Observations and analyses of dislocations and stacking faults in the massive γ{sub m} phase in a...

OBSERVATIONS AND ANALYSES OF DISLOCATIONS

AND STACKING FAULTS IN THE MASSIVE gm PHASE IN

A QUENCHED Ti±46.5 AT.% Al ALLOY

P. WANG1, M. KUMAR2, D. VEERARAGHAVAN1 and V. K. VASUDEVAN1

1Department of Materials Science and Engineering, University of Cincinnati, Cincinnati, OH 45221-0012 and 2Department of Mechanical Engineering, Johns Hopkins University, Baltimore, MD 21218-

2686, U.S.A.

(Received 27 January 1997; accepted 21 July 1997)

AbstractÐThe defect structures in the massively formed gamma (gm) grains in a Ti±46.5 at.% Al alloy,rapidly quenched from the high-temperature a-phase ®eld, have been studied using transmission elec-tron microscopy (TEM). The results reveal that the defect structures are composed of dislocations,stacking faults and antiphase boundaries intimately associated with dislocations or stacking faults. Con-trast analysis indicates that both 1/2<110] and 1/2<101] unit dislocations were present in gm phase,with the latter linked by highly curved non-conservative antiphase boundaries. Comparison of exper-imental and computer simulated TEM images established that wide stacking faults, which are createdby the dissociations of 1/2h101i unit dislocations, lie on {111} planes and are bound by b= 1/6h121iShockley partial dislocations of all possible types. In addition, antiphase boundaries are found to com-mence or terminate on the stacking faults at the partial dislocations with b= 1/6<121], but not thosewith b= 1/6 < 112]. Based on the observations and subsequent analyses, a model for the formation ofthese defectsÐinvolving the occurrence of an intermediate disordered f.c.c. phase during the a4 gmmassive transformationÐis proposed. # 1997 Acta Metallurgica Inc.

1. INTRODUCTION

Since the ®rst report of the a 4 gm massive

transformation in rapidly quenched TiAl alloys

by Wang and co-workers [1, 2], several groups

[3±12], including themselves, have reported the

observation of defects such as dislocations, stack-

ing faults (SFs), microtwins and antiphase

boundaries (APBs) in the massive gm grains.

However, no detailed studies of the dislocations

and SFs in the gm phase have been performed,

although the APB-like defects have been charac-

terized in detail [4±12] and models for their for-

mation mechanisms put forth [6, 9±12]. The aim

of this paper is to present ®rst time TEM obser-

vations elucidating the nature of the SFs and

their bounding dislocations, as well as of other

dislocations found singly or in association with

the APB-like defects in gm; the detailed structural

characteristics of the APB-like defects themselves

are considered in a follow-on paper [13]. The

nature of the stacking faults and the bounding

partials has been established experimentally and

by comparison of experimental and simulated

images. A model for the formation of these

defects during the a 4 gm massive transformation

is proposed based on these observations.

2. EXPERIMENTAL PROCEDURE

The Ti±46.5 at.% Al alloy was obtained as a 250g arc-melted ingot in the cast + hot isostatically

pressed (12008C, 2h, 105 MPa) condition. Wetchemical analysis gave the actual alloy compositionas Ti±46.54 at.% Al, with 0.051 wt% oxygen, 0.005

wt% nitrogen and 0.012 wt% carbon being presentas impurities. Sectioned samples of the alloy were

coated with a protective glass±ceramic [1, 2], thensolutionized in the single phase a regionÐ13808Cfor 20 minÐand subsequently water or iced-brinequenched. Thin foils for TEM were prepared from

the quenched samples using conditions reportedelsewhere [1, 2], and observed in a Philips CM20

TEM operated at 200 kV, utilizing bright ®eld(BF), superlattice dark ®eld (SLDF), weak beam

dark ®eld (WBDF) and selected area di�raction(SAD) modes. The defect structures in the gm and

a2 phases were examined under two-beam con-ditions using a variety of low-index g-vectors. Onlyrepresentative micrographs are presented in this

paper.

Conventional g�b type contrast analysis [14] wasused to determine the Burgers' vectors of unit dislo-cations, as well as of the two individual partial dis-

locations bounding the geometrical faults and thedisplacement vectors associated with the faults. BF

Acta mater. Vol. 46, No. 1, pp. 13±30, 1998# 1997 Acta Metallurgica Inc.

Published by Elsevier Science Ltd. All rights reservedPrinted in Great Britain

1359-6454/98 $19.00+0.00PII: S1359-6454(97)00237-1

13

and WBDF images from di�erent g vectors werecompared and the invisibility criterion applied

(g.b= 0 for unit dislocations and g.b = 0 or 1/3for Shockley partial dislocations). However, themagnitude of the vectors associated with the partial

dislocations bounding stacking faults could not bedetermined entirely in this manner. To establishexactly the direction and magnitude of the individ-

ual partial dislocations, experimental BF imageswere compared with simulated images of variouspossible con®gurations. In particular, the number

and intensity of fringes and the placement of brightand dark lobes along the dislocation lines in theimages were compared. The simulations were car-ried out using CUFOUR [15], a many-beam pro-

gram based on the earlier two-beam program ofHead et al. [16]. The program, CUFOUR [15], wasoriginally written for cubic materials, but sub-

sequently modi®ed for tetragonal crystals [17]. Thegeneralized cross-section formalized by Head et al.[16] was used in the calculations. Use of this cross-

section assumes that the dislocation line direction isacute to the foil normal, and this ®xes the place-ment of the dislocation line (direction) from the left

(bottom of foil) to the right (top of the foil) of thesimulated image. Hence, the leading dislocation isalways at the top of the image and the trailing par-tial at the bottom. Stacking faults, which can be

distinguished from APBs because of the di�erencesin contrast characteristics of the outermost fringesin two-beam BF±DF images (symmetrical±asymme-

trical vs symmetrical±symmetrical, i.e. complemen-tary), can be established to be intrinsic or extrinsicby the method of Gevers et al. [18]. However, as

pointed out by Viguier et al. [17] errors in this esti-mation can arise owing to unusual foil geometries.Hence, the nature of the stacking faults was alsoestablished by image simulations. The elastic con-

stants for Ll0±TiAl were obtained from the work ofTanaka et al. [19], and extinction distances andabsorption coe�cients were calculated using the

EMS program [20]. In all cases three di�ractedbeams in the systematic row of the re¯ection andone transmitted beam (ÿg, 0, +g, 2g) were con-

sidered. The parentheses used with the Miller indi-ces of directions and planes in the text follow theconvention introduced for TiAl by Hug et al.

[21, 22]: in the notation <hkl] or {hkl), only h and kare mutually permutable; l is ®xed and can be posi-tive or negative.

3. RESULTS

Optical microscopy observations of the rapidlyquenched samples of the Ti±46.54 Al alloy revealeda microstructure consisting of darkly etched massive

gm in a light a2 matrix, similar to those reportedpreviously in the same alloy [2]. Figure 1(a) is a BFimage, recorded using g = 002, of a typical regionof the gm phase showing several dislocations, a few

SFs and highly curved APB-like defects. The latter

reveal contrast even when imaged with a fundamen-tal re¯ection [Fig. 1(a)], similar to the observationsof Li and Loretto [7] and Zhang and co-workers

[8±12]. The displacement vector, R, of these APBshas been proposed to be 1/2<101] + DR =05%of R, the latter being responsible for the obser-

vation of contrast with fundamental re¯ections [7].The APB-like defects were often found to be com-

posed of a partially or completely ®lled thin layerof a g variant with c-axis rotated 908 relative to thesurrounding g matrix, somewhat similar to the situ-

ation reported by Zhang and co-workers [8±12].Nevertheless, in the interest of brevity, these APB-like defects are referred to simply as APBs in the

remainder of the text. The structural characteristicsof the APBs are considered in detail in a follow-onpaper [13].

A striking feature is that the APBs commence orterminate at dislocations; pairs of these dislocations

are shown labeled A±B, C±D, E±F and G±H inFig. 1(a) and (b). Details of the contrast analysis ofthe dislocations are given in Table 1, and represen-

tative images are shown in Fig. 1. Since all of thesedislocations are visible in the BF±WBDF image

pair [Fig. 1(a), (b)], recorded using g = 002, and inthe WBDF images recorded using g = 200[Fig. 1(c)], g=111 [Fig. 1(d)] and g = 111 [Fig. 1(e)],

but invisible with g = 111 [Fig. 1(f)] and g = 111(Table 1), their Burgers vector is determined to bebA±H=1/2[101] (in the ordered TiAl lattice this is

one half of the Burgers vector of a [101] superdislo-cation). Stereographic trace analysis revealed thatthese dislocations were mixed (608) or edge in char-

acter, with line directions parallel to the [011] or[121] directions, respectively, and were lying on the

(111) plane. Furthermore, the observations revealthat the APBs run through the foil and the dislo-cations at which they terminate show the oscillatory

contrast that is characteristic of dislocations thatare inclined and running top to bottom through thethickness of the foil [14]. This means that the APBs

are not terminated at the foil surface and the unitdislocations at the ends have not been produced

because of sectioning of the APBs by the foil sur-faces.Examination of many other areas of the sample

revealed the same feature, namely, APBs linked bypairs of b= 1/2<101] dislocations or by a1/2<101] dislocation at one end, with the other

end terminating at a g intervariant domain or high-/low-angle boundary. Examples of the latter are

shown in Fig. 2, where the APBs labeled A and Bare attached to the b= 1/2 [101] dislocationslabeled A1 and B1, respectively, at one end in the

middle of the micrographs and to a g domain/low-angle boundary labeled DB at the other (upper)end. The APB labeled C is attached to one end to

the dislocation labeled C1, which has bC1=1/2[101].These dislocations were determined to be mixed 308

WANG et al.: OBSERVATIONS AND ANALYSES14

(A1, B1) or near-screw (C1) in character and lying

on the (111) plane. In some areas, APBs were

observed as closed loops without any association

with dislocations. The mechanism of formation of

these APBs is considered elsewhere [13].

A few undissociated 1/2<110] unit dislocations

were also observed (those labeled I1, I3, J2, J3 in

Fig. 1, set D in Fig. 2) and in this case no APBs

were found to be associated with them. For

example, the dislocations labeled I1, I3 in Fig. 1

have b=21/2[110], since they are visible with

g= 200 [Fig. 1(c)], g= 111 [Fig. 1(e)] and g = 111

[Fig. 1(f)], but practically invisible with g = 002

[Fig. 1(a),(b)], g=111 [Fig. 1(d)] and g = 111

(Table 1). Similarly, the dislocations labeled J2, J3

have b=21/2[110], since they are visible with

g= 200 [Fig. 1(c)], g=111 [Fig. 1(d)] and g = 111

(Table 1), but practically invisible with g = 002

Fig. 1(a,b).

WANG et al.: OBSERVATIONS AND ANALYSES 15

[Fig. 1(a),(b)], g= 111 [Fig. 1(e)] and g= 111

[Fig. 1(f)]. Stereographic analysis established that

dislocations I1±I3 and J1±J3 were mixed 308 and

458 in character and hence lying on the (111) and

(111) planes, respectively; this is responsible for the

weak, residual contrast seen from these dislocations

in the image with g = 002. The dislocations I2 and

J1 have b =21/2[110] and b =21/2[110], respect-

ively, although they appear to be visible in the

image with g= 002. This is because APBs appar-

ently pass through these dislocations [Fig. 1(a),(b)]

and the ensuing superimposition a�ects image con-

trast. Also, these dislocations appear to have

already dissociated into Shockley partial dislo-

cations with the associated SF. This can be seen

clearly with dislocation I2, but is just discernible

with dislocation J1. For dislocation I2, the dis-

sociation occurs according to:

Fig. 1(c,d).

WANG et al.: OBSERVATIONS AND ANALYSES16

1

2�110� ! 1

6�21�1� � SF� 1

6�121�

Thus, in the image with g= 002, the SF is expected

to be visible (a= 2p�g�R= 2p/3), whereas both the

bounding Shockley partials are expected to be invis-

ible (g�b = 1/3), exactly as observed in Fig. 1(b).

The dislocation set D (Fig. 2) was determined to

have bD=21/2[110], a line direction of [011] and

hence lying on the (111) plane. Owing to their

Fig. 1. TEM micrographs showing dislocations and APBs in the massive gm phase in the quenched Ti±46.54Al alloy. (a), (b) BF and WBDF images recorded using g = 200 (B0[010]), (c)±(f) WBDF micro-graphs recorded using (c) g= 200 (B0[010]), (d) g=111 (B0[110]), (e) g= 111 (B0[011] and (f)g = 111 (B0[121]), respectively. Note the close association between the APBs and 1/2[101] dislocations(labeled A±H). Unit dislocations labeled I1±13 and J1±J3 have b =21/2[110] and b=21/2[110],

respectively.

WANG et al.: OBSERVATIONS AND ANALYSES 17

Table 1. Experimentally observed contrast$ from dislocations labeled A± J in Fig. 1 and deduced Burgers vectors

Operatingre¯ection (g) Beam Direction (B) Dislocations A±H Dislocations I1±I3 Dislocations J1±J3 Figure No.

002 0[010] V W/I W/I 1(a), 1(b)200 0[010] V V V 1(c)111 0[110] V I V 1(d)111 0[011] V V I 1(e)111 0[121] I V I 1(f)111 0[011] I I V Ð022 0[011] V V V Ð202 0[121] V V V Ð

21/2[101](111) 21/2[110](111) 21/2[110](111)

$ W = weak; I = Invisible and V = Visible

Fig. 2(a,b).

WANG et al.: OBSERVATIONS AND ANALYSES18

mixed 608 character, the dislocations in this set

show weak, residual contrast in the images obtained

with g= 002/001 [Fig. 2(a),(b)].

Another remarkable feature observed in the gmphase was the presence of many wide SFs, with

APBs commencing or terminating at the partial dis-

locations bounding these faults, as shown, for

example, in Fig. 3. Consider the SFs labeled A and

B and the partial dislocations at the ends of these

faults labeled A1±A2 and B1. Fault A was deter-

Fig. 2. TEM micrographs showing examples of dislocations and APBs in the massive gm phase in thequenched Ti±46.54Al alloy. (a) BF image g= 002, (B0[010]), (b) SLDF image g= 001 (B0[010]; (c),(d) WBDF micrographs recorded using g=111 (B0[110]) and g=111 (B0[110]). The APBs labeled Aand B are attached to b= 1/2[101] unit dislocations labeled A1, B1 at the bottom ends and to adomain/low-angle boundary, labeled DB, at the upper ends. The APB labeled C is attached at one endto a b = 1/2[101] unit dislocation. The set of unit dislocations labeled D and shown by arrows have

b=21/2[110].

WANG et al.: OBSERVATIONS AND ANALYSES 19

mined to lie on the (111) plane, whereas B, which is

partly attached to the bottom of fault A, was deter-

mined to lie on the (111) plane. The partial A1 is

visible with g= 200 [Fig. 3(c)], but invisible (weak

in contrast) with g = 002 [Fig. 3(a),(b)], g = 111

[Fig. 3(d)] and g = 111 (Table 2). Thus, its Burgers

vector is bA1=21/x[211], where x= 3 or 6. The

partial A2 is visible with g = 002 [Fig. 3(a),(b)] and

g = 111 [Fig. 3(d)], but invisible (weak in contrast)

with g= 200 [Fig. 3(c)] and g = 111 (Table 2),

which gives its Burgers vector as bA2=21/x[112]

(where x= 3 or 6). By stereographic trace analysis,

both partials were determined to have line direc-

tions parallel to [101], which means that the partials

are mixed (308) in character and lie on the (111)

plane. The Burgers vector analysis of these dislo-

cations was self-consistent with the contrast

observed using a number of other g-vectors

(Table 2), but all of the images from these are not

shown herein. Another aspect to note in Fig. 3 is

that APBs are visible with weak contrast; these are

attached to the end of the stacking fault A, but

only at the partial with bA1=21/x[211] and not at

the partial with bA2=21/x[112]. A couple of APBs

can also be seen passing through fault A. The

Burgers vector of partial B1 was determined in a

similar way to be bB1=21/x[112].

Image simulations were carried out to provide

further con®rmation of the nature of the stacking

faults and the Burgers vectors of the partial dislo-

Fig. 3. TEM micrographs showing stacking faults, their bounding Shockley partial dislocations andAPBs in the massive gm phase in the quenched Ti±46.54Al alloy. Faults labeled A and B with boundingShockley partials A1±A2 and B1 are shown. (a), (b) BF±WBDF image pairs recorded using g= 002(B0[010]); (c), (d) WBDF micrographs recorded using g= 200 (B0[010]) and g = 111 (B0[110]), re-

spectively. The dislocation labeled C has b= 1/2[101].

WANG et al.: OBSERVATIONS AND ANALYSES20

cations associated with them. Experimental images

of the fault labeled A in Fig. 3(a), recorded under

two di�erent di�raction conditions, are shown in

Fig. 4(a), (b)Ði. From analysis of the contrast in

these images as discussed above, it appears that the

partial dislocations on the (111) plane have the fol-

lowing Burgers vectors: 21/x[211] and 21/x[112],

where x = 3 or 6. Image simulations were carried

out for all possibilities of the h101i dislocations (i.e.

1/2[101], [101], 1/2[011], [011] and 1/2[110]) disso-

ciating on the (111) plane to introduce a single geo-

metrical fault bounded by two partial dislocations,

and assuming the fault to be extrinsic, intrinsic, or

a complex fault in the ordered lattice. (In the inter-

ests of space only the pertinent images will be pre-

sented.)

Based on the images obtained from computer

simulations, it was possible to narrow down the

Burgers vector of the undissociated dislocation to

be in the [101] direction. Simulated images of the

dissociated dislocations are shown along with the

experimental images obtained with g= [200]

[Fig. 4(a)Ði] and [202] [Fig. 4(b)Ði] for the three

cases where btotal=1/2[101] [Fig. 4(a,b)Ðii], 1/

2[110] [Fig. 4(a,b)Ðiii] or 1/2[011] [Fig. 4(a,b)Ðiv]

with single Shockley partial dislocations bounding

the fault, and the fourth case where btotal=[101]

[Fig. 4(a,b)Ðv] with double Shockley partial dislo-

cations bounding the fault. It is judged from the

number and placement of the lobes along the dislo-

cation line in the simulated images that the closest

match with the experimental images [Fig. 4(a,b)Ði]

is obtained for the ®rst case [Fig. 4(a,b)Ðii], where

the dissociation occurs (on the (111) plane) in the

following manner:

1

2��10�1� ! 1

6��21�1� � SF� 1

6��1�2�1�

These results con®rm the postulated Burgers' vec-

tors of the partial dislocations bounding SF A,

which were deduced from the set of images in Fig. 3

and those recorded using other g-vectors (Table 2).

The mode of dissociation is found to involve the

formation of an intrinsic stacking fault (relative to

the Thompson tetrahedron for an ordinary f.c.c.

lattice), as the simulated images associated with the

extrinsic stacking fault (not shown) caused by dis-

sociation of the dislocation (with respect to a disor-

dered f.c.c. lattice) were not in agreement with the

experimental images. (It is more appropriate to

report the nature of the stacking fault in terms of

the disordered f.c.c. lattice and not the ordered glattice as will be discussed in the next section). The

dislocation labeled C in Fig. 3 was determined to

have bc=21/2[101], a line direction of [211] (308mixed) and hence lying on the (111) plane. An APB

can be seen attached to this dislocation.

Figure 5 shows another region containing a

stacking fault and APBs. Contrast analysis of the

fault labeled A reveals that in this case the fault

plane is (111) and the bounding partials labeled A1

and A2 have Burger's vectors given by bA1=21/

x[211] and bA2=21/x[121], respectively, where

x = 3 or 6; the line directions were determined by

stereographic trace analysis to be [213] and [112],

respectively. For instance, partial A1 is visible with

g= 200 [Fig. 5(d)], but invisible (weak in contrast)

with g = 002 [Fig. 5(a),(c)], g = 111 [Fig. 5(e)] and

g= 111 [Fig. 5(f)], whereas partial A2 is visible

with g = 111 [Fig. 5(e)], but invisible (weak in con-

trast) with g= 002 [Fig. 5(a),(c)], g= 200 [Fig. 5(d)]

and g= 111 [Fig. 5(f)]. This analysis was self-con-

sistent with the contrast observed in the images

recorded utilizing other g-vectors (Table 3).

Further con®rmation of the determined Burgers

vectors of the partials and the SF was obtained by

comparing, as before, experimental images recorded

under two di�erent conditions [Fig. 6(a),(b)Ði]

with simulated images [Fig. 6(a),(b)Ðii±iv], assum-

ing the fault is produced by the dissociation of

1/2[110] [Fig. 6(a), (b)Ðii], 1/2[101] [Fig. 6(a),

(b)Ðiii], and 1/2[011] [Fig. 6(a),(b)Ðiv] unit dislo-

cations. The best match is obtained for the ®rst set

of simulated images [Fig. 6(a),(b)Ðii], correspond-

ing to a dissociated 1/2[110] dislocation leading to

an intrinsic stacking fault (with respect to a disor-

dered f.c.c. lattice), but not the other possible unit

dislocation 1/2[101] [Fig. 6(a),(b)Ðiii] and 1/2[011]

[Fig. 6(a),(b)Ðiv] or superdislocation (not shown)

con®gurations on the (111) plane. The dissociation

in this case proceeds as follows:

Table 2. Experimentally observed contrast$ from fault A and associated partials A1, A2 in Fig. 3 and deduced fault displacement anddislocation burgers vector

Operatingre¯ection (g) Beam Direction (B) Partial dislocation A1 Partial dislocation A2 Stacking fault A Figure No.

002 0[010] W/I V V 3(a), 3(b)200 0[010] V W/I V 3(c), 4(a)Ði111 0[110] W/I V V 3(d)111 0[110] I I I Ð202 0[010] V V I 4(b)Ði220 0[110] W/I V V Ð022 0[011] I V I Ð202 0[010] W/I W/I V Ð

21/x[211] 21/x[112] Rf=1/3[111]

$W= weak; I = Invisible and V = Visible

WANG et al.: OBSERVATIONS AND ANALYSES 21

1

2�110� ! 1

6��211� � SF� 1

6��12�1�

An important feature to note is that unlike the situ-

ation with the stacking faults in Fig. 3, APBs are

seen to be attached to both the partials bounding

the stacking fault in Fig. 5. These APBs can be seen

clearly in the images in Fig. 5(b),(f). The other ends

of these APBs terminate at a grain boundary visible

at the left side of the micrographs in Fig. 5. Notice

Fig. 4. Experimental (i) and simulated (ii±v) BF image pairs of the fault labeled A in Fig. 2 under twodi�erent di�racting conditions: (a) g= 200, B0[010] and (b) g=202, B0[010]. The con®gurations corre-sponding the dissociation of di�erent dislocations (unit and super-) on the (111) plane are shown in thesimulated images: (ii±ii) b= 1/2[101], (iii±iii) b= 1/2[110], (iv±iv) b = 1/2[011] and (v±v) b= [101].The Burgers vector of the bounding Shockley partial dislocations in each set of the simulated imagesare also shown labeled in between the images. The best match with the experimental images (i±i) is

obtained for the set of simulated images in (ii±ii).

WANG et al.: OBSERVATIONS AND ANALYSES22

that in both cases (Figs 3 and 5) residual contrast is

observed outside the bounding partials owing to

APBs that had attached to the partial dislocations.

Attempts were made to simulate this contrast but

were largely unsuccessful as it is quite apparent that

the displacements associated with these non-conser-

vative thermal APBs are not entirely in the plane of

the stacking fault, and moreover, the displacement

vectors may not be exactly 1/2h101i. This aspect

will be addressed in a later communication. The dis-

location labeled B in Fig. 5 was determined to have

bB=21/2[101], a near-screw character and hence

lying on the (111) plane. The end of an APB can

also be seen attached to this dislocation, with the

other end terminating at the neighbouring grain

boundary.

Another point to note is that most of the SFs

observed were monolayer faults, i.e. non-overlap-

ping and not twins. A few overlapping faults were

also observed, but for those that were analyzed, the

Fig. 5(a,b).

WANG et al.: OBSERVATIONS AND ANALYSES 23

partial dislocations bounding the faults were deter-

mined to have the same 1/6<112] Burgers vector,

similar to the situation observed in deformed TiAl

[23]. When h110i SAD patterns were recorded with

these faults edge-on, extra twin re¯ections at the

expected 1/3{111} positions could be seen, together

with streaking along the same h111i direction owing

to the thinness of the twins. Through SAD, it was

also established that no retained a2 layers existed in

the gm grains in regions without and with stacking

faults. Also, the a2 matrix near the gm was observedto contain, on occasion, a few SFs on [0001] planes,

but to be relatively free of dislocations.

4. DISCUSSION

The results of this study have brought to lighta number of new characteristics of the defectstructures in the massive gm phase that have notbeen reported before. Naturally, how these

Fig. 5(c,d).

WANG et al.: OBSERVATIONS AND ANALYSES24

defects originate during the a 4 gm transform-ation requires explanation, and this is considered

below.

According to Hug et al. [21, 22], there are threetypes of SFs that can be de®ned in the L10 structure

of g-TiAl: Complex Stacking Fault (CSF),Superlattice Intrinsic Stacking Fault (SISF) and

Superlattice Extrinsic Stacking Fault (SESF). The

CSF can be thought to be a superimposition ofan SISF + APB. The possible dissociation of a

Fig. 5. TEM micrographs showing another area in the massive gm phase in the quenched Ti±46.54Alalloy containing a stacking fault (labeled A), the bounding Shockley partial dislocations (labeled A1,A2) and APBs. (a); (b) BF/SLDF images pairs recorded using g= 002/001 (B0[110]); (c)±(f) WBDFmicrographs recorded using (c) g= 002 (B0[110]); (d) g= 200 (B0[011]), (e) g = 111 (B0[011]) and (f)

g= 111 (B0[110]). The unit dislocation labeled B has b =21/2[101].

WANG et al.: OBSERVATIONS AND ANALYSES 25

Table 3. Experimentally observed contrast* from fault A and associated partials A1, A2 in Fig. 5 and deduced fault displacement anddislocation burgers vectors

Operating re¯ection (g) Beam direction (B) Partial dislocation A1 Partial dislocation A2 Stacking fault A Figure No.

002/001 0[110] W/I W/I V 5(a,c)/5(b)200 0[011] V W/I V 5(d)111 0[011] W/I V V 5(e), 6(a)±i111 0[110] W/I W/I V 5(f)111 0[110] I I I Ð202 0[111] W/I V V 6(b)±i022 0[011] I V I Ð220 0[110] W/I W/I V Ð

21/x[211] 21/x[121] Rf=1/3[111]

*W= Weak, I = Invisible, and V = Visible

Fig. 6. Experimental (i) and simulated (ii±v) BF image pairs of the fault labeled A in Fig. 4 under twodi�erent di�racting conditions: (a) g= 111, B0[011] and (b) g = 202, B0[111]. The con®gurations cor-responding to dissociation of di�erent unit dislocations on the (111) plane are shown in the simulatedimages: (ii±ii) b= 1/2[110], (iii±iii) b= 1/2[101], and (iv±iv) b = 1/2[011]. The Burgers vector of thebounding Shockley partial dislocations in each set of the simulated images are also shown labeled inbetween the images. The best match with the experimental images (i±i) is obtained for the set of simu-

lated images in (ii±ii).

WANG et al.: OBSERVATIONS AND ANALYSES26

1/2[110] unit dislocation and a [101] superdisloca-

tion [on (111)] are:

1

2��110� ! 1

6��211� � SF� 1

6��12�1��1�

�101� ! 1

2�101� � �APB� � 1

2�101� ! 1

6�112��2�

� �SISF� � 1

6�2�11� � �APB� � 1

6�112� � �CSF�

� 1

6�2�11�

Hug et al. [22] argued that, owing to the extremely

high APB energy in g-TiAl, the [101] dislocations

are expected to be present as pairs of narrowlyspaced 1/2[101] dislocations linked by APBs; each

1/2[101] dislocation may further dissociate intoShockley partials connected by an SISF or CSF as

shown above. Experimental observations [22, 24±27]indicate that the separations between the partial dis-

locations bounding SISFs and APBs are less than 6nm and 4.5 nm, respectively, whereas CSFs have

not been observed to date. The 1/2[110] dislo-cations, which do not disrupt order and hence have

no APBs associated with them, never appear disso-ciated because of the high CSF energy; this feature

is con®rmed by both weak-beam [22, 25±29] andhigh-resolution [29, 30] TEM observations of these

dislocations.

The evidence presented in this study has estab-lished that 1/2<101] dislocations (A±B, C±D, E±F

and G±H in Fig. 1; A1, B1, C1 in Fig. 2, C inFig. 3, B in Fig. 5) are present as widely separated

single dislocations (e.g. projected separations >150nm for the A±B pair in Fig. 1). These 1/2<101]

unit dislocations destroy the order in TiAl andhence are always observed to be connected by

APBs. Furthermore, the dissociations of these 1/2h101i dislocations on {111} planes create wide SFs

bound by 1/6h211i Shockley partial dislocations of

all possible types (Figs 3±6, projected widthsbetween partials >400 nm). For instance, the SF

marked A in Fig. 3 has been created presumably bythe dissociation of an 1/2[101> dislocation on the

(111) plane into 1/6[112] and 1/6[211] Shockley par-tial dislocations, whereas the SF in Fig. 5 has been

created by the dissociation of a 1/2[110] dislocationon the (111) plane into 1/6[211] and 1/6[121]

Shockley partial dislocations. Although the dis-sociation of an 1/2<101] dislocation in ordered

TiAl can create either a SISF or CSF [reaction (2)above], it is quite apparent that the projected separ-

ations between the partials (>400 nm) in theseimages (Figs 3 and 4) for either a SISF or CSF are

extremely large when compared with those forSISFs (3±6 nm) reported in deformed samples

[22, 25±28].

Though dislocation reactions in deformed g-TiAlhave been studied and reported extensively, dislo-

cation dissociations with such high superpartial/superShockley partial separations with an APB/SF

have not been observed before. The exceptions are

stacking fault dipoles [21, 22, 26±28, 31, 32], bound

by single or double Shockley partials having the

same 1/6 < 112] Burgers vector, and overlapping

SFs (mechanical twins) [23]. The high CSF or APB

energy in g-TiAl [33±36] strongly inhibits the cre-

ation of such large partial dislocation separations

with SFs or APBs in deformed samples. Since there

is currently no model that can satisfactorily explain

the possible mechanism for the occurrence of these

``unexpected'' dislocations and their dissociations

observed in ordered TiAl in this study, alternative

explanations for their formation are required, as

discussed in the following.

Two possible scenarios can be initially con-

sidered: the ®rst based on the assumption that these

defects were inherited from the parent, high-tem-

perature a, and the second [3, 5, 6, 10] that they

were created during the formation of ordered gmdirectly from a on cooling. TEM examination of

the remaining a2 in the neighbourhood of gm grains

showed practically no dislocations. A few SFs, lying

only on the (0001) planes, were observed in the a2,whereas in gm the stacking falults were seen on all

four {111} planes. This implies that the defects

observed in the gm phase do not exist in the parent

a phase before transformation. Therefore, the ®rst

scenario which assumes that the observed defects

survive through the a 4 gm massive transformation

and are frozen in the gm grains appears to be inva-

lid. With regard to the second scenario, the high

APB and CSF energies would be a barrier for the

introduction of these widely separated unit dislo-

cations and wide stacking faults, respectively, by

thermal or mechanical means during the formation

of ordered gm directly from the a, as it appears to

be an obstacle for the formation of these defects in

deformed g-TiAl. Thus, it appears unlikely that the

1/2<101] dislocations, the associated APBs and the

SFs produced by the dissociation of the 1/2<101]

and 1/2<110] dislocations were introduced in the

ordered gm phase either during its formation

directly from the a phase or subsequently on cool-

ing after it had already formed. This conclusion is

also supported by the fact that widely extended

defects of the types reported herein have not been

observed previously in ordered g-TiAl alloys

deformed at either low (to ÿ1968C) or high

(>8008C) temperatures [22, 26, 28, 29, 37].

Therefore, a new model for the massive trans-

formation in TiAl alloys is proposed based on the

observations and the discussion detailed above.

Instead of the single step a 4 gm reaction, the

model postulates that a disordered f.c.c.-type phase

is formed as an intermediate phase between the

parent a and the product gm phases, and the trans-

formation then proceeds as follows:

WANG et al.: OBSERVATIONS AND ANALYSES 27

a ÿÿÿ4�massive transformation�disordered f.c.c. phase

ÿÿÿ4�ordering transformation�gm

Based on this model, it can be postulated that the

1/2h101i dislocations and SFs observed in the gmphase are introduced in the ®rst step of the reaction

during which the high-temperature disordered aphase massively transforms to the high-temperature

intermediate disordered f.c.c. phase. Although the

possible methods by which these defects are intro-

duced in the disordered f.c.c. phase have not been

entirely considered yet, it is assumed that they are

created by chance during the change in structure

from h.c.p. to f.c.c. Since this intermediate f.c.c.

phase is disordered, all the 1/2h101i unit dislo-

cations on {111} planes are equivalent. In the sec-

ond step of the reaction, the intermediate

disordered f.c.c. phase undergoes an ordering trans-

formation to the L10g phase, thus making the 1/

2<110] and 1/2<101] unit dislocations non-equiv-

alent in the ordered structure. The 1/2<110] unit

dislocations do not destroy the order and can be

present individually (e.g. those labeled I1, I3 and

J2, J3 in Fig. 1 and set D in Fig. 2), without being

attached to APBs. On the other hand, 1/2 < 101]

unit dislocations will destroy the order, so APBs

are created in their wake as the ordering front

sweeps past them. The other end of these APBs will

either terminate at another 1/2 < 101] unit dislo-

cation (A±B, C±D, E±F pairs in Fig. 1) or at an

interface such as a intervariant g domain or low-

high-angle boundary (Fig. 2) to restore the order.

Several examples, with TEM evidence, can be

found in the literature pertaining to order±disorder

transformations in other alloy systems, of attach-

ment of APBs to unit dislocations present in the

disordered phase [38±41]. In these cases, it is

deduced that the unit dislocations were originally

present in the disordered phase and that during the

subsequent ordering process on aging, APBs

become attached to these dislocations to restore

order.

The same situation holds good with the stacking

faults bound by 1/6h211i Shockley partials, which,

as shown in Section 3, Figs 3±6, are produced by

the dissociation of 1/2h101i dislocations on {111}

planes in the disordered f.c.c. phase. The TEM evi-

dence suggests that the SF energy of the disordered

f.c.c. phase is quite low, especially at the elevated

temperatures where the massive transformation

takes place (1100±11308C [42, 43]). This allows the

easy splitting of the 1/2h101i unit dislocations into

widely separated (>400 nm) Shockley partial dislo-

cations, similar to the situation in other low-SF

energy f.c.c. alloys. For the sake of discussion,

these SFs have been assumed to have been pro-

duced by dissociation of 1/2h101i dislocations,

although a priori there is no reason why they could

not have formed directly by errors in atomic stack-

ing during the a4 f.c.c. massive reaction. The lat-ter appears to be the likely source given the highgrowth rates of the gm phase [43].

The geometry in the simulated images in Figs 4and 6 that gave the best match indicate that the SFsare intrinsic in nature relative to the disordered

f.c.c. lattice. It should be recognized that in theordered TiAl lattice these faults could become high-

energy CSFs. In the case of the SFs in Figs 3 and 4,since the net unit dislocation comprising the twoShockley partials (1/6<112] and (1/6<211]) has

Burgers vector (1/2<101]) that is one half of a<101] superdislocation, order is destroyed whenthe disorder±order interface encounters these

defects. Examination of the atomic stacking on the(111) plane of ordered TiAl, the plane of the SF,

reveals that the 1/6<112] displacement does notchange the nearest neighbor environment aroundthe atoms, whereas the 1/6<211] displacement

does. Consequently, to restore order, APBs are cre-ated at one end of the fault where the 1/6<211]Shockley partial resides, the other end of the APBs

terminating at another Shockley partial associatedwith another fault, at another 1/2<101] unit dislo-

cation, or at an interface such as an intervariant gdomain or other type of boundary. After this pro-cess, the SF would be expected to be a pure geo-

metrical fault (SISF). The random introduction ofdislocations (undissociated or dissociated into SFswith 1/6<211] type Shockley partials) in the disor-

dered f.c.c. phase during the ®rst step of the reac-tion, together with the multi-directional progress ofthe ordering process thereafter could account for

the highly curvaceous and meandering nature of theAPBs that become intimately attached to the dislo-

cations.Interestingly, and unexpectedly, APBs were also

observed to be attached to each of the 1/6[211] and

1/6[121] Shockley partial dislocations bounding theSF produced by the dissociation of a 1/2[110] unitdislocation (Fig. 5), with other partial dislocations

or intervariant domain/grain boundaries. As men-tioned above, the 1/2<110] unit dislocations do

not destroy order in TiAl and can (and should, inprinciple) be present without being linked by APBs(e.g. those labeled I1±I3, J1±J3 in Fig. 1 and set D

in Fig. 2). These undissociated dislocations couldhave been introduced either in the disordered f.c.c.or in the gm phase after the ordering transform-

ation. However, the dissociation of a 1/2[110] unitdislocation into widely separated 1/6[211] and 1/

6[121] Shockley partial dislocations on the (111)plane with the SF in between (Fig. 5) is deduced tohave occurred in the disordered f.c.c. phase prior to

the ordering reaction, since the high CSF energy inthe ordered g phase would prevent wide separationof the Shockley partials.

One possible explanation for the attachment ofAPBs to both the component 1/6<211] partial dis-

WANG et al.: OBSERVATIONS AND ANALYSES28

locations is as follows. Since the wide SFs produced

by the dissociation of the 1/2<110] dislocations inthe disordered f.c.c. phase would become high-energy CSFs [reaction (1) above] after the ordering,

they will have a tendency to reduce their energy.APBs that exist in the vicinity after the orderingreaction may have some mobility at the high tem-

peratures and hence be ``attracted'' to these high-energy faults. Presumably, after the APBs join the

CSF the local atomic displacement of the lattice issuch that the stacking in the faulted region re-sembles a pure geometrical fault (with lower

energy). A better explanation is that the APBs arecreated as the ordering front passes through thesepre-existing wide SFs with 1/6<211] Shockley par-

tial dislocations, around which, otherwise, local vio-lations of the nearest neighbor atomic con®guration

would occur. Since the direction of movement ofthe disorder±order interface is probably randomand not necessarily parallel to the plane of the SF,

the APBs are left attached from the two bounding1/6<211] partial dislocations to restore the nearestneighbor environment and order locally around the

dislocations. In both cases (splitting of 1/2<101] or1/2<110]), the SFs after interaction with the APBs

from the ordering are expected to be lower energy,pure geometrical faults. The only situation wherethe SFs would become CSFs is when the atomic

displacements caused by the motion of the disorder±order interface are exactly parallel to and in theplane of the SF, but this is thought to be a rare

occurrence. Further work is required to fully estab-lish and understand these features.

It should be noted that the occurrence of an in-termediate disordered f.c.c. phase during thea 4 gm transformation has also been hypothesized

recently by other researchers [6, 11, 12]. Forinstance, Denquin and Naka [6] proposed that the

APBs could be produced either by a directa 4 ordered gm transformation or through the me-diation of a disordered f.c.c. phase during this

transformation. They were, however, unable tostate which of these two processes was more likelyto occur. More recently, Zhang et al. [11, 12] expli-

citly proposed the occurrence of an intermediatedisordered f.c.c. phase during the a 4 gm massive

transformation to account for the APBs in the gmphase. However, their proposal that the existence ofAPBs and three gm variants alone is the evidence

for the occurrence of an intermediate disorderedf.c.c. phase prior to ordering is not entirely justi®ed,since these could also be created by impingement of

growing ordered gm domains formed directly in thea phase [3, 6]. Our interpretation of the mechanism

of formation of these APBs di�ers from that ofthese authors [11, 12]. In particular, the existenceand role of widely dissociated 1/2<110] unit dislo-

cations, and 1/2<101] unit dislocationsÐundisso-ciated or widely dissociated into 1/6h211i Shockleypartial dislocations bounding stacking faultsÐwas

not entirely recognized until the present study.Indeed, the occurrence of these unexpected defects

in the gm phase provides, perhaps, better evidencefor the occurrence of an intermediate disorderedf.c.c. phase during the a4 gm transformation. The

APBs, which are intimately attached to these dislo-cations, are then just natural by-products from thesubsequent f.c.c. to g ordering reaction that passes

through these dislocations.Detailed observations and characterization of the

APBs, the structure within them, and the inter-

actions between APBs with dislocations/SFs, as wellas a discussion of the ordering process, will be pro-vided in a follow-on paper [13], which will furthervalidate the new model for the formation and evol-

ution of these defects during the a4 gm massivetransformation in TiAl alloys.

5. CONCLUSIONS

In conclusion, the defect structures in the massivegm phase in a rapidly quenched Ti±46.5 at.% Alalloy have been studied by TEM. The defect struc-

tures are found to be composed of dislocations, SFsand APBs intimately associated with dislocations orSFs. Both 1/2<110] and 1/2<101] unit dislo-cations were present in the gm phase, with the latter

linked by highly curved non-conservative APBs.Comparison of experimental and simulated TEMimages shows that SFs, which are created by the

dissociations of 1/2h101i unit dislocations, lie on{111} planes and are bound by well-separatedb= 1/6h121i Shockley partial dislocations of all

possible types. In addition, APBs are found to com-mence or terminate on the stacking faults only atthe partial dislocations with b =1/6<121], but not

those with b= 1/6<112]. The evidence and argu-ments presented in this study suggest that thesedefects are created through the occurrence of an in-termediate disordered f.c.c. phase during the

a4 gm massive transformation in TiAl alloys.

AcknowledgementsÐSupport for this research from theNational Science Foundation (DMR-9224473, B.MacDonald, Program Monitor) and the WrightLaboratories, Materials Directorate/UES (S-269-000-002/SUB AF, M. Mendiratta, Program Monitor), is deeply ap-preciated. The authors also wish to deeply thank K. J.Hemker at Johns Hopkins University and R. Schaublinand P. Stadelmann of EPFL (Switzerland) for allowingaccess to the simulation program. One of the authors(MK) is thankful for the ®nancial support received as apost-doctoral fellow with K. J. Hemker during the courseof this work.

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WANG et al.: OBSERVATIONS AND ANALYSES30