NBSIR 77-1400

Proof Testing of Hot Pressed

Silicon Nitride

S. M. Wiederhorn and N. J. Tighe

Inorganic Materials Division

Institute for Materials Research

National Bureau of Standards

Interim Report, December 1, 1977

Prepared for United States Air Force

Air Force Materials Laboratory (AFSC)

Aeronautical Systems Division

Wright- Patterson AFB, Ohio

NBSIR 77-1400

PROOF TESTING OF HOT PRESSED

SILICON NITRIDE

S. M. Wiederhorn and N. J. Tighe

Inorganic Materials Division

Institute for Materials Research

National Bureau of Standards

Washington, D C. 20234

Interim Report, December 1, 1977

Prepared for United States Air Force

Air Force Materials Laboratory (AFSC)

Aeronautical Systems Division

Wright-Patterson AFB, Ohio 45433

U.S. DEPARTMENT OF COMMERCE, Juanita M. Kreps, Secretary

Dr. Sidney Harman, Under Secretary

Jordan J. Baruch. Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Acting Director

PROOF TESTING OF HOT PRESSED SILICON NITRIDE

S. M. Wiederhorn and N. J. Tighe

Institute for Materials Research

Washington, D. C. 20234

ABSTRACT

Proof testing was investigated as a method for insuring the reliability

of hot-pressed silicon nitride in high temperature structural applications. The objective of the study was to determine if the strength

distribution of a population of test specimens could be truncated by

proof-testing. To achieve this objective the strength of silicon nitride

was measured at 25°C and 1200°C, both with and without proof-testing.

At 25°C the strength distribution was effectively truncated by prooftesting. At 1200°C, however, the effectiveness of proof-testing as a

means of truncating the strength distribution was determined by the

resistance of the silicon nitride to oxidation. Although oxidation

removes machining flaws that limit the strength of silicon nitride, long

term exposure to high temperature oxidizing conditions resulted in the

formation of surface pits that severely degraded the strength. Provided

the effects of high temperature exposure are taken into account, prooftesting Ls shown to be useful for truncating the strength distribution of

hot-pressed silicon nitride at elevated temperatures.

1 . Introduction

Because they are more refractory and resistant to oxidation than metal

alloys, ceramics, such as silicon nitride and silicon carbide, are being

considered for use as structural components in turbine engines.

Operation of these engines at high temperatures will improve their
efficiency in vehicles and power generation facilities. To be used in

these applications, however, ceramic materials will have to be as reliable

as the alloys that they are to replace, which means that hundreds, or

even thousands of hours of service-free operation will be necessary.

Since ceramic materials are brittle, long-term reliable operation will requir

methods of design particularly suited for brittle materials. Fortunately,

recent developments in materials science have provided new design methods

2-9

that can be used for turbine applications.

Proof-testing is one of the new methods seriously being considered as a

means of improving the reliability of silicon nitride for turbine applications. Proof-testing is used to eliminate weak components before they

are placed in service since it is these components that are most susceptible

to failure 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 components that contain flaws greater than a certain

maximum size so that the strength distribution is truncated, leaving

only the strong components to be put into service. The theory of prooftesting provides a unique relation between the proof-test stress, the

service stress, and the predicted minimum time-to-failure in service.

This can be expressed in terms of design diagrams (figure 1) that can be

used by engineers to establish proof-test levels for specific applications.

As indicated in figure 1, the proof-test stress is estimated from the

r» diagram once the exposure time and the service stress are known. As a

hypothetical example, if 100 hours 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.
In this paper, proof-testing is investigated as a means of improving the

lifetime of hot-pressed silicon nitride, one of the main materials being

considered for high temperature structural applications. Because the

theory of proof-testing assumes that the population of flaws remains

invarient after the proof-test, the effect of temperature on the flaw

population must be considered before proof-testing can be accepted as a

viable screening procedure. High temperature exposure in air is known to

result in surface oxidation which can modify the flaw population and

thus invalidate the life-time predictions made by proof-testing. Therefore,

to establish proof-testing as a screening procedure, it is necessary to

characterize the effect of high temperature oxidation on the flaw population.

This can be accomplished by determining the change in the strength

distribution of the hot-pressed silicon nitride after exposure to high

temperatures

.

2. Experimental Procedure

Specimens were made from two billets (A and B) of hot-pressed magnesia

doped silicon nitride (NC-132) , which were manufactured at different

times from different powder lots, but had the same nominal composition.

Typical impurities (weight %) included: Al(0.2), Ca(0.05), Fe(0. 4-0. 64)

,

Mg(0. 7-. 74) , W(2 . 6-2 . 7)

.

Within these normal levels of impurity, the

original powder for Billet A was reported by the manufacturer to contain

about twice as much calcium and aluminum as the powder for billet B.

Other levels of impurities were about the same for the two billets.

Test bars approximately 4x5x50 mm were cut from the billets and ground

along their length with a 180 grit diamond wheel. The bars were tested
in 4-point bending, in air, at temperatures of 25°C and 1200°C. The

upper temperature was selected because it is believed to be the maximum

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 subsequently loaded to failure after various heat treatments without

altering their position. In this way, the stress distribution during

proof-testing was duplicated 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 a rate of approximately A0°C per

minute and the test specimens were equilibrated for one-half hour at

1200°C before conducting strength measurements.

As will become apparent in the discussion of results, changes in the

flaw population as a result of oxidation significantly affected the test

results. To study the effect of surface oxidation on the flaw population,

strength measurements were made on billet B at 25°C and 1200°C after an

exposure time of 2, 16, 64 and 100 hours at 1200°C. A similar study was

conducted on billet A using an exposure time of 100 hours. In addition

to these studies, a controlled set of flaws was artifically introduced

into the four-point bend specimens from billets A and B by indenting with a 2

Kg load using a Knoop indenter. These specimens were then broken in fourpoint bending at 25°C and 1200°C after exposure times of 1/2, 16 and 100

hours at 1200°C.

Since the strength studies reported in this paper were conducted in air

which oxidized the silicon nitride, microstructural characterization of
the ceramic surface was necessary to clarify our strength results.

Light microscopy , and scanning and transmission electron microscopy were

used to characterize 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 characterize the structure and

chemical composition of the oxide coat. The results of these studies

were then correlated with the results of the strength measurements.

3 . Experimental Results

3.1 Proof Testing

The results of all strength tests are presented in figures 2 through 7

in a form of Weibull diagrams. Figure 2 presents the results of a room

temperature study on billet B, in which a proof stress of 630 MPa (88

Ksi) was applied to a set of 40 specimens. As predicted by theory'*’

(curved line in figure 2) the initial distribution was truncated, all

proof-test survivors breaking at stresses that were greater than the

proof test level. The results of a room temperature study on billet A

were similar to those on billet B. For both billets, fracture initiated

from surface flaws indicating that machining damage was the primary

source of failure at room temperature. From this data we conclude that

proof-testing is a useful procedure for truncating the room temperature

strength distribution of hot-pressed silicon nitride.

For turbine applications* the value of proof testing depends on its

effectiveness in truncating the high temperature strength distribution

of silicon nitride. Furthermore, it is important that the strength
distribution be truncated at room temperature where testing is easy and

11 12 where subcritical crack growth is greatly supressed. * Once the

strength distribution has been truncated at room temperature it can be

13 shown theoretically that the distribution will also be truncated at

high temperatures provided subcritical crack growth from preexisting

flaws is the only cause for fracture. To test this prediction, a group

of 46 specimens from billet A were proof-tested at room temperature and

then broken at 1200°C. Although truncation of the distribution at high

temperature was expected, the results of the study (figure 3) show that

the strength distribution of the proof-test survivors was not truncated

at 1200°C. The strength data after proof-testing fell almost exactly in

line with the data obtained on specimens that had not been proof-tested.

As a consequence, the strength distribution after proof-testing did not

agree with the theoretical curve given by the solid curve in figure 3.

From these results, it is concluded that the flaw population is affected

by the high temperatures and that before proof-testing can be used as a

screening procedure the effect of high temperatures on the flaw population

must be understood.

3.2 Exposure Tests

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

exposure experiments was conducted at 1200°C. In some experiments the

specimens were proof-tested before exposure in order to truncate the

initial distribution, while in others the specimens were exposed without

proof-testing. Strengths were measured after exposure either at room

temperature or at 1200°C. Figure 4 shows the results 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

4A we note the strength distribution after exposure differs from the

initial 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. In figure 4B we see that the initial

normal distribution was also affected by exposure to high temperatures;

only now the specimens broke at somewhat higher loads after exposure. A

comparison of figures 4A and 4B indicates that after exposure, both the

normal and truncated populations gave nearly the same distribution,

suggesting that 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 elevated temperatures.

This suggestion is explored in a later section of the paper.

The results of exposure tests on billet B conducted at 1200°C for 1/2,

2, 16, 64 and 100 hours are shown in figures 5 and 6. In figure 5

strength measurements were made at 1200°C; specimens were broken immediately

after exposure without being cooled to room temperature. In figure 6

strength measurements were made at room temperature after high temperature

exposure. As can be seen from figure 5, the strength of hot-pressed

silicon nitride is continuously modified by exposure to 1200°C. Exposure

times of 2 and 16 hrs. improve the strength distribution over that found

for the 1/2 hour exposure (figure 5A) . The average strength is observed

to increase and the scatter in strength (as indicated by the slope of

the Weibull curve) is observed to decrease. After 64 hrs. of exposure

(figure 5B) , the strength data lay close to the 16 hr. data (with the
exception of one low strength value, which does not seem to be part of

the main body of data) . Results from the 100 hour exposure (figure 5C)

/ suggest that the strength distribution may be starting to become adversely

affected by the temperature, since some of the strengths are somewhat

lower than those obtained after 64 hours, and the scatter in strength is

somewhat greater than that obtained after 64 hours of exposure. Never theles

the 100 hour strength distribution is still better than the distribution

obtained after 1/2 hour of exposure.

Strength measurements made on billet B at room temperature after exposure

to 1200°C, figure 6, differed considerably from those discussed above.

After 1/2 hour exposure the strength distribution is somewhat better

than the strength distribution of specimens that were not exposed to

1200°C (figure 6A) : there is less scatter in the strength data and the

mean strength, ^660 MPa, is somewhat greater than the mean strength,

^640 MPa, of the initial distribution. After 16 hours of exposure, the

mean strength, ^580 MPa, of the distribution has decreased somewhat,

while the scatter in the strength data appears to be less than that of

the initial distribution.* After 100 hours of exposure (fig 6C) the

mean strength of the distribution is about the same as that obtained

after 16 hours of exposure, but the scatter in strength has been reduced

considerably. By comparing figures 5 and 6 we see that there is a major

*Again one data point fell far below most of the other values. This data

point does not seem to be a main part of the strength distribution and was

not used in fitting the line shown in figure 6b.
9

difference in behavior that depends on whether the specimens have been

cooled to room temperature before testing. Tests conducted at 1200°C

without cooling to room temperature show a significant improvement in

mechanical strength with exposure time, whereas those that have been

cooled to room temperature before testing show a slight, but definite

degradation in the mean strength. It is believed that this difference

in behavior is the result of cracking due to a thermal expansion mismatch

between the oxide coat and ceramic substrate.

The results of exposure tests on billet A, conducted at 1200°C for 100

hours, are shown in figure 7. In contrast to the results given above

for billet B, billet A suffered severe strength degradation as a result

of high temperature exposure. After exposure, there is a shift of the

strength distribution to lower values of strength, both for specimens

broken at 25°C (fig. 7A) and for specimens broken at 1200°C (fig. 7B)

.

At 25°C, the mean value of the strength of billet A decreased from ^670 MPa

to ^450 MPa. This decrease in strength is somewhat greater than that observed for

billet B, ^640 MPa to ^590 MPa, for the same test conditions. At 1200°C,

the mean value of the strength of billet A decreased from ^530 MPa for

the 1/2 hour exposure to ^440 MPa for the 100 hour exposure. By contrast

the strength of billet B for the same exposure conditions was observed

to increase from ^520 MPa to ^620 MPa. It is concluded from these

results that after high temperature exposure, billets A and B exhibit

very different mechanical behavior. As is discussed in Section 4.2,

strength degradation of hot-pressed silicon nitride is believed to be
the result of preferential surface attack during oxidation, which occurs

because of impurities in the material. Minor differences in purity

probably account for the differences in the mechanical behavior observed

for billets A and B.

3.3 Indentation Studies

Since proof testing is used to eliminate weak components (i.e. those

with large flawsX it was necessary to evaluate the effect of temperature

on specimens that contained relatively large flaws. A Knoop indenter

was used to introduce flaws into specimens from both billets. The

strength of these specimens was measured after 1/2, 16 and 100 hours

exposure at 1200°C, table 1. At both room temperature and 1200°C (1/2

hour exposure), the strengths of the indented specimens fell well below

the mean strength of the specimens that were not indented. By comparing

table 1 with figure 2, we see that the mean strength of the indented

specimens falls at 0.02 failure probability level of the as received

specimens tested at 25°C, while the mean strength of the indented

specimens tested at 1200°C falls at the 0.03 probability level of the as

received specimens tested at the same conditions. In both cases fracture

occurred through the indentation. Thus, the indentation introduces a

flaw that is more severe than the majority of flaws initially present in

the specimens, and, as a consequence, the indented specimens can be used

to study the effect of temperature on the low-strength portion of the

strength distribution.
11

Results of exposing indented specimens at 1200°C for extended periods

of time are given in table 1. For billet A, fracture originated at the

indentation site in specimens that had been exposed for 16 hours. After

100 hours of exposure, however, fracture in specimens from billet A

always nucleated at some place other than the indentation. In contrast

to this result, fracture in specimens from billet B generally originated

from the site of the indentation, indicating that the cracks that were

introduced into this billet by the indentation process were still effective

nucleating sites for fracture after high temperature exposure. This

difference in behavior is apparently caused by differences in oxidation

behavior of the two billets. In billet A, the oxidation was rapid enough

so that after 100 hours of exposure the indentation was completely

removed by the oxidation process. In addition, accelerated localized

oxidation produced pits in this billet that acted as new sources of

fracture (see Section 4.3). By contrast, the oxidation reaction was

less severe for billet B, so that the cracks introduced by indentation

still served as nucleation sites for fracture after 100 hours of exposure

at 1200°C. Furthermore, fewer and smaller pits formed in billet B after

100 hours of exposure so that these did not act as fracture origins. As

noted in Section 4.4, these results have significant implications with

regard to the proof testing of these materials for high temperature

applications

.

3.4 Microscopic Examination of Oxidized Surfaces

From the strength studies just described, it may be concluded that high

temperature oxidation has a strong effect on the strength and hence on
1

the flaw distribution of the test specimens. In order to gain a deeper

| insight into the effect of oxidation on strength, light and scanning electron

microscopy were used to investigate the surface of test specimens that

were exposed to temperatures of 1200°C for various periods. Light microscopy

studies indicate that heating in air results in the formation of an

oxide coat at the surface of the specimens, the oxide coat thickening

with exposure time. As can be seen from figure 8 , the surface texture

changes significantly with time, so that after 16 hours of exposure,

grinding traces are no longer visible on the specimen surface. Platelets

of oxide are first observed after 2 hours of oxidation and by the time

the specimens have been exposed for 100 hours, the surface is covered

with large plates of oxide. As can be seen from figure 8 c and 8<4 considerable

cracking occurred within the oxide coat after 16 hours of exposure.

After 100 hours of exposure, the oxide coat thickness was approximately

50 pm, and isolated mounds were observed in the oxide coat. These

mounds have a glassy appearance, and, as shown in figure 9, some of the

mounds from billet A were found to have a hole in them. By focusing

through these holes, pits ^50 pm were observed beneath the mounds. The

appearance and structure of the mounds suggest that a liquid phase was

present during exposure. Removal of the oxide coat by etching with an

HF-H2 SO 4 solution revealed pits ^50 pm in diameter beneath the mounds.

These pits were the primary sources of fracture in specimens from

billet B that had been exposed for 100 hours, figure 10.

To further characterize the nature of the oxide surface, x-ray diffraction,

and scanning and transmission electron microscopy studies were conducted
13

on specimens that had been oxidized. X-ray diffraction powder analysis showed

that specimens exposed for as short a time as 1/2 hour contained silicon

oxynitride (Si 2 N 2 0) within the oxide coat. After 100 hours of

exposure enstatite and B-cristobalite (B-Si02 ) as well as siliconoxynitride were found in the oxide coat. In addition to these crystalline

phases, the presence of two amorphous phases were identified in the

oxide coat by transmission electron microscopy. The chemical composition

of the oxide coat and the mounds that formed in this coat were evaluated

by means of a scanning electron microscope with elemental analysis

capability. Specimens that had been exposed for 100 hours contained

relatively large amounts of silicon, magnesium, calcium, and iron in the

oxide coat. Since these last three elements are the major impurities of

the magnesium-doped silicon nitride, their presence in the oxide coat

suggests diffusion of the impurities from the body of the ceramic to the

oxide layer during oxidation. The surface mounds of billet A contained

only silicon and iron, and hence differed in composition from the remainder

of the oxide coat, figure 11. By contrast, the mounds in billet B were

approximately same composition as the oxide coat.

One may conclude from the above results that oxidation of silicon nitride

is a relatively complex process involving both mass transport and the

nucleation of new phases within the oxide coat. The results of this

14 15

study are consistent with earlier studies by Singhal, Tighe and

Freiman et al.

1 ^’ 1 '7 Tighe 1

"*, who used transmission electron microscopy to

characterize the oxide coat, noted that the structure of the coat is

fairly complex, consisting of several polycyrstalline phases that contained
many pores and cracks. Tighe also noted that silicon oxynitride always

formed adjacent to the bulk silicon nitride, apparently nucleating and

growing preferentially along grain boundaries. Other oxide phases such

as enstatite, tridymite and cristobalite were observed forming in a

layer over the silicon oxinitride.'*'

4. Discussion

4.1 Strength

The strength measurements reported in this paper can be compared with

18

earlier work by Singhal and Kossowsky. Their studies were on a grade

of silicon nitride (HS-130) that contained a slightly higher impurity

content than the material used in the present study. They found that

the flexural strength of their materials was reduced considerably by

exposure to high temperatures. After several hundred hours exposure to

1200°C, the strength measured at 1200°C (after first cooling to room

temperature) was found to be approximately 60 percent of the initial

strength. The major amount of strength loss (^75 percent) occurred in

the first 100 hours of exposure. In the present study a comparable loss

(^80 percent) in the high temperature strength was observed on billet A

after 100 hours of exposure. By contrast, the high temperature strength

of billet B was still greater than the initial strength after 100 hours

of exposure. In this case however, a degradation in strength (compared

to the 16 or 64 hour exposure) was becoming evident, and it is probable

that a further decrease would have occurred with longer exposure times.

A greater resistance to oxidation and hence to pit formation would

account for the slower rate of strength degradation for billet B.
15

4.2 Surface Oxidation

From studies reported in sections 3 and 4 of this paper; we conclude that

the integrity and thickness of the oxide coat influences the strength of

silicon nitride specimens that have been exposed to high temperature

oxidizing conditions. During the initial stages of oxidation, the silicon

nitride is strengthened as a result of the elimination of surface flaws

produced during the machining process. Initially only the smaller flaws

are eliminated in this manner, so that the larger flaws still serve as

effective stress concentrators. However, as the oxidation process

continues and the oxide coat penetrates deeper into the specimen,

the largest surface flaws normally present from machining are eliminated

by the oxidation process. After 100 hours of exposure at 1200°C the

oxide coat was deep enough to have removed most of the machining flaws initially

present in the silicon nitride specimens.

4.3 Pit and Crack Formation

If the only result of oxidation were the formation of a uniform coherent

coat on the surface of silicon nitride, then the oxidation process would

be beneficial because it eliminates surface flaws. Unfortunately, a

uniform coherent coat was not formed on the silicon nitride studied in

this paper, and as a consequence its strength was degraded by pits and

cracks that developed during the oxidation process. The most serious

flaws that developed were the pits that formed after long term exposure

at 1200°C (Figure 10). These pits, which were 40-50 pm in size, were

larger than most of the machining flaws that were initially present in

the specimens. As a consequence, these pits were the main sources of
fracture after high temperature exposure. As indicated by the typical

crazed appearance of the oxide surface, (fig. lid) the cracks that were

observed in the oxide coat probably formed as the result of a thermal

expansion mismatch between the oxide coat and the silicon nitride.

However, the low temperature phase transformation of the cristobalite in

the oxide coat may also have contributed to crack formation. Although

the cracks were generally distributed throughout the oxide coat (fig.

11), the ones located at pits were probably most effective in reducing

strength because of the depth of these pits. The narrow size range of

the pits and the possible presence of cracks at the base of the pits

probably accounts for the decrease in the scatter in the strength of

specimens (from billet B) after 100 hours of exposure at 1200°C.

The observation that pits are the main sites for crack nucleation is

consistent with earlier observations by Freiman et al.^*^ and Richerson and

19 Yonushonis in their investigation of the fracture of Si^N^. Freiman et al.

attributed the pit formation to locally severe corrosion, but presented no

14 details of the corrosion process. Singhal , on-the-other-hand , suggests

that pore and crack formation during oxidation is the result of nitrogen

generation at the silicon nitride interface. As nitrogen pressure

builds up at the interface, the oxide coat ruptures locally and cracks

and pits are formed. Another possible explanation of the mechanism of

pit formation involves the oxidation of tungsten compounds which are

known to be present as inclusions in these materials. As the oxidation
interface advances and the tungsten inclusions are uncovered, oxidation

of the inclusions will occur. Because of the high vapor pressure of the

tungsten oxide at 1200°C, tungsten oxide gases will be released from the

surface, leaving pits behind in the silicon nitride surface where the

tungsten inclusions had been located. Although these mechanisms have

been suggested for pit formation, the correct mechanism has yet to be

confirmed experimentally. Since the elimination of pits from the

surface of hot-pressed silicon nitride is a crucial step in improving

the mechanical behavior of this material, ascertaining the correct

mechanism of pit formation is an important topic for future research.

4.4 Proof-Testing

The present study has shown that the initial flaw population is modified

by exposure to high temperatures. The machining flaws that are present

initially are replaced by, or transformed into other flaws that change

slowly with time as a result of high temperature oxidation. Because of

the transient nature of the flaw population over the 100 hour test

period, the effect of proof-testing (or any other non-destructive technique

for that matter) on long-term reliability must be re-examined.

As discussed in the introduction to this paper, proof-testing eliminates

specimens that contain large flaws. The proof— test load determines the

maximum flaw size that can be present in specimens that pass the prooftest, and as long as these flaws remain effective nuclei for crack

growth, the minimum time for failure can be determined from proof—test

theory. If, however, the flaw structure of the solid is modified by

high temperature exposure then the predicted lifetime will differ from
18

the actual lifetime. The proof-test prediction may nevertheless still

be of value if the flaws generated by the high temperature exposure are

less severe than the original ones, for then the predicted lifetime will

be shorter than the actual lifetime and the proof-test prediction will

provide a conservative estimate of the lifetime. If, on-the-other-hand

,

the flaws that replace the initial flaws are more serious stress concentrators,

the predicted lifetime will be longer than the actual lifetime and

proof-testing will be of limited value for lifetime predictions.

In the present study, both possibilities discussed above seemed to have

occurred. Studies on billet A show that cracks formed by hardness

indentation became ineffective as stress concentrators in a period of

somewhere between 16 and 100 hours of high temperature exposure. Furthermore,

the strengths of the specimens after 100 hours of exposure were considerably

less than those measured at the start of exposure, so that after exposure,

the flaws that were generated in specimens from billet A were more

severe than the original machining flaws in this billet. From these

results^ we conclude that proof testing would be of limited value in

assuring the performance of components made from billets similar to

billet A.

In contrast to the above conclusions on billet A, billet B seems to give

much more promise for the application of proof testing as a failure

prediction technique. The strength studies conducted on this billet

suggest that the initial flaw population is replaced by one that is not

as severe after long-term exposure. After 100 hours of exposure, for

example, the strength at 1200°C is greater than the initial distribution
measured at this temperature. Furthermore, after 100 hours of exposure,

the strength distribution determined at room temperature, is better than

the initial strength distribution because the slope of the strength

distribution curve has increased.* Therefore, while the initial flaw

population in billet B is also affected by oxidation, the oxidation rate

is slower and the flaws that are generated by oxidation are not as

serious as the ones originally present in the specimens, which suggests

that proof testing can be used to give a conservative estimate of the

lifetime of components made from billets that are similar to billet B.

The proof test level selected for billet B should be based on both the

design requirements of the components being tested and the severity of

flaws that are generated during oxidation. The latter determine the

maximum load that can be applied during a proof test. The application

of a proof test load greater than this will serve no useful purpose

because flaws will be generated during service that are more severe than

those eliminated during the proof test. From the data generated on

billet B, a maximum proof test load of 400 MPa would seem reasonable for

this billet. The selection of this value is based on the fact that the

flaws formed by hardness indentation were still effective crack nuclei

after 100 hours of exposure to 1200°C. The elimination of mechanical

*Because of this increase in slope, the strength of the low probability

portion of the distribution after 100 hours of exposure is greater than

the strength of the initial distribution Figure 7.
flaws of this magnitude by proof testing, therefore, truncates the

strength distribution so that a minimum lifetime in service can be

assured. This maximum load is indicated by the horizontal line in

figure 1.

For turbine designs currently under consideration, it can be shown that

a maximum proof-test level of 400 MPa is not restrictive to the application of proof testing as a screening procedure. In one vehicular turbine

that is currently under development for example, a maximum stress of 220

MPa is expected in the rotors of this turbine for steady state operation

20

under full-power. To assure an operational lifetime of 25 hours (which

is the current program goal) under these conditions, a proof-stress of

390 MPa is necessary, figure 1. Because it is less than the suggested

maximum proof-stress, this proof-stress should assure satisfactory

operation of the turbine rotors for the expected exposure time. As a

second example, a turbine being designed for electric power generation

is expected to experience extreme stress conditions during emergency

21

shutdown. Thermal transients during shutdown will expose the trailing

edges of the stator blades to a stress of 350 MPa for a period of

approximately 1 second. If a proof-test could be devised for a thermal

shock situation, then a proof-stress of approximately 400 MPa would be

needed for this application, figure 1. Since this value lies at the

proof-test limit shown in figure 1, it too would be acceptable as a

proof-stress. As a final example, a turbine being developed for Naval

22 applications is expected to experience a 98 MPa stress in the first

stage rotor at a temperature of 1040°C . If we assume that the 1200°C
21

data presented in this paper are applicable, then, for a design lifetime

of 50 hours, a proof stress of 170 MPa will be needed.* This stress is

considerably less than the maximum indicated in figure 1. Thus, for the

three examples noted, the proof-stresses are less than the maximum

recommended stress and one may use proof-testing with the assurance that

surface oxidation will not interfere with lifetime predictions.

5 . Summary

Proof-testing has been investigated as a method of improving the lifetime

of hot-pressed silicon nitride in turbine applications. Oxidation of

silicon nitride at 1200°C modified the strength by reacting with surface

flaws. The effect of oxidation on the strength of silicon nitride

appears to depend on the chemical purity of the silicon nitride. Of

the two billets investigated, the one with the higher purity (billet B)

exhibited greater resistance to strength degradation as a result of high

temperature exposure. For short periods of exposure (1/2 hour), the

initial flaw population of both billets was replaced by a second population

of flaws which continued to be modified as the exposure time increased.

For exposure times up to 100 hours, the flaws in billet B became less

severe than the initial set of flaws, and as a consequence, the high

temperature strength of the silicon nitride increased. By contrast,

*Because crack propagation is more difficult at 1040° C than at 1200°C^\

the proof test diagram at 1040°C will give a lower proof stress for

the same operating conditions. Consequently, lifetimes resulting from a

proof stress of 170 MPa will be much greater than 50 hours.
strength measurements on billet A indicate severe strength degradation

as a result of high temperature exposure. After exposure to 1200°C,

specimens tested at room temperature were weak relative to specimens

tested at 1200°C. The apparent weakening was probably due to cracks

that formed on cooling as a consequence of a thermal expansion mismatch

between the oxide coat and the silicon nitride substrate. High temperature

modification of the strength resulted from the formation of a complex

oxide coat, which initially removes machining flaws that were present in

the surface, and later results in the formation of pits that weaken the

material. As a result of this oxidation process and the formation of

pits, there is a possibility of premature failure after proof-testing if

the proof-stress is set too high. Hence, for the particular grade of

hot-pressed silicon nitride investigated, a maximum proof-test stress of

400 MPa was recommended. For the experimental turbines currently under

development, this recommendation does not restrict the application of

proof testing as a screening procedure.

Acknowledgment: The authors greatfully acknowledge the support of the

Air Force Materials Laboratory.
23

References

1. Workshop on Ceramics for Advanced Heat Engines, ERDA Division of

Conservation Research and Technology, Orlando, Fla., Jan. 24 to 26,

1977, CONF-770110.

2. A. G. Evans and S. M. Wiederhorn, "Proof-Testing of Ceramic

Materials — An Analytical Basis for Failure Prediction," Int. J.

Fract, 10, 379-392 (1974).

3. R. W. Davidge, J. R. McLaren and G. Tappin, "Strength-Probability-Time

(SPT) Relationships in Ceramics," J. Mat. Sci 8^, 1699-1705 (1973).

4. J. E. Ritter and J. A. Meisel, "Strength and Failure Predictions for

Glass and Ceramics," J. Am. Ceram. Soc. 59^ 478-481 (1976).

5. S. M. Wiederhorn, E. R. Fuller, Jr., J. Mandel, and A. G. Evans, "An

Error Analysis of Failure Prediction Techniques Derived from Fracture

Mechanics," J. Am. Ceram. Soc. 59^, 403-411 (1976).

6. D. F. Jacobs and J. E. Ritter, Jr., "Uncertainty in Minimum Lifetime

Predictions," J. Am. Ceram. Soc. 59^ 481-487 (1976).

7. P. N. Thorby, "Experimental Errors in Estimating Times to Failure,"

J. Am. Ceram. Soc. _59, 514-517 (1976).

8. S. M. Wiederhorn, "Reliability, Life Prediction and Proof Testing of

Ceramics," in Ceramics for High Performance Applications, J. J. Burke,

A. E. Gorum and R. N. Katz (Editors), Brook Hill Publishing Co. (1974)

Chestnut Hill, Mass. PP 633-663.

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

for Brittle Materials," in Mechanical Properties of Ceramics for

High Temperature Applications, AGARD Report No. 651, pp . 41-55.

10. N. J. Tighe, "The Structure of Slow Crack Interfaces," J. Mat. Sci.

in press.

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

Prediction in Silicon Nitride at Elevated Temperatures," J. Mat. Sci.

9, 270-278 (1974).

12. K. D. McHenry, T. Yonushonis and R. E. Tressler, "Low Temperature

Subcritical Crack Growth in SiC and Si„N, ," J. Am. Ceram. Soc. 59 , 262-3 (1976).

13. E. R. Fuller, Jr., S. M. Wiederhorn, P. B. Oates and J. E. Ritter, Jr.,

"Proof Testing of Ceramics - Theory," To be published, J. Am. Ceram. Soc.
24

14. S. C. Singhal, "Oxidation and Corrosion Erosion of Si^N, and SiC,"

pp 533 to 548 in Ceramics for High Performance Applications Editors:

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

Chestnut Hill, Mass. (1974).

15. N. J. Tighe, "Microstructural Aspects of Deformation and Oxidation

of Magnesia Doped Silicon Nitride," Presented at NATO-ASI Institute,

Canterbury, England. NBSIR Report No. 76-1153, Sept. 1976.

16. S. W. Freiman, J. J. Mecholsky, W. J. McDonough and R. W. Rice,

"Effect of Oxidation on the Room Temperature Strength of Hot Pressed

Si^N^-MgO and Si^N^-Zr02 Bodies," To be published in: Ceramics for

High Performance Applications-II , proc. of the Fifth Army Materials

Technology Conf . , Held at the Sheraton Islander Inn, Newport, Rhode

Island, March 21-25, 1977.

17. S. W. Freiman, J. J. Mecholsky and R. W. Rice, "Fracture of Si^N^

and SiC," to be published in the Proceedings of the Sixth International

Materials Symposium, Ceramic Microstructures 1976, U. of California,

Berkley, CA. Aug. 1976.

18. S. C. Singhal and R. Kossowsky, "Effect of Oxidation on the Flexural

Strength of Hot-Pressed Silicon Nitride and Silicon Carbide," to be

published in J. Am. Ceram. Soc.

19. D. W. Richerson and T. M. Yonushonis, "Environmental Effects on the

Strength of Silicon-Nitride Materials," to be published in the

Proceedings of the ARPA/NAVSEA-Garret/Airesearch Ceramic Gas Turbine

Engine Demonstration Program Review, held at the Maine Maritime Academy,

Aug. 1-4, 1977.

20. A. Paluszney, Private Communication.

21. R. J. Bratton, Private Communication.

22. ARPA/NAVSEA Ceramic Gas Turbine Engine Demonstration Program. Interim

Report No. 2 Sept. 1976, Airesearch Manufacturing Co.
25

Figure Captions

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

diagram is based on crack propagation data obtained in air

at 1200°C (Ref. 10).

Figure 2. Proof test conducted on hot-pressed silicon nitride

in air at room temperature. The proof test level,

650 MPa, is given by the arrow in this figure. The

solid curve gives the theoretical distribution of strength

after proof testing.

Figure 3. The effect of a room temperature proof test on the high

temperature strength of hot-pressed silicon nitride. The

solid curve gives the expected theoretical distribution for

the high temperature strengths after proof testing.

Figure 4. Effect of high temperature exposure on the room temperature

strength of hot-pressed silicon nitride: (a) exposure of the

survivors of a room temperature proof test; (b) exposure of

as-received specimens. The solid curve in (a) gives the

strength distribution after proof testing at room temperature,

while the arrow gives the proof test level, 630 MPa.

Figure 5. Effect of exposure on the strength of hot-pressed silicon

nitride at 1200°C-Billet B: (a) 1/2, 2, and 16 hours of exposure;

(b) 64 hours of exposure; (c) 100 hours of exposure.

Figure 6. Effect of high temperature, 1200°C, exposure on the room temperature strength of hot-pressed silicon nitride-Billet B: (a) 1/2,

hour, (b) 16 hours; (c) 100 hours.

Figure 7. Effect of high temperature exposure on the strength of hot-pressed

silicon nitride-Billet A: (a) room temperature strength; (b) strength

at 1200°C

.

Figure 8. Effect of temperature on the surface structure of hot-pressed

silicon nitride: (a) 1/2 hour exposure; (b) 2 hour exposure;

(c) 16 hour exposure; (d) 100 hours exposure.

Figure 9. Mound formed during the high temperature oxidation of hot-pressed

silicon nitride. The hole in the center of the mound suggests

that gas had issued forth from the mound. Microscopic examination

indicates that a pit lies beneath the surface of the mound.

Figure 10. Fracture source in hot-pressed silicon nitride. The test was conducted on the specimen (from Billet B) at 1200°C. The arrow

at the corner of the specimen indicates the pit that nucleated

fracture

.
26

Figure 11. Microscopic examination of a mound formed during oxidation.

(a) and (b) are low magnification photographs of the mound.

(c) chemical analysis of both the mound and the normal oxidized

surface of the silicon nitride. (d) High magnification micrograph showing crack formation on the mound.
27

Table 1. Strength of Indented Specimens: 2kg load.

Material

Test-Temp

.

°C

Exposure

conditions

Strength

MPa

Fracture

at Indentation

NC132 25 as indented 397 + 12 Yes

Billet A 25 16 hr 1200°C 423 Yes

25 100 hr 1200°C 461 + 42 No

1200 1/2 hr 1200°C 438 + 15 Yes

1200 16 hr 1200°C 396 + 42 Yes

1200 100 hr 1200°C 402 + 12 No

NC132 25 as indented 408 + 17 Yes

Billet B 25 100 hr 1200°C 524 + 47 Yes

1200 1/2 hr 1200°C 441 + 10 Yes

1200 100 hr 1200°C 00 00 + 11 Yes
10

28

PROOF STRESS MPa

70 100 200 300 400 500 600 700 MPa

SERVICE STRESS

Figure 1. Proof test diagram for hotpressed silicon nitride. This diagram

is based on crack propagation data

obtained in air at 1200°C (Ref. 10).

400 450 500 550 600 650

ElEXURE STRENGTH MNm " 2

Figure 2. Proof test conducted on hotpressed silicon nitride in air at room

temperature. The proof test level, 650

MPa, is given by the arrow in this figure.

The solid curve gives the theoretical

distribution of strength after proof

testing.

Figure 3. The effect of a room temperature proof test on the high temperature

strength of hot-pressed silicon nitride.

The solid curve gives the expected

theoretical distribution for the high

temperature strength after proof

testing

.
FAILURE.

PROBABILITY

29

250 350 450 550 650 750 850 250 350 450 550 650 750 850

FLEXURE STRENGTH. MNm

2

Figure 4. Effect of high temperature exposure on the room temperature

strength of hot-pressed silicon nitride: (a) exposure of as-received

specimens; (b) exposure of th>e 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.

Figure 5. Effect of exposure on the strength of hot-pressed silicon

nitride at 1200°C-Billet B: (a) 1/2, 2, and 16 hours of exposure;

(b) 64 hours of exposure; (c) 100 hours of exposure.
30

Figure 6. Effect of high temperature, 1200°C, exposure on the room temperature strength of hot-pressed silicon nitride-Billet B: (2) 1/2,

hour, (b) 16 hours; (c) 100 hours.

Figure 7. Effect of high temperature exposure on the strength of hotpressed silicon nitride-Billet A: (a) room temperature strength;

(b) strength at 1200°C.
31

Figure 8. Effect of temperature 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.

Figure 9. Mound formed during the high temperature oxidation of hot-pressed

silicon nitride. The hole in the center of the mound suggests that gas had

issued forth from the mound. Microscopic examination indicates that a pit

lies beneath the surface of the mound.
32

Figure 10. Fracture source in hot-pressed silicon nitride. The test was conducted on the specimen (from Billet B) at 1200°C. The arrow at the corner of the specimen indicates the pit that nucleated fracture.

Figure 11. Microscopic examination of a mound formed during oxidation,

(a) and (b) are low magnification photographs of the mound. (c) chemical

analysis of both the mound and the normal oxidized surface of the silicon

nitride. (d) High magnification micrograph showing crack formation on

the mound

.
NBS-114A (REV. 7-73)

; U S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

. SHEET

4 . TITLE AND SUBTITLE!

1. PUBLICATION OR RLPORT NO.

NBSIR 77-1400

2. Gov’t Accession

No.

Proof Testing of Hot Pressed Silicon Nitride

3. Recipient’s Accession No.

5. Publication Date December 1977

6. Performing Organization Code

7. AUTHOR(S)

S. M. Wiederhorn and N, J. Tighe

8 . Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No. 3130157

11. Contract/Grant No.

12. Sponsoring Organization Name and Complete Address (Street, City, State, ZIP)

United States Air Force

Air Force Materials Laboratory (AFSC)

Aeronautical Systems Division

Wright-Patterson AFB, Ohio 45433

13. Type of Report & Period Covered

14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

16. ABSTRACT (A 200-word or less (actual summary ot most significant information. If document includes a significant

bibliography or literature survey, mention it here.)

|N Proof testing was investigated as a method for insuring the reliability of

hot-pressed silicon nitride in high temperature structural applications. The objective

of the study was to determine if the strength distribution of a population of test

specimens could be truncated by proof-testing. To achieve this objective the strength

of silicon nitride was measured at 25°C and 1200°C, both with and without proof-testing,

At 25°C the strength distribution was effectively truncated by proof-testing. At 1200°C,

however, the effectiveness of proof-testing as a means of truncating the strength distribution was determined by the resistance of the silicon nitride to oxidation. Although

oxidation removes machining flaws that limit the strength of silicon nitride, long

term exposure to high temperature oxidizing conditions resulted in the formation of

surface pits that severely degraded the strength. Provided the effects of high

temperature exposure are taken into account, proof —testing is shown to be useful for

hot-pressed silicon nitride at elevated temperatures.

17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper name; separated by semicolons)

Ceramic turbines; failure prediction; fracture; oxidation; silicon nitride; strength

,18. AVAILABILITY Dt Unlimited

Q For Official Distribution. Do Not Release to NTIS

1 Order From Sup. of Doc., U.S. Government Printing Office Washington, D.C. 20402, SD Cat. No. Cl 3 E Order Prom National Technical Information Service (NTIS)

Springfield, Virginia 22151

19. SECURITY CLASS 21. NO. OF PAGES

(THIS REPORT)

UNCL ASSIFIED 34

20. SECURITY CLASS 22. Price

(THIS PAGE)

UNCLASSIFIED $4.00