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Effect of temperature and sulfur doping on the plastic deformation of InP single crystals
Materials Science and Engineering, 61 (1983) 167-172
167
Effect of Temperature and Sulfur Doping on the Plastic Deformation of InP Single Crystals D. BRASEN and W. A. BONNER
Bell Laboratories, Murray Hill, N J 07974 (U.S.A.) (Received April 11, 1983)
SUMMARY
In order to assess the mechanical behavior o f single crystals o f InP at high temperatures, deformation studies at various temperatures and dopant levels have been investigated. It is observed that the doped crystals are harder than the u n d o p e d crystals and that their strength depends on the doping level. In addition, deformation-induced substructures o f sulfur-doped crystals have been compared at various temperatures. Lomer-Cottrell locks and well-aligned cell structures are observed at 4 75 °C, whereas these features are absent at 675 °C. This observation implies that the annihilation o f edge dislocations within the cell walls may be occurring by glide and climb.
1. INTRODUCTION
Over the past several years a considerable a m o u n t of work on the plastic flow of semiconducting materials has been carried o u t [1-6]. In particular, "stage I deformation", which includes yield drop, has been studied most extensively. For example, Mahajan et al. [2] have examined the temperature dependence of the upper yield stress and the Li]ders band formation in d e f o r m e d silicon crystals. They have found that, as the temperature is increased, the crystal yields more uniformly and the upper yield stress decreases. Further, the dopant dependence of the upper yield point in GaAs, GaSb and InSb has been studied by Osvenskii et al. [3]. The doping of these III-V c o m p o u n d s with donor impurities appears to increase the upper yield point and the activation energy for dislocation movement, while doping with acceptors causes an 0025-5416/83/$3.00
apparent decrease in these parameters. Unfortunately, Osvenskii et al. did n o t report the density of grown-in dislocations for the differently d o p e d samples. Since Johnston [ 4] has shown that the upper yield point is strongly dependent on the density of mobile dislocations initially present, the conclusions of Osvenskii et al. [3] remain somewhat clouded. In contrast with the many observations made on the early stage of deformation there have been only a few studies of diamond cubic or sphalerite crystals at larger strains. Alexander [ 5] and Schiller et al. [6] have investigated germanium and InSb crystals oriented for single slip. Basically, they have found that a typical resolved shear stressshear strain curve exhibits three additional stages b e y o n d stage I. Stage II is characterized b y a small constant hardening rate, stage III consists of a much higher although still constant hardening rate, whereas in stage IV the curve takes on a concave shape. In the present investigation the deformation behavior of u n d o p e d and sulfur-doped InP single crystals, oriented for single slip, has been examined as a function of temperature. In addition, dislocation distributions present on the primary slip plane at higher strains have been investigated, and their effect on the stress-strain curves is discussed.
2. EXPERIMENTAL DETAILS
Nominally u n d o p e d and sulfur-doped InP single crystals were grown by the liquid encapsulated Czochralski technique. The boule was then cut and shaped into compression samples a b o u t 0.5 cm X 0.5 cm X 1.26 cm in size. However, because of material limitations the u n d o p e d samples were of a © Elsevier Sequoia/Printed in The Netherlands
168
size a b o u t 0.4 cm X 0.5 cm X 1.14 cm. The crystals were oriented so that the compression axis was in the [ i 2 3 ] direction and the side faces were parallel to the ( 1 i l ) and (541) planes. Subsequently, the samples were ground and then chemically polished in a 1.5% solution of Br 2 in methanol to remove the surface damage. Differently d o p e d samples were taken from the same boule in order to ensure that the densities of grown-in dislocations of all the specimens were approximately equal. The chemically polished samples were then deformed to a b o u t 34% strain in compression at a strain rate of 6.7 X 10-5s -i for the sulfurd o p e d specimens and a b o u t 9% strain at a rate of 7.3 X 10 -5 s-1 for the u n d o p e d crystals. Five samples with carrier concentrations of 1.35 X 10 is cm -3 were deformed at 400, 475, 550, 625 and 675 °C. In addition, two samples with carrier concentrations of 3.86 X l 0 is cm -3 were compressed at 475 and 675 °C, while the u n d o p e d samples were deformed at 400 and 500 °C. Deformation was carried o u t in a protective atmosphere of 4% H 2 in N 2 gas. The load was removed at the conclusion of a test and the specimen was allowed to cool to ambient temperature. The deformation behavior was evaluated using a c o m p u t e r program to convert stress versus cross-head speed curves to resolve shear stress-glide strain curves. In addition, samples d o p e d to 1.35 X 10 is cm -3 and deformed at 475, 550 and 675 °C were examined using transmission electron microscopy (TEM). Samples for microscopy were prepared b y cutting slices parallel to the primary slip plane, i.e. the (111) plane, and thinning from the ( l l l ) B face in a 1.5% Br in methanol solution. The thinned samples were then examined in a JEM 200 microscope operating at 200 keV. With the [ 1 0 i ] (111) as the primary slip system, the operating reflections were indexed in an internally consistent manner.
3. RESULTS
3.1. D e f o r m a t i o n characteristics
Figure 1 shows the stress-strain curve of an u n d o p e d crystal deformed at 500 °C. It is evident that the sample shows some nonuniform yielding and exhibits a Liider strain
~7 ~6 I-CD
500°C
3
I
Q
I
5
I
I
I
I
6 9 % GLIDE STRAIN
I
I
12
I
15
Fig. 1. Resolved shear stress-glide strain of a nominally undoped InP single crystal.
of less than 1%. According to Hahn [ 7], the fairly rapid increase in the strain around the lower yield point implies that the material outside the Lfider band is contributing a large part to the total deformation leading to a small Ltider strain. In addition, visual examination of the deformed samples indicates that a small Ltider strain is occurring up to 675 °C. Figure 2 shows the temperature dependence of samples d o p e d with sulfur to carrier concentrations of 1.35 X l 0 is c m -3. It is evident that, as the temperature is increased, there is a general tendency for the lattice to soften. In addition, the easy glide region, typical of stage II, becomes more prominent at higher temperatures. At higher strains the slopes of the curves between 475 and 625 °C remain roughly constant, behavior similar to that observed in stage III. However, at larger strains, at 675 °C the slope of the curve decreases with increasing strain, more closely resembling stage IV. Figure 3 compares two samples d o p e d to different dopant levels b u t deformed at the same temperature (675 °C). The higher d o p e d sample has a consistently greater flow stress than the lower d o p e d sample. This observation is in agreement with the results of several workers [3, 8, 9] who have observed that
169 2P
donor-type impurities increase the strength of III-V compounds with a sphalerite structure. In addition, the slope of the easy glide region is slightly higher for the higher doped sample, while the curvature at the higher strains remains approximately the same. The undoped sample deformed at 400 °C cracked while approaching its lower yield point. However, the sample with a carrier concentration of 1.35 X 1018 cracked at 400 °C before the upper yield point and the higher doped sample showed microcracking at 475 °C and near the upper yield point. Hence, it appears that, as the doping levels increase, the samples become more brittle.
2~
% GLIDE STRAIN
Fig. 2. R e s o l v e d s h e a r stress-glide s t r a i n curves showing temperature dependence of sulfur-doped InP single crystals h a v i n g carrier c o n c e n t r a t i o n s o f 1 . 3 5 x 1018 c m -3.
19 17 15
w9
cr
'0
7
14 21 % GLIDE STRAIN
28
55
Fig. 3. R e s o l v e d s h e a r stress-glide s t r a i n curves
showing t h e e f f e c t o f s u l f u r d o p i n g o n I n P at 6 7 5 °C. T h e u p p e r curve was p r o d u c e d f r o m a s a m p l e h a v i n g a carrier c o n c e n t r a t i o n o f 3.86 X 1018 c m -3, while the l o w e r curve is t h e r e s u l t o f a s a m p l e w i t h a carrier c o n c e n t r a t i o n o f 1 . 3 5 x 1018 c m -3.
3.2. Electron microscopy Figure 4 depicts the substructure observed in a sample deformed to approximately 34% strain at 475 °C. This microstructure is characterized by the presence of well-aligned cell walls which appear to lie along the (1i0) directions. On the assumption that the activated slip vector is + ~ [ 10S], all the dislocations belonging to the primary slip system should be in contrast for all the g values shown in the figure. However, when Figs. 4(a) and 4(b) are compared, it is evident that several dislocations, some of which are labeled L, lying in the ( 0 i l ) direction could have a +21[011] Burgers vector. Most of the dislocations are straight, although some do have kinks. A few other dislocations, some of which are identified by the labeling S, appear to be the result of secondary slip. When Figs. 4(a) and 4(c) are compared, it is evident that their Burgers vector is +~ [110]. Figure 5(a) is a TEM micrograph of the sample deformed at 550 °C. Contrast experiments similar to those described above show evidence of -+½[10S], +½ [110] and +½ [011] dislocations. Further, these dislocations also appear to align parallel to the (011) directions. The TEM micrograph of the sample deformed at 675 °C (Fig. 5(b)) shows a marked difference in the dislocation structure when compared with TEM micrographs of the crystals compressed at the lower temperatures. The first difference is that the dislocation density is much lower. In addition, although a few dislocations such as that marked K have kept their crystallographic directionality, most of the dislocations are randomly aligned. The Burgers vectors of
170
Fig. 4. TEM micrographs illustrating the contrast behavior of substructure features in a sulfur-doped sample deformed at 475 °C. The planes of the micrographs are approximately (111 ). The operating reflections are (a) (202), (b) (022) and (c) (220). The scale bar represents 1.0pro.
Fig. 5. Micrograph showing the mierostructure of a sulfur-doped sample compressed at (a) 550 °C and (b) 675 °C. The planes of the mierographs are approximately (111). The scale bar represents 1.0 pro.
171
these dislocations appear to be identical with those of the samples compressed at the lower temperatures.
4.
DISCUSSION
Several interesting features pertaining to the deformation behavior of sulfur-doped samples have emerged: (a) the general softening of the crystal lattice as the temperature is increased, (b) the dramatic change in yield drop, (c) the emergence of an easy glide region and (d) the shift from linear to approximately parabolic hardening. The general softening of the lattice characterized by the overall lowering of the flow stress as the temperature increases may be rationalized by noting that the frictional drag on dislocations characterized b y the PeierlsNabarro stress is apparently reduced. This could be envisaged as increasing the thermal energy of the dislocations and thereby enabling them to overcome the Peierls potential energy barrier at lower applied stresses. Therefore, dislocations become more mobile at high temperatures and move easily. The observed yield drop behavior could essentially be explained by either of two possible mechanisms. Cottrell and Bilby [10] envisaged that, when dislocations break away from their impurity atmospheres, a yield drop occurs because the stress required to propagate slip is smaller than the break-away stress. Johnston and Gilman [ 11], however, argued that the abrupt yield drop is a consequence of rapid multiplication of dislocations between the upper and lower yield points and the stress dependence of dislocation velocity. At the present time, on the basis of the current work, it is n o t possible to make an assessment of which model is correct. However, Hahn [ 7] has discussed this problem and has concluded that the yield drop, for iron and related b.c.c. metals, is based primarily on dislocation multiplication and velocity, although he has n o t ruled o u t the possibility that the unpinning of dislocations could have a secondary effect. The next striking feature is the temperature dependence of the easy glide region. At 475 °C this region is virtually non-existent. Hence, after the lower yield point the sample could be deforming by multiple slip. However,
as the temperature increases, single slip is able to a c c o m m o d a t e the imposed deformation. At the higher strains the change in character of the work hardening from linear to parabolic may be rationalized by examining the observed microstructures. In order to do this, it must first be realized that strain hardening can be attributed to the progressive introduction, during straining, of barriers to the m o v e m e n t of dislocations. One wellknown barrier is the Lomer-Cottrell (LC) lock. Glide dislocations, as observed in Figs. 4 and 5, could interact with each other according to the following reaction to form the LC locks: a
a
-_
a
2 [10i] + 2 [ 1 1 0 ] - ~ 2 [ 0 i i l In addition, Karnthaler [ 12] has shown that for f.c.c, metals either (1) the LC dislocations are not dissociated and are able to glide on the {001} planes or (2) the locks are sessile and actually split on two intersecting {111} planes. If the locks are sessile, then the movement of primary and secondary dislocations is severely limited. For example, at 475 °C the high directionality of the dislocations such as those marked L appear for the most part to be LC sessiles. Further, it appears that at 475 °C a majority of the work hardening could be due to the pile-up of the 60 ° dislocations. This appears to be the result of the interaction of the dislocations with the LC locks. This can be seen by comparing Figs. 4(a) and 4(b) and noting that in most cases, wherever pile-up occurs, LC locks are present. At 550 °C, most of the LC dislocations still appear to be sessile. However, the absence of cell walls in the [10i] direction seems to indicate that the density of the screw dislocations of the primary slip has decreased as a result of thermally activated cross-glide and annihilation. At 675 °C, very few of the dislocations show directionality, as observed in Fig. 5(b). Although some LC locks such as that labeled K still show sessile characteristics, the vast majority show indications of m o v e m e n t such as the LC lock labeled D. In addition, the low density of dislocations at this temperature implies that annihilation has occurred. Most of the primary and secondary screw dislocations could annihilate each other by glide
172
and cross-slip. However, the edge character of LC and 60 ° dislocations implies that their annihilation occurred as a consequence of both glide and climb. Further, at this temperature the thermal energy is so high that covalent b o n d e d InP acts mechanically in a very similar w a y to f.c.c, metals. This assessm e n t is in agreement with previous observations on deformed germanium [ 1 ]. If this metallic behavior is truly indicative of InP at high temperatures, then it would be reasonable to believe that, as Karnthaler [12] has found in f.c.c, metals, the majority of LC locks, which are split on two intersecting planes at the lower temperatures, form whole dislocations at higher temperatures. These whole LC dislocations are then free to glide on the {001} planes. Hence, at the higher temperatures, n o t only are the 60 ° dislocations annihilated by glide and climb but similarly so could the LC dislocations be. Thus, at high strains the temperature dependence of the flow stress of the InP lattice appears to be determined n o t only b y the Peierls stress b u t also by thermally induced dislocation interaction mechanisms. Consequently, at lower temperatures and moderate strains the work hardening continues to be linear. However, at higher temperatures the annihilation of LC dislocations, coupled with the corresponding decrease in resistance to dislocation motion, has the effect of slowing the work hardening and thus producing the concave shape of the stress-strain curve. The effect of doping level on the flow stress seems to be slightly different from the temperature dependence. In particular, the similar curvature at the higher strains of the graph in Fig. 3 seems to indicate that the mechanism of dislocation annihilation is similar. However, the increased stress required at the higher dopant level suggests that the resistance to the m o v e m e n t of dislocations is
higher at the higher doping levels. The effect of increasing the amount of sulfur apparently increases the drag on the dislocations in a way similar to that proposed by Gilman [13] and used b y Swaminathan and Copley [9] to explain the hardening of GaAs, i.e. there may be a short-range electrostatic interaction b e t w e e n solute-vacancy pairs and dislocations associated with the displacement of the oppositely charged defects relative to each other b y the shearing action of the gliding dislocations. Hence, the major effect of increasing the sulfur doping, b y almost a factor of 3, at a high constant temperature appears to be the increasing of the resistance to dislocation m o v e m e n t whilst not affecting the bonding or annihilation mechanisms.
REFERENCES 1 H. Alexander and P. Haasen, Solid State Phys., 22 (1968) 27-156. 2 S. Mahajan, D. Brasen and P. Haasen, Acta Metall., 27 (1979) 1165-1173. 3 V. B. Osvenskii, O. G. Stolyarov and M. G. Mil'vidskii, Soy. Phys. -- Solid State, 10 ( 11 ) (1969) 2540-2543. 4 W. G. Johnston, J. Appl. Phys., 33 (9) (1962) 2716-2730. 5 H. Alexander, Z. Metallkd., 52 (5) (1961) 344352. 6 S. Schiifer, H. Alexander and P. Haasen, Phys. Status Solidi, 5 (1964) 247. 7 G. T. Hahn, Acta Metall., 10 (1962) 727-738. 8 N. P. Sazhin, M. G. Mil'vidskii, V. B. Osvenskii and O. G. Stolyarov, Soy. Phys. -- Solid State, 8 (5) (1966) 1223-1227. 9 V. Swaminathan and S. M. Copley, J. Appl. Phys., 47 (10) (1976) 4405-4413. 10 A. H. Cottrell and B. A. Bilby, Proc. Phys. Soc. London, Sect. A, 62 (1949) 49. 11 W. G. Johnston and J. J. Gilman, J. Appl. Phys., 30 (1959) 129. 12 H. P. Karnthaler, Philos. Mag. A, 38 (2) (1978) 141-156. 13 J . J . Gilman, J. Appl. Phys., 45 (1974) 508.
167
Effect of Temperature and Sulfur Doping on the Plastic Deformation of InP Single Crystals D. BRASEN and W. A. BONNER
Bell Laboratories, Murray Hill, N J 07974 (U.S.A.) (Received April 11, 1983)
SUMMARY
In order to assess the mechanical behavior o f single crystals o f InP at high temperatures, deformation studies at various temperatures and dopant levels have been investigated. It is observed that the doped crystals are harder than the u n d o p e d crystals and that their strength depends on the doping level. In addition, deformation-induced substructures o f sulfur-doped crystals have been compared at various temperatures. Lomer-Cottrell locks and well-aligned cell structures are observed at 4 75 °C, whereas these features are absent at 675 °C. This observation implies that the annihilation o f edge dislocations within the cell walls may be occurring by glide and climb.
1. INTRODUCTION
Over the past several years a considerable a m o u n t of work on the plastic flow of semiconducting materials has been carried o u t [1-6]. In particular, "stage I deformation", which includes yield drop, has been studied most extensively. For example, Mahajan et al. [2] have examined the temperature dependence of the upper yield stress and the Li]ders band formation in d e f o r m e d silicon crystals. They have found that, as the temperature is increased, the crystal yields more uniformly and the upper yield stress decreases. Further, the dopant dependence of the upper yield point in GaAs, GaSb and InSb has been studied by Osvenskii et al. [3]. The doping of these III-V c o m p o u n d s with donor impurities appears to increase the upper yield point and the activation energy for dislocation movement, while doping with acceptors causes an 0025-5416/83/$3.00
apparent decrease in these parameters. Unfortunately, Osvenskii et al. did n o t report the density of grown-in dislocations for the differently d o p e d samples. Since Johnston [ 4] has shown that the upper yield point is strongly dependent on the density of mobile dislocations initially present, the conclusions of Osvenskii et al. [3] remain somewhat clouded. In contrast with the many observations made on the early stage of deformation there have been only a few studies of diamond cubic or sphalerite crystals at larger strains. Alexander [ 5] and Schiller et al. [6] have investigated germanium and InSb crystals oriented for single slip. Basically, they have found that a typical resolved shear stressshear strain curve exhibits three additional stages b e y o n d stage I. Stage II is characterized b y a small constant hardening rate, stage III consists of a much higher although still constant hardening rate, whereas in stage IV the curve takes on a concave shape. In the present investigation the deformation behavior of u n d o p e d and sulfur-doped InP single crystals, oriented for single slip, has been examined as a function of temperature. In addition, dislocation distributions present on the primary slip plane at higher strains have been investigated, and their effect on the stress-strain curves is discussed.
2. EXPERIMENTAL DETAILS
Nominally u n d o p e d and sulfur-doped InP single crystals were grown by the liquid encapsulated Czochralski technique. The boule was then cut and shaped into compression samples a b o u t 0.5 cm X 0.5 cm X 1.26 cm in size. However, because of material limitations the u n d o p e d samples were of a © Elsevier Sequoia/Printed in The Netherlands
168
size a b o u t 0.4 cm X 0.5 cm X 1.14 cm. The crystals were oriented so that the compression axis was in the [ i 2 3 ] direction and the side faces were parallel to the ( 1 i l ) and (541) planes. Subsequently, the samples were ground and then chemically polished in a 1.5% solution of Br 2 in methanol to remove the surface damage. Differently d o p e d samples were taken from the same boule in order to ensure that the densities of grown-in dislocations of all the specimens were approximately equal. The chemically polished samples were then deformed to a b o u t 34% strain in compression at a strain rate of 6.7 X 10-5s -i for the sulfurd o p e d specimens and a b o u t 9% strain at a rate of 7.3 X 10 -5 s-1 for the u n d o p e d crystals. Five samples with carrier concentrations of 1.35 X 10 is cm -3 were deformed at 400, 475, 550, 625 and 675 °C. In addition, two samples with carrier concentrations of 3.86 X l 0 is cm -3 were compressed at 475 and 675 °C, while the u n d o p e d samples were deformed at 400 and 500 °C. Deformation was carried o u t in a protective atmosphere of 4% H 2 in N 2 gas. The load was removed at the conclusion of a test and the specimen was allowed to cool to ambient temperature. The deformation behavior was evaluated using a c o m p u t e r program to convert stress versus cross-head speed curves to resolve shear stress-glide strain curves. In addition, samples d o p e d to 1.35 X 10 is cm -3 and deformed at 475, 550 and 675 °C were examined using transmission electron microscopy (TEM). Samples for microscopy were prepared b y cutting slices parallel to the primary slip plane, i.e. the (111) plane, and thinning from the ( l l l ) B face in a 1.5% Br in methanol solution. The thinned samples were then examined in a JEM 200 microscope operating at 200 keV. With the [ 1 0 i ] (111) as the primary slip system, the operating reflections were indexed in an internally consistent manner.
3. RESULTS
3.1. D e f o r m a t i o n characteristics
Figure 1 shows the stress-strain curve of an u n d o p e d crystal deformed at 500 °C. It is evident that the sample shows some nonuniform yielding and exhibits a Liider strain
~7 ~6 I-CD
500°C
3
I
Q
I
5
I
I
I
I
6 9 % GLIDE STRAIN
I
I
12
I
15
Fig. 1. Resolved shear stress-glide strain of a nominally undoped InP single crystal.
of less than 1%. According to Hahn [ 7], the fairly rapid increase in the strain around the lower yield point implies that the material outside the Lfider band is contributing a large part to the total deformation leading to a small Ltider strain. In addition, visual examination of the deformed samples indicates that a small Ltider strain is occurring up to 675 °C. Figure 2 shows the temperature dependence of samples d o p e d with sulfur to carrier concentrations of 1.35 X l 0 is c m -3. It is evident that, as the temperature is increased, there is a general tendency for the lattice to soften. In addition, the easy glide region, typical of stage II, becomes more prominent at higher temperatures. At higher strains the slopes of the curves between 475 and 625 °C remain roughly constant, behavior similar to that observed in stage III. However, at larger strains, at 675 °C the slope of the curve decreases with increasing strain, more closely resembling stage IV. Figure 3 compares two samples d o p e d to different dopant levels b u t deformed at the same temperature (675 °C). The higher d o p e d sample has a consistently greater flow stress than the lower d o p e d sample. This observation is in agreement with the results of several workers [3, 8, 9] who have observed that
169 2P
donor-type impurities increase the strength of III-V compounds with a sphalerite structure. In addition, the slope of the easy glide region is slightly higher for the higher doped sample, while the curvature at the higher strains remains approximately the same. The undoped sample deformed at 400 °C cracked while approaching its lower yield point. However, the sample with a carrier concentration of 1.35 X 1018 cracked at 400 °C before the upper yield point and the higher doped sample showed microcracking at 475 °C and near the upper yield point. Hence, it appears that, as the doping levels increase, the samples become more brittle.
2~
% GLIDE STRAIN
Fig. 2. R e s o l v e d s h e a r stress-glide s t r a i n curves showing temperature dependence of sulfur-doped InP single crystals h a v i n g carrier c o n c e n t r a t i o n s o f 1 . 3 5 x 1018 c m -3.
19 17 15
w9
cr
'0
7
14 21 % GLIDE STRAIN
28
55
Fig. 3. R e s o l v e d s h e a r stress-glide s t r a i n curves
showing t h e e f f e c t o f s u l f u r d o p i n g o n I n P at 6 7 5 °C. T h e u p p e r curve was p r o d u c e d f r o m a s a m p l e h a v i n g a carrier c o n c e n t r a t i o n o f 3.86 X 1018 c m -3, while the l o w e r curve is t h e r e s u l t o f a s a m p l e w i t h a carrier c o n c e n t r a t i o n o f 1 . 3 5 x 1018 c m -3.
3.2. Electron microscopy Figure 4 depicts the substructure observed in a sample deformed to approximately 34% strain at 475 °C. This microstructure is characterized by the presence of well-aligned cell walls which appear to lie along the (1i0) directions. On the assumption that the activated slip vector is + ~ [ 10S], all the dislocations belonging to the primary slip system should be in contrast for all the g values shown in the figure. However, when Figs. 4(a) and 4(b) are compared, it is evident that several dislocations, some of which are labeled L, lying in the ( 0 i l ) direction could have a +21[011] Burgers vector. Most of the dislocations are straight, although some do have kinks. A few other dislocations, some of which are identified by the labeling S, appear to be the result of secondary slip. When Figs. 4(a) and 4(c) are compared, it is evident that their Burgers vector is +~ [110]. Figure 5(a) is a TEM micrograph of the sample deformed at 550 °C. Contrast experiments similar to those described above show evidence of -+½[10S], +½ [110] and +½ [011] dislocations. Further, these dislocations also appear to align parallel to the (011) directions. The TEM micrograph of the sample deformed at 675 °C (Fig. 5(b)) shows a marked difference in the dislocation structure when compared with TEM micrographs of the crystals compressed at the lower temperatures. The first difference is that the dislocation density is much lower. In addition, although a few dislocations such as that marked K have kept their crystallographic directionality, most of the dislocations are randomly aligned. The Burgers vectors of
170
Fig. 4. TEM micrographs illustrating the contrast behavior of substructure features in a sulfur-doped sample deformed at 475 °C. The planes of the micrographs are approximately (111 ). The operating reflections are (a) (202), (b) (022) and (c) (220). The scale bar represents 1.0pro.
Fig. 5. Micrograph showing the mierostructure of a sulfur-doped sample compressed at (a) 550 °C and (b) 675 °C. The planes of the mierographs are approximately (111). The scale bar represents 1.0 pro.
171
these dislocations appear to be identical with those of the samples compressed at the lower temperatures.
4.
DISCUSSION
Several interesting features pertaining to the deformation behavior of sulfur-doped samples have emerged: (a) the general softening of the crystal lattice as the temperature is increased, (b) the dramatic change in yield drop, (c) the emergence of an easy glide region and (d) the shift from linear to approximately parabolic hardening. The general softening of the lattice characterized by the overall lowering of the flow stress as the temperature increases may be rationalized by noting that the frictional drag on dislocations characterized b y the PeierlsNabarro stress is apparently reduced. This could be envisaged as increasing the thermal energy of the dislocations and thereby enabling them to overcome the Peierls potential energy barrier at lower applied stresses. Therefore, dislocations become more mobile at high temperatures and move easily. The observed yield drop behavior could essentially be explained by either of two possible mechanisms. Cottrell and Bilby [10] envisaged that, when dislocations break away from their impurity atmospheres, a yield drop occurs because the stress required to propagate slip is smaller than the break-away stress. Johnston and Gilman [ 11], however, argued that the abrupt yield drop is a consequence of rapid multiplication of dislocations between the upper and lower yield points and the stress dependence of dislocation velocity. At the present time, on the basis of the current work, it is n o t possible to make an assessment of which model is correct. However, Hahn [ 7] has discussed this problem and has concluded that the yield drop, for iron and related b.c.c. metals, is based primarily on dislocation multiplication and velocity, although he has n o t ruled o u t the possibility that the unpinning of dislocations could have a secondary effect. The next striking feature is the temperature dependence of the easy glide region. At 475 °C this region is virtually non-existent. Hence, after the lower yield point the sample could be deforming by multiple slip. However,
as the temperature increases, single slip is able to a c c o m m o d a t e the imposed deformation. At the higher strains the change in character of the work hardening from linear to parabolic may be rationalized by examining the observed microstructures. In order to do this, it must first be realized that strain hardening can be attributed to the progressive introduction, during straining, of barriers to the m o v e m e n t of dislocations. One wellknown barrier is the Lomer-Cottrell (LC) lock. Glide dislocations, as observed in Figs. 4 and 5, could interact with each other according to the following reaction to form the LC locks: a
a
-_
a
2 [10i] + 2 [ 1 1 0 ] - ~ 2 [ 0 i i l In addition, Karnthaler [ 12] has shown that for f.c.c, metals either (1) the LC dislocations are not dissociated and are able to glide on the {001} planes or (2) the locks are sessile and actually split on two intersecting {111} planes. If the locks are sessile, then the movement of primary and secondary dislocations is severely limited. For example, at 475 °C the high directionality of the dislocations such as those marked L appear for the most part to be LC sessiles. Further, it appears that at 475 °C a majority of the work hardening could be due to the pile-up of the 60 ° dislocations. This appears to be the result of the interaction of the dislocations with the LC locks. This can be seen by comparing Figs. 4(a) and 4(b) and noting that in most cases, wherever pile-up occurs, LC locks are present. At 550 °C, most of the LC dislocations still appear to be sessile. However, the absence of cell walls in the [10i] direction seems to indicate that the density of the screw dislocations of the primary slip has decreased as a result of thermally activated cross-glide and annihilation. At 675 °C, very few of the dislocations show directionality, as observed in Fig. 5(b). Although some LC locks such as that labeled K still show sessile characteristics, the vast majority show indications of m o v e m e n t such as the LC lock labeled D. In addition, the low density of dislocations at this temperature implies that annihilation has occurred. Most of the primary and secondary screw dislocations could annihilate each other by glide
172
and cross-slip. However, the edge character of LC and 60 ° dislocations implies that their annihilation occurred as a consequence of both glide and climb. Further, at this temperature the thermal energy is so high that covalent b o n d e d InP acts mechanically in a very similar w a y to f.c.c, metals. This assessm e n t is in agreement with previous observations on deformed germanium [ 1 ]. If this metallic behavior is truly indicative of InP at high temperatures, then it would be reasonable to believe that, as Karnthaler [12] has found in f.c.c, metals, the majority of LC locks, which are split on two intersecting planes at the lower temperatures, form whole dislocations at higher temperatures. These whole LC dislocations are then free to glide on the {001} planes. Hence, at the higher temperatures, n o t only are the 60 ° dislocations annihilated by glide and climb but similarly so could the LC dislocations be. Thus, at high strains the temperature dependence of the flow stress of the InP lattice appears to be determined n o t only b y the Peierls stress b u t also by thermally induced dislocation interaction mechanisms. Consequently, at lower temperatures and moderate strains the work hardening continues to be linear. However, at higher temperatures the annihilation of LC dislocations, coupled with the corresponding decrease in resistance to dislocation motion, has the effect of slowing the work hardening and thus producing the concave shape of the stress-strain curve. The effect of doping level on the flow stress seems to be slightly different from the temperature dependence. In particular, the similar curvature at the higher strains of the graph in Fig. 3 seems to indicate that the mechanism of dislocation annihilation is similar. However, the increased stress required at the higher dopant level suggests that the resistance to the m o v e m e n t of dislocations is
higher at the higher doping levels. The effect of increasing the amount of sulfur apparently increases the drag on the dislocations in a way similar to that proposed by Gilman [13] and used b y Swaminathan and Copley [9] to explain the hardening of GaAs, i.e. there may be a short-range electrostatic interaction b e t w e e n solute-vacancy pairs and dislocations associated with the displacement of the oppositely charged defects relative to each other b y the shearing action of the gliding dislocations. Hence, the major effect of increasing the sulfur doping, b y almost a factor of 3, at a high constant temperature appears to be the increasing of the resistance to dislocation m o v e m e n t whilst not affecting the bonding or annihilation mechanisms.
REFERENCES 1 H. Alexander and P. Haasen, Solid State Phys., 22 (1968) 27-156. 2 S. Mahajan, D. Brasen and P. Haasen, Acta Metall., 27 (1979) 1165-1173. 3 V. B. Osvenskii, O. G. Stolyarov and M. G. Mil'vidskii, Soy. Phys. -- Solid State, 10 ( 11 ) (1969) 2540-2543. 4 W. G. Johnston, J. Appl. Phys., 33 (9) (1962) 2716-2730. 5 H. Alexander, Z. Metallkd., 52 (5) (1961) 344352. 6 S. Schiifer, H. Alexander and P. Haasen, Phys. Status Solidi, 5 (1964) 247. 7 G. T. Hahn, Acta Metall., 10 (1962) 727-738. 8 N. P. Sazhin, M. G. Mil'vidskii, V. B. Osvenskii and O. G. Stolyarov, Soy. Phys. -- Solid State, 8 (5) (1966) 1223-1227. 9 V. Swaminathan and S. M. Copley, J. Appl. Phys., 47 (10) (1976) 4405-4413. 10 A. H. Cottrell and B. A. Bilby, Proc. Phys. Soc. London, Sect. A, 62 (1949) 49. 11 W. G. Johnston and J. J. Gilman, J. Appl. Phys., 30 (1959) 129. 12 H. P. Karnthaler, Philos. Mag. A, 38 (2) (1978) 141-156. 13 J . J . Gilman, J. Appl. Phys., 45 (1974) 508.