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Room-temperature deformation mechanisms and the defect structure of tungsten carbide
A<,” Mrrolfurg,~
Vol. 29. pp. 1645 lo 1654. 1981 Prmted in Great Bnlam. All rlghls reserved
~1-6l60/~l/o91~5-10S02.0010 Copyright 0 1981 Pergamon Press Ltd
ROOM-TEMPERATURE AND THE DEFECT
DEFORMATION MECHANISMS STRUCTURE OF TUNGSTEN CARBIDE
M. K. HIBBS
and R. SINCLAIR
Department of Materials Science and Engineering, Stanford University, Stanford, CA 94305, U.S.A. (Received 9 February 1981) Abstract-Defects in undeformed, annealed single crystals and in deformed single crystals of WC are characterized by transmission electron microscopy. Deformation is induced by indentation with a microhardness tester. High densities of stacking faults with atomic displacement vectors R = l/6(1123) are found lying on 1liO0) planes near the indentations. Stacking faults with R = l/6( 1123) and undissociated dislocations with b = (OCOl). l/3( 1120), and l/3( 1123) are observed at distances greater than 2 pm from an indentation. The Burgers vectors of the partial dislocations bounding one fault are found to be identical b, = 1/6[1123], resulting in a total Burgers vector for the extended dislocation b = 1/3[1123]. (Stacking fault energy = 55 mJ/m’). Stacking faults with R = l/6(1123) are described in terms of the crystal structure of WC, illustrating several points: the coordination of carbon atoms by tungsten atoms is preserved at the fault, the number of bonds per atom which are broken during slip may be minimized, and the l/6( 1123) type of partial Burgers vector may be combined to form several types of total Burgers vector.
avons caractkrisk par microscopic Clectronique en transmission les dtfauts dans des monocristaux de WC recuits, ainsi que dans des monocristaux d&form&. La dtformation est produite par indentation &I’aide d’un appareil de microdurett. On observe au voisinage des empreintes une forte densit de c!kfauts d’empilement dans des plans (liOO), avec des vecteurs de d&placement atomique R = l/6(1123). On observe d’autre part des dkfauts d’empilement avec R = l/6(1123) et des dislocations non dissocites avec des vecteurs de Burgers b = (Oool), l/3( 1120) et l/3( 1123) ?t des distances de l’empreinte suptrieures g 2 pm. Le vecteur de Burgers des dislocations partielles bordant une faute est &gal g b, = l/6(1123), ce qui conduit g un vecteur de Burgers total b = l/3(1123) pour la dislocation dissociire (tpergie du dCfaut d’empilement: 55 mJ/m’). Nous discutons les dCfauts d’empilement avec R = l/6( 1123) dans la structure de WC et nous montrons les r&hats suivants: les atomes de carbone conservent le nombre d’atomes de tcngstene voisins sur le dtfaut d’empilement, on peut minimiser le nombre de liaisons par atome rompues au tours du glissement et I’on peut combiner des vecteurs de Burgers de dislocations partielles l/6(1 123) pour obtenir plusieurs types de vecteurs de Burgers pour les dislocations parfaites. RbumGNous
Zusammenfassung-Die Defekte, die in unverformten gegltihten und in verformten WC-Einkristallen auftreten, werden mittels Durchstrahlungselektronenmikroskopie analysiert. Verformt wurde mit dem Stempeleindruck eines Mikrohlrtegerltes. Es wurde eine hohe Dichte von Stapelfchlern aufgefunden, deren Verschiebungsvektor R = l/6( 1123) war und die nahe dem Eindruck auf {llOO}-Ebenen lagen. AuBerhalb einem Abstand von 2 pm vom Eindruck wurden Stapelfehler mit R = l/6( 1123) und nicht aufgespaltene Versetzungen mit Burgersvektoren b = (OOOl), l/3( 1120) und l/3( 1123) beobachtet. Der Burgersvektor der Teilversetzungen war immer b, = l/3( li23), welches zu einem Burgersvektor der vollstlndige_n Versetzung von b = l/3( 1123) fiihrt. (Stapelfehlerenergie = 55 mJ/m*). Stapelfehler mit R = l/6( 1123) werden anhand der WC-Kristallstruktur beschrieben. Einige Punkte werden hervorgehoben: die Wolframumgebund der Kohlenstoffatome bleibt am Stapelfehler erhalten, die Anzahl der bei Gleitung aufzutrennenenden Bindungen kann minimalisiert werden, und die Burgersvektoren der Teilversetzungen vom Typ l/6( 1123) lassen verschiedene Typen von perfekten Burgersvektoren zu.
1. INTRODUCTION Tungsten carbide is the major constituent of cemented carbide cutting tools. Efforts are made to optimize the toughness, hardness and wear resistance of these composite materials by altering their composition and microstructure. However, the basic deformation and wear processes in the various phases are still not well understood. This paper describes the study of plastic deformation in WC and the character of the defects which result from it. Plastic flow is
induced in single crystals at room temperature by micro-indentation; the slip plane, slip direction, and nature of the dissociation of the dislocations are determined by transmission electron microscopy. This information is then related to the crystal structure of the material. 2. BACKGROUND The unit cell of WC is shown in Fig. 1. It is hexagonal with tungsten atoms located at 0, 0, 0 and
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n
OF
DEFORMATION
AND STRUCTURE
WC
n
OF TUNGSTEN
these indents in the final TEM specimen. The slip directions are identified and the nature of the dislocation dissociation is determined by complete, conventional contrast experiments [4]. A preliminary report of this work has appeared previously [S]. 3. EXPERIMENTAL
Glrboll
0
0 = 2.906 % c = 2.037 ii
Tungsten
c,o
_
o.g6
slip plone: (1010) Fig. 1. Unit cell of WC. a corresponds to the l/3( i2iO) directions, the cell edges and the cell face diagonal in the
horizontal plane. c is the [OOOl]direction, the vertical cell edge. The l/3( i2i3) directions correspond to the cell face diagonals. The { lOi0) slip planes are the vertical cell faces and the vertical diagonal plane through the unit cell.
carbon at l/3, 213, l/2. The unit cell dimensions are virtually identical, i.e. c/u _ 1. In terms of MillerBravais indices, a corresponds to the l/3( 1120) vectors and c to (0001). The conclusions of previous studies of WC deformation are conflicting. In order to deduce the slip system, Takahashi and Friese [l] indented single crystals on various faces. From observations of the resulting slip steps they determined the slip plane to be {liOO}. They further suggested that the slip directions were <1120) and (0001) since these are the directions of smallest lattice repeat. However, from a careful analysis of slip steps produced by a similar indentation process, Luyckx [2] proposed a slip direction of the type (1123). Johannesson and Lehtinen [3] performed electron microscope studies on tungsten carbide grains within WC-Co cemented carbides. They observed double dislocations and dislocation networks with rectangular nodes. From diffraction contrast experiments and comparison with calculated images, they concluded that the double dislocations and the nodes were extended (1123) superdislocations with partial dislocations of the type l/6(2023). The results reported in their paper were somewhat ambiguous, however, since only one dislocation disappearance criterion was presented for any particular dislocation, whereas two are necessary to unequivocally identify b. Therefore, while the WC structure is simple, the dislocation mechanism of plastic flow has not been established. The present article addresses this situation. The dislocation structure resulting from deformation is distinguished, from that which is either grown-in or present after sintering, by indenting single crystals grown from a melt and clearly locating
CARBIDE
PROCEDURE
Bulk specimens were prepared from mixtures of WC and Co powders melted in an argon atmosphere at 1923 K and then cooled at rates down to 2.5 K/h. WC single crystals were precipitated in a Co matrix by this process. The Co was then etched away with warm HCl. The best crystals were in the form of triangular plates about 3 mm on a side and 1 mm thick. The large faces of the plates correspond to the basal plane while the remaining three sides are (liO0) planes. Some of the crystals were mounted, polished on two different types of crystal faces (the basal plane and {1100) planes), and indented. The arrangement of slip steps which appeared after indentation was observed by optical microscopy. Both deformed and undeformed samples were prepared for transmission electron microscopy. Discs were spark-cut from the triangular plates, then mechanically ground and polished to a thickness of about 1OOpm. The samples were deformed by diamond indentation with a microhardness tester using a 50gm load. Fifty to seventy-five indents were positioned 1OOpm apart in a square array across the sample. All samples were thinned to electron transparency by ion milling. Undeformed samples were milled from both sides, while the deformed samples were thinned from the non-indented side in order to preserve the defects resulting from deformation. However, the indented sides of these samples were exposed to the ion beam for approximately one hour in order to remove surface deformation resulting from the polishing process. The foils were observed in a Philips EM400 operated at 120 kV. The exact positions of indents in the TEM thin foils were determined by comparing low magnification electron micrographs with optical micrographs of the prepared samples. 4. RESULTS 4.1 Undeformed material In the undeformed material, which had been annealed at 2073 K, undissociated dislocations were observed, often arranged in networks. One such network is shown in Fig. 2. In Fig. 2 (a) g = [1120], and all the dislocations are visible. The same area is imaged under different diffracting conditions, g = [OilO], in Fig. 2 fb), and the dislocations labelled A are invisible. Complete conventional contrast experiments were performed on this network [4] and are summarized in Table 1. All three sets of dislocations, A, B and C, were found to have Burgers vectors of the type l/3( 1120). Although the defect structure in the undeformed material was dominated by single dislocations with
HIBBS
AND
SINCLAIR:
DEFORMATION
AND STRUCTURE
OF TUNGSTEN
CARBIDE
(b) Fig. 2. Dislocation network in undeformed, an_nealed single crystal of WC. In (a) g = [1120], and all the dislocations are visible. In (b) g = [01 lo] and the dislocations labelled A are invisible.
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Table 1. Summary of the contrast experiment performed on the dis-
location network shown in Fig. 2 Burgers vector determination Reflection Visibility 1120 Oil0 O!11 AlOO 1101 lOA0 1011
all dislocations visible type A invisible type A invisible type B invisible type B invisible type C invisible type C invisible
type A: b = 1/3[2ijO]. type B: b = 1/3[1!20]. type C: b = 1/3[1210]. All dislocations were found to have Burgers vectors of the type l/3(1120).
b = l/3( 1120), low-angle grain boundaries, dislocation loops with b = (OOOl), and an occasional stacking fault were also observed.
4.2 Deformed material Figure 3 shows two optical micrographs of indentations in single crystal tungsten carbide. The indentation in Fig. 3(a) was made on the basal plane, while the sample shown in Fig. 3(b) was in the (1010) orientation. As noted by Takahashi and Friese Cl], the slip steps visible on the (Oool) crystal face indicate that some component of the slip direction must lie along [el]. Similarly, the presence of slip steps on the (1010) face indicates that a component of the slip lies out of the (lOf0) plane and therefore along a (1120) direction. It may be noted that (1123) directions fulfill this requirement. The indented side of a TEM specimen is shown in Fig. 4(a). The square arrangement of the indentations made it possible to correlate the hole visible in the low magnification electron micrograph in Fig. 4(b) with the indent indicated by the arrow in Fig. 4(a). A high density of stacking faults, each bound by a pair of partial dislocations, exists around the hole as may be observed from Fig. 5(a). In this orientation the faults are practically end-on, lying on (liO0) planes. The zone axis of the corresponding diffraction pattern, Fig. 5(b), is very close to the c-axis, and the streaks along the (1 iO0) directions arising from the faults confirm that they lie on (IiOO} planes. No undissociated dislocations are found in this region. At distances greater than 2 pm from the indentation a combination of dissociated and undissociated dislocations is observed (e.g. Fig. 6). The faults exhibit contrast similar to those near the indentation and also lie on ( liOO! planes. Undissociated dislocations with three different types of Burgers vector, (OOOl), l/3( 1120) and l/3( 1123), have been identified. Since these defects are located far from regions where macroscopic slip steps are visible it seems unlikely
AND STRUCTURE
OF TUNGSTEN
CARBIDE
that they are part of the plastic deformation resulting from indentation. They may arise from the difference in thermal contraction between the phases during the cooling of the newly formed WC crystals in the Co matrix, or may result from plastic deformation during the sample preparation process and not removed by ion-milling. Fault and partial dislocation identification also followed conventional procedures. A stacking fault imaged in both bright field and centered dark field modes (g = [lOlO]) is shown in Fig. 7. The contrast is symmetrical in both micrographs indicating that the fault is a n-fault, that is, c( = 2ng.R = 71 where R describes the lattice displacement across the stacking fault. Examples of the experiment performed to determine the displacement vector and the Burgers vectors of the bounding partial dislocations of the fault are found in Fig. 8. In Fig. 8 (a) g = [ioll], and the fringe contrast has disappeared while the partial dislocations are still visible. Under the diffracting conditions in Fig. 8(b), g = [liOO], the fringe contrast and both partial dislocations are invisible. A summary of the analysis is given in Table 2. The two partial dislocations were found to have the same Burgers vector, bl = b2 = R = 1/6[1123]. This results in a total Burgers vector for the defect b = 1/3[1123], confirming the macroscopic, slip step analysis of Luyckx [2]. This vector has components in both the [OOOl] and [ 1120] directions as was required by the information derived from optical microscopy. It also satisfies the criterion that a = 27rg.R = B (g = [iOlO], R = 1/6[1123]). Hagege et al. [6] have observed stacking faults with independently R = l/6( 1123) in the carbide grains of a WC-Co composite. Using the Burgers vectors given above for the partial dislocations and assuming a line vector along [OOOl],the stacking fault energy is estimated to be 55 mJ/m’. The particular fault whose analysis is summarized here was located at a distance greater than 5pm from an indentation. Many faults located near indentations were also analyzed and found to have the same type of fault vector, R = l/6(1123). However, the high density of defects in these areas made the determination of the Burgers vectors of the partial dislocations very difficult, but they are presumed to be the same. 5. DISCUSSION The results given above demonstrate that plastic deformation induced by indentation in WC results in a defect structure consisting of dislocations dissociated on {liOO} planes with fault vectors of the type l/6( 1123). Although previous optical studies [l, 21 determined the slip system of WC, they did not reveal the extended nature of the dislocations. Figure 9 contains a diagram of the perfect and faulted structures of WC projected onto a slip plane, (1100). Three layers of atoms are indicated. The bottom layer consists of tungsten atoms, the next of
HIBBS
AND
SINCLAIR:
DEFORMATION
AND
STRUCTURE
OF
TUNGSTEN
CARBIDE
a)
b)
Fig. 3.
A.M. 29/9--c
:a1 micrographs of indentations in single crystal WC. (a) Basal plane orient al orientation. (Courtesy D. J. Rowcliffe and S. Pattanaik).
(b) (lOi0)
I1649
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OF TUNGSTEN
CARBII )E
Fig. 4(a).
Fig. 4(b).
Fig. 4. (a) Optical micrograph of a deformed TEM specimen of WC. Indentations have been placed in a square array. The arrow points to the location of an indentation which appears as a large hole in the : low magnification electron micrograph in (b).
HIBBS
AND
SINCLAIR:
DEFORMATION
AND STRUCTURE
OF TUNGSTEN
CARBIDE
ta)
(b) Fig. 5. (a) A high magnification electron micrograph taken near the hole shown in Fig. 4(b). A high density of stacking faults may be seen lying almost parallel to the electron beam on jlloO) planes. (b) The corresponding diffraction pattern with the electron beam direction very close to the c-axis. The streaks along the (1100) directions confirm that the faults lie on { 1100) planes.
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Fig. 6. Example of the defect structure found at distances greater than 2 pm from an indentation. Both dissociated and undissociated dislocations are observed.
carbon, and the top layer again of tungsten. The faulted structure was formed by moving the top layer of tungsten atoms across the layer of carbon atoms with the displacement vector R = 1/6[1123]. This diagram illustrates several points. The first observation which may be made, and which was also noted by Hagege et al. [6], is that the 1/6[t 1233 fault vector preserves the trigonal prismatic coordination of the carbon atoms by tungsten atoms. In fact, at the fault the structure appears to have been merely rotated by 90” from the perfect structure. Since tung-
sten carbide is highly stoichiometric, with directional bonding between tungsten and carbon atoms [7], the structure of a low energy fault would be expected to preserve the number and the angular relationship of the tungsten-carbon bonds. Figure 9 may also provide an explanation for the fact that the { 1iO0) planes are the slip planes rather than the more closely packed basal planes. The strong covalent bonding implies that it would be desirable to break as few bonds as possible in the slip process. For slip to occur on the basal plane three tungstewar-
7 *..
O.OSpm
~ c
r* P
Fig. 7. (a) Bright field and (b) centered dark field pair of micrographs of a stacking fault taken under the diffracting condition g = [iOlO]. The symmetrical contrast in both micrographs indicate that the fault is a n-fault.
HlBBS
AND
SINCLAIR:
~~FG~~ATION
AND
STRUCTURE
OF ‘TUNGSTEN
CARBIDE
Fig. 8. Exampfes of the contrast experiment performed on the fault shown in Fig. 7. In (a) g = [iOI 11 and the fringe contrast has disappeared while both partial dislocations are still visible. In (b) g = [l MO] and the fringe contrast and both partial dislocations are invisible.
bon bonds must be broken per carbon atom. However, as was noted by Takahashi and FrieseCl] and may be seen from the diagram, the possibility of breaking only two tungsten~arbon bonds per carbon atom exists for slip on { 1iGQ] planes. This will occur if disio~atjons move between planes equivalent to the layer of carbon atoms and the top layer of tungsten; motion between the carbon atoms and the bottom layer of tungsten atoms would result in four broken bonds per carbon atom. The third point which may be made from the diagram is that the 1/6[11?3] partial Burgers vector could give rise to more than one type of totai Burgers vector. The following combinations could exist on the slip plane shown: 1/6[11233 + 1/6fii23]
(h2 = 0.0805 nm’) or t/3(1120)
{b2 = 0.0844nm2) to be energetically favorable and not dissociated into partials b, = l/6(1123> (2bz = 0.0825 nm2). How-
ever, it would be favorable for dislocations with b = l/3( 1123) (b2 = 0.1649 nm’) to dissociate into b, = b2 = f/6(llf3). In fact, an extended dis~~ation with the l/3(1123} type total Burgers vector was demons~ated to exist in the contrast experiment summarized earlier in this paper. Since the b2 criterion for dissociation is an oversimplification, especially in a covalent type material, it is possible that faults whose SIRUCTLIRE
OF TuNGSTEC\I
CARBIDE
= [ooot]
1/6[11?3] + 1;6[11?2?J = 1/3[11?!0] 1/6fl123] + 1/6[ll233 = t,/3[11133 IF the energies of the dislocations are simply taken to be proportional to h*, then one would expect dislocations Burgers vectors with
Projectton
of faulted crystal anto T&p pkme ?4= h rt I231
Table 2. Summary of the contrast experiment performed on the stacking fault shown in Figs 7 and 8 Stacking fault analysis Visibility
Reflection iota
Fringe contrast
ior I Ii00
No
toil
fringe contrast, both partiafs visibte No fringe contrast, partials invisible No fringe contrast, partials invisible
b, -b, = 1/6[lli%J. b = b, + b, = 1/3[liijJ. The Burgers vectors of-both partial disfocations were found to be b, = t/6fl123], resulting in a total Burgers vector for the extended dislocation b = 1,!3[ll?Q.
of
Prajecfm shp
9erfwt
pime
c’jstai (‘?>c‘~
mto
Fig. 9. Projection of the perfect and faulted structures of WC onto the @iOaf slip plane. The fault was created by moving the top Iayer of W atoms across the C atom tayer with the displacement vector R = 1/6[11233.
1654
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DEFOR~ATrON
bounding partials sum to b = (0001) or l/3( 1120) may also exist. The particular total Burgers vector which is found after deformation may depend on the stress state induced in that section of the crystal, The (0001) directions are the only vectors common to more than one {lfOOt slip plane; therefore, cross-slip will occur only for disl~ations with Burgers vectors along the c-axis. Slip will continue on a given plane until the defect density becomes too high; cross-slip onto other slip planes cannot occur for b = l/3(1120) or l/3(1123). If they exist, extended dislocations with total b = (0001) dissociated into l/6( 1123) partial dislocations must be constricted before they can cross-slip. It seems likely, then, that the extended nature of dislocations resulting from indentation and the identification of l/3( 1123) as the total Burgers vector associated with a stacking fault may partially explain the limited extent of plastic flow resulting from deformation. The indentation and subsequent characterization of defects performed on single crystal WC has not yet been repeated on WC-Co composites, so no direct comparison can be made. However, in the study reported by Hagege et al. [6] on sintered WC-Co composites, stacking faults extended across an entire carbide grain with fault vectors l/6( 1123) and rectangular dislocation networks with nodes dissociated into l/6(1123) partials were observed. Although the appearance and arrangement of dislocations in the composite materials differs from that in deformed single crystals, l/6< 1113) fault vectors have been identified in both types of samples. It seems likely then that defects resulting from indentation in the carbide grains in the composite material will also be dissociated into l/6( 1123) partials. The presence of both stacking faults with R = l/6( 1123) and undissociated dislocations with b = l/3(1123), l/3(1120) and (0001) in areas far from an indentation is difficult to understand. The latter two types of dislocations occur in annealed WC and so may just be the remains of the as-grown state. A hypothesis for the co-existence of dissociated and undissociated dislocations with b = i/3( 1123) is that the directional nature of the bonding may allow one of the states to be metastable, requiring additional input of energy to move atoms to the other, more stable configuration.
AND STRUCTURE
OF TUNGSTEN
CARBIDE
6. CONCLUSIONS 1. The defects resulting from plastic deformation induced in WC by indentation are predominantly stacking faults on {1iOO} planes with l/6( 1123) fault vectors. This is consistent with the slip system of WC at room tem~rature being (1iOOJ (1123). 2. Analysis of the partial dislocations bounding one fault determined their Burgers vectors to be identical and of the type l/6(1 123), demonstrating that extended dislocations with the total Burgers vector b = l/3( 1123) exist. A!though it is possibie that these defects move as undissociated dislocations, it seems likely that a partial dislocation mechanism involving l/6( 1123) dislocations is therefore primarily responsible for plastic flow in WC. 3. Away from indentations stacking faults with fault vectors l/6( 1123) and undissociated dislocations with b = (0001), l/3( 1120), and l/3( 1123) have been identified. 4. Undissociated dislocations with b = l/3( 1120) are the predominant type of defect found in undeformed, annealed single crystals of WC. AcknowledgemenrsThe authors wish to thank Drs D. J. Rowcliffe and S. Pattanaik for the preparation of the bulk materials, the optical micrographs of indentations on the same batch of WC material (Fig. 3) and for helpful discussions. They also wish to thank Mr. S. L. Shinde for many useful suggestions regarding the experimental details. Support from the National Science Foundation Grant No. DMR 77-22647 is gratefully acknowledged.
RE~REN~S T. Takahashi and E. J. Friese., Phil. Mag. 12, 1 (1965).
:: S. B. Luyckx, Acta metal!. 18, 233 (1970). 3. T. Johannesson and B. Lehtinen, Phil. Mug. 24, 1079 (1971). 4. P. B. Hirsch, A. Howie, R. B. Nicholson, D. W. Pashley and M. J. Whelan, Electron Microscopy of Thin Crystals, p. 263. Butterworth, London (1965). 5. M. K. Hibbs and R. Sinclair, Proceedings of the 38th Annual Meeting of EMSA (edited by G. W. Bailey) p. 216. CIaitors, Baton Rouge (1980). 6. S. Hagege, J. Vicens, G. Nouet and P. Delavignette, Physica scatus solidi (a) 61, 675 (1980). 7. W. Hume-Rothery and G. V. Raynor, The Structure of ~~~a~s and Alloys, p. 272. The Institute of Metals, London (1956).
Vol. 29. pp. 1645 lo 1654. 1981 Prmted in Great Bnlam. All rlghls reserved
~1-6l60/~l/o91~5-10S02.0010 Copyright 0 1981 Pergamon Press Ltd
ROOM-TEMPERATURE AND THE DEFECT
DEFORMATION MECHANISMS STRUCTURE OF TUNGSTEN CARBIDE
M. K. HIBBS
and R. SINCLAIR
Department of Materials Science and Engineering, Stanford University, Stanford, CA 94305, U.S.A. (Received 9 February 1981) Abstract-Defects in undeformed, annealed single crystals and in deformed single crystals of WC are characterized by transmission electron microscopy. Deformation is induced by indentation with a microhardness tester. High densities of stacking faults with atomic displacement vectors R = l/6(1123) are found lying on 1liO0) planes near the indentations. Stacking faults with R = l/6( 1123) and undissociated dislocations with b = (OCOl). l/3( 1120), and l/3( 1123) are observed at distances greater than 2 pm from an indentation. The Burgers vectors of the partial dislocations bounding one fault are found to be identical b, = 1/6[1123], resulting in a total Burgers vector for the extended dislocation b = 1/3[1123]. (Stacking fault energy = 55 mJ/m’). Stacking faults with R = l/6(1123) are described in terms of the crystal structure of WC, illustrating several points: the coordination of carbon atoms by tungsten atoms is preserved at the fault, the number of bonds per atom which are broken during slip may be minimized, and the l/6( 1123) type of partial Burgers vector may be combined to form several types of total Burgers vector.
avons caractkrisk par microscopic Clectronique en transmission les dtfauts dans des monocristaux de WC recuits, ainsi que dans des monocristaux d&form&. La dtformation est produite par indentation &I’aide d’un appareil de microdurett. On observe au voisinage des empreintes une forte densit de c!kfauts d’empilement dans des plans (liOO), avec des vecteurs de d&placement atomique R = l/6(1123). On observe d’autre part des dkfauts d’empilement avec R = l/6(1123) et des dislocations non dissocites avec des vecteurs de Burgers b = (Oool), l/3( 1120) et l/3( 1123) ?t des distances de l’empreinte suptrieures g 2 pm. Le vecteur de Burgers des dislocations partielles bordant une faute est &gal g b, = l/6(1123), ce qui conduit g un vecteur de Burgers total b = l/3(1123) pour la dislocation dissociire (tpergie du dCfaut d’empilement: 55 mJ/m’). Nous discutons les dCfauts d’empilement avec R = l/6( 1123) dans la structure de WC et nous montrons les r&hats suivants: les atomes de carbone conservent le nombre d’atomes de tcngstene voisins sur le dtfaut d’empilement, on peut minimiser le nombre de liaisons par atome rompues au tours du glissement et I’on peut combiner des vecteurs de Burgers de dislocations partielles l/6(1 123) pour obtenir plusieurs types de vecteurs de Burgers pour les dislocations parfaites. RbumGNous
Zusammenfassung-Die Defekte, die in unverformten gegltihten und in verformten WC-Einkristallen auftreten, werden mittels Durchstrahlungselektronenmikroskopie analysiert. Verformt wurde mit dem Stempeleindruck eines Mikrohlrtegerltes. Es wurde eine hohe Dichte von Stapelfchlern aufgefunden, deren Verschiebungsvektor R = l/6( 1123) war und die nahe dem Eindruck auf {llOO}-Ebenen lagen. AuBerhalb einem Abstand von 2 pm vom Eindruck wurden Stapelfehler mit R = l/6( 1123) und nicht aufgespaltene Versetzungen mit Burgersvektoren b = (OOOl), l/3( 1120) und l/3( 1123) beobachtet. Der Burgersvektor der Teilversetzungen war immer b, = l/3( li23), welches zu einem Burgersvektor der vollstlndige_n Versetzung von b = l/3( 1123) fiihrt. (Stapelfehlerenergie = 55 mJ/m*). Stapelfehler mit R = l/6( 1123) werden anhand der WC-Kristallstruktur beschrieben. Einige Punkte werden hervorgehoben: die Wolframumgebund der Kohlenstoffatome bleibt am Stapelfehler erhalten, die Anzahl der bei Gleitung aufzutrennenenden Bindungen kann minimalisiert werden, und die Burgersvektoren der Teilversetzungen vom Typ l/6( 1123) lassen verschiedene Typen von perfekten Burgersvektoren zu.
1. INTRODUCTION Tungsten carbide is the major constituent of cemented carbide cutting tools. Efforts are made to optimize the toughness, hardness and wear resistance of these composite materials by altering their composition and microstructure. However, the basic deformation and wear processes in the various phases are still not well understood. This paper describes the study of plastic deformation in WC and the character of the defects which result from it. Plastic flow is
induced in single crystals at room temperature by micro-indentation; the slip plane, slip direction, and nature of the dissociation of the dislocations are determined by transmission electron microscopy. This information is then related to the crystal structure of the material. 2. BACKGROUND The unit cell of WC is shown in Fig. 1. It is hexagonal with tungsten atoms located at 0, 0, 0 and
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these indents in the final TEM specimen. The slip directions are identified and the nature of the dislocation dissociation is determined by complete, conventional contrast experiments [4]. A preliminary report of this work has appeared previously [S]. 3. EXPERIMENTAL
Glrboll
0
0 = 2.906 % c = 2.037 ii
Tungsten
c,o
_
o.g6
slip plone: (1010) Fig. 1. Unit cell of WC. a corresponds to the l/3( i2iO) directions, the cell edges and the cell face diagonal in the
horizontal plane. c is the [OOOl]direction, the vertical cell edge. The l/3( i2i3) directions correspond to the cell face diagonals. The { lOi0) slip planes are the vertical cell faces and the vertical diagonal plane through the unit cell.
carbon at l/3, 213, l/2. The unit cell dimensions are virtually identical, i.e. c/u _ 1. In terms of MillerBravais indices, a corresponds to the l/3( 1120) vectors and c to (0001). The conclusions of previous studies of WC deformation are conflicting. In order to deduce the slip system, Takahashi and Friese [l] indented single crystals on various faces. From observations of the resulting slip steps they determined the slip plane to be {liOO}. They further suggested that the slip directions were <1120) and (0001) since these are the directions of smallest lattice repeat. However, from a careful analysis of slip steps produced by a similar indentation process, Luyckx [2] proposed a slip direction of the type (1123). Johannesson and Lehtinen [3] performed electron microscope studies on tungsten carbide grains within WC-Co cemented carbides. They observed double dislocations and dislocation networks with rectangular nodes. From diffraction contrast experiments and comparison with calculated images, they concluded that the double dislocations and the nodes were extended (1123) superdislocations with partial dislocations of the type l/6(2023). The results reported in their paper were somewhat ambiguous, however, since only one dislocation disappearance criterion was presented for any particular dislocation, whereas two are necessary to unequivocally identify b. Therefore, while the WC structure is simple, the dislocation mechanism of plastic flow has not been established. The present article addresses this situation. The dislocation structure resulting from deformation is distinguished, from that which is either grown-in or present after sintering, by indenting single crystals grown from a melt and clearly locating
CARBIDE
PROCEDURE
Bulk specimens were prepared from mixtures of WC and Co powders melted in an argon atmosphere at 1923 K and then cooled at rates down to 2.5 K/h. WC single crystals were precipitated in a Co matrix by this process. The Co was then etched away with warm HCl. The best crystals were in the form of triangular plates about 3 mm on a side and 1 mm thick. The large faces of the plates correspond to the basal plane while the remaining three sides are (liO0) planes. Some of the crystals were mounted, polished on two different types of crystal faces (the basal plane and {1100) planes), and indented. The arrangement of slip steps which appeared after indentation was observed by optical microscopy. Both deformed and undeformed samples were prepared for transmission electron microscopy. Discs were spark-cut from the triangular plates, then mechanically ground and polished to a thickness of about 1OOpm. The samples were deformed by diamond indentation with a microhardness tester using a 50gm load. Fifty to seventy-five indents were positioned 1OOpm apart in a square array across the sample. All samples were thinned to electron transparency by ion milling. Undeformed samples were milled from both sides, while the deformed samples were thinned from the non-indented side in order to preserve the defects resulting from deformation. However, the indented sides of these samples were exposed to the ion beam for approximately one hour in order to remove surface deformation resulting from the polishing process. The foils were observed in a Philips EM400 operated at 120 kV. The exact positions of indents in the TEM thin foils were determined by comparing low magnification electron micrographs with optical micrographs of the prepared samples. 4. RESULTS 4.1 Undeformed material In the undeformed material, which had been annealed at 2073 K, undissociated dislocations were observed, often arranged in networks. One such network is shown in Fig. 2. In Fig. 2 (a) g = [1120], and all the dislocations are visible. The same area is imaged under different diffracting conditions, g = [OilO], in Fig. 2 fb), and the dislocations labelled A are invisible. Complete conventional contrast experiments were performed on this network [4] and are summarized in Table 1. All three sets of dislocations, A, B and C, were found to have Burgers vectors of the type l/3( 1120). Although the defect structure in the undeformed material was dominated by single dislocations with
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(b) Fig. 2. Dislocation network in undeformed, an_nealed single crystal of WC. In (a) g = [1120], and all the dislocations are visible. In (b) g = [01 lo] and the dislocations labelled A are invisible.
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Table 1. Summary of the contrast experiment performed on the dis-
location network shown in Fig. 2 Burgers vector determination Reflection Visibility 1120 Oil0 O!11 AlOO 1101 lOA0 1011
all dislocations visible type A invisible type A invisible type B invisible type B invisible type C invisible type C invisible
type A: b = 1/3[2ijO]. type B: b = 1/3[1!20]. type C: b = 1/3[1210]. All dislocations were found to have Burgers vectors of the type l/3(1120).
b = l/3( 1120), low-angle grain boundaries, dislocation loops with b = (OOOl), and an occasional stacking fault were also observed.
4.2 Deformed material Figure 3 shows two optical micrographs of indentations in single crystal tungsten carbide. The indentation in Fig. 3(a) was made on the basal plane, while the sample shown in Fig. 3(b) was in the (1010) orientation. As noted by Takahashi and Friese Cl], the slip steps visible on the (Oool) crystal face indicate that some component of the slip direction must lie along [el]. Similarly, the presence of slip steps on the (1010) face indicates that a component of the slip lies out of the (lOf0) plane and therefore along a (1120) direction. It may be noted that (1123) directions fulfill this requirement. The indented side of a TEM specimen is shown in Fig. 4(a). The square arrangement of the indentations made it possible to correlate the hole visible in the low magnification electron micrograph in Fig. 4(b) with the indent indicated by the arrow in Fig. 4(a). A high density of stacking faults, each bound by a pair of partial dislocations, exists around the hole as may be observed from Fig. 5(a). In this orientation the faults are practically end-on, lying on (liO0) planes. The zone axis of the corresponding diffraction pattern, Fig. 5(b), is very close to the c-axis, and the streaks along the (1 iO0) directions arising from the faults confirm that they lie on (IiOO} planes. No undissociated dislocations are found in this region. At distances greater than 2 pm from the indentation a combination of dissociated and undissociated dislocations is observed (e.g. Fig. 6). The faults exhibit contrast similar to those near the indentation and also lie on ( liOO! planes. Undissociated dislocations with three different types of Burgers vector, (OOOl), l/3( 1120) and l/3( 1123), have been identified. Since these defects are located far from regions where macroscopic slip steps are visible it seems unlikely
AND STRUCTURE
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that they are part of the plastic deformation resulting from indentation. They may arise from the difference in thermal contraction between the phases during the cooling of the newly formed WC crystals in the Co matrix, or may result from plastic deformation during the sample preparation process and not removed by ion-milling. Fault and partial dislocation identification also followed conventional procedures. A stacking fault imaged in both bright field and centered dark field modes (g = [lOlO]) is shown in Fig. 7. The contrast is symmetrical in both micrographs indicating that the fault is a n-fault, that is, c( = 2ng.R = 71 where R describes the lattice displacement across the stacking fault. Examples of the experiment performed to determine the displacement vector and the Burgers vectors of the bounding partial dislocations of the fault are found in Fig. 8. In Fig. 8 (a) g = [ioll], and the fringe contrast has disappeared while the partial dislocations are still visible. Under the diffracting conditions in Fig. 8(b), g = [liOO], the fringe contrast and both partial dislocations are invisible. A summary of the analysis is given in Table 2. The two partial dislocations were found to have the same Burgers vector, bl = b2 = R = 1/6[1123]. This results in a total Burgers vector for the defect b = 1/3[1123], confirming the macroscopic, slip step analysis of Luyckx [2]. This vector has components in both the [OOOl] and [ 1120] directions as was required by the information derived from optical microscopy. It also satisfies the criterion that a = 27rg.R = B (g = [iOlO], R = 1/6[1123]). Hagege et al. [6] have observed stacking faults with independently R = l/6( 1123) in the carbide grains of a WC-Co composite. Using the Burgers vectors given above for the partial dislocations and assuming a line vector along [OOOl],the stacking fault energy is estimated to be 55 mJ/m’. The particular fault whose analysis is summarized here was located at a distance greater than 5pm from an indentation. Many faults located near indentations were also analyzed and found to have the same type of fault vector, R = l/6(1123). However, the high density of defects in these areas made the determination of the Burgers vectors of the partial dislocations very difficult, but they are presumed to be the same. 5. DISCUSSION The results given above demonstrate that plastic deformation induced by indentation in WC results in a defect structure consisting of dislocations dissociated on {liOO} planes with fault vectors of the type l/6( 1123). Although previous optical studies [l, 21 determined the slip system of WC, they did not reveal the extended nature of the dislocations. Figure 9 contains a diagram of the perfect and faulted structures of WC projected onto a slip plane, (1100). Three layers of atoms are indicated. The bottom layer consists of tungsten atoms, the next of
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a)
b)
Fig. 3.
A.M. 29/9--c
:a1 micrographs of indentations in single crystal WC. (a) Basal plane orient al orientation. (Courtesy D. J. Rowcliffe and S. Pattanaik).
(b) (lOi0)
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Fig. 4(a).
Fig. 4(b).
Fig. 4. (a) Optical micrograph of a deformed TEM specimen of WC. Indentations have been placed in a square array. The arrow points to the location of an indentation which appears as a large hole in the : low magnification electron micrograph in (b).
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ta)
(b) Fig. 5. (a) A high magnification electron micrograph taken near the hole shown in Fig. 4(b). A high density of stacking faults may be seen lying almost parallel to the electron beam on jlloO) planes. (b) The corresponding diffraction pattern with the electron beam direction very close to the c-axis. The streaks along the (1100) directions confirm that the faults lie on { 1100) planes.
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Fig. 6. Example of the defect structure found at distances greater than 2 pm from an indentation. Both dissociated and undissociated dislocations are observed.
carbon, and the top layer again of tungsten. The faulted structure was formed by moving the top layer of tungsten atoms across the layer of carbon atoms with the displacement vector R = 1/6[1123]. This diagram illustrates several points. The first observation which may be made, and which was also noted by Hagege et al. [6], is that the 1/6[t 1233 fault vector preserves the trigonal prismatic coordination of the carbon atoms by tungsten atoms. In fact, at the fault the structure appears to have been merely rotated by 90” from the perfect structure. Since tung-
sten carbide is highly stoichiometric, with directional bonding between tungsten and carbon atoms [7], the structure of a low energy fault would be expected to preserve the number and the angular relationship of the tungsten-carbon bonds. Figure 9 may also provide an explanation for the fact that the { 1iO0) planes are the slip planes rather than the more closely packed basal planes. The strong covalent bonding implies that it would be desirable to break as few bonds as possible in the slip process. For slip to occur on the basal plane three tungstewar-
7 *..
O.OSpm
~ c
r* P
Fig. 7. (a) Bright field and (b) centered dark field pair of micrographs of a stacking fault taken under the diffracting condition g = [iOlO]. The symmetrical contrast in both micrographs indicate that the fault is a n-fault.
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Fig. 8. Exampfes of the contrast experiment performed on the fault shown in Fig. 7. In (a) g = [iOI 11 and the fringe contrast has disappeared while both partial dislocations are still visible. In (b) g = [l MO] and the fringe contrast and both partial dislocations are invisible.
bon bonds must be broken per carbon atom. However, as was noted by Takahashi and FrieseCl] and may be seen from the diagram, the possibility of breaking only two tungsten~arbon bonds per carbon atom exists for slip on { 1iGQ] planes. This will occur if disio~atjons move between planes equivalent to the layer of carbon atoms and the top layer of tungsten; motion between the carbon atoms and the bottom layer of tungsten atoms would result in four broken bonds per carbon atom. The third point which may be made from the diagram is that the 1/6[11?3] partial Burgers vector could give rise to more than one type of totai Burgers vector. The following combinations could exist on the slip plane shown: 1/6[11233 + 1/6fii23]
(h2 = 0.0805 nm’) or t/3(1120)
{b2 = 0.0844nm2) to be energetically favorable and not dissociated into partials b, = l/6(1123> (2bz = 0.0825 nm2). How-
ever, it would be favorable for dislocations with b = l/3( 1123) (b2 = 0.1649 nm’) to dissociate into b, = b2 = f/6(llf3). In fact, an extended dis~~ation with the l/3(1123} type total Burgers vector was demons~ated to exist in the contrast experiment summarized earlier in this paper. Since the b2 criterion for dissociation is an oversimplification, especially in a covalent type material, it is possible that faults whose SIRUCTLIRE
OF TuNGSTEC\I
CARBIDE
= [ooot]
1/6[11?3] + 1;6[11?2?J = 1/3[11?!0] 1/6fl123] + 1/6[ll233 = t,/3[11133 IF the energies of the dislocations are simply taken to be proportional to h*, then one would expect dislocations Burgers vectors with
Projectton
of faulted crystal anto T&p pkme ?4= h rt I231
Table 2. Summary of the contrast experiment performed on the stacking fault shown in Figs 7 and 8 Stacking fault analysis Visibility
Reflection iota
Fringe contrast
ior I Ii00
No
toil
fringe contrast, both partiafs visibte No fringe contrast, partials invisible No fringe contrast, partials invisible
b, -b, = 1/6[lli%J. b = b, + b, = 1/3[liijJ. The Burgers vectors of-both partial disfocations were found to be b, = t/6fl123], resulting in a total Burgers vector for the extended dislocation b = 1,!3[ll?Q.
of
Prajecfm shp
9erfwt
pime
c’jstai (‘?>c‘~
mto
Fig. 9. Projection of the perfect and faulted structures of WC onto the @iOaf slip plane. The fault was created by moving the top Iayer of W atoms across the C atom tayer with the displacement vector R = 1/6[11233.
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bounding partials sum to b = (0001) or l/3( 1120) may also exist. The particular total Burgers vector which is found after deformation may depend on the stress state induced in that section of the crystal, The (0001) directions are the only vectors common to more than one {lfOOt slip plane; therefore, cross-slip will occur only for disl~ations with Burgers vectors along the c-axis. Slip will continue on a given plane until the defect density becomes too high; cross-slip onto other slip planes cannot occur for b = l/3(1120) or l/3(1123). If they exist, extended dislocations with total b = (0001) dissociated into l/6( 1123) partial dislocations must be constricted before they can cross-slip. It seems likely, then, that the extended nature of dislocations resulting from indentation and the identification of l/3( 1123) as the total Burgers vector associated with a stacking fault may partially explain the limited extent of plastic flow resulting from deformation. The indentation and subsequent characterization of defects performed on single crystal WC has not yet been repeated on WC-Co composites, so no direct comparison can be made. However, in the study reported by Hagege et al. [6] on sintered WC-Co composites, stacking faults extended across an entire carbide grain with fault vectors l/6( 1123) and rectangular dislocation networks with nodes dissociated into l/6(1123) partials were observed. Although the appearance and arrangement of dislocations in the composite materials differs from that in deformed single crystals, l/6< 1113) fault vectors have been identified in both types of samples. It seems likely then that defects resulting from indentation in the carbide grains in the composite material will also be dissociated into l/6( 1123) partials. The presence of both stacking faults with R = l/6( 1123) and undissociated dislocations with b = l/3(1123), l/3(1120) and (0001) in areas far from an indentation is difficult to understand. The latter two types of dislocations occur in annealed WC and so may just be the remains of the as-grown state. A hypothesis for the co-existence of dissociated and undissociated dislocations with b = i/3( 1123) is that the directional nature of the bonding may allow one of the states to be metastable, requiring additional input of energy to move atoms to the other, more stable configuration.
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CARBIDE
6. CONCLUSIONS 1. The defects resulting from plastic deformation induced in WC by indentation are predominantly stacking faults on {1iOO} planes with l/6( 1123) fault vectors. This is consistent with the slip system of WC at room tem~rature being (1iOOJ (1123). 2. Analysis of the partial dislocations bounding one fault determined their Burgers vectors to be identical and of the type l/6(1 123), demonstrating that extended dislocations with the total Burgers vector b = l/3( 1123) exist. A!though it is possibie that these defects move as undissociated dislocations, it seems likely that a partial dislocation mechanism involving l/6( 1123) dislocations is therefore primarily responsible for plastic flow in WC. 3. Away from indentations stacking faults with fault vectors l/6( 1123) and undissociated dislocations with b = (0001), l/3( 1120), and l/3( 1123) have been identified. 4. Undissociated dislocations with b = l/3( 1120) are the predominant type of defect found in undeformed, annealed single crystals of WC. AcknowledgemenrsThe authors wish to thank Drs D. J. Rowcliffe and S. Pattanaik for the preparation of the bulk materials, the optical micrographs of indentations on the same batch of WC material (Fig. 3) and for helpful discussions. They also wish to thank Mr. S. L. Shinde for many useful suggestions regarding the experimental details. Support from the National Science Foundation Grant No. DMR 77-22647 is gratefully acknowledged.
RE~REN~S T. Takahashi and E. J. Friese., Phil. Mag. 12, 1 (1965).
:: S. B. Luyckx, Acta metal!. 18, 233 (1970). 3. T. Johannesson and B. Lehtinen, Phil. Mug. 24, 1079 (1971). 4. P. B. Hirsch, A. Howie, R. B. Nicholson, D. W. Pashley and M. J. Whelan, Electron Microscopy of Thin Crystals, p. 263. Butterworth, London (1965). 5. M. K. Hibbs and R. Sinclair, Proceedings of the 38th Annual Meeting of EMSA (edited by G. W. Bailey) p. 216. CIaitors, Baton Rouge (1980). 6. S. Hagege, J. Vicens, G. Nouet and P. Delavignette, Physica scatus solidi (a) 61, 675 (1980). 7. W. Hume-Rothery and G. V. Raynor, The Structure of ~~~a~s and Alloys, p. 272. The Institute of Metals, London (1956).