NBSIR 76-1075

The Structure of Slow Crack

Interfaces in Silicon Nitride

N. J. Tighe

Inorganic Mater!

Institute for Mat

National Bureau

Washington, D.C

Tighe, N. J., The structure of slow crack interfaces

in silicon nitride, J. Mater. Sci. 13, 1455-1463

C1978) .

New

562

5^

NBSIR 76-1075

Interim Report to be published in the Journal Material Science

Prepared for United States Air Force

Air Force Materials Laboratory (AFSC)

Aeronautical Systems Division

Wright- Patterson AFB, Ohio 45433

NBSIR 76-1075

THE STRUCTURE OF SLOW CRACK

INTERFACES IN SILICON NITRIDE

N. J. Tighe

Inorganic Materials Division

Institute for Materials Research

National Bureau of Standards

Washington, D C. 20234

Interim Report to be published in the Journal Material Science

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

3213

THE STRUCTURE OF SLOW CRACK INTERFACES

IN SILICON NITRIDE

N. J. Tighe

National Bureau of Standards

Institute for Materials Research

Washington, D. C. 20234

ABSTRACT

Fracture interfaces formed in silicon nitride at elevated temperatures

were studied by the techniques of light and electron microscopy. The

structure of the fracture interface depended on the type of silicon

nitride fractured. High-purity, reaction-bonded silicon nitride always

formed flat relatively featureless fracture surfaces. Fracture occurred

by a brittle mode even at the highest temperature (1500°C) studied. The

critical stress intensity factor for reaction bonded silicon nitride

3/2

('V'2.2 MN/m ) is relatively low and is insensitive to temperature. By

contrast, hot-pressed silicon nitride gave evidence of plastic flow

during fracture at elevated temperatures. Crack growth in magnesia doped,

hot-pressed silicon nitride occurs by creep, caused by grain boundary

sliding and grain separation in the vicinity of the crack tip. As a

consequence of this behavior, extensive crack branching was observed

along the fracture path. The primary and secondary cracks followed

intergranular paths; sometimes dislocation networks, generated by momentary

crack arrest, were found in grains bordering the crack interface. As a

result of the high temperature, cracks were usually filled with both

amorphous and crystalline oxides that formed during the fracture studies.

Electron microscopy studies of the compressive surfaces of four-point

bend specimens gave evidence of grain deformation at high temperatures

by diffusion and dislocation motion.
The Structure of Slow Crack Interfaces

in Silicon Nitride

N. J. Tighe

' 1 . Introduction

Hot-pressed drid reaction-bonded silicon nitride materials are being

developed for use in gas turbine engines which have service temperatures

as high as 1400°C. The strength and fracture toughness data obtained

for these materials show that most currently available compacts may not

be useful because their strength degrades during exposure to such temperatures [1-7]. In order to improve the high temperature behavior, it is

essential to understand why the strength and fracture toughSnss degrade

during high temperature exposures and why the two types of silicon

nitride behave differently. >

Although the strength oi magnesia doped hot-pressed silicon nitride

(HPSN) is 2-3 times greater than reaction-bonded silicon nitride (RBSN)

at 25°C, the strength of reaction bonded silicon nitride changes little

with exposure to high temperature while that of hot-pressed silicon

nitride, degrades at temperatures above '\^1200°C. At 1400°C the strength

of both materials is 'V'250 MPa. From measurements of crack velocity and

fracture toughness Lange [5], and Evans and Wiederhorn [4] correlated

the strength degradation of magnesia doped hot-pressed silicon nitride

with the onset of slow crack growth. They suggested that slow crack

growth occurred as a result of plastic deformation or stress corrosion

in the neighborhood of the crack tip. Reaction bonded silicon nitride

did not exhibit slow crack growth under similar test conditions.

2
In this paper we characterize fracture and slow crack growth and

the strength degradation in terms of the microstructural changes which

occurred in 4-point bend and double torsion specimens during experiments

at 1400°C. The fracture interfaces and the regions near fracture interfaces in these specimens were examined by light and electron microscopy.

Earlier microscopy studies [8,9] showed that hot-pressed silicon nitride

was composed primarily of blocky a ^i^N^ grains ('^^1 ym diam) and columnar

6 Si^N^ grains ('^^IxS ym) . In addition, it contained micron sized grains

of a and B SiC, and Si2N20 . Furthermore, glass was occasionally found

at triple junctions and as micron sized regions. The microstructure is

therefore heterogeneous with respect to composition. The reaction

bonded silicon nitride [10] was composed of ultra fine a Si^N grains

('^'0.02 nm) and larger B ^-^3^4 grains ('v^l-S ym) . The SiC phases

occurred in both materials because they were hot-pressed or reaction

bonded in a carbonaceous atmosphere. Both materials oxidized at the

test temperature; and, a complex crystalline and amorphous oxide layer

was formed on the surface [11]

2 . Experimental Procedure

2.1 Materials

The hot-pressed and the reaction bonded silicon nitride materials

examined in this study are currently under consideration for turbine

engine applications. Two hot-pressed compositions* (HS130 and NC132) of

magnesia-doped Si^N^ were obtained in the form of billets 75.5 by 75.5

by 2.5 cm. The manufacturer listed the following impurities (wt pet):

0.20 Al, 0.05 Ca, 0.40-0.64 Fe, 0.70-0.74 Mg, 2.6-2.7 W. Two sets of

reaction-bonded silicon nitride specimens (NC350) were obtained as

*The Si N materials were obtainea from tne worton Co., iioston, nass.

3 4
plates 7.6 by 2.5 by 0.2 cm. Test specimens were cut from the billets

and ground into plates 7.6 by 2.5 by 0.2 cm for double-torsion tests and

into bars 5.5 by 1.0 by 5.5 cm for 4 point bend tests. The doubletorsion specimens were grooved along their length by sawing through

approximately 1/2 the plate thickness. The specimens were notched at

one end of the groove.

2.2 Mechanical Testing

Stress intensity factors and crack velocity data were obtained on

the double-torsion specimens using the procedure for propagating a

controlled crack at high temperature which was used by Evans [7]. In

this method, the grooved and pre-cracked double- torsion specimens (Figure

1) were held in a silicon carbide jig with the grooved side up; and, the

load was applied at a constant displacement rate to a ram, which straddled

the central notch at (P) . The load was increased until the critical

stress for crack initiation occurred and the crack propagated along the

-5 -2

groove at constant load. Displacement rates from 5x10 to 5x10 cm/min

were used. The stress intensity factors were calculated from the load

and specimen dimensions. The crack velocity was obtained from measurements of the crack length as a function of time at constant load. Very

low crack velocities were obtained by applying a constant load for fixed

time intervals up to 48 hours. Crack lengths that were measured using a

fluorescent dye penetrant were in good agreement with measurements of

the crack tip displacement after the plate was broken into two parts.

2.3 Microscopy

Transverse and parallel sections of the test specimens were examined

using light and electron microscopy. Specimens were prepared for transmission electron microscopy by standard ion thinning procedures [12]. In

4
Fig. 1. Schematic diagram of a

double-torsion specimen

showing the saw groove

(Wg) and the notch at one

end. The load is applied

at P and the crack is drawn

as a dashed line.

Fig. 2. Strength vs stressing time

for hot-pressed silicon

nitride and reaction bonded

silicon nitride tested in

4-point bending, o hot-pressed

at 1400°C, a reaction bonded

at 1500°C.

TIME HOURS

10

10"

10

-1 HOT PRESSED SI^N^

3 - •

^ UOO'C

1200*C

10

10"

10"

10

10

-7

10

-9

10 -10

Fig. 3. K, V data for silicon nitride

tested in air at 1200°C and

1400°C.

1 3 5 7 9

STRESS INTENSITY FACTOR HN/bVj

5
this method, disks 3 mm in diameter were cut from polished thin plates

50-100 ym thick, using a diamond core drill. These disks were then

polished to a thickness of approximately 50 ym and ion-thinned using

argon ions. Copper single hole grids were cemented onto the specimens

with epoxy cement to strengthen and to support the thin foil. The thin

foils were examined in an electron microscope operating at 200 kV. The

double-torsion specimens and the bend specimens were examined with a

light microscope in order to establish the macroscopic characteristics

of the deformation, and to compare the electron microscopy results with

these features. '

' - '

' 3. Results

3.1 Mechanical Properties

The hot-pressed silicon nitride and reaction-bonded silicon nitride,

behaved differently during the mechanical tests. In the A-point bend

tests, hot-pressed silicon nitride deformed plastically and exhibited

secondary creep at 1400°C; whereas, reaction-bonded silicon nitride

_2 deformed elastically to failure at approximately 212 MNm at 1500°C.

Curves showing the calculated stress vs time for these two materials are

shown in Figure 2. The load was applied at the displacement rate of

8.5x10"^ m/s.

The stress intensity factor obtained for hot-pressed silicon

nitride showed distinct temperature dependence. The stress intensity vs

crack velocity results obtained for NC 132 at 1400°C and at 1200°C are

plotted in Fig 3. The variation in with crack velocity shows that

extensive slow crack growth occurred at 1400°C . At 1200°C slow crack

6
Fig. 4. Light micrographs of the

tension sides of doubletorsion specimens tested at

1400°C: (a, b) hot-pressed

silicon nitride, (c) reaction bonded silicon nitride,

Cracks propagated in the

direction of the arrow.

inn m

Fig. 5. Light micrograph of a hotpressed silicon nitride

specimen deformed in 4-point

bending at 1400°C: (a)

polished section of specimen,

(b) details around crack tips(X]

7
growth was minimal as indicated by the change in slope of the K-V plot.

Reaction-bonded silicon nitride did not exhibit slow crack growth at

3/2 1400°C and at this temperature was '^'?..2 MN/m over the

range of loading rates used for the tests.

3.2 Fracture Interfaces

3.2.1 Light Microscopy

The hot-pressed silicon nitride double-torsion specimens bent

plastically during crack growth experiments at lAOCC. This deformation

and subsequent oxidation of the surfaces kept the crack interfaces

separated along the tension surface after the stresses were removed,

therefore, the path of the primary crack was outlined by oxide and was

easy to see. In Fig. 4a, it is seen that the primary crack oscillated

along the groove and, at successive intervals, the tips of the crack

segments overlapped. The oscillations occurred as the crack propagated

and were not related to the crack arrest at the end of each experiment.

In specimens, which were broken into 2 pieces the overlapping crack tips

were torn apart as illustrated in Fig. 4b.

In contrast with the hot-pressed material the reaction bonded

silicon nitride specimens did not deform at temperatures as high as

1500°C. Slow crack growth did not occur and the fracture interfaces

were nearly featureless. Some disruption in the surface occurred at the

crack origins as at x in Fig. 4c.

Further evidence for the tearing apart of the structure of the hotpressed silicon nitride during deformation was found by examining 4-

point bend specimens after polishing away the surface oxide. The unnotched

specimen shown in Figs. 5a and b was bent uniformly across the gage length.

8
Fig. 6. Electron micrograph of Fig. 7. Secondary cracks arrayed

a section from double- approximately parallel to

torsion specimen showing: the primary crack,

oxide along the primary

crack.

9
Cracks which grew along the machined edges of the tension side of the

specimen penetrated up to 1mm deep. However, near the middle of the

tension surface of the bar, cracks were only 0.1 mm deep thus demonstrating the strong corner effect of the crack nucleation. As seen in the

figure the surface and sub-surface cracks were filled with oxide which

enhanced the crack visibility in polished sections.

It is clear that the crack branching characteristics of slow cracks

in hot-pressed silicon nitride and the smooth interfaces of rapid cracks

in reaction bonded silicon nitride can be distinguished by light microscopic examination of the specimens. However, the grain size in both

materials is small and cannot be resolved by light microscopy. Therefore,

the relations between the cracks and the microstructure can be determined

only by electron microscopy.

3.2.2 Electron Microscopic Examination

Areas around the primary cracks in hot-pressed silicon nitride

specimens tested at 1400°C were examined by electron microscopy using

thin foils made from sections cut parallel to the length of the cracks.

As expected from the light microscopy, the primary crack shown in Fig. 6

had oxide along its interfaces. Near the primary crack tips many intergranular cracks several micrometers long were found which were arranged

approximately parallel to the primary crack. From the array shown in

Fig. 7 it is apparent that the compact was pulled apart along grain

boundaries. The fine structure of cracks is similar to the crack

branching macrostructure seen in the light micrographs.

. The oxides on the crack surfaces can form by oxidation of the crack

interfaces during fracture. Further the cracks which are held open

10
Fig. 9. Tip of grain boundary crack

in Fig. 8a showing dislocation

arrays extending into adjacent

grains

.

g. 10. Reaction bonded silicon nitride showing: (a) fine, grains of aSigNt, and large grains of 3-Si3Ni+ and 3-SiC, (b) a grains.

11
because of the deformation can be filled by migration of the oxide

formed on the specimen surfaces, and from the silicate phase in the

matrix. Examples of cracks which have oxide along the interfaces are

seen in Fig. 8. An amorphous phase was found also in pores and in open

cracks of this specimen. These amorphous phases had electron diffraction

patterns similar to the patterns taken of Si02 in other oxidized silicon

nitride specimens [11].

When the grain boundary crack shown in Fig. 8a was followed to its

tip, dislocation arrays were found in the grains adjacent to the crack

tip (Fig. 9). It appears that this crack propagated along the grain

boundaries and, when it arrested, the strain release produced dislocation

pile-ups on both sides of the crack tip. In grain A, the dislocations

have formed in a network approximately parallel to the basal plane; and,

their presence indicates that plastic deformation occurred within these

grains at 1400°C.

The microstructure of the reaction-bonded material as shown in Fig.

10 was considerably different from the microstructure of the hot-pressed

material. The structure was bi-modal with a fine-grained a-phase 10-20 nm

in diameter, and a 3-phase. Some large grains of silicon carbide were

found also; and, these grains and the 3-silicon nitride grains usually

were faulted. The images of the small a-grains in Fig. 11 overlapped

and gave a very confused picture; however, individual grains can be

distinguished by the presence of moir^ fringes from overlapping grains.

There was little difference in microstructure between as-received and

tested specimens except at the surface where a thin layer of 3-cristobalite

and glass had formed during the experiments. This observation suggests

that the microstructure of the reaction bonded material was not affected

by the stress.

12
Fig. n. Regions near compression side of specimen bent at 1400°C:

(a and b) dislocations and faults interacting at grain boundaries,

(c) array of pores forming "ghost" boundary, (d) recrystal 1 ized

grains.

13
In the hot-pressed specimens which were bent at 1400°C, grains near

the tension side were pulled apart at grain boundaries in a manner

similar to grains in the double-torsion specimens. Extensive fragmentation occurred near the silicon nitride-oxide interfaces and the cracks

near this interface were filled with oxide. Thin foil specimens which

were made from the tension side of the bend bars were extremely fragile

and some areas were lost in handling.

In specimens taken from the compression side of the bend bar it was

found that most grains had extensive dislocation arrays which were

associated with grain boundaries as in Fig. 11. When stereo-pairs were

taken of the area in Fig. 11a it appeared that the dislocations left the

grain boundary, interacted with small particles, and then broke up into

loops. Since these dislocation interactions were not found in asreceived specimens, this dislocation motion must have occurred in the

grains as a result of bending. Extensive arrays of bubbles near grain

boundaries also were found only in foils made from the compression side

of the bent bars (Fig. lib) . These bubbles were found in undulating

arrays which were displaced slightly from the grain boundaries in some

grains and in the center of other grains. Such arrays are suggestive of

diffusion along the grain boundary or movement of the grain boundary

during compression. Such bubble formation was not found in the silicate

regions shown in Figs. 6 and 8. The bubbles were observed in most

grains and could not be associated with any degradation resulting from

electron beam irradiation. Clusters of small defect-free grains, as

shown in Fig. lid, also were found in foils made from the compression

side. These microstructural changes are manifestations of plastic

deformation, diffusion recrystallization and grain growth, which occurred

on the compression side of bars bent at 1400°C.

14
4 . Discussion

The microscopic analysis presented above suggests that the differences

in fracture toughness between the hot-pressed silicon nitride and the

reaction bonded silicon nitride can be related to the differences in

microstructure between these two materials. The MgO and CaO based

additives promoted sintering and densif ication during hot pressing at

temperatures near lyOO^C. As a result, the grain size distribution is

more homogeneous but more crystalline and amorphous second pores are

present in hot pressed material than in reaction bonded silicon nitride.

The reaction bonded material did not have additives and was porous and

had a bi-modal grain size distribution. A crack in reaction bonded

silicon nitride traversed the weakly sintered matrix easily because it

encountered only ultrafine grains, pores and large faulted grains. A

crack in hot-pressed silicon nitride encountered well-sintered grains

and propagated intergranularly indicating that grain boundaries presented

easy but more tortuous fracture paths. The micrographs taken from

specimens fractured at 1400°C showed that slow crack growth was enhanced

by plastic deformation in regions surrounding the crack tip and that the

grain boundaries were pulled apart over a considerable distance. The

dislocations that were observed along grain boundaries and within grains

in MgO doped hot-pressed silicon nitride are evidence for some localized

plastic deformation. Oxides were observed along fracture interfaces and

along some grain boundaries in the neighborhood of the cracks and could

be evidence for grain boundary sliding.

The strength degradation that is observed in hot-pressed silicon

nitride has been attributed to grain boundary sliding involving shear

deformation of a thick ('^100 nm) glass phase between grains [5, 6].

15
This type of grain boundary sliding could occur without plastic deformation

of the grain. Although micro-probe analysis by Kossowsky [13] showed

that magnesium and calcium segregated near grain boundaries a significant

amount of glass phase (> 3 nm thickness) has not been found by electron

microscopy at grain boundaries in undeformed commercial materials. In

most of the micrographs shown in this paper, the silicon nitride grains

were sintered together and an amorphous phase could not be resolved

between grain boundaries before fracture. Glass phases were detected at

triple junctions and as distinct separate regions several hundred nanometers

in diameter. Recently, Clarke and Thomas [14], using high resolution

electron microscopy obtained lattice images of grains and detected a

phase "^1/2 nm thick between some grains in MgO doped hot-pressed silicon

nitride and at triple junctions.

The contribution of the crystalline or amorphous silicate phases

to high temperature deformation can be illustrated by considering that

at 1400°C, magnesium silicate has the viscosity of water; and, although

pure silica glass has a higher viscosity, many of the possible Mg and Ca

silicate phases can melt and thus be very mobile at such temperatures.

With a liquid phase present, grain boundary sliding could cause the

dislocation arrays observed here by a mechanism akin to abrasion and

wear on lubricated surfaces. Thus grain boundary sliding does not

necessarily require a 100 nm thick phase between grains when the magnesium

based sintering aid can melt or otherwise decompose and when deformation

can occur within the grains and along grain boundaries. Grains under a

tensile or shear stress can simply pull apart and grains under compressive

stress can accommodate the strain by diffusional creep. The bubbles and

dislocation arrays observed near grain boundaries support this concept

16
of localized deformation in the bent specimens. Rapid diffusion of

impurities at 1400°C is evidenced by the presence of Mg and Ca in the

outer oxide layer. The exact composition of the oxide phase is a

function of the exposure time [11].

The reaction-bonded silicon nitride did not deform plastically in

the present experiments and analysis of this deformation mode in reaction bonded silicon nitride could not be made. Thin foil specimens

made from a series of extensively deformed samples of both hot-pressed

and reaction bonded silicon nitride must be examined before the deformation mode can be defined completely. Silicon nitride dissociates

without melting, and oxidizes or reduces depending on the surrounding

atmosphere. It will be necessary, therefore, to do tests in vacuum and

nitrogen as well as in air.

5 . Conclusions

By examining specimens from double-torsion, fracture mechanics

tests and from 4-point bend tests, we have shown that in hot-pressed

silicon nitride: (1) extensive crack branching occurred in specimens

fractured or bent at 1400°C; (2) slow crack growth observed in hightemperature tests resulted from plastic deformation; (3) this plastic

deformation occurred by grain boundary sliding which involved both slip

along grain boundaries and within grains, and separation of grains

through reaction with the silicate phase. Reaction bonded silicon

nitride deformed elastically under the same conditions. The difference

in high-temperature strength and fracture toughness between hot-pressed

and reaction-sintered silicon nitride was related to compositional and

microstructural differences between the materials.

Acknowledgment

The author greatfully acknowledges the support of the Air Force

Materials Laboratory.

17
References

[I] A. G. Evans and R. W. Davidge, J. Mat. Sci. 5, (1970) 314-325.

[2] W. Ashcroft, in Ceramics for Turbines and Other High-Temperature

Engineering Applications , Br. Ceram. Soc, Stoke-on-Trent (1973),

pp. 169-179.

[3] M. L. Tortl, R. A. Alllegro, D. W. Richerson, M. E. Washburn and

G. Q. Weaver, in Ceramics for Turbines and Other High-Temperature

Engineering Applications , Br. Ceram. Soc, Stoke-on-Trent (1973),

pp. 129-146.

[4] A. G. Evans and S. M. Wiederhorn, J. Mat. Sci. 9_, (1974) 270-278.

[5] F. F. Lange, J. Am. Ceram. Soc. 57, (1974) 84-89.

[6] R. Kossowsky, D. G. Miller, and E. S. Diaz, J. Mat. Sci. 10, (1975)

983-997.

[7] A. G. Evans, in Ceramics for Hlgh-Performance Applications , ed.:

J. J. Burke, A. E. Gorum and R. N. Katz, Brook Hill, Mass. (1974)

pp. 373-396.

[8] P. Drew and M. H. Lewis, J. Mat. Sic. % (1974) pp. 261-269.

[9] N. J. Tighe, in "Nitrogen Ceramics," F. Riley Ed., (in press).

[10] A. E. Evans and J. V. Sharp, J. Mat. Sci. 6^, (1971) 1292-1302.

[II] N. J. Tighe, Proc. Electron Microscopy Society of American, Vol. 32

Claitors Publ. Co., Baton Rouge, La. 1974 pp 470-471.

[12] N. J. Tighe, in "Electron Microscopy in Mineralogy", H.-R. Wenk Ed.,

Springer-Verlag. Berlin, 1976, pp. 144-171.

[13] R. Kossowsky, J. Mat. Sci. 8, (1973) 1603-1615.

[14] D. R. Clarke and G. Thomas, J. Amer. Ceram. Soc. (in press).

18
References

A. G. Evans and R. W. Davidge, "The strength and oxidation of

reaction sintered silicon nitride", J. Mat. Sci. 5^, 314-325

(1970).

W. Ashcroft, "Mechanical properties of silicon nitride at elevated

temperatures," in Ceramics for Turbines and Other High-Temperature

Engineering Applications , Br. Ceram. Soc, Stoke-on-Trent (1973),

pp. 169-179.

M. L. Torti, R. A. Alliegro, D. W. Richerson, M. E. Washburn and

G. Q. Weaver, "Silicon nitride and silicon carbide for hightemperature engineering applications," in Ceramics for Turbines

and Other High-Temperature Engineering Applications , Br. Ceram. Soc.

,

Stoke-on-Trent (1973), pp. 129-146.

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).

F. F. Lange, "High- temperature strength behavior of hot-pressed

Si^N^: evidence for sub-critical crack growth," J. Am. Ceram. Soc.

57_, 84-89 (1974).

R. Kossowsky, D. G. Miller, and E. S. Diaz, "Tensile and creep

strengths of hot-pressed Si^N^," J. Mat. Sci. 10, 983-997 (1975).

A. G. Evans, "High-Temperature slow crack growth in ceramic materials,"

in Ceramics for High-Performance Applications , ed. : J. J. Burke,

A. E. Gorum and R. N. Katz, Brook Hill, Mass. (1974), pp. 373-396.

P. Drew and M. H. Lewis, "The microstructure of silicon nitride ceramics

during hot-pressing transformations," J. Mat. Sci. 9, 261-269 (1974).

19
[9] N. J. Tighe, "Mlcrostructural Aspects of Deformation and Oxidation

of Magnesia Doped Silicon Nitride," in Nitrogen Ceramics, F. Riley

Ed. , (in press)

.

riOl A. E. Evans and J. V. Shart>. "Microstructural studies on silicon nitride,"

J. Mat. Sci. 6, 1292-1302 (1971).

[11] N. J. Tighe, "Microstructure of oxidized silicon nitride;" Proc. Electron

Microscopy Society of America, Vol. 32 Claitors Publ. Co., Baton Rouge,

La. 1974 p. 470-471.

[12] N. J. Tighe, "Experimental techniques", in "Electron Microscopy in

Mineralogy", H.-R. Wenk Ed., Springer-Verlag. Berlin, 1976, p. 144-171.

[13] R. Kossowsky, "The microstructure of hot-pressed silicon nitride,"

J. Mat. Sci. 8, 1603-1615 (1973).

[14] D. R. Clarke and G. Thomas, "Grain boundary phases in a hot-pressed

MgO) silicon nitride," J. Amer. Ceram. Soc. (in press)

.

20
ImBS-IUA (REV. 7-73)

U.S. DEPT. OF COMM.

BIBLIOORAPniC DATA

SHEET

1. PUBLICATION OR REPORT NO.

NBSIR 76-XQ75

2. Gov't Accession

No.

3. Recipient's Accession No.

' 4. TITLE AND SUBTITLE

THE STRUCTURE OF SLOW CRACK INTERPACES IN SILICON NITRIDE

5. Publication Date May 1976

6. Performing Organization Code

7. AUTHOR(S)

N. J. Tiehe

0. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAN4E AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No. 313Q459

11. Contract/Grant No.

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

Department of the Air Force

Headquarters Aeronatucial Systems Division (AFSC)

jj Wright-Patterson Air Force Base, Ohio 45433

13. Type of Report & Period

Covered

Interim, 1/I-3/31/76

14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

16. ABSTRACT (A 200-word or less factual suwmary ol most si^iticant information. If document includes a significant

bibliography or literature survey, mention it here.)

Cracks in hot-pressed silicon nitride (HPSN) and in reaction bonded silicon

nitride were studied by light and electron microscopy. Slow crack growth,

resulting from plastic deformation, occurred in HPSN at 1400°C. However, the

RBSN exhibited brittle fracture at this temperature. This difference in

behavior can be explained by the effects of compositional and microstructural

differences between the two materials. The MgO additions to HPSN promoted

sintering but its presence contributed to plastic deformation. The RBSN was

made without added phases. It does not sinter as completely as the HPSN and

the grain size is an order of magnitude smaller. The small grained regions

offered no resistence to fracture.

17. KEY VORDS (six to twelve entries; alphabetical order; capitalize only the first letter of Ae first key word unless a proper name; separated by aemicolons) ^ , , ^ , _ Cracks; electron microscopy; fracture interfaces; plastic

deformation; silicon nitride; turbine materials

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