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Dislocation structures and in situ observations of dislocation motion in InSb
Physica 116B (1983) 641.645
Paper presented at ICDS-12 Amsterdam, August 31 - September 3, 1982
North-Holland PublishingCompany
DISLOCATION STRUCTURES AND IN SITU OBSERVATIONS OF DISLOCATION MOTION IN InSb M. FNAIECH* and F. LOUCHET**
* IMpartement de Physique, Facultd des Sciences, Tunis, Tunisia ** Laboratoire d'Optique Electronique du C2V.R.S., Toulouse, France Dislocation mobility in InSb is investigated both by a microscopic analysis of dislocation structures at lO0 kV and by in situ straining in a high voltage microscope. The observed differences in deformation processes between bulk specimens and thin foils is discussed. Dislocation pileups observed in thin foils indicate the existence of localized and recurrent dislocation sources.
I. INTRODUCTION
II. EXPERIMENTAL PROCEDURE
In III-V compounds of sphalerite structure, two types of 60 ° dislocations (labelled ~ and B) can be imagined, depending on the nature of the atoms which terminate the extra half planes in the dislocation cores. A perfect hexagonal loop is shown in fig. l, the sides of which lie parallel to the directions. The screw segments separate the two ~ 60 ° segments (on the right) from the two ~ 60° segments (on the left). The different electronic structures associated with the cores of ~, B and screw dislocations are expected to result in different mobilities for these three types of dislocations.
InSb specimens strained up to the lower yield point by compression along [132] were provided by D. Ferr6 and R. Kesteloot (3). The geometry of these specimens is shown in fig. 2. The orientation of the compression axis favours the slip system [IT0] (Ill). One of the (5~4) surfaces was scratched in order to select in compression the nucleation of s-type half loops, i.e. ~ dislocations connected at both ends to screw dislocations.
/,lj
$ Figure I - Perfect hexagonal loop. ~ and B segments, of 60 ° character, lie on opposite sides of the screw segments.
I
BTol
\
/
//1~\
l f (1111 I
\\
S
15t"41 Indirect measurements of dislocation mobilities in InSb performed previously by etching techniques (], 2) showed indeed a large difference in mobilities of ~, 8 and screw dislocations. More recently, formation of long screw dipoles by the faster propagation of ~ dislocations from the surface has been shown by X-ray topography (3). However, 8 dislocations were not observed by these authors. In the present work, dislocations structures produced previously (3, 4) by compression of bulk specimens were observed by 100 kV electron microscopy. Some results of in situ deformation of InSb thin foils in a high voltage microscope are also reported. The latter specimens were also analysed in a conventional transmission microscope.
0378-4363/83/0000-0000/$03.00 © 1983 North-Holland
[TTI] Figure 2 - Orientation of bulk compression specimens. Thin foils for TEM observations were sliced parallel to the (Ill) plane. Thin speimens for in situ experiments were parallel to_the (514) plane, with a tensile axis along [132]. In order to observe the resulting dislocation structures, we sliced the specimens along the primary glide plane (Ill). The micro-tensile specimens for in situ experiments, on the opposite, were cut from (514) slices. This latter orientation is a compromise which allows good visibility of dislocations in the primar~ slip plane (Ill) in single slip conditions (132 tensile axis).
642
M. Fnaiech, F. Louchet / Dislocation motion in lnSb
Figure 3 - Dislocation structures in specimens strained by compression. The foil plane is the primary plane (Ill). The screw dislocations emerging at the surface have been probably built up by glide of the 60 ° segment. A closed hexagonal loop, which necessarily contains a B segment, is shown in A. (Note that half-loops are elongated along the Burgers vector direction, which suggests that 60 ° dislocations are faster than screws).
Figure 4 - Same foil as in fig. 3, showing dislocation dipoles resulting from mutual interaction of loops with the same Burgers vector a/2 [ll0] lying on parallel (Ill) planes.
M. Fnaiech, F. Louchet
/Dislocationmotion
The slices were first thinned down mechanically, and then chemically polished. The micro-tensile specimens were shaped and thinned down by preferential chemical etching, using wax coating as a protection. The in situ straining experiments were performed in the 3 MeV microscope of Toulouse, operating at I MeV. The heating straining stage (5) was used at about 250°C. The dynamical observations were recorded on a video tape recorder.
III. OBSERVATIONS OF DISLOCATION STRUCTURES IN SPECIMENS STRAINED BY COMPRESSION. - Figure 3 shows dislocation half-loops, of hexagonal shape. The dislocation segments lie along the three directions of the primary (Ill) plane. The screw segments intersect the surface. These half-loops are similar to those observed by Kesteloot on a larger-scale in the same specimens (3). However, a complete closed loop is also observed in fig. 3 which demonstrates the existence of B dislocations.
in lnSb
643
in a 100 kV microscope (fig. 6 and fig. 7) showed that the screw dislocations mentioned above belong to two families, the Burgers vectors of which are a/2 [110] and a/2 [0Jl]. The slip planes have been unambiguously determined, and the results are summarized in the following table :
Burgers vector
Slip plane
Observation frequency
a
(711)
frequent frequent
0,36 0,12
a
(ITs)
frequent less frequent
0,29 0
a
(lll)
not observed not observed
0,48 (primary) 0
a
(111)
not observed not observed
0,36 0,18
7 [11o] (~TI)
T [o~] (TT1) T [17o] (TT~) T [To1] (iTl)
Theoretical Schmid's factor
- More complicated configurations were also frequently observed : figure 4 shows a dipole structure resulting from interaction of loops with the same Burgers vector, lying on neighbouring (Ill) planes.
IV. IN SITU DEFORMATION
\
The specimens were constrained at 2 5 ~ C . Fast motion of straight dislocations was observed. Frequent cross-slip events indicate that the mobile dislocations have a screw character (fig. 5). Motion of 60 ° dislocation was not observed. Further analysis of the dislocation structure performed
Figure 6 - Stereographic projection of the specimen during in situ deformation. N = normal to the foil plane. T.A = tensile axis.
V. DISCUSSION
Figure 5 - The path of a short screw dislocation aa' is evidenced by the slip traces labelled st. A kink on these slip traces shows that the dislocation has undergone a double cross-slip. The direction of the Burger's vector is given directly by the line connecting the two kinks. The picture is taken from dynamical video recordings.
We shall mainly discuss here the results obtained in situ, which can be surmnarized by three main points : (i) all the mobile dislocations are of screw character. (ii) their movement is rather fast. (iii) they do not belong to the primary system. We would like to discuss these three points one by one:
644
M. Fnaiech, 1~ Louchet /Dislocation motion in InSb
Figure 7 - Dislocations of the [II0] (I~I) system observed after in situ straining. The total extinction in (d) confirms the pure screw character.
M. Fnaiech, F. Louchet / Dislocation motion in lnSb
(i) The first point seems to be a simple thin foil effect : if one assumes, following (3), that half-loops are nucleated at the surface, and if the mobility of ~ dislocations is larger than that of screws, the former dislocations will readily propagate through the thin specimen, leaving a pair of opposite screw dislocations across the sample (fig. 8). A similar result would be obtained by operation of dislocation sources in the bulk, if both ~ and B dislocations are able to reach the surface readily. We are not able at the moment to decide between these two possibilities. However, the pile-ups of screw dislocations shown in fig. 7 suggest the existence of recurrent localized sources, which are easier to imagine as bulk sources than as surfaces ones.
w_
/11 (
ll¢,
I I
Figure 8 - Cross section of the thin foil during straining. Sources like E produce fast ~ dislocations which annihilate readily on the opposite surface. The deformation must then be accomodated by the remaining screw dislocations labelled S.
645
slowest type of dislocation. Although this criterion has been worked out for BCC metals at low temperatures, and cannot be applied directly here, a similar effect cannot be excluded at the moment. CONCLUSION Observations by electron microscopy of dislocation structures after bulk deformation confirm X-ray results and demonstrates the existence of B dislocations and the formation of dipoles by loop interaction. The mobile screw dislocations observed in thin foils during in situ deformation are thought to be left in the specimen by annihilation of fast 60 ° dislocations at the surfaces. These dislocations seem to be responsible for the major part of plastic deformation, and are probably produced by a small number of localized sources. We hope that further in situ experiments on specimens orientated for double slip will allow a direct observation of source operation and of ~ and B dislocations motion. AKNOWLEDGEMENTS The authors are indebted to Dr D. Ferr~ and coworkers for providin~ the strained specimens and for very helpfuldiscussions, to Mr L. Bernard for skilful specimen preparation and to the staff of the 3 MeV microscope for assistance during in situ experiments. REFERENCES
(I) Steinhardt H. and Sch~fer S., Acta Met. 19 p. 65, (1971). (ii) The dislocation structures in bulk specimen mentioned in § 3 (screw dipoles) suggest that ~ dislocations are faster than screws. Here, as shown in (i), all the 60 ° dislocations have disappeared, and the remaining screw dislocations have to accomodate the applied strain rate. This results in a fast movement of these dislocations, whatever their mobility might be, owing to an eventual adjustment of the applied stress. (iii) The shape of the specimen (non-uniform thickness, holes, ...) is probably responsible for the apparent deviation from Schmid's law. However, it has been shown by Vesel# (6) that in thin foils Schmld's factor might be modified by a term involving the angle between the Burger's vector and the foil surafce, which accounts for the ability of the surface to annihilate the
(2) Mihara M. and Ninomiya T., Phys. Stat. Sol. (a) 32, p. 43, (1975).
(3) Kesteloot R., Th~se de SpecialitY, Universit~ des Sciences et Techniques de Lille
(1981).
(4) Farvacque J.L. and Ferr~ D., J. Phys. 40-6, p. 157, (1979). (5) Valle R. and Martin J.L., 8th Int. Congress on Electron Microscopy, Vol. I. Eds. J.V. Sanders and D.J. Goodchild (Canberra, Australian Academy of Sciences) p. 180, (1974). (6) Vesely D., Phys. Stat. Sol. 29, 685 (1968).
Paper presented at ICDS-12 Amsterdam, August 31 - September 3, 1982
North-Holland PublishingCompany
DISLOCATION STRUCTURES AND IN SITU OBSERVATIONS OF DISLOCATION MOTION IN InSb M. FNAIECH* and F. LOUCHET**
* IMpartement de Physique, Facultd des Sciences, Tunis, Tunisia ** Laboratoire d'Optique Electronique du C2V.R.S., Toulouse, France Dislocation mobility in InSb is investigated both by a microscopic analysis of dislocation structures at lO0 kV and by in situ straining in a high voltage microscope. The observed differences in deformation processes between bulk specimens and thin foils is discussed. Dislocation pileups observed in thin foils indicate the existence of localized and recurrent dislocation sources.
I. INTRODUCTION
II. EXPERIMENTAL PROCEDURE
In III-V compounds of sphalerite structure, two types of 60 ° dislocations (labelled ~ and B) can be imagined, depending on the nature of the atoms which terminate the extra half planes in the dislocation cores. A perfect hexagonal loop is shown in fig. l, the sides of which lie parallel to the
InSb specimens strained up to the lower yield point by compression along [132] were provided by D. Ferr6 and R. Kesteloot (3). The geometry of these specimens is shown in fig. 2. The orientation of the compression axis favours the slip system [IT0] (Ill). One of the (5~4) surfaces was scratched in order to select in compression the nucleation of s-type half loops, i.e. ~ dislocations connected at both ends to screw dislocations.
/,lj
$ Figure I - Perfect hexagonal loop. ~ and B segments, of 60 ° character, lie on opposite sides of the screw segments.
I
BTol
\
/
//1~\
l f (1111 I
\\
S
15t"41 Indirect measurements of dislocation mobilities in InSb performed previously by etching techniques (], 2) showed indeed a large difference in mobilities of ~, 8 and screw dislocations. More recently, formation of long screw dipoles by the faster propagation of ~ dislocations from the surface has been shown by X-ray topography (3). However, 8 dislocations were not observed by these authors. In the present work, dislocations structures produced previously (3, 4) by compression of bulk specimens were observed by 100 kV electron microscopy. Some results of in situ deformation of InSb thin foils in a high voltage microscope are also reported. The latter specimens were also analysed in a conventional transmission microscope.
0378-4363/83/0000-0000/$03.00 © 1983 North-Holland
[TTI] Figure 2 - Orientation of bulk compression specimens. Thin foils for TEM observations were sliced parallel to the (Ill) plane. Thin speimens for in situ experiments were parallel to_the (514) plane, with a tensile axis along [132]. In order to observe the resulting dislocation structures, we sliced the specimens along the primary glide plane (Ill). The micro-tensile specimens for in situ experiments, on the opposite, were cut from (514) slices. This latter orientation is a compromise which allows good visibility of dislocations in the primar~ slip plane (Ill) in single slip conditions (132 tensile axis).
642
M. Fnaiech, F. Louchet / Dislocation motion in lnSb
Figure 3 - Dislocation structures in specimens strained by compression. The foil plane is the primary plane (Ill). The screw dislocations emerging at the surface have been probably built up by glide of the 60 ° segment. A closed hexagonal loop, which necessarily contains a B segment, is shown in A. (Note that half-loops are elongated along the Burgers vector direction, which suggests that 60 ° dislocations are faster than screws).
Figure 4 - Same foil as in fig. 3, showing dislocation dipoles resulting from mutual interaction of loops with the same Burgers vector a/2 [ll0] lying on parallel (Ill) planes.
M. Fnaiech, F. Louchet
/Dislocationmotion
The slices were first thinned down mechanically, and then chemically polished. The micro-tensile specimens were shaped and thinned down by preferential chemical etching, using wax coating as a protection. The in situ straining experiments were performed in the 3 MeV microscope of Toulouse, operating at I MeV. The heating straining stage (5) was used at about 250°C. The dynamical observations were recorded on a video tape recorder.
III. OBSERVATIONS OF DISLOCATION STRUCTURES IN SPECIMENS STRAINED BY COMPRESSION. - Figure 3 shows dislocation half-loops, of hexagonal shape. The dislocation segments lie along the three
in lnSb
643
in a 100 kV microscope (fig. 6 and fig. 7) showed that the screw dislocations mentioned above belong to two families, the Burgers vectors of which are a/2 [110] and a/2 [0Jl]. The slip planes have been unambiguously determined, and the results are summarized in the following table :
Burgers vector
Slip plane
Observation frequency
a
(711)
frequent frequent
0,36 0,12
a
(ITs)
frequent less frequent
0,29 0
a
(lll)
not observed not observed
0,48 (primary) 0
a
(111)
not observed not observed
0,36 0,18
7 [11o] (~TI)
T [o~] (TT1) T [17o] (TT~) T [To1] (iTl)
Theoretical Schmid's factor
- More complicated configurations were also frequently observed : figure 4 shows a dipole structure resulting from interaction of loops with the same Burgers vector, lying on neighbouring (Ill) planes.
IV. IN SITU DEFORMATION
\
The specimens were constrained at 2 5 ~ C . Fast motion of straight dislocations was observed. Frequent cross-slip events indicate that the mobile dislocations have a screw character (fig. 5). Motion of 60 ° dislocation was not observed. Further analysis of the dislocation structure performed
Figure 6 - Stereographic projection of the specimen during in situ deformation. N = normal to the foil plane. T.A = tensile axis.
V. DISCUSSION
Figure 5 - The path of a short screw dislocation aa' is evidenced by the slip traces labelled st. A kink on these slip traces shows that the dislocation has undergone a double cross-slip. The direction of the Burger's vector is given directly by the line connecting the two kinks. The picture is taken from dynamical video recordings.
We shall mainly discuss here the results obtained in situ, which can be surmnarized by three main points : (i) all the mobile dislocations are of screw character. (ii) their movement is rather fast. (iii) they do not belong to the primary system. We would like to discuss these three points one by one:
644
M. Fnaiech, 1~ Louchet /Dislocation motion in InSb
Figure 7 - Dislocations of the [II0] (I~I) system observed after in situ straining. The total extinction in (d) confirms the pure screw character.
M. Fnaiech, F. Louchet / Dislocation motion in lnSb
(i) The first point seems to be a simple thin foil effect : if one assumes, following (3), that half-loops are nucleated at the surface, and if the mobility of ~ dislocations is larger than that of screws, the former dislocations will readily propagate through the thin specimen, leaving a pair of opposite screw dislocations across the sample (fig. 8). A similar result would be obtained by operation of dislocation sources in the bulk, if both ~ and B dislocations are able to reach the surface readily. We are not able at the moment to decide between these two possibilities. However, the pile-ups of screw dislocations shown in fig. 7 suggest the existence of recurrent localized sources, which are easier to imagine as bulk sources than as surfaces ones.
w_
/11 (
ll¢,
I I
Figure 8 - Cross section of the thin foil during straining. Sources like E produce fast ~ dislocations which annihilate readily on the opposite surface. The deformation must then be accomodated by the remaining screw dislocations labelled S.
645
slowest type of dislocation. Although this criterion has been worked out for BCC metals at low temperatures, and cannot be applied directly here, a similar effect cannot be excluded at the moment. CONCLUSION Observations by electron microscopy of dislocation structures after bulk deformation confirm X-ray results and demonstrates the existence of B dislocations and the formation of dipoles by loop interaction. The mobile screw dislocations observed in thin foils during in situ deformation are thought to be left in the specimen by annihilation of fast 60 ° dislocations at the surfaces. These dislocations seem to be responsible for the major part of plastic deformation, and are probably produced by a small number of localized sources. We hope that further in situ experiments on specimens orientated for double slip will allow a direct observation of source operation and of ~ and B dislocations motion. AKNOWLEDGEMENTS The authors are indebted to Dr D. Ferr~ and coworkers for providin~ the strained specimens and for very helpfuldiscussions, to Mr L. Bernard for skilful specimen preparation and to the staff of the 3 MeV microscope for assistance during in situ experiments. REFERENCES
(I) Steinhardt H. and Sch~fer S., Acta Met. 19 p. 65, (1971). (ii) The dislocation structures in bulk specimen mentioned in § 3 (screw dipoles) suggest that ~ dislocations are faster than screws. Here, as shown in (i), all the 60 ° dislocations have disappeared, and the remaining screw dislocations have to accomodate the applied strain rate. This results in a fast movement of these dislocations, whatever their mobility might be, owing to an eventual adjustment of the applied stress. (iii) The shape of the specimen (non-uniform thickness, holes, ...) is probably responsible for the apparent deviation from Schmid's law. However, it has been shown by Vesel# (6) that in thin foils Schmld's factor might be modified by a term involving the angle between the Burger's vector and the foil surafce, which accounts for the ability of the surface to annihilate the
(2) Mihara M. and Ninomiya T., Phys. Stat. Sol. (a) 32, p. 43, (1975).
(3) Kesteloot R., Th~se de SpecialitY, Universit~ des Sciences et Techniques de Lille
(1981).
(4) Farvacque J.L. and Ferr~ D., J. Phys. 40-6, p. 157, (1979). (5) Valle R. and Martin J.L., 8th Int. Congress on Electron Microscopy, Vol. I. Eds. J.V. Sanders and D.J. Goodchild (Canberra, Australian Academy of Sciences) p. 180, (1974). (6) Vesely D., Phys. Stat. Sol. 29, 685 (1968).