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Glide Dislocations Dissociation in Inversion Domain Boundaries of Plastically Deformed Aluminium Nitride
Virginia Feregotto, Jean-Pierre Michel
To cite this version:
Virginia Feregotto, Jean-Pierre Michel. Glide Dislocations Dissociation in Inversion Domain Bound- aries of Plastically Deformed Aluminium Nitride. Journal de Physique III, EDP Sciences, 1996, 6 (9), pp.1261-1269. �10.1051/jp3:1996184�. �jpa-00249523�
Virginia Feregotto and Jean-Pierre Michel (*)
LMPSM (**), kcole des Mines, Parc de Saurupt, 54042 Nancy Cedex, France
(Received 16 February1996, re&ised 9 May1996, accepted 3 June 1996)
PACS.61,16.Bg Transmission, reflection and scanning electron microscopy PACS.81.40.Lm Deformation, plasticity and creep
Abstract. A ten per cent plastic deformation of polycrystalline aluminium nitride, at a
temperature ranging from 1500 to 1650 °C creates a new kind of intragranular defect. Observed by transmission electron microscopy, they look like torsion subboundaries created by dislocations
with 1/3(11%0) Burgers vectors and so nodes are dissociated into Shockley partials. They are located in the basal plane. In fact, these defects appear only in the plane areas of grown-in defects, the inversion domain boundaries. The formation of these faulted networks is interpreted
as being the ultimate stage of the interactions between inversion domain boundaries and glide
dislocations.
R4sumd. Une d4formation plastique de 10 % de nitrure d'aluminium polycristallin, entre
1500 et 1650 °C introduit
un nouveau type de d4fauts intragranulaires. Au microscope 41ectro-
nique par transmission, ils apparaissent comme des sous-joints de torsion cr44s par des dislo- cations de vecteurs de Burgers 1/3(ll10) dont les nceuds triples sont dissoc14s en partielles
de Shockley ils sont situ6s dans le plan de base. En fait, ces d6fauts ne se produisent que
sur les parties planes de d6fauts originels, les parois de domaines d'inversion. La formation de
ces r6seaux faut6s est analys6e comme l'ultime stade des interactions entre parois de domaines d'inversion et dislocations de glissement.
1. Introduction
Aluminium Nitride (AIN) is a stoichiometric compound with a wurzite-type hexagonal struc-
ture. This structure is built from a hexagonal closed-packed (hcp) lattice with a basis of two
atoms, Al and N, at a distance of 0.385c along [0001j. It could equally be described as two
interpenetrating hcp lattices, the one for cations (Al), the other for anions IN) displaced from each other by the vector 0.385[0001]. Each atom is within a tetrahedron whose summits are
occupied by four first neighbours of the other species. AIN is non-centrosymmetric so two ar-
rangements of atoms can be considered (Figs. la and b) for an identical Al sublattice. N atoms
give rise to upward and downward-pointing tetrahedrons. In a same grain, the two configu-
rations can exist, being separated by a surface called Inversion Domain Boundary (IDB). An
(*) Author for correspondence (e-mail: michel@mines.u-nancy.fr) 1'*) assoc16 au CNRS
no. 155
Q Les kditions de Physique 1996
1262 JOURNAL DE PHYSIQUE III N°9
/ '
' /
', / ',
,
N
a a
a) a b) a
Fig. I. Schematics of the two AIN polymorphs showing the two different tetrahedral site orienta- tions. These two orientations are related to one another by an inversion operation.
IDB is an orientable surface (it is always possible to define unambiguously inside and outside
areas) which cannot be formed or removed by shearing operations.
IDBS and dislocations are the only extended defects observed so far in AIN by transmission electron microscopy.
Dislocations have a 1/3(ll10) Burgers vector and appear generally undissociated when ob- served by the weak beam method. Nevertheless some cases of dissociation were mentioned:
ii) In thin single crystals grown from the vapour phase where triple nodes are bounded by Shockley partials [1,2j, the stacking fault energy ~ is evaluated at 4 mJ m~~
iii) In some bulk materials obtained by sintering, triple nodes and individual dislocations appear to be dissociated with a stacking fault ribbon width of eight nanometers (~ Gt 7.5 mJ m~~) [3].
Observed IDBS are constituted by a succession of adjacent surfaces alternatively ii) curved, often with an orientation close to (loll) and iii) planar, parallel to (0001) [4,5j. These
interfaces are characterized by a displacement vector, R.
Notice that, by definition, R represents the displacement between two neighbouring domains
separated by a bidimensional defect. An IDB separates regions which cannot be derived one from the other by means of a parallel displacement. Nevertheless an equivalent R vector
can be defined unambiguously [6j, its value usually being deduced from contrasts observed by
transmission electron microscopy.
For planar portions:
Rp = (1010) + a[0001j
3
Two a values have been proposed: 0.39 [7j and close to 0.15 [8,9j, this is consistent with the contrasts observed by different authors.
For curved portions, Rc is not unique. For some authors it is parallel to c:
Rc = fl[0001j
with fl
= 0.5 [8j or 0.05 < fl < 0.43 [10j. For other authors it has an another small component,
not determined [4j or equal to 1/3(1010) [9j.
[7]. The proposed structures are AION [13j and A1203 whose misfit with AIN is particularly
low [14].
Interactions between dislocations and IDBS were rarely observed after sintering. The em- phasis was put on "hinge" dislocations located at the border of two successive portions of
IDB. The determination of their Burgers vector by the invisibility criterion is difficult. It is necessary that the IDBS themselves be invisible and also in that case, residual contrast is often observed for these dislocations. This is probably due to the "decoration" by oxygen. Burgers
vectors (hh01) [5], and more precisely, )(1100) [12] or )(1100) + t[0001] where measured and
calculated t values, 0.157 and 0.164 respectively [7], have been suggested.
A plastic deformation increases the dislocation density but changes neither the number,
about one per grain, nor the characteristics of IDBS [5j. After a few per cent deformation,
some cases of dislocations piercing an IDB and dislocations trapped by a plane or a curved
IDB were observed [5j.
The aim of this work is to study interactions between dislocations and IDBS during an important plastic deformation. We chose the maximum value permitted by fracture of the
material: 10%.
2. Experimental Procedure
Aluminium nitride was obtained from ESK (Electro Schmelzwerk Kempten, Gmbh, Germany).
Two types of material, sintered (S-AIN) and hot-pressed (HP-AIN), were studied. Both were
polycrystals with a mean grain size of five microns. They contained a sintering additive Y203
or La203 with a volume fraction of 5%. The oxygen concentration was measured to be 2%
weight by infrared absorption. Compression tests were conducted on 8 x 3 x 3 mm~ samples,
either at a constant strain rate of
= 5 x 10~6 s~~ or, by creep at constant stress (150, 200 or
250 MPa) under an argon and helium atmosphere. Temperature was chosen between 1500 and 1650 °C so that the test duration varied between 6 and 80 hours. At the end of the test, samples
were cooled down with the final load applied in order to freeze in dislocation configurations.
Thin foils were prepared by ion milling and observed in a JEOL 200 CX microscope operating
at 200 kV.
3. Results and Discussion
Few dislocations were found in the grains of the material on receipt (zero to 10~ cm~~ depending of the grain). After 10% deformation, dislocations density strongly increases. Every grain
contains an homogeneous density of10~ cm~~ on both sides of the IDB if one is present. Burgers
vectors are more or less distributed between the three 1/3(1110). Most of the dislocations are
undissociated but some large stacking fault ribbons have been observed. The number of IDBS is unchanged but a new type of interactions between IDBS and dislocations appeared in nearly
one fifth of the grains. These structures are planar, lying in (0001) and resemble the usual networks of extended nodes. The configuration of Figure 2 is representative of all observed
cases in both materials IS and HP-AIN). It is made up of triangles, alternatively bright and
dark. Figure 3 describes the usual mechanism of formation of these networks. Faulted nodes of one of two types, for example F1, spread whereas F2 nodes shrink into points. Finally, the
configuration of Figure 3d is obtained. The size of the triangles varies between 80 and 130 nm.
1264 JOURNAL DE PHYSIQUE III N°9
o.3 pm
9 -
Fig. 2. Partial dislocations network in planar IDB. Stacking faults are in contrast (a curved IDB is also visible). HP-AIN deformed to e = 9.5$l at constant strain rate (1650 °C). Weak beam micrograph
g, 2g g = i103, beam direction [till].
2
3
2
a) b)
n n
uF
uF fi
w w uF
c) d)
Fig. 3. Sketch of the network formation. a) Two basal dislocations I and 2 interact attractively to form dislocation 3, b) extended Fl and contracted F2 nodes result from the dissociation of dislocations into partials, c) generalization to several nodes; UF: unfaulted zone and d) shrinkage of F2 nodes and realignment of dislocations give this schematic configuration.
Fig. 4. Same
area as in Figure 2. The planar IDB and the three partial dislocations families are in
contrast. Weak beam micrograph g, 5g g = 1010, beam direction [1%13].
Generally, networks are a consequence of the increase of the dislocation density during plastic
deformation whereas dissociation is less common. The radius of curvature of partial dislocations
bounding Fl areas are too large to be accurately measured. As ~ is inversely proportional to this radius, we can only conclude that the stacking fault energy is very low. Low stacking fault
energy of the dislocation ribbon is not the only parameter in effect since matrix dislocations in the vicinity are undissociated. In fact, micrographs obtained with other reflection vectors
show that, on each network, is superimposed a plane IDB. In Figure 4 obtained by weak beam technique, faults are out of contrast while IDB and partial dislocations are visible. IDB are apparent because of narrow thickness fringes. It can be noted that the curved IDB, visible in Figure 2 is here out of contrast; curved IDBS are invisible for any reflection in the basal
plane, clearly visible with 10il
or 1013 and either weakly or invisible with 1112 reflections, which agree more or less with [8j. Determination of the partial Burgers vector was done using invisibility criterion when stacking faults were out of contrast. Partials can be divided into
three families, each corresponding to a line of direction. Each family has its own Burgers
vector 1/3(1010) parallel to the lines. These are therefore screw Shockley partials resulting
from dissociation of1/3(1110) perfect dislocations. In Figure 5, obtained with g =
ii10, the b
= 1/3[01I0j family whose line is perpendicular to g, is invisible. If the Burgers vector of the partials can be determined unambiguously, it is not the case for Rj, the displacement of the faults. The superposition "IDB plus fault" is clearly out of contrast for the three 11§0 diffraction vectors. For others reflections, the IDB is visible and the relative fault's contrast is strong for 1013 only; it is weak with 10I1
or 10I2 and difficult to distinguish with 11§2 or 10I0. Dislocation density is greater in the IDB than in the surrounding matrix. We can thus conclude that an attractive interaction exists between these two types of defects. No network,
whether dissociated or not, has been observed either in the matrix or in the curved IDBS.
In our opinion, these results can be interpreted as a continuity to the interactions between dislocations and IDBS as observed so far and hence justify a growth model for IDBS, a problem
which is actually debated. IDBS are created at ahigh temperature (G£ 180o °C) during sintering.
Our hypothesis is that IDBS generally nucleate at a grain boundary and develop in the grain
as a curved surface, initially unique, leaning on the grain boundaries. The presence of oxygen in solution in the matrix is necessary for growth [15]. The additional energy due to the IDB'S
surface increase is balanced by the decrease of the oxygen content of the volume swept by the
1266 JOURNAL DE PHYSIQUE III N°9
fi O.3 pm'~? ".~ '"
Fig. 5. Same
area as in Figure 2. The planar IDB and one partial dislocations family (b = 1/3[0i10]) are out of contrast (curved IDB is also invisible). Weak bean~ micrograph g, 2g g = il10,
beam direction [0ill].
IDB. A small dissymmetry of oxygen contents on both sides of the IDB has been detected [7].
During and after growth, trapping of dislocations can occur. There is an attractive interaction between dislocations and curv~d IDBS because dislocation density is always higher inside the
two-dimensional defects than in their vicinity [5]. Some trapped dislocations dissociate into
Shockley partials in the (0001) plane creating planar surfaces. This mechanlm explains the sequence:
grain boundary curv.ed IDB dislocation plane area -~dislocation curved area grain boundary which agrees with our observed cases, more than a thousand. The displacement
vector of a planar IDB is therefore the sum of the displacement vectors relative to curved IDB
(Rc) and stacking fault ribbon (RF)i
Rp=Rc+RF
RF which can differ by a period of the lattice is equal to the Burgers vector of one of the two
Shockley partials leading to a Rp value compatible with the above relationships. It agrees too with:
b=Ri-R2
established in SiC [16] where b is the Burgers vector of a hinge dislocation and RI and R~ the
displacement vectors of two adjacent IDBS.
For some authors, IDBS may nucleate inside the grains [9,17]. In this case; they would be constituted by one curved surface like a dome and one flat surface connected together by a dislocation. This process is probably rare because only very few of such nucleus were really observed. On the other hand, many "D shaped" IDBS were seen IS,9,17]. These configurations
can be interpreted as either the result of the growth of intragranular nucleus or as the cutting of a part of one multinapped IDB by the thin foil planes. Only the observation of materials
with very fine grains, completely included in the thin foil, could give an answer.
After sintering, a 10% plastic deformation created a high density of dislocations, this increase
being typically two orders of magnitude. Figure 6 shows the different stages of a network for- mation. The number of interactions between dislocations and IDB increases from the left to the right of the micrograph. In the left part, fault ribbons are wide and have no particular shape. in the right one, a network of triangles with regular shapes is formed. In the middle,
a)
- + 1/3
(1210j
+ 1/3 (l100j
+ 1/3 (1010j
+ 1/3 (0110j
b)
Fig. 6. a) 1/3(11%0) perfect dislocations trapped by a planar IDB. Matrix segments are undis- sociated whereas segments in the IDB are widely split. S-AIN deformed to e = 9.I$l at constant
stress (150 MPa) (1550 °C). Dark field micrograph g
= i103, beam direction [I%11]. b) Sketch of the micrograph center. A perfect dislocation (b = 1/3[1§10]) is dissociating into partials (bi = 1/3[l100] + b2
= 1/3[0i10]).
a dislocation with a 1/3[1§10] Burgers vector can be seen having partially reacted with other dislocations in the IDB. In the matrix its line becomes perpendicular to (0001) when it ap-
proaches the IDB, confirming an attractive interaction. It can be noted that the dissociation of this dislocation exists only in the IDB. The process can be described in this way: first dislo-
cations in the IDB dissociate separately whereas the following ones with other Burgers vectors react with the first ones to build a regular array of triangles. So the formation of subboundaries in plane IDBS occurs by a slightly different way of the mechanism described in Figure 3 even if the result is the same.
In hexagonal structures the only faults bordered by Shockley partials are intrinsic ones, called12. When the sequence of partials is reversed the most probable configuration is an other fault 12, with the same energy but located on an adjacent basal plane (see for example
Ref. [18]). In our case dislocation nodes are alternatively spreaded and shrunk implying very different stacking fault energies. A rough explanation can be given considering the variation of bounding energy between atomic layers introduced by the creation of the stacking fault.
Consider the two nearest planes above the IDB where dissociation can occur (planes I and II in Fig. 7) (Notice that faults happen on glide planes contrary to planar IDBS which take
place on shuffle planes). For the Al layers, called 1 and 2, respectively above planes I and
1268 JOURNAL DE PHYSIQUE III N°9
IDB
Fig. 7. Wurtzite structure projected along (ll10) with a planar IDB in the basal plane (0001).
Large white circles and small black ones represent respectively Al and N atoms in the plane of the
drawing, grey ones at a/2 behind this plane. Planes I (and II) and layer I (and 2) represent stacking
fault planes above the IDB and the first induced translated Al layers.
II, interatomic distances with lower layers are modified only from the third one. The faults increases distances between neighbouring atoms of layers I or 2 and their corresponding third
layer a or b constituted of Al or N atoms, respectively. So to the habit fault planes I and II could correspond to lower and higher stacking fault energy.
This model is only a first approach because it does not take into account the influence of oxygen. In fact the structure is complex, constituted byii) one IDB, defect which can be created
only by an inversion operation and which contains oxygen in an undetermined structure iii) stacking faults created by shear containing probably oxygen too. The resulting stacking fault is weak because dissociation widths can reach several tenths of a micron. It cannot be described
from measurements because configurations are out of equilibrium. The applied stress create
different forces on the partial dislocations inducing fault ribbons. Dissociation is afterwards modified and stabilized by reaction with other dislocations.
It was not detected in the studied domain that temperature influenced either the number or the geometry of stacking faults. To separate the possible effects of heat treatment from those due to plastic deformation, observations were made on a sample heated at 1550 °C during 80
hours; no effects were seen, the observed results thus need first glide dislocations and applied
stress.
4. Conclusion
After a 10$l plastic deformation many interactions between perfect dislocations and plane IDBS
were observed by TEM. IDBS contain more dislocations than the surrounding matrix. Plane IDBS being parallel to the main glide plane of dislocations, an attractive interaction between these defects is expected. The first dislocation segments caught by the IDBS are individually
dissociated in 1/3(1010) widely separated partials; neighbouring segments located in the matrix
are undissociated. Dislocation splitting occurs in IDBS although these defects cannot be created
or annihilated by shearing. When new dislocations are trapped by the plane IDB, they react with the first dislocations building twist subboundaries with nodes alternatively extended and contracted. The corresponding stacking fault asymmetry is attributed jointly to different fault
planes and to oxygen the main impurity.
