Materials Science and Engineering, B I I (1992) 325-330 325

Transmission electron microscopy studies of dislocation mechanisms in as-sintered a-SiC and after creep experiments at high temperature

S. J. Lee and J. Vicens LERMA T, ISMRa, URA CNRS 1317, Boulevard du Mar~chal Juin, F-14050 Caen Cedex (France)

Abstract

Hot-pressed SiC with small additions of aluminium has been deformed at 1600 °C by three-point bending and studied by transmission electron microscopy. This paper reports observations of Shockley partials ~ (10 i 0). These dislocations result from the activation of dissociated dislocation sources located in the basal plane. Evidence of dissociated dislocation climb has been found and two different climb mechanisms are proposed. Defects in the as-sintered material have been studied and two examples are presented.

1. Introduction

Many of the technological problems as well as much of the theoretical interest associated with SiC lie in its numerous polytypic forms. A lot of studies have been devoted to this problem and several theories have been proposed in order to explain the occurrence of polytypism in SiC. They have been reviewed recently [1-3]. Many of them (screw dislocation theory [4], faulted matrix model [5], periodic slip process [6], etc.) are based mainly on interactions between crystallographic defects in SiC crystals. On the other hand, the precise nature of these defects and their prop- erties are not well known. Furthermore, SiC is a compound with tetrahedral coordination of its atoms and it is of great theoretical interest to test the dislocation mechanisms already developed for semiconducting materials.

The high mechanical strength of SiC makes it one of the most important engineering ceramics for high temperature and high strength applica- tions. Many studies have been performed on the plasticity of polycrystalline SiC at high tempera- ture. For a long time, results of creep experiments have been explained mainly by diffusion mecha- nisms. However, more recent publications have shown a contribution of dislocation glide and/or climb in high temperature creep experiments [7-9] as well as below 1000 °C [10]. Furthermore, Fujita et al. [11] and Maeda et al. [12] have studied the plasticity of 6H single crystals up to 1650 °C,

showing the important role of dislocation glide. These authors have proposed a microplasticity mechanism in order to explain the single-crystal deformation. The aim of this work is to identify dislocation mechanisms occurring during creep experiments performed at 1600 °C in hot-pressed a-SiC sintered with a small amount of aluminium. In the first part of the paper, defect analyses in as- sintered a-SiC will be presented.

2. Materials and experimental techniques

Hot-pressed a-SiC samples with 0.3% AI- based additives (supplied by Elektroschmelz- werk, F.R.G.) have been deformed (three-point bending) in the temperature range 1300-1760 °C under vacuum. Slices of the as-sintered material as well as high temperature (1600°C) deformed specimens were mechanically ground and pol- ished down to a thickness of 80/zm. Discs 3 mm in diameter were ion milled (Ar ÷, 6 kV) for transmission electron microscopy (TEM) obser- vations (Jeol 120 CX and 200 CX).

3. Defect analyses in the as-sintered material

3.1. S t ruc tura l cons idera t ions

SiC exhibits a number of one-dimensional ordering sequences without any variation in stoichiometry. All SiC structures are made up of single basic units of tetrahedra SiC 4 or C4Si.

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Successive layers can be arranged in different ways, giving rise to cubic (C), hexagonal (H) or rhombohedral (R) unit cells. The most common polytypes observed are 4H and 6H (Ramsdell notation [13]), often reported as the a-SiC hexagonal form, and 3C, the fl-SiC cubic form. Polytype stacking sequences can be described by the a, b, c notation, where each plane (a, b or c) is a double atomic plane of silicon and carbon. The corresponding stacking sequences of 4H and 6H polytypes can be written following this notation as: 4H . . . . a b c b a . . . ; 6H . . . . a b c a c b a . . . . Both polytypes are mixed cubic-hexagonal stack- ing sequences (or mixed zinc blende-wurtzite structures).

Electron diffraction studies performed on our samples show that the carbide grains have crystal- lized as 4H and 6H polytypes. Only deformed specimens show a thin layer (approximately 10 nm) of 3C polytype in some areas located near grain boundaries [14]. In as-sintered specimens, many defects can be seen inside grains as a conse- quence of stresses induced during preparation and/or polytype transformations. Two defect analyses are presented in the following.

3.2. Extended dislocation network In the f.c.c, structure, constricted-extended

node pairs result from interactions between dis- sociated dislocations in the same glide plane [15] or in successive glide planes [16]. They have also been observed in the diamond structure [16]. Because cubic sequences are a part of 4H and 6H polytype stacking sequences, such configurations can be found in a-SiC. One example is shown in Figs. l(a) and l(b) and drawn schematically in Fig. l(c). The dislocation network of Fig. l(a) forms a twisted low angle grain boundary located in the (0001) plane. Contrast analysis of the total configuration [14] has revealed the existence of three families of screw dislocations, oA, o13 and oC (Fig. l(d)), with triangular meshes of faulted (black areas, Fig. l(b); F, Fig. l(c)) and unfaulted (white areas, Fig. l(b)) zones. Three types of nodes are indicated in Figs. l(b) and l(c) (denoted 1, 2 and 3). Type 1 and 3 nodes are due to ribbon interactions in the same (0001) plane located in the cubic sequence of the polytype, whereas two successive (0001) planes are respon- sible for type 2 nodes. Type 2 node formation is illustrated in Fig. l(d). The same mechanism has already been published for the f.c.c, structure [16]. Modifications of 4H stacking sequences are

shown in Figs. l(e) and l(f) after l(e) one or l(f) two successive shears have occurred in a 4H crystal. Consequently, as in cubic structures, triple ribbons consisting of three Shockley par- tials with the same Burgers vector can exist in a-SiC. They have been observed in deformed a-SiC crystals [14].

3.3. Stacking fault analysis" A high number of extended growth faults in the

basal plane can be imaged in our crystals. They can be caused by movements of Shockley partial dislocations. Thus in this case, two parts of the crystal are sheared by a ~(lOiO) fault displace- ment vector. This is not the case for the stacking fault (4H crystal) shown in Fig. 2, which has a displacement vector R = ~[0001]. This fault is in contrast in Figs. 2(a) and 2(b) (with a = 2~g 'R = - ~ ) and out of contrast in Figs. 2(c)-2(f), where only the two partials denoted 1 and 2 are (2(c), 2(d)) in strong contrast and (2(e), 2(f)) in residual contrast (see Table 1 ). It can be deduced that this growth fault (with interstitial or vacancy character) lying in the (0001) plane is located in the cubic stacking sequence. The same fault can exist in the hexagonal part of the sequence but with an addi- tional displacement vector R '=}{10]0) in order to avoid high energy fault formation. An equiva- lent defect has already been found in a 6H crystal with a fault displacement R"=~[2021] [17]. Furthermore, a stacking fault with R" = ), [0001] (6H crystal) has been identified recently after neutron irradiation experiments [ 18].

4. Defect analysis in deformed specimens

4.1. Activation of the (0001) glide plane Shear slip along basal planes was found to

occur after deformation at 1600 °C. Figure 3 is a weak beam image of Shockley partials (denoted 2 and 3) observed in their glide plane. They result from dissociated dislocation source activation located in (0001) planes. Dislocation lines are frequently straight or slightly rounded and have a tendency to be aligned close to (1120) or (1010) directions. Thus Shockley partials in a-SiC can be classified into four different types according to their character (0°-60 °, 60°-60 °, 30°-30 °, 30°-90°). This suggests the presence of a high Peierls potential in the crystal. The stacking fault energy (approximately 8 mJ m-2) estimated from the measured separation distance between par- tials seems to be higher than the value in a

327

,Q2__~_,~m

#,- i

, e.| l im..

: : c I c

b .I. c j . b

c ~ Co" a o'C a a no-

a b

/ , c / 2 3 . ~ ."2- c c'ot " *'*" "~ " OB

I I %.o* a

=11~" = o ' C

Fig. 1. (a) Dislocation network of a twisted low angle grain boundary. The boundary plane is parallel to the basal plane of the 4H crystal. (b) Enlargement of part of (a) showing extended meshes (black areas) of the dislocation network. Extended nodes (denoted 1 and 2) and constricted nodes (denoted 3) are visible. Some cavities (P) have been trapped inside the dislocation network. (c) Schematic drawing of the outlined part of (b). Faulted areas (denoted F) and the three different types of nodes are indicated. (d) Formation of a type 2 node by interaction of two dissociated ribbons in two successive glide planes. A similar process has already been reported for extrinsic-intrinsic node pairs in the f.c.c, structure [ 16]. (e) Shockley partial dislocations oC and aB in the basal plane of a 4H crystal. Each letter of the stacking represents a double plane of silicon and carbon. The dis- sociation has occurred in the slip system. (f) A_ rray of dislocations in two successive glide planes occurring in the type 2 node of (d); oA=~[1 i00], oB =~[[010] and crC =~[0110].

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Fig. 2. Contrast analyses of a stacking fault lying in the basal plane of a 4H crystal. (a) Bright field image with g = 1 i02. The fault is in contrast. (b) Corresponding dark field image of (a). (c) Bright field image with g = 0004. In this orientation the fault plane is parallel to the electron beam and the fault i_s out of contrast. (d)-(f) Bright field images with g = 1214, 1210 and 2 i i 0 respectively. Only the partials are in contrast.

TABLE 1

Contrast analyses of a stacking fault with b = J[0001 ]; g b × u

and g . b values are given

g g ' b I × u I g ' b 2 × u 2 g ' b a g'b2

a, b 1 |()2 0.5 1.3 - 0.5 - 0.5 c 0004 0 0 1 I d 1212, 1.4 - 2 . 2 3 - 1 - 1 e ]2 ]0 1.4 - 2 . 2 0 0 f 2110 0 - (I.8 0 0

u, = [21101. u 2 = l i i 2o], b, = bz = J(00011.

deformed single crystal (2.5 mJ m 2) [12]. More- over, there is no large change in the shape of a dislocation line with respect to its character as has been observed in a deformed single crystal [12]. These differences between single-crystal and polycrystalline SiC behaviour might be attributed to impurities (aluminium) present in our speci- mens. Figure 3 also shows partials (denoted 4) recombined along a part of their line. This phe- nomenon will be analysed below.

4.2. Dissociated dislocation climb via jog pairs The dislocation configuration of Fig. 4(a)

imaged with g= 1120 consists of two 300-90 ° Shockley partials which have recombined along their lines in parts A, B and C. Contrast analyses have revealed that these three segments are in the same (10]0) plane. They are respectively parallel to [3632] (parts A and B) and [2421] (part C). Moreover, we can see a redissociation of partials in an opposite sense between segments A and C. This configuration has been drawn in Fig. 4(b). Constriction and inversion of Shockley partials clearly indicate a climb motion of basal disloca- tions. In this climb mechanism the dissociated dislocations recombine, climb over the length of the constriction in a prismatic plane and dissoci- ate again. This reaction is a pure climb by jog pair nucleations. In this case there are no residual dislocations associated with the climb motion.

4.3. Climb mechanism by vacancy loop nucleation Another climb mechanism has been found in

a-SiC after high temperature deformation [19]. It is illustrated in Fig. 5. A high density of loops (15-50 nm in size) is observed around disloca- tion lines of the two partials 1 and 2 (0°-60°). A strong interaction exists between these small loops and dislocation lines. Small loops (denoted L) can be seen to have nucleated on only one partial (partial 2). Some loops (T) have taken the form of a triangular projection with sides parallel to (1120). In another area, loops (R) detach from the dislocation line. These observations can also be explained by a model involving a climb mechanism based on reactions between vacancies and dislocations during high temperature defor- mation. The first step is to assume the nucleation of a perfect prismatic loop with vacancy character on the dislocation line. Three Burgers vectors are possible for these loops, namely AB=~[2110], BC=~[l l20] and CA=~[1210]. Reactions between prismatic loops and the four types of

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Fig. 3. Weak beam image of Shockley partial dislocations observed in a 4H crystal after a creep experiment at 1600 °C. The Burgers vector directions are indicated for each of the partials: 1, o13; 2, aA; 3, oC; 4, CA. Partial 1 is out of contrast.

b

Fig. 4. (a) Bright field image of two partials in a 6H crystal. The partials have. recombined along parts of their lines denoted A, B and C. The Burgers vector directions are indicated. A redissociation of partials in an opposite sense is observed. (b) Schematic representation of (a).

partials (00-60 ° , 300-30 ° , 300-90 ° , 600-60 ° ) have been studied assuming that loop nucleation is only possible if it corresponds to a minimum acti- vation energy [20]. This model shows that partials 0°-60 ° and 60°-60 ° are the most suitable for prismatic loop nucleation.

Figure 6 illustrates the different steps of this climb process, which has already been applied to cubic crystals [20-22]. Loops AB and CB which

have nucleated on partial 2 are indicated in Fig. 6(a). Parts of loops situated in the basal plane are dissociated into Shockley partials. Accordingly, these loops form a chain of dipoles (Fig. 6(b)) which can react with partial 1 (Fig. 6(c)). Closed loops can be formed by this climb motion, leaving a complex loop in the matrix (Fig. 6(d)).

5. Conclus ions

(1) An extended dislocation network has been analysed in the as-sintered specimen. It contains constricted-extended nodes. An analysis of this configuration in a 4H crystal assumes interactions between Shockley partials situated in two succes- sive layers of the cubic stacking sequence. A stacking fault with R=¼[0001] has also been found in this specimen.

(2) TEM observations of high temperature deformed a-SiC (1600 °C) indicate activation of the basal glide plane. !

(3) The dissociated dislocations are Shockley partials which have four possible characters, i.e. 0°-60 °, 30°-30 °, 30°-90 ° and 60°-60 °.

(4) Climb motion of basal dissociated disloca- tions occurs. A mechanism has been proposed which assumes pure climb via jog pair formation in prismatic planes.

330

Fig. 5. Weak beam image of two partials 1 and 2 (00-60 °) in a 4H crystal with a high density of small loops along the dislocation lines. Some loops (T, L and R) are arrowed.

, 1 , o c

~A o,e

A'O (~O BO BO q Ba

a b c Fig. 6. Schematic drawing of the climb mechanism and loop nucleations on partials 1 and 2 (0°-60°).

| oC

B o

d

(5) Another climb mechanism has also been found. In this case dissociated dislocations can act as loop sources. This same process has already been applied to cubic crystals.

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