NBSIR 75-764

PYRAMIDAL SLIP AND BASAL ^- TWINNING IN ALUMINU MJDXjPE fcl;C!;-V

Hockev, B. J., Pyramidal slip on {il23j

^ IICO and basal twinnins in Al-^0^,

B. J. Hockey - ^ 3'

(Froa. Syrjp. on Plastic DeforiEation of

Ceracic i-.at^rials, Pennsylvania State

Inorganic Materials Division

Institute for Materials Research Univ., univsrsity Park, PA, July 17-19,

National Bureau of Standards

Washington, D. C, 20234 197'^).- Cdapter in Deforrration of Cerarr-ic

Materials, R. C. Bradt and H. S. Trsssls

August, 1975 Sis., Sec. II, do. 167-"'"179 (Plenum Press

New York, TTY, 1975).

This paper has been formerly published as follows:

(Proc. Symp. on Plastic Deformation of Ceramic

Materials, Pennsylvania State Univ., University

Park, PA, July 17-19, 1974), Chapter in Deformation

of Ceramic Materials, R. C, Bradt and R. E. Tressler, Eds,,

Sec. II, pp. 1 67-1 79 (Plenum Press, New York, NY, 1 975).

Prepared for Department of the Army

U. S. Army Research Office

Box CM

Duke Station

Durham, North Carolina

U.S. DEPARTMENT OF COMMERCE, Rogers C. B. Morton, Secretary

James A. Baker, lit. Under Secretvy

Dr. Betsy Ancker-Johnson, Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS, Errvett AmWer, Acting Director

NBSIR 75-764

Pyramidal Slip and Basal

Twining in Aluminum Oxide

B. J. Hockey

Inorganic Materials Division

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

August, 1975

This paper has been formerly published as follows:

(Proc. Symp. on Plastic Deformation of Ceramic

Materials, Pennsylvania State Univ., University

Park, PA, July 17-19, 1974), Chapter in Deformation

of Ceramic Materials, R. C. Bradt and R. E. Tressler, Eds.,

Sec. II, pp. 167-179 (Plenum Press, New York, NY, 1975).

Prepared for Department of the Army

U.S. Army Research Office

Box CM

Duke Station

Durham, North Carolina 27706

PYRAMIDAL SLIP ON {1123; <1100> AND BASAL WINNING IN Al^O^

5 . J , Hockey

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

ABSTRACT

Plastic deformation of Al^O- hy slip and twinning has been

investigated by exar.ining -he regions surrounding a microhardness

indentation using transr.ission electron microscopy (TE:-!) . The

results establish: {1) the occurrence of pyrarriidal slip on {1123}<11CC> , and (2) the nature of basal twins in this material.

The observations on basal twins, in particular, have led to a completely different description for the twinning process,

which is briefly described.

1 . INTRODUCTION

Previous observations by TEl-i [1] have shcsn that generalized

plastic flow involving both slip and "cwinning occurs d'uring the rcorn

terr.perat'ure indentation of Al^O^ . Indentation of various surface

orientations —for exa-ple, by tne placement of indentations within

individual grains of polycr\'stals —thus provides a sim.ple experir.ental

means for causing and subsequently studying the possible r.odes of

deforr.ation by TZM.

In this paper, the general usefulness of these techniques,

which are particularly applicable zc "hard, brittle materials" [2]

,

is illustrated and specific results (1) on the occurrence of pyramidal

slip on the '.1123/ <110C> system and (2) on the structure of

basal ru.crotwins in Al_C. are presented.
2. EXPERIMENTAL DESCRIPTION

The methods used in the initial specimen preparation, indentation, and the preparation of thinned electron microscope specimens

have been previously published [1]. In the present study, the

electron microscope was operated at an accelerating voltage of 200

kv and a double-tilt (+ A5°) specimen holder was used to tilt the

specimen to various two-beam conditions.

All observations presented here were obtained from the region

surrounding a 200 gram Vickers hardness indentation, made at room

temperature within a single grain of sintered polycrystalline alumina.

Fig. 1 is a sterographic projection* onto the plane of the foil and

shows that the normal to the original crystal surface was close

to the (1123) pole.^^The other poles listed correspond to the

various reflections used in the analysis of slip and twinning

described in the following sections.

3. RESULTS AND DISCUSSION

A. Slip on {1123} <il00>

Observations on the indented specimen revealed numerous arrays

of dislocations which extended from the indentation in various directions. Fig. 2 shows the configuration of dislocation found in one area. In this micrograph, the beam is nearly normal to the

foil, and it is clear that the region contains arrays of short dislocation segments (E) situated on inclined planes that intersect

the foil along A-B, together with long dislocation segments (S)

which are nearly parallel to A-B. The regions of dark contrast, F, are apparently faults which also intersect the foil along A-B. Although not apparent in" this projected view, stereomicroscopy revealed

that all of the isolated dislocations (S and E) lie on the same set

of parallel planes, which is also the plane of the faults, F.

The presence of curved dislocations on parallel planes strongly suggests that the observed configuration is a result of slip.

Miller-Bravais indices corresponding to the structural unit cell

c/a = 2.730, are used throughout.

Indices of corresponding reciprocal_space vectors satisfy the

reflecting conditions: - H+k+Jl =3n; hh2h£ : (Jl=3n) ; hho£ :h+£=3n,

i+2n [3] .

^
3

Figure 1. Stereographic projection normal to the indented specimen

surface. With the exception of 1123, all poles listed

correspond to the various reflections used in the analyses

described in Section 3.

This conclusion was confirmed by determining both the apparent slip

plane and the Burgers vector of the isolated dislocations. The

lack of direct evidence for the various slip systems proposed for

Al^O^ [4] , make it worthwhile to consider the details of the results.

(a) Slip Plane . With reference to Fig. 1, the intersection of

the apparent slip planes and faults, F, with the surfaces of the

foil was determined to be along [1100] . Accordingly, the plane of

the dislocations must belong to the [1100] zone, which is indicated

by the dotted great circle in Fig. 1. Identification of this plane

was then made by tilting the foil about [1100] until the dislocations and faults were edge-on (i.e. parallel to the electron beam).

As shown in Fig. 3, this position of the foil was determined by

electron diffraction to coincide with the 1123 reflecting position.

Since 1123 lies on the [1100] zone, this uniquely defines the slip

(and fault) plane as (1123).
4

Figure 2. Arrays of dislocations associated with 200 g Vickers

hardness indentation which were determined to be a result of slip on (1123) [1100] . See text for description.

(b) Slip Direction . The region of Fig. 2 was examined under various

two-beam reflecting conditions_ (see Fig. 1) to determine the nature of

the dislocation arrays on (1123) . These observations confirmed that

the isolated dislocations, S and E, were total, slip dislocations. The

nature of the faults, F, and their bounding partial dislocations also

present on (1123), however, could not be clearly established.

As will be seen, the determination of the Burgers vector for the

total dislocations was complicated by the fact that strong residual

contrast was produced for nearly all of the reflections used. As a result, a unique solution for b could not be made sol'ely on the basis

of the 5 • b = 0 invisibility criterion, but instead required a lengthy

comparison of the observed contrast with that expected for edge and

screw dislocations of different Burgers vectors .^^ »^-' All possible

lattice vectors for the AI2O3 structure of reasonable magnitude (i.e.

<12 A) were included in this comparison. On the basis of this analysis - which for the sake of brevity is not detailed here - the only reasonable

agreement between observed and expected contrast was obtained for

B = [1100]

.
Figure 3. Dislocations and_faults are seen "edge-on" with specimen

tilted to the 1123 reflecting position.

In agreement, comparison of Figs. 4a and 4b with Fig. 3 shows

that the long dislocation segments (S) lying close to [1100] are either out of contrast or show only a weak residual contrast for

g = 1123 and g = 1126. Under both conditions, g* [1100] = 0 and it can be concluded that these dislocations axe screw or near screw

segments with b = [1100] . From their line orientation, those dislocations which are steeply inclined to the foil (E) appear to

correspond to edge or near edge [1100] dislocation segments. Although these dislocations are in contrast, particularly in Fig. 4b,

the nature of the line contrast is characteristic of that expected

for edge dislocations when g-b = 0, but m = l/8g-b_ x u_ 5 0.2, where

u is a unit vector tangent to the dislocation line (see e.g. [5]

and [6] ) . Since b x u_ is simply a vector of magnitude |b' along

the normal to the (1123) slip planes, this expression for m can be

reduced to

1/8 Igj |b| cos - = 1/8 -——

hkiJl ^Tikii

Ji
for pure_edge dislocations, where 6 is the angle between ^j-^j^^^^'^^

the (1123) pole. Using this expression, calculated values of m

for (1123) [1100] edge dislocations are: m = 0.61 for g = 1126,

Fig. 4b, and m = 0.49 for g = 1123, Fig. 4a. These values are quite large and the strong contrast observed experimentally is

expected.

Figure 4. Image contrast_of dislocations_on (1123) planes produced

for: a) g = 1123 and b) g = 1125. For both reflections,

g-b = 0 for [1100] dislocations. Only [1100] screw

segments are invisible; edge segments show strong con- trast due to large values of m = 1/8 g'b x _u.

In further agreement, strong double line contrast for both

screw (S) and edge (E) segments was observed in reflections for

which g* [1100] = 2 at s = 0 (i.e., at the Bragg reflecting posi- . tion) . This is illustrated in Fig. 5, where g = 1102.
Figure 5. Edge and screw [1100] dislocation segments in doiible

image contrast (g-b = 2) for g = 1102 at s = 0.

The_present results thus provide clear evidence that slip on

{1123} <1100> occurs during the room temperature indentation of

AI2O3. As yet, no explanation can be given regarding either the

choice or activity of this particular system. In particular, it

should be noted that 1/3 [011l2_and not [1100] corresponds to the

shortest lattice vector in (1123) . Conceivably, the extremely

high stresses developed during point loading may result in

the activation of certain slip systems which may never be

energetically favorable modes of bulk deformation under uni- or bi-axial loading. The assxamption then that {1123} <1100> slip occurs during the high temperature bulk deformation of AI2O3 cannot

be made solely on the basis of these results. However, KlassenNeklyudova et_ al_. [7] have reported impurity segregation along

{1123} planes in r\iby after deformation above 1850°. The present

results, thus, appear to justify their conclusion that these bands
are the result of slip. In addition, the present observations provide the first positive identification of glissile <loIo> dislocations in ^^2^3' ^^i'^h until now, could only be inferred from etch

pits or the macroscopic displacements associated with {1210} <1010>

slip [8,9,10].

Unlike {1123} <1100> slip, there is ample evidence for deformation twinning on {0001} [11,12,13]. The macroscopic elements of

basal twinning-first determined by morphological examination and

later confirmed by X-ray diffraction-are known to be:*

S = 0.635

Crystallographically , twinning on (0001) can be described by a rotation of 180° about [0001] , which requires that the macroscopic

shear direction be [1010] , [1100] , or [0110] and not the reverse

directions.

However, as pointed out by Kronberg, [14] uniform shear displacements along th^ observed twinning shear direction, o-^, cannot

possibly lead to a twinned structure. As a consequence, Kronberg

proposed a rather complicated shearing mechanism involving the

synchronized motions of "quarter-partial" twinning dislocations.

In this section, the structure of basal microtwins, as determined by TEM, is presented. The results identify the actual displacements associated with the twinned structure and have led to

the development of a completely different description of the twinning process.

Figure 1 agains specifies the orientation of the indented specimen which contained numerous thin lamallar regions that often extended far beyond the zone of intense slip deformation. The identification of these lamallae as basal microtwins was readily established

by selected area diffraction and the contrast observed under different diffracting conditions. Figs. 5a and 6b. Typically, these

twins were quite thin (< 2iJm) and essentially wedge-shaped with

arrays of dislocations present within the twin-matrix boundary.

These features are best illustrated in Fig. 6c, which shows the same twin viewed edge-on. Here the twin boundaries are delineated

B. Basal Twinning

(0001);

{1011};

<1010>

<1012>

*It is important to emphasize that indices based on the true structiaral cell are used here.
by the residual contrast associated with the boundary dislocations.

Previous observations [15] on basal twins produced by indentation

or surface grinding have revealed similar characteristics.

a

Figure 6. Basal microtwin under various diffracting conditions:

a) Bright field g = (0114) . , b) Dark field matrix

g = (1014) . , c) Bright field g = 000.12.

twxn

In the present study, the peculiar nature of the twin boundaries was established by unambigously determining the Burgers

vectors of the boundary dislocations by application of the g-b = 0

invisibility criterion [5] . The essential results are contained

in Figs. 7a-f , which show that, for this orientation, one face of

the wedge boundary contains only +_[1010] dislocations, while the

other face contains only [0110] dislocations. Since similar contrast is observed at the points where the dislocations emerge from

the foil, the dislocations within each face must be of the same

sign. Although the magnitude of b was not determined, it can be

shown that the boundary dislocations cannot be actual twinning

partials, but instead must correspond to accommodation dislocations
Figure 7. Basal twin boundary dislocations in different diffracting

conditions. Dislocations in both faces of wedge-shaped

twins are in contrast for a) g = 1126 and b) g = 1123.

Dislocations in one face are invisible for c) g = 2116

and d) g = 2113; g-b = 0 for b = x[0110]. Dislocations

in opposite face are invisible for e) g =1216 and f)

g = 1213; g.b = 0 for b = x[loIo] .

For the purposes of discussion, it is asi^jmed that the unequal

densities of dislocations present within the two boundaries simply

indicates that the boundaries are not equally inclined to (0001) - the subtended angle at the tip was too small (>2° in Fig. 6c) to

allow experimental verification. Fig. 8a, thus, schematically

represents the configuration of boundary dislocations at the tip

of an ideal fully symmetric basal twin.
11

Figure 8. a) Schematic representation cf configuration of boundary dislocations at the tip of a syrraetric wedgeshaped basal twin based on present observations.

Twin is viewed parallel to (0001) along [1120] direction as in Fig. 6c.

b) Vector diagram illustrating nature of the relative

shear displacements on (0001) planes through crosssection of basal twin, e.g., ABCDE in (a).

The significance of the results can be understood by noting

that, since the boundary dislocations are glissile on {OOOl), each

face of the twin corresponds to a simple shear boundary. The observed structure of the twin boundaries thus directly suggests a twinning mechanism involving uniform shear on (0001) along +_[1010]

above the mid-plane of the twin and along +_[0110] below the midplane. Such a process is clearly imcompatible with a 'onidirectional

shear process normally used to describe twinning. It is also

incompatible with Kronberg ' s proposed dislocation model for basal

twinning [14] .
The observed structure can, however, be explained on the basis

of a different and relatively simple dislocation model for basal

twinning. The details of this model are included in a formal treatment of basal twinning which is currently being prepared for publication elsewhere. Only those aspects relating to the present

results and admittedly idealized by Fig. 8a will be described here.

Basic to this model, is the realization the displacements of

1/3 [1010], 1/3 [Olio], or 1/3 [1100] , i.e., opposite to the macroscopic twinning shear directions, on every other basal lattice

plane will result in a twinned structure. As a result, a symmetric

twin, whose orientation and boundary displacements are depicted in

Fig. 8a, can be described by the passage of 1/3 [1010] partial dislocations on every other (0001) lattice plane situated above the

mid-plane of the tv;in and the passage of 1/3 [0110] partial dislocations on every other (0001) lattice plane below the mid-plane.

This process can be envisioned by considering the action of a hypothetical double-ended pole dislocation source emitting

1/3 [1010] dislocations at one end anc 1/3 [0110] dislocations at the

other end under a uniform applied stress. The resulting displacements through any cross-section of the twin, e.g. ABCDE in Fig. 8a,

will then be represented by a simple vector diagram. Fig. 8b. As

seen, the relative shear displacements above and below the midplane correspond to the experimentally determined directions (the

sign of which cannot be uniquely specified) . Moreover, the resultant matrix displacement across the twin is along [1100] . Propagation of the twin through the crystal would thus eventually result

in a maximiam surface offset along [1100] , which corresponds to the

required macroscopic_twinning shear direction, o_^. Since the

passage of each 1/3 [1010] and 1/3 [0110] dislocations produces a net resolved displacement of 1/5 [1100] on every other (0001) lattice

plane, the resultant shear along [1100] is simply

l^ff^l= /3 ./c. o. 0.634,

which corresponds to the experimentally verified required value.

The proposed model thus satisfies all the macroscopically

observed aspects of twinning. Unlike Kronberg ' s model, the synchonous displacements of adjacent oxygen and aluminum lattice

planes in different direction under a uniform stress is not required

Moreover, the resulting twin interface on (0001) is a true mirror

plane, not a shear (faulted) interface.

The author wishes to acknowledge the support of the Army

Research Office.
13

References

B. J. Hockey, J. Am. Ceram. Soc, 54_, 223 (1971).

B. J. Hockey, pp. 21-50 in The Science of Hardness Testing and

Its Research Applications. Edited by J. H. Westbrook and

H. Conrad. American Society for Metals, Metals Park. Ohio 1973.

International Tables for X-ray Crystallography, Vol. 1; No. 167,

p. 275. Kynoch Press, Briiningham , England, 1962.

J. D. Snow and A. H. Heuer, J. Am. Ceram. Soc, 56_, 153 (1973).

P. B. Htrsch, A. Howie, R. B. Nicholson, D. W. Pashley, and

M. J. Whelan. Electron Microscopy of Thin Crystals. Butterworth, Inc., Washington, D.C., 1965.

J. M. Silcock and W. J. Tunstall , Phil. Mag. 10, 361 (1964).

M. V. Klassen-Neklyudova, V. G. Govorkov, A. A. Urusovskaga,

M. N. Voinova, and E. P. Kozlovskaga, Phys. Stat. Solidi, 39 , 679 (1970)

.

M. V. Klassen-Neklyudova, J. Tech. Phys. (USSR), 12, 519, 535

(1942)

.

R. Scheuplein and P. Gibbs, J. Am. Ceram. Soc, 43_, 458 (1960).

D. J. Gooch and G. W. Grotfes, J. Am. Ceram. Soc, _55, 105 (1972)

K. Veit, Neues Jahrb. Mineral., Geol . Palaeontol., Beilageband,

45 , 121 (1921)

.

E. Stofel and H. Conrad, Trans. AIME, 227 , 1053 (1963).

A. H. Heuer, Phil. Mag., _13, 379 (1966).

M. L. Kronberg, Acta Met., _5, 507 (1957).

B. J. Hockey, Proc Brit. Ceram. Soc, 20, 95 (1972).

NBS-n4A (REV. 7-73)

1 U .S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO.

NBSIR-75-764

2. Gov't Accession

No.

3. Recipient's Accession No.

4. TITLE AND SUBTITLE

ryraimaaj. oiip on \±±d.jj <nuu> and. r>asal i winning

in AI2G3

5. Publication Date

6. Performing Organization Code

7. AUTHOR(S)

B. J. Hockey

8. Performing Organ. Report No.

9. PERFOR.MING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/Work Unit No. 3130152

11. Contract/Grant No.

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

Department of the Army

U. S. Army Research Office

Box CM, Duke Station, Durham, North Carolina 27706

13. Type of Report & Period

Covered

14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

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

bibliography or literature survey, mention it here.)

Plastic deformation of Al20^ by slip and twinning has been investigated

by examining the regions surrounding a microhardness indentation using

transmission electron microscopy ( jM) . _The results establish: (1) the

occurrence of pyramidal slip on {1123) <1100> and (2) the nature of basal

twins in this material. The observations on basal twins, in particular,

have led to a completely different description for the twinning process,

which is briefly described.

17. KEY U'ORDS fsix to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper name; separated by semicolons)

Aluminum oxide; microhai^ess; plastic deformation; slip; transmission electron

microscopy; twinning

18. AVAILABILITY [^^Tnlimited 19. SECURITY CLASS

(THIS REPORT)

21. NO. OF PAGES

i'^^^'^For' Official Distribution. Do Not Release to NTIS

UNCL ASSIFIED

15

1

' Order From Sup. of Doc, U.S. Government Printing

Washington. D.C. 20402. SD Cat. No. C13

Office 20. SECURITY CLASS

(THIS PAGE)

22. Price

Order From National Technical Information Service

Springfi-eld, Virginia 22151

(NTIS)

UNCLASSIFIED

USCOMM-DC 29042-P74

I

II

I