NBSIR 88-3710

Structural Reliability and Damage

Tolerance of Ceramic Composites

for High-Temperature Applications

Semi-Annual Progress Report for

the Period Ending September 30,

1987

E. R. Fuller, Jr., T. W. Coyle, and R. F. Krause, Jr.

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

Ceramics Division

Gaithersburg, MD 20899

October 1987

Issued February 1988

Prepared for

U.S. Department of Energy

Fossil Energy

Advanced Research & Technology Development

Fossil Energy Material Program

Oak Ridge National Laboratory

Oak Ridge, TN

NBSIR 88-3710

STRUCTURAL RELIABILITY AND DAMAGE

TOLERANCE OF CERAMIC COMPOSITES

FOR HIGH-TEMPERATURE APPLICATIONS

SEMI-ANNUAL PROGRESS REPORT FOR

THE PERIOD ENDING SEPTEMBER 305 1987

E. R. Fuller, Jr., T. W. Coyle, and R. F. Krause, Jr.

U.S. DEPARTMENT OF COMMERCE

National Bureau of Standards

Ceramics Division

Gaithersburg, MD 20899

October 1987

Issued February 1988

Prepared for

U.S. Department of Energy

Fossil Energy

Advanced Research & Technology Development

Fossil Energy Material Program

Oak Ridge National Laboratory

Oak Ridge, TN

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

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director

STRUCTURAL RELIABILITY AND DAMAGE TOLERANCE OF

CERAMIC COMPOSITES FOR HIGH -TEMPERATURE APPLICATIONS

Semi-Annual Progress Report for the Period Ending September 30, 1987

Prepared for US Department of Energy, Fossil Energy

Advanced Research & Technology Development

Fossil Energy Material Program

Oak Ridge National Laboratory

Oak Ridge, TN

Interagency Agreement DE-AI05-800R206179

E. R. Fuller, Jr., T. W. Coyle, and R. F. Krause, Jr.

National Bureau of Standards

Institute for Materials Science and Engineering

Ceramics Division

Gaithersburg, MD 20899

INTRODUCTION

The achievement of higher efficiency heat engines and heat recovery

systems requires the availability of high- temperature , high-performance

structural materials. Structural ceramics, and more recently, ceramic

matrix composites have received particular attention for these

applications due to their high strength, and resistance to corrosion and

thermal shock. Even with these positive attributes, improved reliability

and extended lifetime under service conditions are necessary for

structural ceramics to gain industrial acceptance. The problems with

these materials are mechanical and chemical in nature and are enhanced by

the fact that they are subjected to high temperatures, reactive

environments and extreme thermal gradients.

With an objective of improved performance for heat engine/heat

recovery applications, the NBS program on structural ceramics and ceramic

composites addresses these problems through the determination of critical

factors that influence mechanical and microstructural behavior. The

activities of the program are grouped under two major subtasks, each

designed to develop key data, associated test methods and companion

predictive models. The status of the subtasks are detailed in the

following sections.
-
DISCUSSION OF CURRENT ACTIVITIES

Subtask A: Interfacial and Microstructural Composite Properties

Measurement of Fiber/Matrix Interfacial Forces

Summary . An instrumented hardness machine was employed to

characterize the mechanical behavior of the fiber/raatrix interface in two

types of composite specimens. Using the technique, described in past

reports [1], several specimens of a Nicalon fiber-reinforced SiC

composites provided by R.A. Lowden of Oak Ridge National Laboratory were

examined. In addition, a variation of this technique was used to provide

an independent analysis of the fiber/matrix interface of DCDC specimens,

whose development and testing were reported in the previous semi-annual

report [1]. Due to the much larger diameter of the AVCO SIC fibers used

in the DCDC specimens, changes in the experimental technique and data

analysis procedure were required from those previously developed for

Nicalon fibers. Due to configurational changes in the displacement

sensors of the instrumented hardness machine, modifications to the data

analysis procedure were also necessary for specimens with Nicalon fibers.

Interfacial Forces in DCDC specimens . Discussions in previous

progress reports on measurements of the fiber/matrix interfacial forces

with the Instrumented hardness machine dealt with testing of bulk

composites containing small diameter (10-15,um) fibers. In these

instances the specimen thickness was much larger than the distance along

the fiber over which slippage occurred. The fibers used in the DCDC

specimens are much larger in diameter (140/im). Since the force required

to cause fiber slippage increases with the interfacial area between fiber

and matrix, the specimen thickness must be reduced in order to limit the

applied force required to an accessible experimental range. To simplify

the analysis, the limiting condition of specimen thickness less than the

distance, over which slippage occurs, was adopted. In this case the

shear strength of the interface (r) is simply r=F/(27rRt) , where F is the

axial force applied to the fiber, R is the fiber radius and t is the

specimen thickness.

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. Load-displacement

curves for two fibers are shown in Fig. 1. The tests typically show a maximum in the applied load followed by an abrupt drop to a plateau load

which is nearly constant over several micrometers of displacement. 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 ).

Following this simple analysis the observed values of r db range from 0.8

to 1.0 MPa and those of r

f from 0.4 to 0.7 MPa.

Nicalon fiber - SiC matrix composites . The materials examined were

2-D reinforced composites produced at Oak Ridge National Laboratory by a chemical vapor infiltration (CVI) process. The goal of the testing is to

relate the strength of the interface to the various interfacial

structures produced during the CVI fabrication process. The requisite

flat, optically polished surface had been prepared on samples

approximately 10mm thick. Indentations were placed in several individual

fibers for each specimen and the load and total indenter displacement

recorded.

The total displacement measured during the test includes not only

the fiber displacement but the displacement due to the penetration of the

diamond indenter into the fiber and displacement due to the compliance of

the sample support and stage as well. Recent changes in the

configuration of the displacement sensor have made the latter

contributions much larger and therefore the simple correction based on

the fiber hardness (see [1]) is no longer adequate. A new procedure was

developed to experimentally determine that part of the total displacement

which was not due to the displacement of the fiber relative to the

matrix. Indentations were placed in fibers laying parallel to the sample

surface and therefore unable to move relative to the matrix. The

load-displacement curve determined for these indentations can be

subtracted from the curves recorded during indentation of fibers whose
'
axis is perpendicular to the surface, thus correcting for both the

compliance of the sample support and penetration of the diamond into the

fiber. When this analysis is completed the results will be correlated

with those of other characterization techniques to be performed at Oak

Ridge National Laboratory.

Characterization of Matrix Crack - Fiber Interactions

Summary . The double cleavage drilled compression (DCDC)

configuration has been used to study the effects of individual SiC fibers

on crack propagation in glass specimens as previously reported [1].

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 in

representative specimens were recently evaluated by measuring the force

required to push the fibers through thin sections of the glass matrix

using the indentation technique discussed above. These results have been

included in further efforts to quantify the effects of the fibers on

crack propagation in these specimens.

Analysis and discussion . An optical micrograph of the fracture

surface of a DCDC specimen in the region about a fiber is shown in

Fig. 2. The light and dark bands visible in the micrograph arise from

the periodic undulations on the fracture surface by virtue of the

differential interference contrast. These bands represent the shape and

position of the crack front at periodic intervals of time [1]. Thus it

can be seen that as the crack propagated through the specimen, it bowed

around the fiber, then again formed a continuous front as it moved beyond

the fiber, leaving the fiber bridging the crack.

The quantitative information obtained from the tests is based on

stress intensity factor analyses of the DCDC specimen [2-6] which

indicate that for a wide range of crack lengths the stress intensity

factor (Kapp ) due to the remotely applied compressive load (<7 app ) as a function of crack length (c) is well described the expression:

Fibers C3, C4

Figure 1. 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.

Figure 2. 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

.
.
?«PPJL ( 1 ) K*pp = a — +

r

where r is the radius of the center hole in the DCDC specimen. The

constants a and y3 can be evaluated empirically by testing specimens of

known KIC . The stress intensity at the crack tip is assumed to be the critical

stress intensity for the glass matrix (KIC ) and is given by the sum of

the remotely applied loading and the closure forces exerted by the

bridging fiber (Kfib ):

( 2 )

(3)

As seen in Fig. 2 the crack front is highly curved in the region near the

fiber rendering detailed, rigorous analysis of these results extremely

difficult. The following development does, however, provide some useful

insight on factors affecting crack growth. Assume that the crack front

develops a profile (Fig. 3) such that Eq. 2 is satisfied at all points

along the front. The expression for Kapp depends only on the crack

length, c, or on x, the distance beyond the fiber, and so cannot account

for the curvature. Thus the curvature must arise from the form of the

expression for Kfib . 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 by [7]:

Kt i p Kic Kapp + Kfib

or

KOPP KIC Kfib

K

I Ji f

(tt x) 3/2

r

v

2 x

2 x +• 2

(4)

where z describes the position along the crack front. Substitution of

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

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

contours in the x-z plane (the fracture surface) where IC i = KIC for

given values of cr app and F:
-
f Jl F

'v o app

r*

K . = — tip x

Or-r

(5)

Or— + Cr~ +

a

r

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 significantly overestimate Kfib . 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 IC. ip , the error should be modest. With an appropriate value of

F, it should therefore be possible to rationalize the observed crack

front profiles and to account for the range of interaction between the

fiber and crack front.

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, the instrumented

hardness machine was used to push fibers through thin sections fabricated

from the broken halves of the DCDC specimens. Measured loaddisplacement curves for two fibers are shown in Fig. 1, as described

above. An estimate of the magnitude of F is obtained by multiplying r by

the interfacial area, i.e. F = (n D w/2) r. Contours of Ktip = KIC given by Eq. 5 with F = 0 . 7N for several

levels of applied stress are shown in Fig. 4, where the point force is

located at the origin. The shape of the calculated, cons tant-K*. L contours have the same general shape as the crack front profiles shown in

Fig. 2, 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 magnitude of the trial bridging force is of the correct order.

However, in the region near the fiber, where the interactions are
-

.
Figure 3. 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.

x (bb)

0.0 0 5 1.0 1.5 2.0

Figure 4. Contours of constant K4 i

D

on the x-z plane (the fracture

surface) calculated using Eq . 5 for several values of a app with F = 0.7N.

Each contour, from left to right, corresponds to a larger a 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

„
‘
strongest, the calculated contours do not describe the experimental

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

F = constant used in this calculation.

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., F cc u

1/2 ). Assuming

u <x x

1

1

2 would then give F a x

1 1

4

. Substitution of such an expression

for F in Eq. 5 may yield better agreement with the observed crack front

contours. However, several other factors could influence the shape of

the contours in the region near the fiber as well. It is felt that the

errors involved in applying expressions such as Eq . 4 to this situation

preclude rigorous quantitative evaluation of the bridging force from the

present results. Additional analysis, probably including finite element

treatments of the system would be required. The qualitative

understanding obtained from such experiments is valuable none- the- less

,

particularly in conjunction with independent evaluation of the

fiber/matrix interfacial properties.

Sub task B: High-Temperature Damage Modes and Failure Mechanisms

Fracture Toughness Behavior

Summary. The fracture toughness behavior of a 25 wt percent silicon

carbide whisker-reinforced alumina ceramic has been further characterized

at room temperature to include a material with a porosity of 0.6 percent,

thus supplementing previous measurements [1] at porosities of 4.9 and

11.5 percent. The composite materials were provided as billets through

the courtesy of J. F. Rhodes of then ARCO Chemical Company [8]. The

intrinsic, flexural strength was determined, and noted to increase as a function of decreasing porosity. A varying flaw size, controlled by

Vickers - indentation loading between 2 and 200 N, was placed on the

previously polished (9 pm finish) , prospective tensile surfaces of fourpoint flexure specimens. The flexural strength was measured as a function of indentation load, and the results indicate that the fracture

resistance of these materials increases as a function of crack extension.

Furthermore, the fracture resistance is influenced by the bulk mass

density. An increase in bulk mass density (a decrease in porosity) of

the ceramic raises the level of fracture resistance, but has negligible

effect on the susceptibility for the fracture resistance to increase with

crack extension (i.e„, rising R-curve behavior).

Intrinsic strength . The open circles which are arbitrarily plotted

at log(P) = 0 in Fig. 5, correspond to the strength, S 0 , for non- indented

specimens. The other open circles in Fig. 5 show strengths of indented

specimens, which did not rupture at the indentation flaw, also indicative

of the intrinsic strength. Table 1 gives the mean S Q indicated by all

the open circles. The results reported earlier [1] for other porosities

are included for comparison. Figure 6 shows S Q as a function of the

porosity, p. This function may be characterized by

S Q = SB (l-p) f

» (6)

where S m and tj are constants. As the porosity approaches zero, which

corresponds to a theoretical dense material, the intrinsic strength

approaches its maximum value, S m

. A least squares fit, the solid line in

Fig. 6, gives S m = 694 ± 8 MPa and rj = 4.12 ± 0.14, where the

uncertainties are standard deviations.

Indented beam strength . Figure 5 also shows the bending strength,

S, as a function of indentation load, P. This straight line function,

which is the logarithm of

S = S 0 (P/Po )-'5 (7)

was fitted by a method of least squares to evaluate the constants: a =

S Q P 0

/3 and P at each porosity. Table 1 gives the values of p and ?Q thus
:

' v’.
1000

CO

03

o

Indentation Load, P in N

2 5 10 20 50 100 200

iog (P)

03 CL

cn

.c

•*-> 03

C

<©

*-» CO

Figure 5. Flexural strength, S, as a function of indentation load, P,

for a 25 wt percent silicon carbide whisker-reinforced alumina ceramic at

0.6 percent porosity. The solid line is a least-squares -fit for the

solid points.

Table 1. Values of parameters that define crack-extension in a 25 wt

percent silicon carbide whisker-reinforced alumina ceramic at selected

porosities. The uncertainty for each value is its standard deviation.

Parameter Billet (a) Billet (b) Billet (c)

A. Porositv

p, percent 11.48 + 0.16 4.89 + 0.14 0.60 + 0.03

No. of measurements 24 24 24

3. Intrinsic Strength

S Q , MPa 421 i 11 559 12 681 it 15

No. of measurements 11 6 6

C. Indented Beam Strength

£ 0.276 i 0.015 0.266 + 0.009 0.262 it 0.014

N 8.5 ± 1.3 4.7 i 0.5 4.9 + 0.7

No. of measurements 13 18 18

Percent <loge (S)>

Student's t for 95 %

+ 1.7 i 1.3 + 2.0

confidence interval 1.80 1.75 1.75

D. Indented Crack Length

c 0 , /xm 28.9 + 2.4 18.1 + 1.2 18.8 ± 1.6

No. of measurements

Percent loge <c I >

6

i 1.5

6

+ 1.3

6

± 2.3

E

.

Fracture Resistance

m 0.068 + 0.018 • 0.080 + 0.012 0.084 ± 0.018

K0 , MPa •

m

1 ' 2 5.17 + 0.25 5.35 0.17 6.61 ± 0.32

cT /cx 2.66 0.04 2.69 0.03 2.70 ± 0.04
•-

'

*
determined at each of the selected porosities. Only those indentation

loads which correspond to the fractured surface having passed through a radial crack of an indentation flaw are considered valid points for the

fit. The dashed line is a line of -1/3 slope, which a material with a

fracture toughness that is invariant with crack size would indicate.

Fracture resistance . To facilitate analysis of the observed

indentation-strength data, a parametric representation of the fracture

resistance, KR , as a fractional power function in crack extension, Ac, is

used [9,10], namely,

KR ^ K0 (Ac/c0 ) m (8)

where K0 , cQ , and m are constants. The exponent m measures the

susceptibility to R- curve behavior, that is, the susceptibility for

fracture resistance to increase with crack extension. R-curve behavior

arises because additional energy is required besides that needed at the

crack tip to propagate the crack. The additional energy is required to

overcome other energy consuming processes, such as tractions of aggregate

interlock and/or pullout, or restraining forces of ligaments in the wake

of the crack. In general, Ac = c - c

p

, where c is the total crack length

and c

p is a traction-free portion, such as exists in a prenotched

specimen.

The lengths of surface cracks produced by indentation loading were

measured by optical microscopy. Only values in the large crack region

were used to evaluate the constant c Q of Eq . (8). When indentation

equilibrium is attained [10], the constant c 0 may be characterized by

c 0 = Cl (P0 /P) (1 + ^2 (9)

where c r is the half-length of surface crack. Figure 7 shows c r observed

as a function of P. Notice the bias in the small crack region of

indentation loads due to the failure to observe complete crack length by

optical microscopy. Table 1 gives the mean value of c 0 determined at

each of the selected porosities.
i

-
Percent Porosity, 100 p

12 10 86420

700

600

500

400

<o a,2

c

o

cn

•—o cnc

©

ka CO

iog (1-p)

Figure 6. Intrinsic, flexural strength, S 0 , as a function of porosity

for a 25 wt percent silicon carbide whisker-reinforced alumina ceramic.

The solid line is a least-squares-fit

.

Indentation Load, P in N

2 5 10 20 50 100 200 500

500

200

100

50

20

10

cnc

©

o

03

~ O

Figure 7. Half-length of surface crack, c = c r

, as a function of

indents uin load, P , for a 25 wt percent silicon carbide whiskerreinforced alumina ceramic at 0.6 percent porosity. The dashed line is

least-squares-fit , and the slope of the solid line is derived from the

strength measurements.

Solution of the fracture mechanics equations [10] enables an

evaluation of the constants K0 and m of Eq. 8. The indentation flaw is

considered to have extended to an onset of instability due to

superposition of the stress-intensity factor for flexural loading on the

residual stress intensity factor of indentation loading. The onset of

crack-extension instability, cT , is related to the indentation flaw size,

cx . b7

cT /Cl = [4/(l-2m)

]

2/ (3+2m) (10)

When the solution is expressed in the form of Eq. (7), the constant /? is

defined by

m = (1-3/?)/ (2+2/3) (11)

and the constant S Q is defined by

K0 = YSt,c 0

1 ' 2 (l+£)‘ 1 +

<’>/,S'J (12)

The reasonable fit of Eq. (7) to the observed indentation- strength data

does not invalidate the supposition of Eq. (8). Table 1 gives the values

of m and KQ deduced from Eqs. (11) and (12), respectively, using Y = 1.174 [9], Figure 8 shows the fracture resistance, KR , as a function of

Eq . (8). The shaded region at each porosity in Fig. 8 describes the 95

percent confidence bands of the data and the range of experimental values

of the critical stress-intensity factor, Kc , at the onset of crackextension instability, cT = Ac. Decrease in the porosity of the

composite raises the level of fracture resistance, but has negligible

effect on the susceptibility for fracture resistance to increase with

crack extension, which indicates that tractions in the wake of the crack

are hardly affected by porosity.
•-

.
re CL2

cr *

reo

c

re

**> as

03

re QC

re

zs

*-o

rew

U.

Figure 8. Fracture resistance as a function of crack extension of a 25

wt percent silicon carbide whisker-reinforced alumina ceramic at selected

porosities. The shaded regions denote the range for the onset of crackextension instabilities deduced in the experiments, and the dashed lines

describe the 95 percent confidence intervals of KR .

Temperature, T,° C

Figure 9. Flexural strength versus temperature for a 25 wt percent

silicon carbide whisker-reinforced alumina ceramic of 4.9 percent

porosity

.
.

'

'
Creep and Crerg Rupture Behavior

Summary . The intrinsic strength of a 25 wt percent silicon carbide

whisker-reinforced alumina ceramic with 4.9 percent porosity was

determined by rapid loading of four-point flexure specimens at several

temperatures between 25 and 1300 C. While the strength remained

approximately invariant up to 900 C, it regressed above this temperature,

« becoming almost half as strong at 1300 C.

The creep strain and creep rupture of this material was measured in

four-point flexure as a function of flexural stress (corresponding to

applied stresses between 98 and 374 MPa) and temperature (from 1000 to

1300 C) . At 1000 C a specimen displayed approximately 0.2 percent strain

without rupture, following a period of 1230 h, during which it was under

a stress of 300 MPa. At 1120 C another specimen displayed approximately

0.5 percent strain without rupture, following a period of 1590 h, during

which it was under a stress of 149 MPa. Creep strain to failure ranged

between 0.1 and 1.6 percent, decreasing with increasing stress at a given

temperature and increasing with increasing temperature at a given stress.

The time required to rupture a specimen ranged from 0.3 to 865 h.

High-temperature strength . Flexure specimens, 3 mmx4mmx 50 mm,

were cut and ground the same as those used in the fracture toughness

study [1] except these did not receive the 9 ,um finishing polish. Using

nitrocellulose cement, they were mounted in place on silicon carbide

platens with fixed 3 mm diam loading bars. Constant loading on a specimen was applied through a pneumatic cylinder with a lubricated

rubber diaphragm. A controlled- leak, pressure regulator was capable of

maintaining the applied load within ± 2 N. The load was transmitted

through 19 mm diam rods of 99 percent recrystallized silicon carbide,

bearing on a hemi-spherical compression tool of the same material. Three

sets of loading rods protruded through the roof and floor of an aluminainsulated furnace cavity, 150 mm x 200 mm x 150 mm high, to accomodate

three test specimens simultaneously . Two such furnace assemblies were

used in these tests. A pre-load of approximately 10 N maintained the

allignment of a specimen while it was heated to the test temperature,

during which the nitrocellulose cement burned away. The furnace cavity
-

'
was heated to the test temperature in approximately one hour by

molybdenum disilicide heating elements, radiating from the side walls

within the furnace cavity. The temperature was controlled and measured

by calibrated platinum-rhodium thermocouples. The loading rods were

water-cooled where they were connected to air cylinders and load

transducers, which were supported in the same frame above and below the

furnace, respectively. The load transducers were calibrated with known

weights and were read through a chart recorder. After the specimens were

held at an elevated temperature for approximately one hour, the loads

were increased at a uniformly rapid rate to rupture the specimens in

succession by increasing the pressure in the air cylinders. Rubber shock

absorbers prevented an interaction amongst the loading rams.

The flexural strength, S, was characterized by

S = 3F(i 1 -i2 )/2bw2 (13)

where F is the rupture load; = 40 mm and i

2 = 10 mm are the outer and

inner loading spans, respectively; and b = 4 mm and w = 3 mm are the

nominal dimensions of a specimen. Figure 9 shows the results of the

strength tests as a function of temperature.

Green strain and rupture time . The same procedure described above

for the high- temperature strength tests were used to conduct creep

experiments except that only two specimens were placed in a furnace

cavity and the stress applied at the elevated temperature was something

less than approximately half the material strength. Also, the loadingpoint displacement, y, was measured as a function of time, t. The creep

strain, e, was characterized by [11]

e = 6 yw/(i 1 2 +f 1 i 2 -2i 2 2

) (14)

for which the neutral axis of bending was assumed to be at the geometric

center of the specimen and the stress exponent was assumed as unity.

The loading-point displacement during creep was monitored by a linear voltage displacement transducer that was attached to each watercooled holder of an upper ram of the test equipment. The transducers
_

'

"
were calibrated with known gauge blocks, and the displacement was measured within ±5 ^m. To enhance the long-term accuracy of the displacement

readings, a block of silicon carbide, the same as that used for the rams,

was used in the middle test station of a furnace. Any displacement measured on the block is believed to reflect changes in the test facility

due mainly to fluctuations in the ambient temperature. These displacement fluctuations were substracted from the displacements measured on the

specimens in the other two stations of the same furnace so that errors

due to differential thermal expansion of the test facility were minimized. The failure of one specimen in a furnace was observed to have negligible effect on the other specimen, mainly because the pneumatic pistons and rubber shock absorbers were used to restrain the propagation of

stress-waves

.

Two examples of typical curves of creep strain versus time are shown

in Fig. 10. They exhibit a work-hardening of the material right up to

the time of rupture. Table 2 gives a summary of the experiments

conducted to date. Strain valves with a greater- than-symbol refer to

specimens that had not as yet ruptured up to the time indicated. Creep

strain was hardly perceptible at 1000 C, but the considerable creep

strain at 1300 C indicates a specimen to specimen variation by at least a

factor of two. The apparent work-hardening characteristic of creep

strain was represented by

£ = Bta (15)

where a and B are constants. Leas t- squares fits of the logarithm of Eq.

(15) are shown in Fig. 11. The time exponent, a, appears to remain

fairly uniform over the range of stresses shown at a given temperature.

Future creep studies . The apparent work-hardening characteristic of

creep strain may be in part a manifestation of a shift in the neutral

axis. We plan to examine this phenomena by comparing strains calculated

from loading point displacement with those calculated from the radius of

curvature. Also, an examination of the distribution of strain along the

specimen height would indicate any possible shift in the neutral axis.

The observed creep strain data will be analysed to relate strain rate as
.
0.5

0 50 100 150 200

Time, t f hour

Time, t, hour

Figure 10. Typical curves of creep strain versus time for flexure

specimens of a 25 wt percent silicon carbide whisker-reinforced alumina

ceramic of 4.9 percent porosity under the given applied stress at

selected temperatures. The solid lines are least-squares fits.

Table 2. Creep rupture behavior of a 25 wt percent silicon carbide whiskerreinforced alumina ceramic, having 4.9 wt percent porosity.

Temperature Applied Stress Percent Strain Failure Time

°C MPa at Failure h

r 202® >0.04 > 360

301 >0.23 >1230

1000

249® Ao o > 360

374 0.01 87

r 126 >0.32 >1590

149 >0.52 >1590

1120 < 200 0.53 . 865

s. 250 vT> o m 149

124 >0.45 > 290

149 1.04 61

1200 < 200 0.73 7.5

< 249 0.15 0.3

r 98 1.61 14

1300 -

W

124 0.87 16

a This is the preliminary stress on a specimen. Following a period of 360 h

the load on the same specimen was increased to that shown on the subsequent

line

.

Time, t, hour

10 30 100 300 1000

Time, t, hour

0.1 0.3 1 3 10 30 100

log (t)

Figure li. Straight line, least-squares fits of the logarithm of

tq. (15), showing creep strain versus time for flexure specimens of a -5 wt percent silicon carbide whisker-reinforced alumina ceramic of 4.9 percent porosity under given applied stresses of constant loading ar selected temperatures.

Percent

Strain,

100 €

functions of applied stress, temperature, and time to failure,

tion of the fracture resistance at elevated temperatures would

the extent of crack-extension to be expected up to the time of

for a specimen under a given stress.

De terminaindicate

failure

REFERENCES

1. E.R. Fuller, Jr., T.W. Coyle, R.F. Krause, Jr., and T.J. Chuang,

"Structural Reliability and Damage Tolerance of Ceramic Composites for

High-Temperature Applications," Semiannual Progress Report ORNL/FMP87/1, Fossil Energy Materials Program . R.R. Judkins, Manager, Oak

Ridge National Laboratory, Oak Ridge, TN, July 1987, pp . 115-140.

2. 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, pp . 23-30.

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

Intermediate Range of Crack Velocities", Strength of Inorganic Glass , ed. C.R. Kurkjian, Plenum Publishing Corp., NY, 1985, pp . 219-229.

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

Percee d'un Trou Soumise en Compression', Mechanics Research

Communications 10 (5) 253-257 (1983).

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

Under Compressive Stress States", Act Me tall. 34 511-526 (1986).

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

Compression Specimen", Int. Journal of Fract. 3

3

223-325 (1987).

7. H. Tada, P.C. Paris, and G.R. Irwin, The stress analysis of Cracks

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

8. J.F. Rhodes, ARCO Chemical Company, Greer, SC, May 14, 1986.

9. R.F. Krause, Jr., "Rising Fracture Toughness from the Bending Strength

of Indented Alumina Beams," submitted to the J. Amer. Ceram. Soc

.

. August 1986

.

10. R.F. Krause, Jr. and E.R. Fuller, Jr., "Fracture Toughness Behavior of

a Silicon Carbide Whisker-Reinforced Alumina Ceramic at Selected

Porosities," in Fossil Energy Materials Program Conference

Proceedings . ed. by R.R. Judkins, Oak Ridge National Laboratory,

Oak Ridge, TN , Aug. 1987, pp . 38-55.

11. G.W. Hollenberg, G.R. Terwilliger, and R.S. Gordon, "Calculation of

Stresses and Strains in Four-Point Bending Creep Tests," J

.

Amer

.

Ceram. Soc. 54 196-199 (1971).
v";

' -

- ^

1

>
KIBS-114A (REV. 2.ac) U.S. DEPT. OF COMM. 1. BIBLIOGRAPHIC DATA

SHEET (See instructions)

4. TITLE AND SUBTITLE

PUBLICATION OR

REPORT NO.

NBSIR 88 3710

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

FEBRUARY 1 5 8 8

STRUCTURAL RELIABILITY AND DAMAGE TOLERANCE OF CERAMIC COMPOSITES FOR HIGHTEMPERATURE APPLICATIONS

Semi-Annual Progress Report for the Period Ending September 30, 1987.

5. AUTHOR(S)

E. R. Fuller, Jr., T. W. Coyle, and R. F. Krause, Jr

6. PERFORMING ORGANIZATION (If joint or other than NBS. see instructions) 7. Contract/Grant No.

9 .

NATIONAL BUREAU OF STANDARDS

U.S. DEPARTMENT OF COMMERCE

GAITHERSBURG, MD 20899

8. Type of Report & Period Covered

SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street. City. State. ZIP

US Dept, of Energy, Fossil Energy

Advanced Res. & Technology Development

Fossil Energy Materials Program

Oak Ridge National Lab., Oak Ridge, TN

10.

SUPPLEMENTARY NOTES

F 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

bibliography or literature survey, mention it here)

The achievement of higher efficiency heat engines and heat recovery systems requires

the availability of high-temeprature , high-performance structural materials. Structure

ceramics, and more recently, ceramic matrix composites have received particular

attention for these applications due to their high strength, and resistance to corrosioji

and thermal shock. Even with these positive attributes, improved reliability and

extended lifetime under service conditions are necessary for structural ceramics to

gain industrial acceptance. The problems with these materials are mechanical and

chemical in nature and are enhanced by the fact that they are subjected to high

temperatures, reactive environments and extreme thermal gradients.

With an objective of improved performance for heat engine/heat recovery

applications, the NBS program on structural ceramics and ceramic composites addresses

these problems through the determination of critical factors that influence mechanical

and micros tructural behavior. The activites of the program are grouped under two

major subtasks, each designed to develop key data, associated test methods and

companion predictive models. The status of the subtasks are detailed in the following

sections

12.

KEY WORDS (Six to twelve entries ; alphabeti cal order; capitalize only proper names; and separate key words by semicolon s) composites; double-cleavage-drilled-compression; indentation; R-curve; SiC fibers;

SiC whiskers; shear resistance; strength

13. AVAILABILITY

E Unlimited

For Official Distribution. Do Not Release to NTIS

Order From Superintendent of Documents, U.S. Government Printing

20402.

Office, Washington, D.C.

14. NO. OF

PRINTED PAGES

26

15. Price

pyj Order From National Technical Information Service (NTIS), Springfield, VA. 22161 $11.95

U SCOMM-DC 8043-P80
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