the bridging fiber (K,rac).

K;ip = K,c = Kapp = Ktrac ( -

)

The stress intensity factor due to a point force acting at a distance 6 behind the

crack tip is given by: 3

where the point force, F, may be a function of 6. Thus, information regarding

the forces exerted by the bridging fiber as a function of crack length can be obtained from data such as shown in Fig. 6. Summary

Sample fabrication methods and experimental techniques have been developed

which allow crack-fiber interactions to be directly examined in a ceramic matrix.

Qualitative information regarding fiber/matrix debonding, fiber bridging, and pinning and bowing of the crack front at the fiber can be obtained. Possibilities for

quantifying the effects of bridging fibers on crack extension in terms of the trac- tions exerted across the crack by the fiber were demonstrated.

Acknowledgments

Support for this work has been provided by the Department of Energy AR&TD

Fossil Energy Materials Program (DE-AI05-800R20679) and by the SDIO/IST program under ONR contract No. N00014-86-F-0096.

References

'C. Janssen. "Specimen for Fracture Mechanics Studies on Glass;" pp. 23-30 in Proceedings of

the Xth International Congress on Glass. Kyoto. Japan. July 8-13. 1974. TheCeranuc Socien of Japan.

Kyoto. Japan.

: H. Richter. "Crack Propagation in Glass Under Liquids in an Intermediate Range of Crack

Velocities;" pp. 219-29 in Strength of Inorganic Glass. Ed. by C. R. Kurkjian. Plenum Publishing

Corp.. 1985.

J H. Tada. P. C. Paris, and G. R. Irwin. The Stress Analysis of Cracks Handbook. Del Research

Corp.. Hellertown. PA.. 1973.

or

( 3 )

Krrac =-sfTF/(TW/2 (4)

41
Fig. 1. Double-cleavage drilled-compression fracture

mechanics specimen. Compressive load opens cracks above

and below centered hole.

Fig 2. View perpendicular to fiber axes of DCDC specimen

under load The cracks are seen edge-on with the crack tips

located between the inner and outer fibers Strain contours are

visible by virtue of the transmitted cross-polarized illumination.

42
Fiq 3 Higher magnification views in cross-polarized transmitted light of region near reinforcing fiber Crack-tip is seen approaching and then passing the fiber as the remotely applied load is increased. Field of view is approximately 2.7

mm wide

Fig 4 Optical micrograph of a region of the fracture surface centered about

a reinforcing fiber Nomarsky contrast reveals the periodic surface undulations

caused by the applied acoustic waves The crack was propagating from right

to left The field of view ib approximately 1 mm wide

43
Fig. 5. Scanning electron micrograph of area shown in Fig. 4 The remnant

of the ligament formed as the crack front swept around the fiber can be seen

to the left of the fiber. (Field of view approximately 1.1 mm wide )

Crack Length (C/R)

Fig 6. Plot of remotely applied K (calculated from Eq. (1)) vs crack length

showing relative position of the reinforcing fiber. Increase in the applied K re- quired to propagate the crack reflects the closure forces exerted by the bridging

fiber.

44
Proceedings of the 1987 Northeast

Regional ASM Meeting

45
CRACK- FIBER INTERACTIONS IN CERAMIC MATRIX COMPOSITES

T. W . Coyle, T. R. Palamides* , S. W. Freiman,

E. R. Fuller and U. V. Deshmukh*

Ceramics Division

Institute for Materials Science and Engineering

National Bureau of Standards

Gaithersburg, MD

*Drexel University

Philadelphia, PA

Abstract

Crack- fiber interactions are the key to the unique properties of

continuous fiber reinforced ceramic matrix composites. This paper describes

the development of a fracture mechanics technique for examining crack- fiber

interactions in model SiC fiber reinforced glass specimens. Direct

examination of the closure force imposed by a bridging fiber on a crack was

accomplished using the double -cleavage drilled-compression sample. The

stress intensity at the crack tip is described in terms of the applied load

and the closure force. The closure force is a function of the interfacial

bonding and/or frictional forces between the fiber and the matrix.

Evaluation of these interfacial forces by indentation experiments in which

the indentation load and the displacement of the fiber are measured directly

is described.

46
Introduction

The fracture behavior of brittle matrix/brittle fiber composite systems

is strongly dependent on the mechanical performance of the fiber/matrix

interface. For a given set of fiber and matrix properties, the fracture

characteristics can range from the desired fibrous, non-catastrophic failure

to a completely brittle, low toughness fracture, depending on the interfacial

properties. The first condition which must be met to avoid a completely

brittle failure mode is for the interface to be weak enough to allow

debonding to occur between the fiber and matrix ahead of the crack tip or to

cause the matrix crack to deflect along the interface. Fracture mechanics

treatments of cracks at bi-material interfaces 1,2 provide some guidance

regarding the relative strength levels required among matrix, fiber, and

interface to achieve the desired interfacial failure but direct experimental

studies are lacking.

When the interface has debonded, the fiber can remain bridging between

the crack faces after the matrix crack has advanced. Propagation of the

matrix crack is then opposed by the closure forces exerted on the crack by

the bridging fibers. The nature of these closure forces is controlled by the

shear properties of the interface and the fiber modulus and strength. Recent

modelling of this situation by Marshall, Evans and coworkers3,4 indicates

that the bridging fibers can either remain intact behind the crack tip

leading to non-catastrophic failure or can break at some distance behind the

crack tip resulting in a failure which is primarily brittle in nature,

although with a higher toughness than the unreinforced matrix. Experimental

results suggest that the trends predicted by the model with changes in

interfacial character are correct, but detailed studies over the range of

possible behavior are not yet available.

Direct experimental examination of the interactions between matrix

cracks and reinforcing fibers is therefore valuable in furthering the

understanding of the influence of interfacial properties on fracture

behavior. Develop of techniques to investigate these issues are described in

the following along with observations on a model ceramic matrix system, SiC

fiber reinforced glass. The double cleavage drilled compression (DCDC)

configuration5 was used to study the effects of individual SiC fibers on

crack propagation in glass specimens. Qualitative observations of the shape

of the crack front and the fracture path in the vicinity of the fiber were

made and some quantitative information regarding the closure forces exerted

by bridging fiber was obtained. The fiber/matrix interfacial shear

properties were evaluated independently by measuring the force required to

push the fibers through thin sections of the glass matrix using an

indentation technique. The results of these tests are discussed in terms of

the current views of fracture in ceramic matrix composites.

Experimental Procedure

Transparent DCDC specimens were produced by sandwiching SiC fibers*

between borosilicate glass plates# which were then fused together by heating

in air at approximately 850°C for 20 to 30 minutes. The samples were ground

to size using standard diamond machining techniques. The four sides were

then polished with the final surface finish produced by 0.05/zm Al2 03 powder.

The dimensions, as defined in Fig. 1, were nominally t = w = 6mm and

SCS-6 Fibers, Avco Corp., Wilmington, MA.

47
h = 60mm, with a hole radius of 0 . 8mm. The fibers were placed symmetrically

about the central hole with the inner fibers located approximately 9mm from

the hole. The specimens were loaded in compression using a screw-driven

mechanical testing machine* at a cross -head displacement rate of 0.05 mm/min.

Figure 1

- Compressive

Double -cleavage drilled-compression fracture mechanics

load opens cracks above and below centered hole.

specimen.

Transmitted cross -polarized illumination was used perpendicular to the

fiber axis to produce visible strain contours in the specimens. The contours

enabled the uniformity of the applied loading to be evaluated, the crack

length to be determined, and local strain concentrations associated with the

crack- fiber interactions to be detected.

The shape of the crack front and the relative velocity of the growing

crack along the specimen were monitored by stress wave fractography

.

6 Acoustic waves were propagated through the specimen during crack extension by

use of a piezoceramic transducer coupled to the specimen through silicon

grease. The acoustic waves produce ripples on the fracture surface which can

be observed by optical microscopy in reflected light using differential

interference contrast. Each "ripple" marks the position of the crack front

at periodic intervals set by the excitation frequency of the transducer,

therefore the spacing of the ripples is a measure of the local crack

velocity. The transducer was driven with bursts of an 80-90 KHz sine wave at

a maximum amplitude of 30 V peak to peak. The burst frequency was 5 Hz; the

burst duration was 25% of the low frequency period.

To characterize the mechanical behavior of the fiber/matrix interface in

shear, the fibers were pushed through thin sections of the matrix by applying

an axial load on the fiber using a Vicker's hardness machine. The hardness

tester was instrumented to allow the applied load and the displacement of the

diamond indenter to be measured continuously during the testing cycle.

Samples were prepared from the broken halves of the DCDC specimens by diamond

grinding them to a thickness of approximately 0.85mm perpendicular to the

fiber axis and then polishing.

48
Results and Discussion

A view of a DCDC specimen under load in cross -polarized transmitted

light which illustrates the type of qualitative observations that could be

made is shown in Fig. 2. The strain contours are symmetric about the center

line of the specimen indicating that the specimen is well aligned. The crack

emanating from the central hole is seen edge-on with the crack tip located at

the center of the hour-glass shaped pattern formed by the strain contours.

At this point in the test the crack tip has already moved past the first

fiber and is now between the inner and outer fibers. A change in contrast

indicating a region of localized strain could be observed in the area where

the fiber intersects the crack, demonstrating that the fiber was bridging the

crack and exerting a closure force across the crack faces. This feature

disappeared following further crack growth, suggesting that the fiber broke

at some distance behind the crack tip.

Figure 2 - View perpendicular to fiber axes of DCDC specimen under load.

Strain contours are visible by virtue of the transmitted cross -polarized

illumination.

An optical micrograph of a fracture surface in the region about the

inner fiber, taken with differential interference contrast, is shown in Fig.

3. The fiber axis is perpendicular to the plane of the photograph. The

light and dark bands visible under these imaging conditions represent

"snapshots" of the position and shape of the crack front at periodic

intervals. With a constant excitation frequency applied to the piezoceramic

transducer, closely spaced bands indicate slow crack growth, while widely

spaced bands reflect regions of more rapid crack propagation. It can be seen

that the crack front bowed around the fiber which was bridging the crack

before the center portion broke away and accelerated to reform a more nearly

straight front. A scanning electron micrograph of the same area (Fig. 4)

shows that after bowing around the fiber the two sides of the crack did not

meet on the same plane. Fracture of the ligament which formed between the

two crack planes was necessary before a continuous crack- front could reform.

The scallop -shaped chips around the fiber occurred during fiber pullout.

From the indications of bridging which could still be observed in transmitted

polarized illumination after the crack had propagated veil beyond the fiber,

49
Figure 3 - Optical micrograph of a region of the fracture surface centered

about a reinforcing fiber. Differential interference contrast reveals the

periodic surface undulations caused by the applied acoustic waves.

Figure 4 - Scanning electron micrograph of area shown in Fig. 6. The remnant

of the glass ligament formed as the crack swept around the fiber can be seen

to the left of the fiber.

50
The quantitative information obtained from the tests is based on stress

intensity factor analyses of the DCDC specimen5

' 9 which indicate that for a wide range of crack lengths the stress intensity factor (Kapp ) due to the

remotely applied compressive load (a app ) is well described as a linearly

decreasing function of crack length (c)

.

The expression

aa_p-p_ fr

Kapp

( 1 )

where r is the radius of the center hole in the DCDC specimen, was used to

calculate Kapp from the applied load and crack lengths measured from

photographs such as shown in Fig. 2. The constants a and were evaluated

empirically by testing specimens of known KIC . In Fig. 5 the calculated Kapp , with a=0.151 and /?=2.158, is plotted versus c/r for several specimens.

Assuming that the crack is growing at a constant stress intensity, Ktip = KIC

for the glass matrix (0.77 MPam1/2

) 5, the increase in the calculated Kapp for

crack lengths beyond the fiber position reflects the contribution to the

total stress intensity of the closure force exerted by the bridging fiber

or

^tip -KIC - Ka pp

+ Kfib

Kapp = KIC Kfib

( 2 )

(3)

Crack Length (C/R)

Figure 5 - Plot of applied K due to remotely applied compressive load versus

crack length showing relative position of the reinforcing fiber.

The measured crack lengths used to construct Fig. 5 are through

thickness averages of the crack length along the crack front. In

unreinforced specimens such measurements are adequate since the crack front

is nearly straight. However, as seen in Fig. 3 the crack front is curved in

the region near the fiber, therefore analysis of these results should

incorporate this feature. A£ r me that the crack front assumes a profile such

51
curvature must arise from the form of the expression for Kf ib . The stress

intensity factor along a straight crack front due to a point closure force,

F, acting at a distance, x, behind the crack tip is given by10

:

K

I

72 F

(tt x) 3/2

f

2 } x

2

X + z

2 (4)

where z describes the position along the crack front (Fig. 6). Substitution

of Eq. 4 for Kfib , with F taken as the bridging force due to the fiber, and

Eq. 1 for Kapp in Eq. 2 results in an expression which describes contours in

the x-z plane (the fracture surface) where Ktip = KIC for given values of

aapp and F >

or- +

a

Or-r

72 7x F

3/2 . 2

(x + z (5)

where a is the distance from the edge of the center hole to the fiber. It is

difficult at this time to estimate the error involved in applying Eq. 4 to a

curved crack front. Near the fiber, where the crack front curvature is

large, Eq. 4 may overestimate Kfib by as much as a factor of two or three.

The approximation of F as a point force also becomes inadequate within a few

fiber radii of the fiber location. Farther from the fiber, where the change

in x becomes small with respect to z along the contours of constant Kti , the

error should be modest. With an appropriate value of F, it should therefore

be possible to account for the observed crack lengths and crack front

profiles as a function of the applied stress.

Figure 6 - Schematic illustration of the bowed crack- front profile showing

the x-z coordinate system defined on the fracture surface with origin located

at the position of the closure force, F.
The closure force is transferred from the fiber to the matrix through

‘shear tractions acting across the fiber/matrix interface. Therefore the

closure force exerted across the crack surfaces is controlled by the

interfacial properties. To independently evaluate the response of the

interface to applied shear stresses an instrumented hardness machine was used

to push the fibers through thin sections fabricated from the broken halves of

the DCDC specimens. Measured load-displacement curves for two fibers are

shown in Fig. 7. The load-displacement records for these tests typically

showed a maximum in the applied load followed by an abrupt drop to a plateau

load nearly constant over several micrometers of displacement. The increase

in load beyond the plateau results from the indenter making contact with the

matrix surrounding the fiber. The customary interpretation of this type of

result is that the peak applied load corresponds to the applied shear stress

required for debonding (r db ) and the plateau load to the stress for

frictional sliding (r f ). Observed values of r db ranged from 0.8 to 1.0 MPa

and of r

f from 0.4 to 0.7 MPa. An estimate of the magnitude of F is obtained

by multiplying r by the interfacial area, i.e. F = (tt D w/2) r.

Displacement (microns)

Figure 7 - Curves showing axial load applied to the end of a fiber by the

Vicker's indenter versus the displacement of the indenter. The abrupt load

drop occurs as the fiber begins to slide through the glass matrix.

Contours of Ktip = KIC given by Eq . 5 with F = 0 . 7N for several levels

of applied stress are shown in Fig. 8, where the point force is located at

the origin. The shape of the calculated constant Ktip contours have the same

general shape as the crack front profiles shown in Fig. 3, suggesting that

the character of the bridging interaction is fairly well described by Eq . 4.

The spatial extent of the calculated interaction is also in fair agreement

with that observed, indicating that the me aitude of the trial bridging force

is of the correct order. However, in the region near the fiber, where the

interactions are strongest, the calculated contours do not describe the

experimental observations very well. This may be due to the assumption of

53
From the results of the indentation experiments it is clear that the

'force exerted by the fiber varies with the relative displacement of the fiber

and the matrix. The bridging force will therefore depend on the debonded

length and the crack opening displacement (u) . In principle the dependence

on crack opening displacement could be introduced into Eq . 4 by taking

F = f(x). For example, when a simple frictional sliding resistance

(r f = constant) can be used to describe the resistance to shear of the

interface, F increases with Ju (i.e. Foc(u1 1 2

) ) 3 * 4

. Assuming uoc(x 1/2

) would

then give Fa(x1/A

). Substitution of such an expression for F in Eq . 6 may

yield better agreement with the observed crack front contours. However,

several other factors influence the shape of the contours in the region near

the fiber. The importance of debonding relative to frictional slippage, the

debonded or slippage length, and the errors involved in applying Eq. 4 to

this situation can also affect the agreement between the calculated and

observed contours. The work now in progress is concentrated on assessing the

relative importance of these factors.

X (mm)

Figure 8 - Contours of constant Ktip on the x-z plane (the fracture surface)

calculated using Eq. 5 for several values of aapp with F = 0.7N. Each

contour, from left to right, corresponds to a larger <7 app . The point force

(the fiber) is located in the lower left corner (x = 0, z = 0) . The area

represented would extend from the center line of the DCDC specimen at z = 0

to the specimen edge at z = 3.0.

54
References

1. T. S. Cook and F. Erdogan, "Stresses in Bonded Materials with a Crack

Perpendicular to the Interface", Int . J . Engng . Sci

.

. 10 (1972) 677.

2. J. R. Rice and G. C. Sih, "Plane Problems of Cracks in Dissimilar Media",

J , Appl . Mech

.

. 32 (1965) 418.

3. D. B. Marshall, B. N. Cox, and A. G. Evans, "The Mechanics of Matrix

Cracking in Brittle -Matrix fiber Composites", Acta Metall

,

. 33 (1985)

2013-2021.

4. D. B. Marshall and A. G. Evans, "Tensile Failure of Brittle -Matrix Fiber

Composites", Proceedings of the 5th International Conference on Composite

Materials. ICCM-V . ed. W. C. Harrigan, J. Strife, and A. K. Dhingra

(Warrendale, PA: The Metallurgical Society, 1985), 557-568.

5. C. Janssen, "Specimen for Fracture Mechanics Studies on Glass",

Proceedings of the Xth International Congress on Glass (Kyoto, Japan: The

Ceramic Society of Japan, 1974), 23-30.

6. H. Richter, "Crack Propagation in Glass Under Liquids in an Intermediate

Range of Crack Velocities", Strength of Inorganic Glass , ed. C. R.

Kurkj ian (Plenum Publishing Corp . , 1985), 219-229.

7. J. H. Henry, "Energie d'Avancement de Fissure, G, pour une Plaque Percee

d'un Trou Soumise en Compression' , Mechanics Research Communications . 10

(5) (1983) 253-257.

8. C. G. Sammis and M. F. Ashby, "The Failure of Brittle Porous Solids Under

Compressive Stress States", Act Metall

.

. 34 (1986) 511-526.

9. W. E. Warren, "Theoretical Analysis of the Double Cleavage Drilled

Compression Specimen", Int. Journal of Fracture . 33 (1987) 223-325.

10. H. Tada, P. C. Paris, and G. R. Irwin, The Stress Analysis of Cracks

Handbook (Hellertown, PA: Del Research Corp., 1973).

55
SECURITY CLASSIFICATION OF THIS PACE fWhan Data Entered)

REPORT DOCUMENTATION PAGE READ INSTRUCTIONS

BEFORE COMPLETING FORM

1 . report number 2. GOVT ACCESSION NO. 3. RECIPIENT’S CATALOG NLMBER

4. TITLE (end Subtitle)

Mechanical Property Enhancement in Ceramic Matrix

Composites

5. TYPE OF REPORT & PERIOD COVERED

Interim Report

*

6. PERFORMING ORG. REPORT NUMBER

7. AUTHOR/#;

S. W. Freiman, T. W. Coyle, E. R. Fuller Jr.,

P. L. Swanson, D. C. Cranmer, and W. Haller

8. CONTRACT OR GRANT NUMBER/ a)

N00014-86-F-0096

9. PERFORMING ORGANIZATION NAME AND ADORESS

National Bureau of Standards, Gaithersburg, MD

Drexel University, Philadelphia, PA

10. PROGRAM ELEMENT. PRCJEC”. TASK

AREA 4 WORK UNIT NUM3ERS

11. CONTROLLING OFFICE NAME AND ADORESS

SDI0/I3T

Washington, D.C.

12. REPORT DATE

April 30, 1988

13. NUMBER OF PAGES

55

14. MONITORING AGENCY name & AOORESS('// different from Controlling Office)

Office of Naval Research

800 N. Quincy Street

Arlington, VA 22217

IS. SECURITY CLASS, (of UUa report)

1S«. DECLASSIFICATION/ DOWNGRADING

SCHEDULE

16. DISTRIBUTION STATEMENT (of this Report)

17.

DISTRIBUTION STATEMENT (of the atatract entered In Block 20, li different from Report)

18.

SUPPLEMENTARY NOTES

19.

KEY WORDS (Continue on reverse aide If necaeaery and Identify by block number)

ceramics; composites; fibers; fracture; interfacial strength

20.

ABSTRACT (Continue on reverse aide It neceeeery end Identify by block number)

The interfacial strength between AVCO SCS— 6 filaments and a borosilicate glass

matrix was measured using an indentation push— in technique. Crack-filament

interactions were investigated using a fracture mechanics technique (DCDC). An

ultraviolet curing technique, coupled with hot pressing was developed to prepare

SiC f ilament-borosilicate glass composites.

DD J j AN ^73 1473 EDITION OF 1 NOV 65 is OBSOLETE

S/N 01 02- LF -014-6601 SECURITY CLASSIFICATION OF THIS PACE (*1ten Data Entered)
NBS-114A (REV. 2-80

U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET (See instructions)

1. PUBLICATION OR

REPORT NO.

NBSIR 88-3798

2. Performing Organ. Report No. 3. Publication Date

APRIL 1988

4. TITLE AND SUBTITLE

Mechanical Property Enhancement in Ceramic Matrix Composites.

5. AUTHOR(S) S. W. Freiman, T. W. Coyle, E. R. Fuller Jr., P. L. Swanson, D. C. Cranmer,

and W. Haller

6. PERFORMING ORGANIZATION (If joint or other than NBS. see instructions)

NATIONAL BUREAU OF STANDARDS

U.S. DEPARTMENT OF COMMERCE

GAITHERSBURG, MD 20899

7. Contract/Grant No.

S. Type of Report & Period Covered

9.

SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City . Stote, ZIP)

10.

SUPPLEMENTARY NOTES

j Document describes a computer program; SF-185, FIPS Software Summary, is attached.

11.

ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant

bi bl iography or literature survey, mention it here)

The interfacial strength between AVCO SCS-6 filaments and a borosilicate glass matrix

was measured using an indentation push-in technique. Crack-filament interactions

were investigated using a fracture mechanics technique (DCDC) . An ultraviolet curing

technique, coupled with hot pressing was developed to prepare SiC f ilament-borosilicate

glass composites.

12.

KEY WORDS (Six to twelve entries ; alphabetical order; capitalize only proper names; and separate key words by semicolon s)

ceramics; composites; fibers; fracture; interfacial strength

13. AVAILABILITY

(j£] Unlimited

[ 1 For Official Distribution. Do Not Release to NTIS

; Order From Superintendent of Documents, U.S. Government Printing Office, Washington, D.C.

20402.

[XU Order From National Technical Information Service (NTIS), Springfield, VA. 22161

14. NO. OF

PRINTED PAGES

60

15. Price

$13.95

USCOMM-DC 6043-P80



IR 88-3799

CANCELLED 03/16/90

UNAVAILABLE FOR BINDING