AFML-TR-78-83

-tit- 7$ -/*?'*"

FRACTURE OF BRITTLE MATERIALS

AT HIGH TEMPERATURE

N. J. TIGHE

S. M. WIEDERHORN

FRA CTURE AND DEFORMA TION DIVISION

NA TIONAL MEASUREMENT LABORA TOR Y

NA TIONAL BUREA U OF STANDARDS

WASHINGTON, D C. 20234

JULY 1978

TECHNICAL REPORT AFML-TR-78-83

Final Report for Period January 1976 — December 1977

Approved for public release; distribution unlimited.

AIR FORCE MATERIALS LABORATORY

AIR FORCE WRIGHT AERONAUTICAL LABORATORY

AIR FORCE SYSTEMS COMMAND

WRIGHT-PATTERSON AIR FORCE BASE, OHIO 45433
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data, is not to be regarded by implication or otherwise as in any manner licensing the holder or any other person or corporation , or conveying any rights or permission to manufacture , use , or sell any patented invention that may in any

way be related thereto.

hnical report has been reviewed and —

-

lENRY C/ GRAHAM

Project Engineer

is approved for publication.

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AIR FORCE/56780/9 January 1979 — 135
UNCLASSIFIED

SECURITY CLASSIFICATION OF THIS PAGE (When Data Entered) REPORT DOCUMENTATION PAGE READ INSTRUCTIONS

BEFORE COMPLETING FORM

1. REPORT NUMBER b. GOVT ACCESSION NO. 3. RECIPIENT'S CATALOG NUMBER

AFML-TR-7 8-83

4. TITLE (and Subtitle)

Fracture of Brittle Materials at High

Temperature

5. TYPE OF REPORT & PERIOD COVERED

Final Report

January 197 6-December 1977

6. PERFORMING ORG. REPORT NUMBER

7. AUTHORfsj 8. CONTRACT OR GRANT NUMBERfs;

N. J. Tighe and S. M. Wiederhorn F33 615-7 6-F-67 52

9. PERFORMING ORGANIZATION NAME AND ADDRESS Fracture and Deformation Division

National Measurement Laboratory

National Bureau of Standards

Washington, D. C. 20234

10. PROGRAM ELEMENT. PROJECT. TASK AREA & WORK UNIT NUMBERS

2306P203

61102F

11. CONTROLLING OFFICE NAME AND ADDRESS Air Force Materials Laboratory

Wright-Patterson AFB, Ohio 45433

12. REPORT DATE July 1978

13. NUMBER OF PAGES

45

14. MONITORING AGENCY NAME & ADDRESS^// different from Controlling Office) 15. SECURITY CLASS, (ot this report)

Unclassified

15a. DECLASSI FI CATION/ DOWN GRADING SCHEDULE

16. DISTRIBUTION STATEMENT (ol this Report)

Approved for Public Release: Distribution Unlimited.

17. DISTRIBUTION STATEMENT (of the abstract entered in Block 20, if different from Report)

18. SUPPLEMENTARY NOTES

19. KEY WORDS (Continue on reverse side if necessary and identify by block number)

Proof Testing; silicon nitride; ceramic turbine; strength; high temperature;

flaw generation; oxidation; strength degradation; hot-pressed Si N^;

reaction bonded Si^N^

20. ABSTRACT (Continue on reverse side if necessary and identify by block number)

DD 1 j an

M

73 1473 EDITION OF 1 NOV 65 IS OBSOLETE

UNCLASSIFIED

SECURITY CLASSIFICATION OF THIS PAGE (When Data Entered)

FOREWORD

This report is the final technical report on "Fracture of Brittle

Materials at High Temperatures" sponsored by the Air Force Materials

Laboratory, AFSC under Contract F33& 1

- 76-F-6752 . The program was

funded for two years.

The principal investigators were Drs. Nancy J. Tighe and

Sheldon M. Wiederhorn. Mr. Lewis Russell assisted in obtaining the

test results. The technical monitor was Dr. Norman Tallan of AFML

iii
TABLE OF CONTENTS

INTRODUCTION 1

EXPERIMENTAL PROCEDURE A

EXPERIMENTAL RESULTS 1 1

DISCUSSION OF RESULTS 29

CONCLUSIONS 35

REFERENCES 37

TABLES 39

PUBLICATIONS A
LIST OF TABLES

TABLE

2

MEAN STRENGTH OF TEST MATERIALS

STRENGTH OF INDENTED SPECIMENS: 2kg LOAD

PAGE

7

17

v
LIST OF ILLUSTRATIONS

FIGURE PAGE

1 DESIGN DIAGRAM, PROOF STRESS VS. SERVICE STRESS 3

2 HYDRAULIC LOADING APPARATUS 9

3 WE 1 BULL PLOT SHOWING EFFECT OF PROOF TESTING ON

STRENGTH DISTRIBUTION

10

4 EFFECT OF PROOF TESTING ON STRENGTH DISTRIBUTION

25°C

AT 12

5 EFFECT OF PROOF TESTING ON STRENGTH DISTRIBUTION

1 200° C AT 14

6 EFFECT OF 1200°C EXPOSURE ON ROOM TEMPERATURE

STRENGTH

15

7 EFFECT OF 1200°C EXPOSURE ON STRENGTH OF NC 1 32

BILLET A

18

8 EFFECT OF 1200°C EXPOSURE ON STRENGTH OF NCX34 19

9 EFFECT OF 1200°C EXPOSURE ON STRENGTH OF NC 1 32

BILLET B

20

10 EFFECT OF 1200°C EXPOSURE ON STRENGTH OF NC350 21

1 1 EFFECT OF EXPOSURE AT 1200°C ON SURFACE 22

12 EFFECT OF 1200°C EXPOSURE ON STRENGTH OF NC 1 32

BILLET B 24

13 SCANNING ELECTRON MICROSCOPIC EXAMINATION OF

SURFACE PIT IN OXIDE LAYER

25

1 A MICROGRAPH OF PIT AS A FRACTURE SOURCE IN HOT-PRESSED SILICON NITRIDE

26

15 MOUND FORMED DURING 1200°C EXPOSURE OF NC 1 32

BILLET B 27

16 SCHEMATIC DRAWING ILLUSTRATING EFFECT OF

OXIDATION ON FLAWS IN Si 0 N,

3 4

28

17 ROOM TEMPERATURE STRENGTH DISTRIBUTION FOR

HOT-PRESSED AND REACTION BONDED MATERIALS

34

18 HIGH TEMPERATURE STRENGTH DISTRIBUTION FOR

HOT-PRESSED AND REACTION BONDED MATERIALS

34
SECTION I

I n t roduct ion

Program Plan

Ceramic materials offer the potential for significantly increased

operating temperatures, increased efficiency and lower costs for gas

turbine engines. However, the brittle nature and resulting catastrophic

mode of failure of ceramic materials requires a design approach based on

consideration of crack nucleation and growth [1-11]. If both the initial

size of pre-existing cracks and the rate at which they grow under an

applied stress are known then the time to reach critical size at which

failure occurs can be predicted. In principle, the initial size can be

determined by non-destructive test techniques. However, state-of-theart NDT equipment does not have the required sensitivity to detect the

flaws in question, which are generally 100 ym or less. An upper limit

on tolerable flaw size can be obtained by proof testing, but the validity

of this approach needs to be demonstrated. There is little information

on the strength of turbine ceramics which can be used for component

design and the relationship between material behavior and microstructure

is not well understood.

The major objective of the proposed research was to demonstrate

that proof testing could be used to assure the reliability or lifetime

of turbine components made from ceramic materials. The secondary

objectives were: (l) the accumulation of critical stress intensity and

crack growth data on new or improved turbine ceramics in order to generate

proof test design diagrams and (2) the determination of the relationship

between the microstructure and the observed fracture behavior.
The technical requirements of the project were as follows: Fourpoint bend tests were conducted on two billets of Norton N

C - 1 3 2 , one

billet of NCX~3^ hot pressed silicon nitride and 5 billets of Norton NC350 reaction bonded silicon nitride. For each material, one set of

specimens was used to determine fracture strength, a second set was used

to determine whether proof testing truncated the fracture strength

distribution, and a third set was used for time to failure measurements

after a proof test was applied. The lifetimes were compared to components

that received no proof tests. A statistical analysis was conducted to

assess the reliability of the proof test technique. The crack growth

rate and the critical stress intensity data required for the time to

failure prediction were obtained by the doub 1 e-tors ion technique developed

by Evans and Wiederhorn [4], In addition, microscopic techniques were

used to study microstructure and the fracture surfaces in order to

determine the mechanisms responsible for crack growth and fracture. The

studies were done both at room temperature and at the temperature expected

in the gas turbine environment, that is 1200°C. The area of primary

emphasis was the validation of the proof test approach to the life

prediction/life assurance of turbine ceramics.

Proof Test Analysis

Proof testing is one of the methods being considered as a means of

improving the reliability of silicon nitride for turbine applications.

Proof testing is used to eliminate weak components, which are most

susceptible to failure in service, before they are placed in service.

In a proof test, a load or proof test stress that is higher than that

expected in service is applied to every component. This load breaks

2
components that contain flaws greater than a certain maximum size and

truncates the strength distribution of the surviving components.

The theory of proof testing provides a unique relation between the

proof test stress, the service stress, and the predicted minimum time to

failure in service [5,7]. This relationship can be expressed in design

diagrams prepared from crack velocity data. These diagrams can be used

by engineers to establish proof test levels for specific materials. In

the diagram shown in Figure 1, the proof test stress is selected for the

70 100 200 300 400 500 600 MPa

SERVICE STRESS

Figure 1. Proof test diagram for hot-pressed silicon nitride NC-132.

This diagram is based on crack propagation data obtained

in air at 1200°C (Ref. 18).

desired exposure time and the service stress. As a hypothetical example,

if 100 hour service at 300 MPa were required, then a proof stress of 560

MPa is determined, from Figure 1, as the stress that would be necessary

to assure full design lifetime of the components.

3
The theory of proof testing assumes also that the population of

flaws remains invarient after the proof test. Therefore if the flaw

population changes during high temperature exposure, the effects of the

exposure must be assessed before proof testing can be used as a viable

screening procedure. High temperature exposure of silicon nitride in

air is known to cause oxidation and strength degradation [12-17]. The

present study has shown that oxidation can modify the flaw population

and thus invalidate the lifetime predictions made by proof testing

[18,19]. Therefore, to establish proof testing as a screening procedure, it is necessary to evaluate the proof test in an inert environment

and to characterize the effect of high temperature oxidation on the flaw

population by determining the change in the strength distribution of the

silicon nitride after exposure. This report will be presented in the

following scheme. In Section II the materials and test procedures will

be presented. In Section III the results of proof tests, strength

tests, and the effects of oxidation on the strength will be presented.

In Section IV we discuss the results, the effects of the changing flaw

population on strength, the statistical significance of the results and

give some recommendations for further experimental work.

SECTION II

Experimental Procedure

Mater i a 1

Specimens were made from two billets of hot-pressed, magnesia-doped

silicon nitride, one billet of yttria-doped silicon nitride, and five

billets of reaction bonded silicon nitride. The specific specimen

details are as follows:

4
Hot Pressed Si 3 N 4 NC-132 - Two bill ets of magnes i a doped

hot pressed silicon nitride ( 15 . 2 x 15 . 2 x 2.5 cm) were

tested. These billets were manufactured at different

times using different powder lots and are designated

billet A and billet B in this report. Both billets had

the nominal manufacturers* analysis in wt % of: 0.20 A 1 ;

0.05 Ca; 0.40-0.64 Fe ; 0.70-0.74 Mg; 2. 6-2. 7 W. Billet A

was manufactured from powder lot HN5 which had the

following manufacturers analysis in wt %: 86 . 5 aSi 3 N 4 ;

0.28 A1 ; 0.06 Ca ; 0.39 Fe. Billet B was manufactured from

powder lot HN 8 which had the following manufacturers analysis

in wt %: 93.3aSi 3 N 4 ; 0.14 A1 ; 0.04 Ca ; 0.36 Fe ; 1.340. The

magnesium addition was not given.

Thirty doubl e-tors ion specimens cut from Billet A were used to

obtain the K-V data in reference [17]; and 1 6 1 bars (5x4x50 mm) were cut

and used in the proof test program. Billet B was cut into 322 bars

(4x5x50 mm) for the proof test program. These bars were numbered with

respect to their position in the billet, the numbers were randomized and

the bars were tested in the order of the random numbers.

Reaction Bonded Si 3 N 4 NC~350 ~ Five billets of reaction

bonded material 15 . 2 x 15 . 2 x 0.6 cm, from the same powder

lot and firing schedule were used for the tests. Each

of the billets was cut into ^72 bars; the bars were

numbered with respect to their position in the billet

and the numbers randomized. The bars from each billet were

tested in groups of 6 specimens so that a total of

Materials were obtained from the Norton Co., Worchester, Mass.

5
30 specimens was used to obtain a complete distribution

of all billets. The strength distribution for the

individual billets differed as shown by comparing the

mean strengths listed in Table 1. The strength

distribution for all 5 billets was used to determine

the proof test load.

Hot-Pressed Si^N 4 NCX-34 - The billet of V2O3 doped

Si 3N4 (15.2x15.2x1 .5 cm) was cut to obtain 90 bars

4x3x50 mm for the proof test program and doub 1 e- tors i on

specimens for crack velocity data. The proof test

program was not completed before the project ended.

All test bars were ground along their lengths with a 180 grit

diamond wheel and the edges were retained. The test bars were prepared

in this way to ensure obtaining a broad strength distribution for each

material. Thus the mean strengths which are listed in Table 1 may be

lower than values obtained in other studies on bars with chamfered edges

or polished surfaces. The bars were numbered as they were cut according

to a scheme so that the location of each bar in the billet was known.

The bar numbers were radomized from a computer program which generated

random numbers by swapping numbers for each set of bars 30,000 times.

Thus each billet had a specific set of random numbers which was based on

the number of bars per billet.

Test Program

The bars were tested in 4-point bending, in air, at 25°C and 1200°C.

The SiC bend fixtures had upper and lower spans of 10 and 40 mm respectively. The upper temperature was selected after tests on the

doub 1 e- tors i on specimens in Billet A showed that considerable plasticity

and slow crack growth occurs above 1200°C. Thus 1200°C is the maximum

6
TABLE 1. MEAN STRENGTH OF TEST MATERIALS

Flexure Strength, MPa

Material 25°C 1200°C (1/2 hr.)

NCI 32 Billet A 670 + 99 534 + 44

NC 1 32 Bi 1 let B 635 + 108 524 + 71

NCX-34 880 + 52 785 + 58

NC350 Billet 6 214 + 33 297 + 53

Billet 7 206 + 41 278 + 50

Bi 1 let 8 259 + 34 274 + 65

Billet 9 181 + 34 206 + 38

Billet 10 262 + 71 342 + 68

Average 6-10 225 + 52 278 + 68

temperature at which this material can be used in high stress applications.

In all tests, the bars were loaded at a cross head speed of 0.02 cm per

minute. After proof testing, the bars were left on the test fixture and

were loaded to failure after the various heat treatments without altering

their position. In this way the stress distribution after proof testing

was not changed during strength measurements. Proof test stresses were

chosen to break approximately 50% of the test specimens for each experimental condition.

For high temperature tests, the furnace temperature was raised at

the rate of rv40°C per minute and the test specimens were equilibrated

for 1/2 hour at 1200°C before conducting strength measurements. It was

found that oxidation significantly affected the test results presumably

by changing the flaw population which controlled the strength. To study

the effect of surface oxidation on the flaw population, strength measurements were made on specimens at 25°C and at 1200°C after different

exposure times up to 100 hours. Additional strength measurements were

7
made on specimens which had an artifical flaw introduced by indentation

before exposure. Specimens were indented using a Knoop indenter with a

2 kilogram load and were broken at 25 and 1200°C after exposure times of

up to 100 hours at 1200°C.

The strength studies at 1200°C were conducted in air and the silicon

nitride surfaces oxidized during the test. It was necessary to characterize

these oxidized surfaces in order to clarify our strength results. Light

microscopy, and scanning and transmision electron microscopy were used

to characterize the morphology of the surface oxide layer. In support

of these studies, x-ray diffraction, electron diffraction, and energy

dispersive x-ray analysis were used to identify the structure and

chemical composition of the oxide coat. The microstructura 1 changes

which occurred as a result of exposure were correlated with the results

of the strength measurements.

Testing Apparatus

Most of the experiments were carried out using a standard screw

operated testing machine, (the Instron type). However, in order to

perform experiments more ef f i c i ent 1 y, a special hydraulic loading test

apparatus was designed and built so that three specimens could be broken

independently in one furnace. Two furnaces (D) were set-up as in Figure 2a.

Each test apparatus has 3 hydraulic cylinders (A) mounted below the

furnace and 3 load cells (E) mounted above the furnace. Silicon carbide

rods (C) attached to each load cell and cylinder with a water-cooled

fixture (B) , extend through the top and bottom of the furnace. Figure 2b

shows a set of three test bars in our standard ^4-point bend jigs assembled

inside the furnace.

8
The load is applied to each specimen individually and the output

is recorded. The loading rate was established by adjusting the oil

pressure to obtain a displacement rate which was equivalent to the

displacement rate used in the standard machine. The data obtained

with this apparatus compared well with that obtained in the standard

machine and the results were reproducible.

(a) (b)

Figure 2. Hydraulic loading apparatus used to test 3 specimens in one furnace (a) furnace and loading arrangements; (b) interior of furnace showing 3 specimens in place.

9
Proof Testing Analysis

The results of the proof test will be displayed in the form of

Weibull diagrams in which the strength values are ordered and plotted as

a function of the failure probability [7,21]. The diagram in Figure 3

shows the scheme of initial and final strength distributions which should

be obtained in proof testing. In our experiments, the initial or inert

population is obtained by breaking specimens at 25°C. This population

is "inert" when no slow crack growth or plasticity occurs. After obtaining

Figure 3. Weibull type probability plot demonstrating the effect of

proof testing on strength distribution: - initial inert

distribution; proof test in an inert environment;

... proof test in an environment in which subcritical

crack growth can occur.

this initial inert strength distribution a stress is chosen that will

break approximately 50% of the population. When the survivors of the

proof test are broken the strength of the suviving specimens should be

higher than this proof stress. Thus when these strength values are

10
ordered and plotted on the Weibull diagram the strengths appear as in

the dashed line of the Figure; that is, the strength distribution is

truncated. If the proof test survivors are affected by the testing

conditions then they will not follow the truncated curve and the strength

distribution can fall somewhere between the ideal proof test distribution

curve and the initial, inert distribution as shown by the dotted line in

Figure 3.

SECTION I I I

E xperimental Results

Proof Testing

The effect of proof testing on the strength distribution at 25°C is

shown in Figure 4A , B for hot pressed silicon nitride and in Figure AC

for reaction bonded silicon nitride. As can be seen in these Weibull plots of

the data, proof testing truncates the strength distribution for both types of

silicon nitride because the breaking strengths of the survivors were

greater than the proof stress. This finding indicates that proof

testing can be used to eliminate weak specimens and thus improve the

structural reliability of silicon nitride in room temperature applications.

However, for turbine applications, proof testing must be able to

truncate the high temperature strength distribution to assure component

reliability. In principle, the easiest way to obtain such a distribution

is to proof test the components at room temperature before using them at

high temperature. Then, provided the high temperatures do not affect

the flaw distribution, the truncated population formed at room temperature

will be truncated also at high temperatures. Alternatively, components must

be proof tested at high temperatures. To test these assumptions, sets of

specimens from both hot pressed silicon nitride and reaction bonded silicon ni

were proof tested at the 50% failure probability level at room temperature

the survivors were unloaded and then broken at 1200°C after 1/2 hour of high
FAILURE

PROBABILITY

.990

Figure A. Effect of proof testing on

the strength distribution of

silicon nitride at 25°C:

(A) hot pressed silicon nitride

billet A; (B) hot pressed

silicon nitride billet B;

(C) reaction bonded silicon

nitride.

12
temperature exposure. Another set of hot-pressed specimens was proof

tested at 1200°C after 1/2 hour at 1200°C and the survivors were unloaded and then

broken at 1200°C. The results of this study are shown in Figure 5.

Included in Figure 5 are the initial strength distribution obtained

after 1/2 hr at 1200°C, and the theoretically predicted values of the

1200° strength distribution (the curved lines) after proof testing. We

note from this figure that none of the strength distributions determined

after proof testing agreed with the theoretical curves. The strength

distributions were not truncated because some proof survivors broke

below the proof test stress. The Weibull slopes for the proof test

survivors are close to those for the initial distribution. It is noteable

that as a result of the 1200°C exposure, specimens which were broken

below the 50% probability level were weaker than predicted and those

broken above the 50% probability level were stronger than predicted.

Further, the 1200°C proof test in Figure 5B did not truncate the strength

distribution much more effectively than the 25°C proof test. Thus it

appears that the proof test predictions which were based on an inert

strength distribution must be modified to account for non-inert service

conditions, i.e. a 1200°C air environment.

The proof test stress determines the largest flaw that can be

present in the region of a specimen that is subjected to the test. As

long as this critical flaw is still an effective nucleus for fracture

after exposure, the minimum time to failure can be estimated by the

proof testing procedure. However, if the critical flaw is modified by

The predictive strength calculations we re done by Sheldon M. Wiederhorn using

empirical formulations being developed for the reliability of ceramic materials.

13
i

mLunc

rnubABILITY

FAILURE

PROBABILITY

350 400 450 500 550 600 650

FLEXURE STRENGTH, MPa FLEXURE STRENGTH, MPa

Figure 5- Effect of proof testing on

the high temperature (1200°C)

strength distribution:

(A) hot pressed silicon nitride

billet A, proof test at 25°C;

(B) hot-pressed silicon nitride

billet B, proof test at 1200°C;

(C) reaction bonded silicon

nitride, proof test at 25°C.

Solid curves are predicted

strengths after proof test.

1A
the high temperature exposure, then proof testing will not be a viable

screening procedure. The results of the tests shown in Figure 5 indicate

that specimens behaved differently from theoretical expectations which

means that the initial flaw population of the materials was modified by

the high temperature exposure. These data show that proof testing will

be of doubtful value at the 50% probability level used to obtain the

data .

To test the effect of temperature on the flaw population, a series

of exposure experiments was conducted at 1200°C. Specimens were prooftested before exposure in order to truncate the initial distribution,

and other specimens were exposed without proof-tes t i ng . Strengths were

measured after exposure at room temperature. Figure 6 shows the results

Figure 6. Effect of high temperature exposure on the room temperature

strength of hot-pressed silicon nitride, billet B: (A) exposure

of as-received specimens; (B) exposure of the survivors of a room temperature proof test. The solid curve in (B) gives

the strength distribution after proof testing at room temperature, while the arrow gives the proof test level, 630 MPa.

15
of a 1/2 hour exposure on the room temperature strength of two sets of

specimens from billet B, one of which had been truncated prior to exposure.

In figure 6A we see that the initial, inert distribution was affected by

exposure to high temperatures, in that the specimens broke at somewhat

higher loads after exposure. In Figure 6B we note the strength distribution after exposure differs from the truncated distribution, which is

given in the Figure by the solid curve. Approximately 50% of the specimens

tested broke at loads that were less than the proof-test load. A

comparison of Figures 6A and 6B indicates that after exposure, both the

normal and truncated populations gave nearly the same distribution. Thus

both populations were altered in such a way that the flaw population

after exposure was independent of the initial population. A reasonable

explanation for this observation is that the initial population of flaws

is replaced by one produced by oxidation at 1200°C in air.

To determine whether proof testing is effective in eliminating

specimens with large flaws, strength measurements were made on specimens

having a flaw that was introduced using a Knoop hardness indenter.

Strength measurements at room temperature (Table 2) indicated that the

indented specimens were comparable to those specimens at the low end of

the strength distribution plot, which broke at flaws of the same magnitude

i.e. 100-200 pm. For proof testing to be useful by eliminating specimens

with flaws of this magnitude, it must be demonstrated that such flaws

are still fracture nucleation sites after exposure to high temperatures.

Strength measurements on indented specimens after exposure to 1200°C

for 100 hours indicated that N C 1 3 2 billet B, the NCX3^ billet, and the

NC350 billets fractured from the site of the indentation, Table 2.

Thus the artifically introduced flaws were still effective nucleation

sites in these materials and proof testing could be used to truncate the

16
TABLE

2.

STRENGTH

OF

INDENTED

SPECIMENS:

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7
FAILURE

PROBABILITY

strength distribution by eliminating large flaws. By contrast, specimens

from N C 1 3 2 billet A failed from origins that were not located at the

site of the indentation, which indicates that proof testing would be of

little value for billet A because specimens would be eliminated

unnecessarily. The indentations were in fact removed during the

oxidation.

The findings obtained in the indentation studies were confirmed in

strength studies conducted on specimens that were not indented. Strength

measurements on billet A after 100 hours of exposure at 1200°C indicated

that the strength at 25°C and 1200°C decreased and the strengths were

considerably below the initial (as ground) strengths shown in Figure 7-

Thus, flaws generated by high temperature oxidation are far more severe

FLEXURE STRENGTH, MPa FLEXURE STRENGTH, MPa

Figure 7. Effect of high temperature exposure in air for 100 hr on the strength of hot-pressed silicon nitride (billet A):

(A) strength at 25°C; (B) strength at 1200°C.

13
than the machining flaws originally present in these specimens. The

flaws also are more severe than the large cracks introduced by indentation.

The strength of the Y 2 O 3 doped silicon nitride (NCX3^) also degraded

after exposure at 1200°C. The data plotted in Figure 8 show that the

mean strength at 25°C decreased after 100 hours exposure from 880 MPa to

692 MPa and the Weibull slope decreased from 16.3 to 5 9 • This 20%

decrease in strength was significantly greater than that observed for

billet A. The indentation results in Table 1 show that the indentation

was still the effective flaw after 100 hr at 1200°C but was not the

effective flaw in specimens broken at 25°C after the same exposure.

Figure 8 . Effect of high temperature exposure for 100 hr on the

strength of Y 2 O 3 doped silicon nitride: (A) strength

at 25°C; (B) strength at 1200°C.

19
The strength behavior of N C 1 3 2 billet B differs from that obtained

on billet A. Strength measurements at 25°C after 100 hours of exposure

to 1200°C, Figure 9a, indicate that the mean strength decreased from 640

MPa to 580 MPa, whereas the Weibull slope increased from 5-9 to 14. A.

The net effect of the increase in the Weibull slope is to increase the

strength of specimens that fail at low levels of failure probability,

which means that large strength impairing flaws are not generated in

billet B by high temperature exposure. As a consequence of this behavior,

FLEXURE STRENGTH, MPa

FLEXURE STRENGTH, MPa

Figure 9- Effect of high temperature exposure for 100 hr on the

strength of hot-pressed silicon nitride (billet B)

(A) strength at 25°C; (B) strength at 1200°C.

20
FAILURE

PROBABILITY

components that contain large flaws can be eliminated by proof testing

without fear that new, more serious flaws will be generated by high

temperature oxidation. This conclusion is supported by the data obtained

at 1200°C after 100 hours of exposure at 1200°C. For this condition of

testing, the measured strengths (figure 7b) were greater than those

obtained after 1/2 hour of exposure at 1200°C. Therefore the natural

flaws generated by high temperature exposure are less severe than those

that were introduced by indentation, and as a consequence, proof testing

can be used to truncate the strength distribution by eliminating large

flaws that might be present in the specimen. We note from Figure 10

and Table 2 that the data on the reaction bonded billets were qualitatively

100 200 300 400

FLEXURE STRENGTH, MPa FLEXURE STRENGTH, MPa

Figure 10. Effect of high temperature on the strength of reaction

bonded silicon nitride; 100 hrs at 1200°C: (A) strength

at 25°C; (B) strength at 1200°C.

21
the same as for those obtained on billet B; that is oxidation improved

the strength by healing or otherwise reducing the size of large flaws.

Therefore, it is concluded that proof testing can be used to improve the

performance of components made from reaction bonded material.

Flaw Healing and Generation

To determine the reason for strength changes in hot pressed silicon

nitride, a study of the oxidized surfaces and of the fracture surfaces

was conducted using light microscopy, scanning electron microscopy and

transmission electron microscopy. Light micrographs shown in Figure 11

present a time sequence of surface oxidation of hot pressed silicon

nitride billet B. This time sequence corresponds with the strength

Figure 11. Effect of exposure at 1200°C on the surface structure

of hot-pressed silicon nitride: (a) 1/2 hour exposure;

(b) 2 hour exposure; (c) 16 hour exposure; (d) 100 hour

exposure

22
measurements shown in Figure 12. As can be seen from Figure 11, the

surface texture changed with time as the oxides formed and after 16 hours

the grinding marks were no longer visible. Oxide filled the grinding

grooves after 1/2 hour, distinct platelets were seen after 2 hours of

oxidation and by the time the components were oxidized for 6k and

100 hours, the specimen was covered with a thick polycrystalline oxide.

The oxide film formed on both billet A and billet B specimens after 100

hr exposure had a glassy appearance and had a dispersion over the surface

of smooth and uneven mounds and pits, 50-100 ym in diameter.

X-ray diffraction studies indicated that silicon oxynitride [ S

2 N 2 0

formed within the scale after 1/2 hours of exposure. While after 100

hours of exposure enstatite [ Mg S

0

] , 3 cristobalite [ 6

_ S

0

] and silicon

oxynitride were found. Transmission electron microscopy studies of oxide

scales demonstrated that at least two amorphous phases (probably magnesium

silicate and silica glasses) can occur within the oxide [21]. The light

microscopy and transmission electron microscopy studies showed that the

silicon oxynitride was located at the silicon nitride-oxide interface and

can be amorphous or crystalline whereas the silicate glasses and other

crystalline materials were located toward the oxide-air interface [14,21]

The chemical composition of the oxide scale and the mounds and the

pits that formed in this scale were evaluated by means of a scanning

electron microscope with elemental analysis capability. Specimens that

were exposed for 100 hours contained silicon, magnesium, calcium, iron

and manganese in the oxide. Since these last four elements are the

major impurities of the magnesium doped silicon nitride, their presence

23
FAILURE

PROBABILITY

.990

.900

.500

.300

DO§ .100

.010 -

.001 300

FLEXURE STRENGTH, MPa

Figure 12.

t i i i i r

1200°C, NC 132-B

* 1/2 HR

64 HR

1 i i l i l i i 400 500 600 700 800

FLEXURE STRENGTH, MPa

Effect of exposure at

1200°C on the strength of

hot-pressed silicon nitride,

billet B: (A) 1/2, 2 and 16

hours; (B) 64 hours; (C) 100

hours

2*4
in the oxide suggests diffusion of some of the impurities from the body

of the ceramic to the oxide layer during oxidation. The surface mounds

and pits found in billet A contained only the silicon and iron, and

hence differed in composition from the remainder of the oxide scale as

seen in Figure 13. By contrast, the mounds and pits in billet B were

approximately the same composition as the oxide scale.

When the oxide scale was removed from N C 1 3 2 specimens oxidized 100

hrs at 1200°C (by etching in Hf+b^SOtJ the silicon nitride surface was

Figure 13- Scanning electron microscopic examination of a pit formed during

exposure at 1200°C for 100 hours: (a,b) low magnification photographs; (c) X-ray element analysis of the pit and the nearby

surface; (d) details of the cracks on the surface. A indicates

same area in a, b and d.

25
found to be heavily pitted. The pits were ^50 ym in diameter and occurred

over all surfaces. By comparing etched and unetched fracture interfaces,

it was evident that fracture originated at the surface pits as shown in

Figure 14. Specimens from Billet A were more heavily pitted than those

from Billet B. Examination of polished cross-sections of the specimens

also showed penetration of the oxide into the silicon-nitride along pits

as well as along cracks.

Figure 1A. Pit as a fracture source in hot-pressed silicon nitride. The

specimen was broken after exposure of 100 hrs at 1200°C. The

surface oxide was removed by etching.

Pits appeared to form under some of the mounds that were observed

in the oxide scale. Some mounds have a glassy appearance and as shown

in Figure 15, sometimes contained a hole in them. By focusing through

these holes [using light microscopy], pits approximately 50 ym in diameter

were observed beneath the mounds. The appearance and structure of some

mounds suggest that a liquid phase was present at the site of the pit

during oxidation. When the indented specimens were examined after 100

26
hr exposure it was found that the indentations were smaller in Billet B

and completely gone in Billet A. Thus the oxidation process had reduced

the effectiveness of the indentation in these specimens and had created

new flaws in the form of surface pits.

Figure 15. Mound formed during exposure of hot-pressed silicon nitride,

billet B, for 100 hrs. at 1200°C. The hole in the center of

the mound suggests that gas had issued forth from the mound.

These microscopy studies on N C 1 3 2 show that the oxidation is a

complex process which involves mass transport of impurities from the

matrix, nucleation and growth of new phases within the oxide scale,

penetration of the silicon nitride interface along cracks and surface

and subsurface inclusions. The diagram in Figure 16 illustrates the

interaction observed between the oxide scale and flaws. Two flaws

27
are indicated: a crack and a subsurface inclusion or pore. Because the

oxide scale grows into the matrix, the original, silicon nitride surface

EFFECT OF OXIDATION ON

FLAWS IN Si 3 N4

1 4 OXIDE—w

)

SI3N4 -Si3 N4—

a b C

Figure 16. Effect of exposure at 1200°C on flaws in silicon nitride:

(a) surface crack and subsurface inclusion before exposure;

(b) after exposure crack size is reduced and subsurface flaw

has reacted with the oxide; (c) oxide growth has removed the

surface crack and enlarged the subsurface flaw to produce

a pit in the silicon nitride.

as indicated by the dotted line recedes. The crack is reduced in size

and the inclusion or pore interacts with the oxide. If the oxidation

continues long enough the crack is removed completely as in (c) and the

inclusion or pore is replaced completely by oxide. The severity of the

oxidation reaction varies with the material. The effect of the oxide scale

on strength degradation can be related to the form of the oxidation

process for each material.

28
Oxidation is less severe for reaction bonded (NC350) and for yttriadoped silicon nitride (NCX3^) specimens than for the magnesia doped (NCI 32 )

materials. An oxide film approximately 10 microns thick formed on the

reaction bonded specimens that were oxidized for 100 hours. The x-ray

powder diffraction patterns from the 100 hour specimens and 1/2 hour

specimens showed that the primary crystalline phase was a-cr i s toba 1 i te

This silica film appeared to be continuous over the surface. The hardness

indentations were visible both before and after acid removal of the

oxide film. The initial porosity of the material makes it difficult to

distinguish additional oxide related pitting. The continuous oxide film

has a strengthening effect on the machined test specimens because it forms a

protective surface layer and fills or heals cracks.

After 100 hours of exposure the oxide film on Y 2 O 3 doped silicon

nitride specimens was approximately 20 ym thick. The X-ray diffraction

patterns of the oxide layers showed that Y 2 Si 2

C>7 and a-Si 02 were the

primary crystalline phases. When the oxide was removed by etching, the

silicon nitride surface showed traces of the machining grooves and had

shallow pits. The absence of deep pits at fracture interfaces indicates

that the strength degradation could not be attributed to the oxide pit

formation as found in N

1 3

and is more likely related to exposure

induced internal oxidation and phase changes in the intergrain phase.

SECTION IV

Discussion of Results

Evaluation of Proof Tests

The results of this study indicate that there are limits to the

application of proof testinq as a method of assurinq the lonq term reliability

of silicon nitride in turbine appl i cat i ons . Proof testinq is

29
of value only if the flaws eliminated during the proof test are in fact

the flaws that will cause failure during service. Modification of the

flaw population by oxidation can invalidate a proof test procedure,

either by eliminating the machining flaws that cause failure at room

temperature, or, by changing their severity. In the present study, it

was shown that high temperature exposure [1200°C] modifies the flaw

population sufficiently so that truncation of the strength distribution

of machined specimens is no longer observed after exposure times as

short as 1/2 hour. Therefore, although the effects of the flaws that

result from machining operations can be reduced by proof testing, this

truncation of the strength distribution will serve little practical

value if those machining flaws probably are not the ones that cause

failure after high temperature exposure. From these results it is

apparent that there is an upper bound to the proof test load above which

no positive benefits can be gained by proof testing. This upper-bound proof

test load depends on the material and on the condition of use. Subjecting

components to loads in excess of the upper-bound proof test level will

merely break an excessive number of components and will not limit the maximum

size of a flaw that can be present during service without causing failure.

The maximum proof test load for a given application can be determined

when an artifical flaw of known size is introduced into a test specimen and

is still effective as a crack origin after high temperature exposure.

In the present study, such flaws were introduced by hardness indentations.

The results of the surface indentation studies for the reaction bonded

silicon nitride show that bend specimens broke at the indentations

after 100 hours of exposure. Since these specimens would have been

broken by a load of ^115 MPa prior to high temperature exposure, the

30
results indicate that a proof load of 115 MPa would have truncated the

strength distribution of the population. In the proof test conducted on

reaction bonded silicon nitride, however, it was shown that a load of

225 MPa broke A5% of the specimens but did not result in truncation of

the strength distribution. From these two results we conclude that the

maximum proof test load for the reaction bonded material should be

between 115 MPa and 225 MPa.

The strength studies on the hot pressed silicon nitride also can be

used to obtain an estimate of the maximum proof test stress for this

material. Indentations in billet B, which reduced the strength to

approximately A00 MPa, were effective flaw nucleation sites after 100

hours of exposure at 1200°C. However, a 630 MPa proof test on this

billet did not result in a truncated population at high temperatures.

From these results we conclude that the maximum proof test load for

billet B lies somewhere between A00 and 630 MPa. In contrast with this

finding on billet B, results on billet A indicate that the maximum

effective load for proof testing must be less than ^A00 MPa, since the

indentations in specimens from this billet were not effective nucleation

sites after high temperature exposure. Therefore, proof testing would

be more effective in truncating the strength distribution of billet B

specimens than billet A specimens. The limited results obtained for

yttr i a-doped , material showed that indentations were effective flaw

nucleation sites for specimens that were broken at 1200°C after 100

hours without cooling but not for specimems that were cooled to room

temperature before breaking. The indentations reduced the inert strength

to %A90 MPa, which indicates that the maximum effective proof test load

should be less than A90 MPa. The cause of the large strength loss which

31
occurs when this material is cooled from 1200°C to room temperature

before testing is more likely related to mi crostructura 1 changes than to

the oxide-flaw interactions that produced the pits in Figure 15-

Effects of Oxidation on Flaws

It is apparent from the results that the oxidation which occurs

during exposure at 1200°C has a significant effect on the strength

distribution. The flaw population which controls the strength can be

changed in the manner shown in Figure 16. Two types of flaw are considered

in Figure 1 6a , a crack at the surface and an internal flaw which can be

an inclusion, a large pore or a crack. When oxidation occurs the crack

interfaces oxidize at least as fast as the ground surface and the crack

can be blunted. The oxide interface grows into the silicon nitride

matr i x, removes surface imperfect ions, and reduces the size of cracks.

Internal inclusions can oxidize and change their volume and composition.

The mounds in the oxide and pits in the matrix, shown in Figures 1^ and 15

respectively, are equivalent to ^50 grain diameters. The accelerated

oxidation in areas this size could be related to a concentration of

impurities in a mass of incompletely sintered grains or to the rapid

oxidation of large tungsten-based inclusions. Such reactions produce an oxide

filled pit in the matrix. Until the microstructure is examined more

closely, the mode of pit formation cannot be described completely.

These oxidation or corrosion pits act as strength limiting flaws in

specimens exposed for long times. It is therefore essential to determine

why they form and how to control their formation.

32
The effects of exposure at 1200°C on the strength distributions of

the four batches of silicon nitride are summarized in Figure 17 and 18. In

these figures, it can be seen that the 25°C and the 1200°C strengths of

reaction bonded material and billet B of N

C - 1 3 2 are increased by exposure

at 1200°C whereas the strengths of the Y 2 O 3 doped-mater i a 1 (NCX-3^)and

billet A of NC-132 are decreased. These differences are statistically

significant and are related to the interaction between the oxide formation

and the material. The mi crostructura 1 differences between the 3 hotpressed materials are related to the composition. The specific interactions could not be pursued during the present program but should be

investigated in order to identify the specific cause of the degradation

and to guide further development of materials.

Suggestions for Future Programs

The effects of exposure on the strength of silicon nitride are

related to compositional variations in the materials. In the present

program the exposure was limited to 1200°C in air for a maximum of 100

hours. The effects of longer exposure and different temperatures on

strength must be documented. It will be necessary also to test these

materials in the service atmosphere. The proof-test procedures for lifetime predictions must be modified to account for the change in flaw

population with exposure. An evaluation of the strength distribution

under static load will be necessary. In order to investigate the effect

of exposure on a wide range of flaw sizes and flaw size distributions,

specimens should be tested with known flaw sizes such as those introduced

with a Knoop indenter. Extensive microstructura 1 analysis should be

carried out on the test specimens as well as on materials prepared with

known compositional and structural variations.

33
FLEXURE STRENGTH, MPa

Figure 17. Comparison of strength of reaction bonded and 3 hot-pressed silicon

nitride compositions: (a) as-received; (b) after 100 hours exposure

at 1 200° C

FLEXURE STRENGTH, MPa

Figure 18. Comparison of the high temperature strength of reaction bonded and

3 hot-pressed composition of silicon nitride: (a) after 1/2 hour

exposure; (b) after 100 hours exposure at 1200°C.
Silicon nitride materials development must parallel the development

of component designs. The materials have not reached the optimum composition

and microstructure for the turbine application. Many long-term use

applications are dependent on having a reasonable expectation of long

service lives and considerable materials analysis and development must

be done to reach this goal.

SECTION V

CONCLUSIONS

Proof testing has been investigated as a means of assuring reliability

and lifetimes of hot-pressed and reaction bonded silicon nitride for

turbine engine applications. It was found that the proof test truncated

the strength distribution at 25°C as predicted by theory. The proof

test was not as successful in truncating the strength distribution of

specimens exposed at 1200°C, because the flaw population changed during

exposure. This change in flaw population occurred because some flaws

healed or were reduced in size, and other flaws were formed when the

specimens oxidized.

Reaction bonded silicon nitride (NC-350) showed an increase in

strength as a result of exposure at 1200°C. The oxidation formed a

protective layer which healed surface cracks. A proof test scheme could

be applied to this material.

The two billets of magnes i a- doped , hot-pressed silicon nitride ( N

C

- 1 32)

behaved differently. The strength distribution of billet A (of earlier

manufacture) decreased as a result of exposure, whereas, the strength

distribution of billet B increased as a result of exposure. Thus, proof

testing could be used for one billet but not for the other. Oxidation

pits were produced in magnesia-doped, hot-pressed material and these

pits become the strength limiting flaws. As a result of oxidation,

35
flaws related to machining healed and could be removed eventually as

the surface oxidized.

The yttria-doped (NCX~3^) silicon nitride proof tests were not

comp 1 eted; however the strength distribution decreased as a result of

exposure and it is unlikely that proof testing could be used for this

billet. Examination of the microstructure before and after exposure

showed that the oxidation produced significant mi crostructura 1 changes

in the near surface areas and these changes could be related to the

strength changes.

36
References

I. Wiederhorn, S. M., Evans, A. G., Fuller, E . R., and

Johnson, H., "Application of Fracture Mechanics to

Space Shuttle Windows," J. Am. Ceram. Soc . 57, 31

- 32

(1974). —

2.. Ritter, J. E., Jr., and Meisel, J. A., "Strength and

Failure Prediction for Glass and Ceramics," J. Am. Ceram.

Soc. 59, 1112 (1976) .

3. Evans, A. G., and Wiederhorn, S. M., "Proof Testing of

Ceramic Materials - An Analytical Basis for Failure

Prediction". Int. J. Fracture, ]_0_, 379~392 (1974).

4. Evans, A. G., and Wiederhorn, S. M., "Crack Propagation and

Failure Prediction in Silicon Nitride at Elevated Temperatures," J. Mat. Sci. % 270-278 (1974).

5. Evans, A. G., and Fuller, E. R. , Jr., "Crack Propagation

in Ceramic Materials Under Cyclic Loading Conditions,"

Met. Transactions 5_, 27 - 33 (1974).

6. Davidge, R.W., McLaren, J. R., and Tappin, G., "StrengthProbability-Time (SPT) Relationships in Ceramics," J. Mat.

Sci 8, 1699-1705 (1973)

7. Wiederhorn, S. M., Fuller, E. R., Jr., Mandel, J., and

Evans, A. G., "An Error Analysis of Failure Prediction

Techniques Derived from Fracture Mechanics," J. Am. Ceram.

Soc. 59_, 403-41 1 (1976) .

8. Jacobs, D. F. , and Ritter, J. E., Jr., "Uncertainty

in Minimum Lifetime Predictions," J. Am. Ceram. Soc. 59 , 481-487 (1976)

9. Thorby, P. N., "Exper i menta 1 Errors in Estimating Times to

Failure," J. Am. Ceram. Soc. 59_, 514-517 (1976).

10. Wiederhorn, S. M., "Reliability, Life Prediction and Proof

Testing of Ceramics," in Ceramics for High Performance

Appl icat ions , J. J. Burke, A. E. Gorum and R. N. Katz

(Editors), Brook Hill Publishing Co. (1974), Chestnut Hill,

Mass. pp 633-663.

II. Wiederhorn, S. M., Tighe, N. J., and Evans, A.G., "New

Design Techniques for Brittle Materials," in Mechanical

Properties of Ceramics for High Temperature Applications,

AGARD Report No. 651, pp. 41-55, 1976.

12.

Davidge, R. W. , Evans, A. G., Gilling. D, and Wilyman, P.R.,

"Oxidation of Reaction Sintered Silicon Nitride and Effects

on Strength," in Special Ceramics Vol . 5, 329-44 (1972).

37
13. Lange, F. F., and Iskoe, J. L., "High Temperature Strength

Behavior of Hot-Pressed S i 3N4 and SiC: Effects of Impurities,"

i n Ceramics for High Performance Applications , ed . by J . J

Burke, A. E. Gorum, and R. N. Katz, Brook-Hill Publishing Co.,

pp 223-38 (1975).

1A. Tighe, N. J., "Mi crostructura 1 Aspects of Deformation and

Oxidation of Magnesia Doped Silicon Nitride," NBSIR Report

No. 76-1153, Sept. 1976, in "Nitrogen Ceramics" F. R. Riley

ed . Nordhoff International Publ., 1977 p AA1-A8.

15. Freiman, S. W. , Mecholsky, J. J., McDonough, W. J., and

Rice, R. W. , "Effect of Oxidation on the Room Temperature

Strength of Hot Pressed Si 3 N 4 -MgO and Si 3 N 4 -Zr02 Bodies,"

To be published in: Ceramics for High Performance Applications-! I, proc. of the Fifth Army Materials Technology Conf.,

Newport, Rhode Island, March 21-25, 1977.

16. Singhal, S. C., and Kossowsky, R., "Effect of Oxidation

on the Flexural Strength of Hot-Pressed Silicon Nitride

and Silicon Carbide," to be published in J. Am. Ceram.

Soc

17. Richerson, D. W. , and Yonushonis, T. M. , "Environmental

Effects on the Strength of Silicon-Nitride Materials,"

to be published in the Proceedings of the ARPA/NAVSEAGarrett/Ai research Ceramic Gas Turbine Engine Demonstration Program Review, held at the Main Maritime Academy,

Aug. 1-A, 1977.

18. Wiederhorn, S. M., and Tighe, N. J., "Application of

Proof Tests to Silicon Nitride," ERDA/Proceed i ngs of

Workshop on Ceramics for Advanced Heat Engines, Orlando,

Fla.

19. Tighe, N. J., "The Structure of Slow Crack Interfaces in

Silicon Nitride," J. Mat. Sci., ]_3 1^55-63 1978 (NBSIR 76-1075).

20. Weibull, W. , "A Statistical Theory of the Strength of

Materials," Ingeniors Vetenskaps Akademian Handlinaar,

No. 151, 1931 .

21. Tighe, N. J., "Microstructure of Oxidized Silicon Nitride

Proc. EMSA p

A 7 O-A 7 I, Claitors Publishing Co., Baton

Rouge , La . ( 1 97A) .

38
Technical Talks Resulting from this Program

N. J. Tighe and S. M. Wiederhorn

Examination of Fracture Interfaces in Silicon Nitride,

Electron Microscopy Society of America, Los Vegas, Nevada,

August 1975- N. J. Tighe

Structure of Fracture Interfaces in Silicon Nitride, American

Ceramic Sodiety Fall Meeting, Indianapolis, I nd . , Sept.

1975. N. J. Tighe

Relationship between Crack Growth

Silicon Nitride, American Ceramic

Indianapolis, Ind., Sept. 1975-

and M i crostructure of

Society Fall Meeting,

N. J. Tighe

Microstructure of Silicon Nitride - Effects of Fracture and

Oxidation, Joint Meeting American Society of Metals and

American Institute of Meta 1 1 urg i ca 1 Engineers, Cincinnati,

Ohio, November 1975. N. J. Tighe

Proof Testing of Silicon Nitride, NATO - Institute on Nitrogen

Ceramics, Canterbury, England, August 1976. N. J. Tighe

Microstructural Aspects of Deformation and Oxidation of

Silicon Nitride, NATO - Institute on Nitrogen Ceramics,

Canterbury, England, August 1976. N. J. Tighe

Proof Testing of Silicon Nitride, USAF Meeting, Dayton,

Ohio, October 1976. N. J. Tighe

Application of Proof Tests to Silicon Nitride, American

Ceramic Society Fall Meeting, October 1976. N. J. Tighe

New Design Techniques for Brittle Materials, 43rd AGARD

Meeting, Structures Panel, London, England, Sept. 1976.

S. M. Wiederhorn

Application of Proof Tests to Silicon Nitride, ASME , Gas

Turbine Conference, Philadelphia, PA., March 1977- N. J. Tighe

Effect of Flaw Generation on Proof Testing of Silicon Nitride,

American Ceramic Society, Annual Meeting, Chicago, Illinois,

April 1977. N. J. Tighe

Characterization of Flaws Produced by Mechanical Grinding of

Si 3 N 4 , Electron Microscopy Society of America, Annual Meeting,

Boston, Mass., August 1977. N. J. Tighe

39
Is Proof Testing Applicable for High Temperature Survival of

S

3 N 4 , American Ceramic Society Fall Meeting, Hyannis,

Mass., Sept. 1977. N. J. Tighe

Effect of Service Conditions on Proof Testing of S

3 N 4 , American Society of Mechanical Engineers, Annual Meeting,

Atlanta, Ga . , Sept. 1977. N. J. Tighe

Application of Proof Testing to Silicon Nitride, Workshop on Ceramics for Advanced Heat Engines, Orlando, Fla., Jan.

1977. S. M. Wiederhorn

Effect of Flaw Generation on Proof Testing, Ceramic Gas

Turbine Engine Demonstration Program Review, Castine, Maine,

August 1977. S. M. Wiederhorn

Application of Proof Testing to Lifetime Predictions in

Silicon Nitride, American Ceramic Society, Metal Systems

Division, Cocoa Beach, Fla., Jan. 1978. S. M. Wiederhorn

40
Technical Publications Resulting from this Program

S. M. Wiederhorn, N. J. Tighe and A. G. Evans, "New Design

Techniques for Brittle Materials," in Mechanical Properties

of Ceramics for High Temperature Application AGARD Report

No. 651 , pp 41-55, 1976.

N. J. Tighe, "M i crostructura 1 Aspects of Deformation and

Oxidation of Magnesia Doped Silicon Nitirde," (NBSIR 76-1153

Sept. 1976) in N i trogen Cerami cs , F. R. Riley ed . , Noordhoff,

Leyden 1977.

N. J. Tighe, S. M. Wiederhorn and L. R. Russell, "Application

of Proof Tests to Silicon Nitride," NBSIR 77-1202, March

(1977) •

N. J. Tighe, "The Structure of Slow Crack Interfaces in

Silicon Nitride," NBSIR 76-1075, (1976), J. Mat. Sci., 1

, 1455-1463, 1978.

S. M. Wiederhorn and N. J. Tighe, "Application of Proof

Testing to Silicon Nitride," ERDA Proceedings of Workshop on Ceramics for Advanced Heat Engines, Orlando, Fla., (Jan.

1977) .

S. M. Wiederhorn and N. J. Tighe, "Effect of Flaw Generation

on Proof Testing," (NBSIR 77“ 1 408) Proceedings Ceramic Gas

Turbine Engine Demonstration Program Review, Castine, Maine

S. M. Wiederhorn and N. J. Tighe, "Proof Testing of Silicon

Nitride," J. Mat. Sci., 13, 1781-1793, (1978).

41

(1977) .

6U.S. Government Printing Office: 1978 — 657-002/320