NBSIR 75-658

Electron Microscopic Observations

of l\llicrocracl(ing About Indentations

in Aluminium Oxide and Silicon Carbide

B. J. Hockey and B. R. Lawn

Inorganic Materials Division

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

January 1975

Interim Report for Period July 1, 1974 through June 30, 1975

Prepared for Departtrent of the Navy

Office of Naval Research

Arlington, Virginia 22217

NBSIR 75-658

ELECTRON MICROSCOPIC OBSERVATIONS

OF MICROCRACKING ABOUT INDENTATIONS

IN ALUMINIUM OXIDE AND SILICON CARBIDE

B. J. Hockey and B. R. Lawn

Inorganic Materials Division

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

January 1975

Interim Report for Period July 1, 1974 through June 30, 1975

Prepared for Department of the Navy

Office of Naval Research

Arlington, Virginia 22217

U. S. DEPARTMENT OF COMMERCE. Frederick B. Dent. Seaetary

NATIONAL BUREAU OF STANDARDS, Richard W. Roberts. Director

1. Introduction

Indentation hardness testing is finding an increased usage in the

scientific investigation of damage processes in highly brittle solids [1]

.

A suitably sharp indenter tip concentrates very high levels of stress,

particularly in shear and hydrostatic compression, and thus induces

irreversible deformation in conveniently localized regions of a test

surface. In many brittle materials the mean contact pressure can

be as much as a tenth of an elastic modulus [2] , indicating that

the intrinsic bond strength of the lattice itself must surely be

exceeded places. Despite a recent proliferation of indentation

damage studies, the exact nature of such irreversible deformation

processes in brittle solids remains something of a contentious issue.

One aspect of indentation damage which has been largely overlooked

is that of indentation- induced cracking. In the general indentation field

a component of tension, however small, is unavoidable [3], and there

is a growing body of evidence suggesting that cracking may be more

prevalent than one might have been led to believe from earlier studies.

An understanding of indentation fracture is important for many practical

as well as academic reasons (for a review see [4]): a pertinent

example is strength degradation as a result of particle-surface contact,

where individual damage events may be usefully simulated in a standard

hardness testing device.
- 2 -

The chief object of this work is to report on observations of

microcrack patterns about small-scale point indentations in aliominiiixn

oxide and silicon carbide, two highly brittle materials currently of

great interest in ceramics engineering. The results represent part of

a detailed study of residual crack interfaces using transmission electron

microscopy. Our emphasis here is directed to geometrical features of

the observed crack interfaces. The observed crack geometry is discussed

in terms of an earlier proposal for crack growth beneath sharp indenters

based on optical observations [3]. The study complements an analysis

of diffraction contrast effects at large-scale remnant cone cracks in

sphere- indented silicon using X-ray topography [5]

.

2. Preparation of Indentation Specimens for Electron Microscopy

The method of specimen preparation was similar to that previously

described [6-8] . Single crystals of sapphire and a-silicon carbide

were cut into slabs and polished, first mechanically and then chemically,

to a thickness of about 100 ym. Indentations were made with either a

Knoop or Vickers diamond pyramid at a load of 200 g on each test surface.

An optical examination of the surfaces showed clear evidence of microcracking in most cases, especially about the Vickers indentations. The

crystals were then thinned to a size suitable for 200 kV transmission

electron microscopy by ion bombardment [7] . Most of the material removed

was from the surface below the indentation, although a little («lym) was

removed from the top surface as well to eliminate any spurious damage

incurred in specimen handling.

The bulk of the foils examined were prepared at room temperature.

However, a few were subjected to prescribed theinnal or mechanical
-3--

treatments between the indentation and thinning stages, as described in

[8]. About one-hundred indentations in sapphire, and about twenty-five

in silicon carbide, were investigated in this way.

3. Electron Microscopy of Cracks

3 . 1 General crack morphology

Intense diffraction contrast was observed in the intmediate vicinity

of all impressions. At the edges of the impressions, the images of

dislocations, and sometimes of twins as well, could be clearly resolved.

An analysis of this deformation has been given in the case of sapphire

in an earlier paper [7] , and similar deformation patterns have recently

been analysed in indented silicon [9,10]. The diffraction spot pattern

was maintained at all locations in the foils , although considerable

broadening of the spots was evident in the most heavily deformed regions.

No diffraction evidence was found for any crystallographic phase tra,ns formation

resulting from the indentation process.

Cracking was observed about all Vickers impressions examined, and

about most Knoop impressions. With the Vickers specimens, cracks were

always observed to extend radially outward from the corners of the

indentations, their plane oriented very nearly normal to the foil surfaces,

as seen clearly in Fig. 1 (silicon carbide) . In addition, cracks lying

nearly parallel to the foil surfaces were often observed between the

radial cracks, again as in Fig. 1, but the geometry and indeed the

frequency of appearance of this type of crack were sporadic In the ca,se

of Knoop specimens, for example Fig. 2 (sapphire), the incidence of
-4-

detectable radial cracking was considerably less than for Vickers

specimens , but for the lateral type of cracking the incidence was almost

the same

.

These observations may be readily explained in terms of the scheme

for crack growth beneath sharp indenters put forward by Lawn and Swain

[3]. Basically, two distinct stages of crack formation are apparent.

The first occurs on indenter loading; at some (low) threshold level of

stress a crack initiates at the deformation zone immediately below the

sharp point of the indenter and extends downward on a plane of symmetry

containing the contact axis. Fig. 3 depicts schematically the shape

of this crack, the so-called "median vent crack," in its well-developed

form. Diagram (a) indicates how the thinned foil samples the crack about

the deformation zone. The median vents accordingly appear as ribbon-like

segments along the lines of greatest stress concentration, as depicted in

(b) for the Vickers and (c) for the Knoop specimens respectively. It is

clear that the shape of the remnant crack will not be too sensitive to

the exact location of the foil with respect to the original indented

surface. The second stage of crack formation occurs on unloading ; just

prior to removal of the indenter new cracks initiate at the deformation

zone and extend sideways into a shallow saucer-shaped configuration, modified

some extent by the presence of the median vents. This crack, the

"lateral vent", is depicted in Fig. 4. In this case the location of the

foil becomes a critical factor; small changes in the relative amounts

of material removed from the upper and lower surfaces of the indented
-5-

crystal during thinning might result in the loss of a portion, if not all,

of the lateral vent system. One further, very important consideration

in connection with the geometry of remnant lateral vents is the axis

of loading; in practice it is no easy matter to ensure near-perpendicularity

between indenter axis and specimen surface, and preferred cracking on one

side of the indentation tends to be the rule rather than the exception.

In Fig. 5 we indicate how such asymmetry accounts for the apparently

complex crack morphology seen in the electron micrograph of Fig. 2.*

3.2 Inter facial fringe and dislocation contrast

An examination of the diffraction contrast of the remnant cracks

provides quantitative evidence for residual lattice mismatch across the

interfaces. Essentially, the crack images are typified by fringe and

dislocation networks. Figs. 6 and 7, showing median vents about Vickers

indents in silicon carbide and sapphire respectively, and Pigs. 8 and 9,

showing lateral vents about a Vickers indent in silicon carbide and a Knoop

indent in sapphire respectively, are examples. The patterns in all cases

are largely characteristic of mismatch contrast [11, 12] , generated by

interference between slightly mismatched, simultaneously diffracting

portions of crystal across the crack interfaces, although there is some

indication of thickness fringe modulation where the interface is inclined

to the foil surface (e.g. Fig. 7). This diagnosis was confirmed by

observing the spacing of the fringes to remain invariant with voltage

of the electron beam in the microscope (e.g. by lowering from 200 kV to

100 kV) [11]

.

*A distinct example of such asymmetry in crack pattern due to inclined loading

of a Knoop indenter on a quartz surface may be seen in Fig. 7 of Ref . [3]

.
-6-

Fringe patterns of the mismatch type associated with residual cavities

or cracks may be categorised into either displacement or moire systems

[12]. The displacement system is generated when the opposing crystal

portions are separated by a rigid-body displacement £(r) , where r is a

position vector contained within the interface and measured from the

crack tip: if the diffraction vector (i.e. reciprocal vector defining

orientation and spacing of lattice planes) in both crystallites is g the

fringes are loci of

' g . c = N (1)

where N is the order of the fringe [12]. Thus one characteristic of the

displacement fringe system is that the observed pattern should remain

geometrically similar for all reflections. In Figs. 6 and 8, where the

effects of systematically varying g are investigated, the geometry of the

fringe pattern is seen to change markedly. Moreover, the reflections in

Fig. 8 are such that we would have g . c «0 if the mutual crack-wall

displacements were, as anticipated, to be closely normal to the plane of

the interface (hence of the foil) . We must generally conclude that the

displacement fringe system can be of no more than secondary ' importance

in the micrographs observed in this work.

The moire fringe system, on the other hand, is generated when the

overlapping crystal portions have a small mismatch in lattice periodicity

such that there is a small difference in diffraction vector,

6g = g^ -g ^ 0. The fringes are then loci of [13, 14]

6g . r = const. (2)
Then since 6g will generally vary with g, we would expect such a

pattern to be reflection dependent, consistent with the observations of

Figs. 6 and 8. This diagnosis in favour of the moire system was also

reached in the earlier X-ray analysis of remnant cone cracks in silicon

[5] . In the event that the opposing crack walls are able to recontact

and heal, interfacial relaxation of the mismatched crystal portions results

in a dislocation network [8, 15]. Examples of spontaneous crack healing

are seen in Figs. 7 (near-tip region) and 9.

It is of interest to determine the degree of mismatch needed to

explain the network patterns seen in the micrographs. Characterising

the fringe system associated with the non-healed crack interfaces by

the moire" vector (i.e. reciprocal vector defining orientation and spacing

of fringe system) [13]

G = 6g = g - g , (3)

we distinguish between two basic moire configurations: parallel moires,

in which g, and g„ differ slightly in magnitude (G parallel to g)

,

and rotation moires, in which g and g differ slightly in direction ~1 ~2

(G perpendicular to g) . In most cases examined here the system appeared

to be predominantly of the second type; note in particular the comparative

absence of fringe contrast in Fig. 8c, corresponding to G^ having a large component

parallel to g. Such a lattice-plane rotation is in fact consistent with

a simple shear displacement field u = u (r)z (z being a unit vector ~ z ~ ~

parallel to the crack front) for the crack walls at the interface (i.e.

g

as the spacing of the moire fringes, we obtain

"mode III" field). Writing d = -1

as the spacing of the lattice planes

and D =

D ^ 1 (rotation) , ... d £
-8-

with e the angular misorientation of the diffracting planes.

Typically, we have ^-^-^20 ~ ^'-^^ ^

(a-SiC) , Dw 0.1 ym (Fig. 8), from

-3

which we compute a "lattice mismatch" d/D w 1 x 10 ; this corresponds

to a mutual lattice-plane misorientation e « u (r)/^ 1 ^ 10 rad.

^ z Our treatment here is necessarily oversimplistic : a more accurate

account of fringe pattern details would require a rigorous diffraction

contrast analysis, taking into account such complications as inclinations

of the interface to the foil surface [14] , superposition of components

u (r)y and u (r)x (y and x being unit vectors parallel to crack normal and y^^ x~~~

crack direction respectively) onto the residual displacement field

(i.e. adding "modes I and II" to "mode III"), etc., in the general

reflection situation. . .. ••

3.3 Other interfacial features in the crack images

Some other contrast features were evident at the residual crack

interfaces. The crack front itself usually showed up by enhanced diffraction contrast, implying residual crack-tip strain due to incomplete

closure of th- interface. This residual-strain effect was mere evident

in some cases (e.g. Fig. 8) than others (e.g. Fig. 7); in general,

those interfaces indicating some tendency to healing showed the least

crack-tip contrast. In addition, the contrast visibility increased in

intensity from zero as the g vector was rotated away from alignment

with the crack front, in accordance with a state of plane strain at the

tip: a particularly clear example of this contrast variation is shown

in Fig. 10. .
-9-

The existence of a residual crack opening points to the operation of

some closure-prevention mechanism. While there was ample indication of

dislocation generation about the central, grossly deformed regions of

the indentations, in no case was there found any evidence for dislocation

activity in association with crack growth, thus ruling out plastic flow

as a significant factor.* The present observations merely reinforce

similar conclusions previously drawn concerning the reversibility of

fracture in highly brittle solids [5, 8, 15]. Closure resistance may

be identified with purely mechanical obstruction, predominantly from

cleavage steps at the fracture surfaces [5] . These are seen clearly in many

of the micrographs, notably in Figs. 8 and 9. A close study of the mechanics

of formation of such steps [16] indicates a process in which crack

segments propagating on closely adjacent planes first overlap and subsequently link up tip- to-plane to effect separation. Segmentation of an

initially planar crack into an array of partial fronts occurs when a

"twist-mode" disturbance is suddenly encountered in the crack propagation

field [17]. Fig, 11 depicts the phenomenon schematically, and illustrates

via the fringe pattern in the accompanying micrograph the significant

influence the stepped regions have on the residual crack opening. The

*The healing dislocations do not, of course, form via a glide process,

but rather as a direct consequence of crack closure; they could

therefore hardly contribute to a residual crack opening.
-10-

smallest of lateral displacements across the crack interface, corresponding

to the lattice mismatch of order one part in a thousand typified by the

fringe spacings, could be quite sufficient to prevent the opposing walls

from keying together in perfect registry upon removal of the indenter

-load. Moreover, these same lateral displacements would account for a

predominance of a mode III type configuration in the moire system.

One further detail which was observed in isolated cases was the

trace of a slightly retracted crack front. This is especially noticeable

in Fig. 10, where protuberant cleavage-step damage beyond the residual

interfacial opening gives some indication as to how far the crack front

must once have extended. The exact source of this detail is not clear,

but its contrast parallels that of dislocations and thus probably arises

from residual displacements on the atomic scale.

3.4 Cracks subjected to special treatments

Some of the indented sapphire specimens were siibjected to a mechanical

shock treatment prior to thinning. This was done simply by delivering

a small impulse, insufficient to rupture the specimen, to the surface

opposite to that containing the indentations. The chief result of this

treatment was to re-propagate the existing median vent cracks (the

lateral vents suffering effectively zero tension in the flexural impulse)

,

to anything up to an order of magnitude increase in length, rather than

to change the nature of the patterns. These extended crack portions also

showed some tendency to partial spontaneous healing [8],
-11-

Other specimens were given an anneal treatment. In this case

diffusion and sintering processes led to a marked enhancement in the

healing [8] , Fig. 12 shows a remnant interface at which the delineation

between restored and non-restored interfacial regions is distinct.

Elongated "pipes" are seen to form along lines of greatest residual

mismatch, namely along network dislocations, but these tend ultimately to

break up into small polyhedral voids in the regions of more prolonged

sintering away from the open interface.

4. Discussion

The observations described here provide useful semi-quantitative

information on the geometry of residual cracks about micro- indentations

.

It should be emphasised, however, that this information relates only to

the history of the propagation, and not the initiation , of the cracks.

That is to say, nothing in the electron micrographs revealed any

indication as to how the cracks originally nucleated and formed. It is

known from earlier section- and-etch examinations [3] that both median-^ •

and lateral-vent initiation is tied up intimately with events within

the gross deformation zone immediately surrounding the indenter;

unfortunately, this is the region of lost resolution and diffractionpattern blurring (Sect. 3.1), and, while it is clear that the theoretical

strength limits of the lattice are undoubtedly approached (at least in

the most brittle of solids) , the nature of indentation- induced deformation

remains obscure [9, 10, 18, 19]. This is one aspect of the indentation

problem v^ich calls for a good deal more attention.
-12-

Although we have given explicit attention here to only sapphire and

a-silicon carbide, the crack patterns described appear to be quite typical

of highly brittle solids in general. In particular, the vent-crack systems

show the same broad geometrical features evident in such structurally diverse

materials as monocrystalline silicon and amorphous silicates [3, 20], thus

emphasising the dominant role of the indentation stress field in determining

the fracture paths. Nevertheless, preferred cleavage tendencies do

usually have some effect on indentation-cracking geometry [4, 21], and

crystallographic tendencies were indeed noted in the present observations.

In this respect the materials studied in the present work provide an interesting contrast in that a-silicon carbide has a distinct tendency to basal

cleavage while sapphire does not, the reason being that although (0001)

may represent the plane intersecting the least number of bonds in the

hexagonal structure a relatively large component of polarity in the

bonding in the case of sapphire renders this plane unfavourable on the

grounds of electrostatic attraction effects [22] . One would therefore

anticipate a greater incidence of lateral venting nearly parallel to (0001)

foils in the fon.j^x material; this trend was in fact apparent in the

micrographs (note extensive crack plane in Figs. 8 and 10, not generally

seen in sapphire) . This raises the question of anisotropy in the

microfracture, a factor which has already been recognised as one of vital

importance in the interpretation of deformation processes in hardness

testing [23]. One striking manifestation of anisotropy in a (cumulative)

indentation chipping process occurs in the abrasion of diamond; by changing
-13-

the abrasion direction on a given surface the wear rate may alter by as

much as three orders of magnitude [24] . Anisotropy is therefore a second

topic which warrants further consideration in relation to indentation

fracture.

The present results also bear on the process of crack healing, and

on the agencies which tend to obstruct it. In those special instances

where spontaneous healing was observed, evidence of cleavage steps and

"debris" at the interface was conspicuously absent [15]. (It is possible,

of course, that the very act of thinning the foil removes, at least

partially, some of these obstructions.) Here again the two materials

examined showed contrasting behaviour, for the only sighting of a spontaneously healed area in silicon carbide was that shown in Fig. 10.

The nature of the bonding thus appears to manifest itself in the mechanics

of the closure as well as of the opening of the cracks; the explanation

is once more to be sought in the long-range electrostatic attractive

forces that tend to develop across fracture interfaces in ionic structures

[25, 26]. The second type of healing reported here, that arising from mass

transport at elevated temperatures, occurs in both materials. On the other

hand, while indentation- induced cracks can be made to heal, so can they be

made to extend still further, as in the event of mechanical shock.

The existence of residual cracks of this sort has obvious implications

in the strength of structural ceramics: an incompletely restored lattice

across any separation interface represents a potential source of weakness

in the material. A more detailed investigation into electron microscopic
-14-

images of fracture interfaces is accordingly under way, and will be

reported at a later date [27]. For the present, it is sufficient to

point out that the strengths of brittle materials can be seriously

degraded by the most minute of contact events , and that this degradation

may vary according to sxibsequent mechanical and thermal history.

Acknowledgements

The authors are indebted to S. M. Wiederhorn and M. V. Swain for

valuable discussions on the work. They also acknowledge the sponsorship

by the Office of Naval Research under Contract No. NR-032-535.
-15-

References

1. "The Science of Hardness Testing and its Research Applications,"

Symposiiom Proceedings, Eds. J. H. Westbrook and H. Conrad (American

Society for Metals, 1973).

2. A. Kelly, "Strong Solids" (Clarendon, Oxford, 1966), Ch. 1.

3. B. R. Lawn and M. V. Swain, J. Mater. Sci . (in press)

.

4. B. R. Lawn and T. R. Wilshaw, J. Mater. Sci . (in press).

5. J. S. Williams, B. R. Lawn and M. V. Swain, Phys. Stat. Sol, (a)

2 (1970) 7.

6. B. J. Hockey, Ref. [1], Ch. 3.

7. B. J. Hockey, J. Amer. Ceram. Soc. 54 (1971) 223.

8. S. M. Wiederhorn, B. J. Hockey and D. E. Roberts, Phil. Mag. 28 (1973)

783.

9. V. G. Eremenko and V. I. Nikitenko, Phys. Stat. Sol, (a ) 14 (1972) 317.

10. M. J. Hill and D. J. Rowcliffe, J. Mater. Sci. 9(1974) lSfi9.

11. P. B. Hirsch,' A. Howie, R. B. Nicholson, D. W. Pashley and M. J.

Whelan, "Electron Microscopy of Thin Crystals" (Plenum, New York,

1965)

.

12. U. Bonse, M. Hart and G. H. Schwuttke, Phys. Stat. Sol. 33 (1969)

361.

13. R. Gevers, Phil. Mag. 7 (1962) 1681

14. R. Gevers, J. van Landuyt and S. Amelinckx, Phys. Stat. Sol. 18

.(1966) 325.
-16-

15. B. R. Lawn and T. R. Wilshaw, "Fracture of Brittle Solids" (Cambridge

University Press, Cambridge, 1975 ), Ch. 2.

16. M. V. Swain, B. R. Lawn and S. J. Burns, J. Mater. Sci. 9 (1974) 175.

17. E. Sommer, Eng. Fract. Mech. 1^ (1969) 539.

18. I. V. Gridneva, Yu V. Milman and V. I. Trefilov, Phys. Stat. Sol, (a)

14 (1972) 177.

19. F. M. Ernsberger, Ann. Rev. Mat. Sci. 2 (1972) 529.

20. D. J. Rowcliffe, unpublished work.

21. B. R. Lawn, J. Appl. Phys. 39 (1968) 4828.

22. G. A. Wolff and J. D. Broder, Acta Cryst. 12 (1959) 313.

23. C. A. Brookes, J. B. O'Neill and B. A. W. Redfern, Proc. Roy. Soc. Lond.

A322 (1971) 73. '

24. E. M. Wilks and J. Wilks, Phil. Mag. 4 (1959) 158.

25. P. J. Bryant, L. H. Taylor and P. L. Gutshall, Trans. Tenth National

Vacuum Sympos. (Macmillan, New York, 1963) , p. 21.

26. J. A. Ryan and J. J. Grossman, Science J. 4_ (1968) 41.

27. B. J. Hockey, to be published.
Figure Captions

1. Transmission electron micrograph showing typical configuration of

dislocariv^ris and c:racks associated with a room temperature 200 g

Vickers indentation on the (0001) plane of a-SiC (4H polytype structure)

.

Two distinct crack configurations are evident: A, radially directed

cracks extending from corners of pyramidal impression, and B,

cracks lying parallel to foil either partially or fully encircling

central deformation zone (see also Figs. 8 and 10)

.

2. Transmission electron micrograph showing room temperature 200 g

Knoop indentation on the (lolo) plane of Al^O^. Note crack remnants

extending from sides of central deformation zone. Geometry of cracks

depicted in Fig. 5.

3. Schematic presentation of median vent cracks, formed during indenter

loading . Cracks are indicated by full, heavy lines, central

deformation zone by dark region. (a) Section view, for both Vickers

and Knoop indenters; location of eventual thin foil indicated by

dashed lines. (b) Plan view, for Vickers indenter; crack tends to

extend well beyond central zone, along both diagonals. (d) Plan

view, for Knoop indenter; crack tends to be obscured below central zone,

and to grow only along major diagonal.

4. Schematic representation of lateral vent cracks, formed during indenter

unloading. Cracks are indicated by full, heavy lines (preceding median

vent cracks by full, light lines), central deformation zone indicated

by dark region. (a) Section view, for both Vickers and Knoop indenters;

location of eventual thin foil indicated by dashed lines. (b) Plan
-18-

view, for Vickers indenter. (c) Plan view, for Knoop indenter.

Note lateral vents tend to form as lobes between the median vents.

5. Effect of skew loading on crack pattern for Knoop indenter. Depicted

to match loading conditions for the indentation illustrated in Fig. 2.

(a) Section view. (b) Plan view.

5. Median vent type crack associated with 200 g Vickers indentation on

(0001) plane of a-SiC, seen under different diffracting conditions.

Crack front indicated as CC. (a) and (b) satisfy Laue reflection

. conditions, diffraction vectors shown. (c) is view of crack obtained by

tilting specimen, multiple reflections. Note change in mismatch fringe

pattern in (a) and (b) with diffraction vector. Specimen indented at

room temperature and annealed in air at 1200°C for four hours prior to

thinning.

7. Portion of median vent crack associated with room temperature 200 g

Vickers indentation on (0001) plane of Al^O^. Note complex fringe

pattern over greater part of crack interface, with interfacial dislocation network near crack front CC. Multiple reflections operating.

8. Enlargement of lateral vent segment B in Fig. 1, under different

reflecting conditions. Note tendency for fringes to run nearly

parallel to the crack front, with modification in patterns at steps,

S, which terminate at crack front. Diffraction vectors indicated.

9. Enlargement of lateral vent at left in Fig. 2. Note interfacial

dislocation network between steps, S, which terminate at crack front.

10. Enlargement of lateral vent segment B in Fig. 1, under different

reflecting conditions. Note disappearance of crack-tip contrast
-19-

for g parallel to front. Note also presence of features exhibiting

dislocation-like contrast ahead of residual crack front. Step _

contrast, S, indicates extent to which crack front must once have

propagated.

11. (a) Portion of remnant crack front associated with room temperature

200 g Vickers indentation on (0001) a-SiC.

(b) Schematic representation, showing how crack breaks up into

partial fronts upon encountering some shear disturbance.

Arrows indicate direction of propagation of main crack.

12. Median vent crack remnant associated with room temperature 200 g

Vickers indentation on (0001) Al^O^. Indented specimen annealed at

1200°C in air for 8 hr prior to thinning. Annealing results in

recession of crack front to C'C^, with healed (sintered) region

characterised by interfacial dislocations and voids, and \inhealed

(open) region characterised by mismatch fringe contrast.

1.
3
(b)

(c)
(a)

(b)

(c)
(a)
9» 1 2 I 0


0.5u m

(a)

* 4

(b)

Modes I +III

1 1 ! 1 1 ! 1 1 1 n

Mode I

II.

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USCOMM NHS [)C
nD3-l iaA (REV. 7-73)

U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO.

NBSIR 75-658

2. Gov't Accession

No.

3. Recipient's Accession No.

4. TITLE AND SUBTITLE

Electron Microscopic Observations of Microcracking About

Indentations in Aluminium Oxide and Silicon Carbide

5. Publication Date January 1975

6. Performing Organization Code

1

7. AUTHOR(S)

B. J. Hockey and B. R. Lawn

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. 3130453

11. Contract/Grant No.

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

Department of the Navy

Office of Naval Research

Arlington, Virginia 22217

13. Type of Report & Period Covered Interim

7/1/74—6/30/75

14. Sponsoring Agency Code

ONR

15. SUPPLEMENTARY NOTES

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

bibliography or literature survey, mention it here.)

Transmission electron microscopy is used to examine the nature of microcracking

about small-scale indentations in two highly brittle solids, sapphire and carborundum.

The observed crack geometry is discussed in terms of an earlier model of indentation

fracture beneath a point force, in which both loading and unloading half-cycles

contribute to the crack growth. The residual interfaces are characterised mainly

by moire patterns, sometimes by dislocation networks. These observations are discussed in relation to spontaneous closure and healing mechanisms, and the "lattice

mismatch" necessary for their production estimated at about one part in a thousand.

It is shown that cleavage steps comprise the main source of obstruction to lattice

restoration across the interfaces. Mechanical and thermal treatments of the indented

surfaces are found to influence the extent of the residual cracking. Some practical

implications of the observations are discussed.

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

)

Brittle solids; dislocation networks; electron microscopy; healing; indentations;

microcracking; moire patterns

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