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Thermal activation of glide in InP single crystals
OOOI-6160/87 $3.00 + 0.00 Copyright 6 1987 Pergamon Journals Ltd
Acra merall. Vol. 35, No. 1, pp. 143-148, 1987 Printed in Great Britain. All rights reserved
THERMAL
P. GALL,’
ACTIVATION OF GLIDE SINGLE CRYSTALS
IN InP
J. P. PEYRADE,’
R. COQUILLI?; F. REYNAUD,J S. GABILLET’ and A. ALBACETE’ ‘Laboratoire de Physique des Solides, INSA, Av. de Rangueil, 31077 Toulouse-Cedex, %.N.E.T., Route de Trkcastel, BP 40, 22301 Lannion-Cedex and ‘Laboratoire d’optique Electronique, CNRS, B.P. 4347, 3 1055 Toulouse-Cedex, France (Received 27 November 1985; in revised form 31 March 1986)
Abstract-The study of the plasticity of undoped indium phosphide between 573 and 1023 K by means of uniaxal compression (i 2: 10e4 s-l), stress relaxation and transmission electron microscopy shows that the deformation is controlled by the motion of screw dislocations in the stress field of the lattice according to a thermally activated Peierls mechanism. Between 573 and 673 K, this mechanism is pure; it is assisted by the internal stress between 673 and 1023 K and the lattice friction has no effect on the screw dislocations above 1023 K: the deformation is then controlled by the forest mechanism. To move freely in the crystal, a screw dislocation needs an energy below 2.8 eV. Rt&um&--Nous avons ttudi& la plasticit du phosphure d’indium non dopi, entre 573 et 1023 K, B I’aide d’essais de compression uniaxale (i N lOA s-‘) et de relaxation de contrainte, ainsi que par microscopic blectronique en transmission. Cette 6tude montre que la d&formation est contr&e par le mouvement des dislocations vis dans le champ de contrainte du riseau, suivant un mkcanisme de Peierls activi thermiquement. Entre 573 et 673 K, ce mkanisme est pur; il est aid& par la contrainte interne entre 673 et 1023 K; au dessus de 1023 K la friction de rkseau n’a pas d’effet sur les dislocations vis: la d&formation est alors contr%e par le m&canisme de la for& Pour pouvoir se d&placer librement dans le cristal, une dislocation vis a besoin d’une bnergie infkrieure g 2,8 eV. Zusammenfassung-Die Untersuchung des plastischen Verhaltens undotierten Indiumphosphids zwischen 573 und 1023 K mittels einachsigen Druckversuchs (e r 10m4s-l), Spannungsrelaxation und Durchstrahlungselektronenmikroskopie zeight, da5 die Verformung durch die Bewegung bon Schraubenversetzungen im Spannungsfeld des Gitters entsprechend einem thermisch aktivierten Peierlsmechanismus gesteuert wird. Zwischen 573 und 673 K lauft dieser Mechanismus in reiner Form ab. Zwischen 673 und 1023 K wird der Mechanismus von der inneren Spannung unterstiitzt; die Gitterreibung beeinflul3t den Mechanismus oberhalb 1023 K nicht mehr. Dann wird das Verformungsver halten durch Waldversetzungen gesteuert. Fiir eine freie Bewegung der Schraubenversetzungen im Kristall ist eine Energie unterhalb von 2.8 eV notwindig.
1. INTRODUCTION Numerous works have shown that the electrical and optical properties of semiconductors are modified at the vicinity of dislocations [l-3]. This localized degradation does not disturb large junctions but is becoming critical with the increasing miniaturization of electronic components. The development on an industrial scale of these new devices, on substrates the size of which is continuously increasing, is linked to the growth of large single crystals with a very low density of as grown dislocations and to the design of structures that do not contain areas favourable to the multiplication of dislocations. A better knowledge of the structure of the core of dislocations and the conditions governing their mobility and multiplication in bulk material should contribute to the solution of these problems in InP. The plastic deformation of undoped InP has been much less investigated than that of the tetravalent semiconducting elements Ge [4] and Si [5-71 and of other III-V compound such as InSb [&lo] and GaAs A.M. W--J
[lO-121. As far as we know, few publications have been devoted to it: the first presents measurement of the behaviour during uniaxial compression [14-l 61 for restricted experiental conditions. In this article, the standard techniques of plasticity [17, 181 have been applied to undoped single crystals of InP. After a detailed description of the experimental procedure, the results of uniaxial compression, stress relaxation and electron microscopy are presented. We have tried to identify the microscopic mechanisms controlling the deformation and to estimate their activation parameters. 2. EXPERIMENTAL The (lTl> single crystal was grown by the liquid encapsulated Czochralski technique at the CNET in Lannion. Its characteristics are presented in Table 1. The parallelepipedal compression samples were cut from a single (171) slice extracted from the middle of the ingot. The samples are ground with Carborundum and then mechanically polished to the final dimension 143
144
GALL
et al.: Table
THERMAL
I.
ACTIVATION
Physical characteristics
OF GLIDE
of the InP single crystal
300K n
(at cm-j) Undoped (ITI) 1x1~
1.5 to 3 x 10’6
IN InP
Mean density of etch pits
71 K p (cm*V-‘s-l)
P
(at CL-“) 0.2 to 2 x IO’6
3500
8 x 3 x 3 mm3. With our polishing apparatus, the thickness is constant to within 0.05 mm. The orientation [T23] is the usual one that favours simple glide of the dislocations with Burgers vectors b = f [iOl] lying in the (111) glide plane (Fig. 1). The deformation is performed by uniaxial compression along [T23], on an Inston machine equipped with an inverted set-up designed and built in the laboratory [19]. The speed of the mobile cross-bar is fixed at 50 pm/min (i N 10e4s -I). A differential transformer monitors the real deformation of the sample with a precision of 0.5 pm and a high sensitivity Instron cell measures the force with a precision of 1 N. The temperature is held constant at a value between 0 and 800°C to within 1°C. Furthermore, when the temperature is higher than 450°C the phosphorus evaporation is restricted by a B,O, encapsulation liquid joint. The stress relaxation tests have been performed by immobilizing the cross-bar during a compression test at a definite temperature. The parasitic relaxation of the compression machine is vey low; it is less than 1 N for the worst conditions. This enables us to measure precisely the strain rate sensitivity. The residual phosphorus evaporation enabled us to observe neither slip lines nor cracks. A slice [(T23) plane, thickness: 0.5 mm] is then spark cut in the middle of each deformed sample for observation by transmission electron microscopy. It is thinned by grinding and subsequently chemically in a solution of bromine in methanol. Such an orientation of the slice has been chosen because the chemical polishing of a (T23) slice is better than that
(cm4W)
(cm-*)
11500 to 8000
0.3 to 2x IO’
of a { 111) slice. Finally, irradiation effects have not been observed, although the accelerating voltage of the microscope (200 kV) was higher than the threshold for P displacement.
3. EXPERIMENTAL
RESULTS
3.1. Uniaxial compression 3.1.1. The predeformation. The samples contain different densities of as grown dislocations, which vary according to a W-shaped curve as a function of the diameter of the ingot [20]. To homogenize the initial structures of the samples, the latter are prestrained at TP = 0.77 T,,,, before the deformation at T < T,. At TP, undoped InP shows the standard semiconductor behaviour (Fig. 2). A yield drop precedes the usual three stages associated with c.f.c. crystals. A predeformation of about 2% has been chosen; this corresponds to the middle of stage 1 and ensures primary simple glide. Furthermore, this predeformation initiates surface and bulk dislocation sources and reduces the brittleness of the crystal at low temperature. 3.1.2. Compression curves. In Fig. 3(a), the compression curves r (yr) are shown at various temperatures, where r is the resolved shear stress, and yp the plastic shear strain in the primary (111) [iOl] glide system. The characteristic yield drop of semiconductors occurs at each temperature. Its magnitude and width increases when the temperature decreases.
ci231
T- 1023
a
Ai-
K
50 ,um/mn
I
/
6-
/rc,
0
I
I
I
100 AL
Fig. 1. Compression
sample. Primary [TOl].
glide system = (I 11)
Fig.
2. Strength-elongation
I
200
300
~-
undoped
InP
(pm)
curve 1023 K.
of
at
(4
t b 62
30
GALL et al.: THERMAL
ACTIVATION
decrease of the temperature involves an expansion of the curves from the origin; the transition between stages is more and more blurred and the curves become parabolic, (iii) InP becomes brittle at about 523 K.
(tAPa)
i
I
145
OF GLIDE IN InP
@r’
523K
In Fig. 3(b), the existence of a microplastic deformation stage is observed, the width of which at the upper yield point increases when the temperature decreases. 3.2. Stress relaxation The stress relaxation tests are performed during the compression. They are analysed by the technique of Guiu and Pratt [21] and by assuming that the crystal is submitted to the apparent state equation
Gap,(71 yapp = j. exp - A ~ kT
10
0
20
30
7,
40
(%)
t 62 MPa
(b)
523K i_lB4r’
4c
3c
573K I-
2(
623K
k II
In this equation, japp is the apparent strain rate resulting from all the microscopic deformation processes, v,, a pre-exponential factor, AC,,,(t) the apparent enthalpy of activation, k Boltzann’s constant and T the absolute temperature. From it we deduce the strain rate sensitivity 1 = (87/a lnv)T and the activation volume V = (kT)/1 that characterises the deformation from the macroscopic point of view. The relaxations become close to logarithmic as the temperature is lowered and they have been performed for a time longer than the characteristic constant. Further details of this argument, which is far from obvious in III-V compounds, are to be found in [19]. At low temperature, the activation volume shows a very high value in the microplastic stage, a minimum at the upper yield point and a weak parabolic variation beyond (Fig. 4). Its value at the lower yield point is very large at high temperatures; it then decreases rapidly when the temperature decreases and it becomes very small at low temperature ( 5 20b3 at 573 K) (Fig. 5).
h
673K T = 623
-723K
1023 s
K
c_ 625 e23K I 2
,
K I 4
I 6 e-8
rp
K
.
(%I
Fig. 3.(a) Resolved shear stress-resolved shear strain curves for undoped InP between 573 and 1023 K. (b) Resolved shear stress-resolved shear strain curves up to the lower yield point for undoped InP between 573 and 1023 K.
1oc
,-
,” l.
LYP
‘.
.-.
-.-s-
r-
We have observed that: (i) the lower yield point is thermally activated, (ii) at 1023 and 923 K, the curves show the three stages of the c-f-c- crystals, but a further
0
10
20 7
30
(%I
Fig. 4. Variation of the activation volume with the resolve shear strain at 623 K.
146
GALL ef al.:
Fig, 5. Activation
TIiERMAL
ACTIVATfON
OF GLlDE IN InP
volume at tie lower yield paint vs
temperature,
The structure of the dislocations produced by deformation in the most characteristic conditions (Fig, 6) has been observed *‘post mortem” by transmission electron microscopy, after unloading and cooling to room temperature. In all the samples, the density of dislocations is about lO*cm/cm3 and more than 80% of the dislmations Ixlong to the primary glide system (111) [TOI]. The predeformation at 1023 K generates mostly primary edge dislocations (Fig, 7) This is the initia1 rn~c~o~t~~ture of all the samples a&r predeformation. A subsequent deformation at low temperature up to the microplastic stage produces homogeneous structure of primary screw dislocations (Fig. 8). At the lower yield point, the typical configuration is still
Fig+ 7. B = [Oll]: undoped InP deformed at 1023 K in stage one. Primary edge dislocations are mainfy observed isolated (a), as toops (b), dipoles (c) or arrays of dipois (d), There are very few CO”dislocations (e>.
1023 K
Fig. 6. Deformation conditions of the undoped InP samples observed bv transmission electron microscouy.
Fig, 8. B = [Ol 11: undoped InP deformed at 573 K in microplastic stage.
GALL
er al.:
147
THERMAL ACTIVATION OF GLIDE IN InP
0
I
, 500
, 1000
Fig. 10. Variation of rZyp vs temperature.
Fig. 9. B = [Oll]: undoped InP deformed at 573 K at the lower yield point. formed by primary screw dislocations, but walls appear (Fig. 9) in which the structure is more complicated. 4. DISCUSSION 4.1. Nature of the dislocations participating in the deformation At low temperature, there is a microplastic deformation stage. In this stage, the edge dislocations created during predeformation have moved leaving a large majority of screw dislocations. It is likely that the microstrain arises from the exhaustion of the edge dislocations and the moving of 60”-dislocations. As in InSb [8,22], these dislocations leave behind screw dislocations that are less mobile and thus control the deformation. At the lower yield point, some strengthening begins to appear and at high temperature all types of dislocations seem to be very mobile. The interaction between dislocations and the dragging of superjogs [7, IO] may explain the presence of edge dislocations: this is the standard high temperature behaviour of semiconductors [7]. 4.2. The dynamical parameters of the deformation Two macroscopic parameters have been chosen to characterize the deformation: (i) the stress at the lower yield point corrected for the weak strengthening due to the increasing density
of dislocations during the yield drop, as calculated by Escaig et al. [23] (this correction is about 0.04 MPa at IO23 K and 2.5 MPa at 573 K), (ii) the activation volume at the lower yield point V. In Fig. 10, the variation of the corrected resolved shear stress at the lower yield point rzvLpis presented as a function of temperature. The results obtained on undoped InP by Brown et al. [14] and by Brasen and Bonner [I5] are in agreement with ours (Fig. 10). The curve may be divided into three parts: .a high temperature part, where the thermal activation is nearly equal to zero and the stress very low, @a low temperature part with an important thermal activation and where the stress is very high, @a transition zone between these two parts. In terms of the simple model where the stress can be split into thermal (7**) and athermal (7~) components [ 171, the athermal stress 7~ that enables us to deform the crystal beyond the athermal temperature r, N 1023 K (at i w 10e4 s-l) is estimated at I .5 MPa. By substraction, the thermal component of the stress is deduced for each temperature. From these results and from the activation volumes, the variation of the free enthalpy of activation is calculated as a function of a thermal stress 7**. The variation of 7** as a function of temperature being known, AG(T) can be deduced (Fig. 11): (i) AG(T) shows the expected linear variation. The slope depends on the logarithm of the strain rate. (ii) extrapolation of the AG(T**) curve to zero stress gives the apparent free enthalpy of activation of the obstacles that impede the motion of the dislocations: it is of the order of 2.8 eV (iii) the asympotic region of 7**(T) at low temperature tells us that the thermal stress
148
GALL et al.: THERMAL ACTIVATION OF GLIDE IN InP
3-
\
t
l\*
3 u
‘\
2
‘\
dislocations is of the order of 2.8 eV; this value is an overestimation because of the influence of high temperature phenomena on the measurement of the activation volume.
. : l*.
2-
*\.
-0
5.2. At high temperature
_b-
I\ I
I
loo0
500
I
I
T(K) 500
I
T**
( MPo)
_-_.-‘-
Y I.
20
10
(i) The stress is no longer thermally activated. (ii) The activation volume becomes very large, so that the Peierls hills have become transparent to the motion of screw and other dislocations. The thermally inactivated long range obstacles (forest and interaction) control the deformation. 5.3. At medium temperature
100 f
1
I-
Fig. 11. Cross-diagram in undoped InP: AC(T), AG(r**) and r**(T) are reported. required for overcoming large.
the obstacles is very
Finally, Fig. 12 shows that the free enthalpy of activation that we have calculated is of the same order as that found in Si [6,7], GaAs [lo] and InSb [IO]. Furthermore it seems that for a given stress the free enthalpy of activation of III-V compounds increases with their bandgap, as show in [24]. 5. CONCLUSION: MECHANISMS
THE DEFORMATION OF UNDOPED InP
5.1. At low temperature
The deformation is thermally activated, but the activation volume is too high to reveal a Peierls mechanism; it is transition zone between the high and low temperature mechanisms of deformation. We think that in such a temperature range, where 2~ is no longer negligible compared with z **, the athermal stress helps the nucleation and migration of double kinks. There may also be an increased contribution from inverse jumps. Acknowledgements-The experimental part of the study was performed in the Laboratoire d’optique Electronique du CNRS, in Toulouse, and we are grateful to Dr. B. Jouffrey for supporting the project. The thin foils for electron microscopy were prepared by J. Crestou.
REFERENCES 1. P. Petroff and R. L. Hartman, J. appl. Phys. 45, 3899
(1974). 2. 0. Ueda, S. Isozumi, S. Yamakoshi and T. Kotani, J. appl. Phys. 50, 765 (1979).
(i) The stress is thermally activated and the activa-
tion volume is small; this means that the microscopic mechanism controlling the deformation is localized like the nucleation and migration of double kinks on a Peierls potential. (ii) Transmission electron microscopy shows that the screw dislocations control the deformation: thus, the parameters that have been determined in this study are those of the motion of screw dislocations. The activation enthalpy for the motion of the screw
3. Y. Nanishi, S. Miyazawa and Y. Matsuoka, Symposium DRIP, Montpellier, p. 217. Elsevier, Amsterdam (1985).
4. H. Schaumburg, Phil. Mug. 25, 1429 (1972). 5. A. George, C. Escaravage, G. Champier and W. Schroter, Physica status solidi (b) 53, 483 (1972). M. Omri, Th&e I.N.P. Nancy (1981). A. Oueldennaoua, Thtse, I.N.P. Nancy (1983). R. Kesteloot, Thtse, Lille 1981) D. Ferre and J. L. Farvacque, Physica status solidi (a) 49, 737 (1978). 10. M. Karmouda, Thkse, Lille 1984). 11. D. Laister and G. M. Jenkins J. Muter. Sci. 8, 1218
6. 7. 8. 9.
(1973).
3-
aSi 15)
2~~~__,:_z~~~~~_.
;I -.
l-
0
-.-_
..*-*-+
1
t
1
20 T**
( MPa
(1983). G. Muller, R. Rupp, J. Vokl, H. Wolf and W. Blum, J. Cryst Growth 71, 771 (1985).
L. P. Kubin Thtse, Paris (1974). J. P. Peyrade, These Toulouse (1978). P. Gall, Thdse, Toulouse (1985). M. Tajima, Jap. J. uppl. Phys. 21, 227 (1982). F. Guiu and P. L. Pratt, Physica status solidi 6, 111 (1964). 22. J. Di Persio and R. Kesteloot, J. Physique Suppl. 44 18. 19. 20. 21.
.-
10
12. H. Steinhardt and P. Haasen Physica stutus solidi (a) 49, 93 (1978). 13. H. Nagai, Jap. J. appl. Phys., 20, 793 (1981). G. T. Brown, B. Cockayne and W. R. Macewan, J. Matter. Sci. 15 1469 (1980). D. Brasen and W. A. Bonner, Mater Sci. Engng 61, 167
I
I
30
40
1
Fig. 12. Comparison of the AG(r **) curves in Si, GaAs, InSb and InP.
C4, 469 (1983).
23. B. Escaig, J. L. Farvacque and D. Ferre, Physica status solidi (a) 71, 329 (1982). 24. S. K. Choi, M., Mihara and T. Ninomiya, Jap. J. uppl. Phys. 16, 737 (1977).
Acra merall. Vol. 35, No. 1, pp. 143-148, 1987 Printed in Great Britain. All rights reserved
THERMAL
P. GALL,’
ACTIVATION OF GLIDE SINGLE CRYSTALS
IN InP
J. P. PEYRADE,’
R. COQUILLI?; F. REYNAUD,J S. GABILLET’ and A. ALBACETE’ ‘Laboratoire de Physique des Solides, INSA, Av. de Rangueil, 31077 Toulouse-Cedex, %.N.E.T., Route de Trkcastel, BP 40, 22301 Lannion-Cedex and ‘Laboratoire d’optique Electronique, CNRS, B.P. 4347, 3 1055 Toulouse-Cedex, France (Received 27 November 1985; in revised form 31 March 1986)
Abstract-The study of the plasticity of undoped indium phosphide between 573 and 1023 K by means of uniaxal compression (i 2: 10e4 s-l), stress relaxation and transmission electron microscopy shows that the deformation is controlled by the motion of screw dislocations in the stress field of the lattice according to a thermally activated Peierls mechanism. Between 573 and 673 K, this mechanism is pure; it is assisted by the internal stress between 673 and 1023 K and the lattice friction has no effect on the screw dislocations above 1023 K: the deformation is then controlled by the forest mechanism. To move freely in the crystal, a screw dislocation needs an energy below 2.8 eV. Rt&um&--Nous avons ttudi& la plasticit du phosphure d’indium non dopi, entre 573 et 1023 K, B I’aide d’essais de compression uniaxale (i N lOA s-‘) et de relaxation de contrainte, ainsi que par microscopic blectronique en transmission. Cette 6tude montre que la d&formation est contr&e par le mouvement des dislocations vis dans le champ de contrainte du riseau, suivant un mkcanisme de Peierls activi thermiquement. Entre 573 et 673 K, ce mkanisme est pur; il est aid& par la contrainte interne entre 673 et 1023 K; au dessus de 1023 K la friction de rkseau n’a pas d’effet sur les dislocations vis: la d&formation est alors contr%e par le m&canisme de la for& Pour pouvoir se d&placer librement dans le cristal, une dislocation vis a besoin d’une bnergie infkrieure g 2,8 eV. Zusammenfassung-Die Untersuchung des plastischen Verhaltens undotierten Indiumphosphids zwischen 573 und 1023 K mittels einachsigen Druckversuchs (e r 10m4s-l), Spannungsrelaxation und Durchstrahlungselektronenmikroskopie zeight, da5 die Verformung durch die Bewegung bon Schraubenversetzungen im Spannungsfeld des Gitters entsprechend einem thermisch aktivierten Peierlsmechanismus gesteuert wird. Zwischen 573 und 673 K lauft dieser Mechanismus in reiner Form ab. Zwischen 673 und 1023 K wird der Mechanismus von der inneren Spannung unterstiitzt; die Gitterreibung beeinflul3t den Mechanismus oberhalb 1023 K nicht mehr. Dann wird das Verformungsver halten durch Waldversetzungen gesteuert. Fiir eine freie Bewegung der Schraubenversetzungen im Kristall ist eine Energie unterhalb von 2.8 eV notwindig.
1. INTRODUCTION Numerous works have shown that the electrical and optical properties of semiconductors are modified at the vicinity of dislocations [l-3]. This localized degradation does not disturb large junctions but is becoming critical with the increasing miniaturization of electronic components. The development on an industrial scale of these new devices, on substrates the size of which is continuously increasing, is linked to the growth of large single crystals with a very low density of as grown dislocations and to the design of structures that do not contain areas favourable to the multiplication of dislocations. A better knowledge of the structure of the core of dislocations and the conditions governing their mobility and multiplication in bulk material should contribute to the solution of these problems in InP. The plastic deformation of undoped InP has been much less investigated than that of the tetravalent semiconducting elements Ge [4] and Si [5-71 and of other III-V compound such as InSb [&lo] and GaAs A.M. W--J
[lO-121. As far as we know, few publications have been devoted to it: the first presents measurement of the behaviour during uniaxial compression [14-l 61 for restricted experiental conditions. In this article, the standard techniques of plasticity [17, 181 have been applied to undoped single crystals of InP. After a detailed description of the experimental procedure, the results of uniaxial compression, stress relaxation and electron microscopy are presented. We have tried to identify the microscopic mechanisms controlling the deformation and to estimate their activation parameters. 2. EXPERIMENTAL The (lTl> single crystal was grown by the liquid encapsulated Czochralski technique at the CNET in Lannion. Its characteristics are presented in Table 1. The parallelepipedal compression samples were cut from a single (171) slice extracted from the middle of the ingot. The samples are ground with Carborundum and then mechanically polished to the final dimension 143
144
GALL
et al.: Table
THERMAL
I.
ACTIVATION
Physical characteristics
OF GLIDE
of the InP single crystal
300K n
(at cm-j) Undoped (ITI) 1x1~
1.5 to 3 x 10’6
IN InP
Mean density of etch pits
71 K p (cm*V-‘s-l)
P
(at CL-“) 0.2 to 2 x IO’6
3500
8 x 3 x 3 mm3. With our polishing apparatus, the thickness is constant to within 0.05 mm. The orientation [T23] is the usual one that favours simple glide of the dislocations with Burgers vectors b = f [iOl] lying in the (111) glide plane (Fig. 1). The deformation is performed by uniaxial compression along [T23], on an Inston machine equipped with an inverted set-up designed and built in the laboratory [19]. The speed of the mobile cross-bar is fixed at 50 pm/min (i N 10e4s -I). A differential transformer monitors the real deformation of the sample with a precision of 0.5 pm and a high sensitivity Instron cell measures the force with a precision of 1 N. The temperature is held constant at a value between 0 and 800°C to within 1°C. Furthermore, when the temperature is higher than 450°C the phosphorus evaporation is restricted by a B,O, encapsulation liquid joint. The stress relaxation tests have been performed by immobilizing the cross-bar during a compression test at a definite temperature. The parasitic relaxation of the compression machine is vey low; it is less than 1 N for the worst conditions. This enables us to measure precisely the strain rate sensitivity. The residual phosphorus evaporation enabled us to observe neither slip lines nor cracks. A slice [(T23) plane, thickness: 0.5 mm] is then spark cut in the middle of each deformed sample for observation by transmission electron microscopy. It is thinned by grinding and subsequently chemically in a solution of bromine in methanol. Such an orientation of the slice has been chosen because the chemical polishing of a (T23) slice is better than that
(cm4W)
(cm-*)
11500 to 8000
0.3 to 2x IO’
of a { 111) slice. Finally, irradiation effects have not been observed, although the accelerating voltage of the microscope (200 kV) was higher than the threshold for P displacement.
3. EXPERIMENTAL
RESULTS
3.1. Uniaxial compression 3.1.1. The predeformation. The samples contain different densities of as grown dislocations, which vary according to a W-shaped curve as a function of the diameter of the ingot [20]. To homogenize the initial structures of the samples, the latter are prestrained at TP = 0.77 T,,,, before the deformation at T < T,. At TP, undoped InP shows the standard semiconductor behaviour (Fig. 2). A yield drop precedes the usual three stages associated with c.f.c. crystals. A predeformation of about 2% has been chosen; this corresponds to the middle of stage 1 and ensures primary simple glide. Furthermore, this predeformation initiates surface and bulk dislocation sources and reduces the brittleness of the crystal at low temperature. 3.1.2. Compression curves. In Fig. 3(a), the compression curves r (yr) are shown at various temperatures, where r is the resolved shear stress, and yp the plastic shear strain in the primary (111) [iOl] glide system. The characteristic yield drop of semiconductors occurs at each temperature. Its magnitude and width increases when the temperature decreases.
ci231
T- 1023
a
Ai-
K
50 ,um/mn
I
/
6-
/rc,
0
I
I
I
100 AL
Fig. 1. Compression
sample. Primary [TOl].
glide system = (I 11)
Fig.
2. Strength-elongation
I
200
300
~-
undoped
InP
(pm)
curve 1023 K.
of
at
(4
t b 62
30
GALL et al.: THERMAL
ACTIVATION
decrease of the temperature involves an expansion of the curves from the origin; the transition between stages is more and more blurred and the curves become parabolic, (iii) InP becomes brittle at about 523 K.
(tAPa)
i
I
145
OF GLIDE IN InP
@r’
523K
In Fig. 3(b), the existence of a microplastic deformation stage is observed, the width of which at the upper yield point increases when the temperature decreases. 3.2. Stress relaxation The stress relaxation tests are performed during the compression. They are analysed by the technique of Guiu and Pratt [21] and by assuming that the crystal is submitted to the apparent state equation
Gap,(71 yapp = j. exp - A ~ kT
10
0
20
30
7,
40
(%)
t 62 MPa
(b)
523K i_lB4r’
4c
3c
573K I-
2(
623K
k II
In this equation, japp is the apparent strain rate resulting from all the microscopic deformation processes, v,, a pre-exponential factor, AC,,,(t) the apparent enthalpy of activation, k Boltzann’s constant and T the absolute temperature. From it we deduce the strain rate sensitivity 1 = (87/a lnv)T and the activation volume V = (kT)/1 that characterises the deformation from the macroscopic point of view. The relaxations become close to logarithmic as the temperature is lowered and they have been performed for a time longer than the characteristic constant. Further details of this argument, which is far from obvious in III-V compounds, are to be found in [19]. At low temperature, the activation volume shows a very high value in the microplastic stage, a minimum at the upper yield point and a weak parabolic variation beyond (Fig. 4). Its value at the lower yield point is very large at high temperatures; it then decreases rapidly when the temperature decreases and it becomes very small at low temperature ( 5 20b3 at 573 K) (Fig. 5).
h
673K T = 623
-723K
1023 s
K
c_ 625 e23K I 2
,
K I 4
I 6 e-8
rp
K
.
(%I
Fig. 3.(a) Resolved shear stress-resolved shear strain curves for undoped InP between 573 and 1023 K. (b) Resolved shear stress-resolved shear strain curves up to the lower yield point for undoped InP between 573 and 1023 K.
1oc
,-
,” l.
LYP
‘.
.-.
-.-s-
r-
We have observed that: (i) the lower yield point is thermally activated, (ii) at 1023 and 923 K, the curves show the three stages of the c-f-c- crystals, but a further
0
10
20 7
30
(%I
Fig. 4. Variation of the activation volume with the resolve shear strain at 623 K.
146
GALL ef al.:
Fig, 5. Activation
TIiERMAL
ACTIVATfON
OF GLlDE IN InP
volume at tie lower yield paint vs
temperature,
The structure of the dislocations produced by deformation in the most characteristic conditions (Fig, 6) has been observed *‘post mortem” by transmission electron microscopy, after unloading and cooling to room temperature. In all the samples, the density of dislocations is about lO*cm/cm3 and more than 80% of the dislmations Ixlong to the primary glide system (111) [TOI]. The predeformation at 1023 K generates mostly primary edge dislocations (Fig, 7) This is the initia1 rn~c~o~t~~ture of all the samples a&r predeformation. A subsequent deformation at low temperature up to the microplastic stage produces homogeneous structure of primary screw dislocations (Fig. 8). At the lower yield point, the typical configuration is still
Fig+ 7. B = [Oll]: undoped InP deformed at 1023 K in stage one. Primary edge dislocations are mainfy observed isolated (a), as toops (b), dipoles (c) or arrays of dipois (d), There are very few CO”dislocations (e>.
1023 K
Fig. 6. Deformation conditions of the undoped InP samples observed bv transmission electron microscouy.
Fig, 8. B = [Ol 11: undoped InP deformed at 573 K in microplastic stage.
GALL
er al.:
147
THERMAL ACTIVATION OF GLIDE IN InP
0
I
, 500
, 1000
Fig. 10. Variation of rZyp vs temperature.
Fig. 9. B = [Oll]: undoped InP deformed at 573 K at the lower yield point. formed by primary screw dislocations, but walls appear (Fig. 9) in which the structure is more complicated. 4. DISCUSSION 4.1. Nature of the dislocations participating in the deformation At low temperature, there is a microplastic deformation stage. In this stage, the edge dislocations created during predeformation have moved leaving a large majority of screw dislocations. It is likely that the microstrain arises from the exhaustion of the edge dislocations and the moving of 60”-dislocations. As in InSb [8,22], these dislocations leave behind screw dislocations that are less mobile and thus control the deformation. At the lower yield point, some strengthening begins to appear and at high temperature all types of dislocations seem to be very mobile. The interaction between dislocations and the dragging of superjogs [7, IO] may explain the presence of edge dislocations: this is the standard high temperature behaviour of semiconductors [7]. 4.2. The dynamical parameters of the deformation Two macroscopic parameters have been chosen to characterize the deformation: (i) the stress at the lower yield point corrected for the weak strengthening due to the increasing density
of dislocations during the yield drop, as calculated by Escaig et al. [23] (this correction is about 0.04 MPa at IO23 K and 2.5 MPa at 573 K), (ii) the activation volume at the lower yield point V. In Fig. 10, the variation of the corrected resolved shear stress at the lower yield point rzvLpis presented as a function of temperature. The results obtained on undoped InP by Brown et al. [14] and by Brasen and Bonner [I5] are in agreement with ours (Fig. 10). The curve may be divided into three parts: .a high temperature part, where the thermal activation is nearly equal to zero and the stress very low, @a low temperature part with an important thermal activation and where the stress is very high, @a transition zone between these two parts. In terms of the simple model where the stress can be split into thermal (7**) and athermal (7~) components [ 171, the athermal stress 7~ that enables us to deform the crystal beyond the athermal temperature r, N 1023 K (at i w 10e4 s-l) is estimated at I .5 MPa. By substraction, the thermal component of the stress is deduced for each temperature. From these results and from the activation volumes, the variation of the free enthalpy of activation is calculated as a function of a thermal stress 7**. The variation of 7** as a function of temperature being known, AG(T) can be deduced (Fig. 11): (i) AG(T) shows the expected linear variation. The slope depends on the logarithm of the strain rate. (ii) extrapolation of the AG(T**) curve to zero stress gives the apparent free enthalpy of activation of the obstacles that impede the motion of the dislocations: it is of the order of 2.8 eV (iii) the asympotic region of 7**(T) at low temperature tells us that the thermal stress
148
GALL et al.: THERMAL ACTIVATION OF GLIDE IN InP
3-
\
t
l\*
3 u
‘\
2
‘\
dislocations is of the order of 2.8 eV; this value is an overestimation because of the influence of high temperature phenomena on the measurement of the activation volume.
. : l*.
2-
*\.
-0
5.2. At high temperature
_b-
I\ I
I
loo0
500
I
I
T(K) 500
I
T**
( MPo)
_-_.-‘-
Y I.
20
10
(i) The stress is no longer thermally activated. (ii) The activation volume becomes very large, so that the Peierls hills have become transparent to the motion of screw and other dislocations. The thermally inactivated long range obstacles (forest and interaction) control the deformation. 5.3. At medium temperature
100 f
1
I-
Fig. 11. Cross-diagram in undoped InP: AC(T), AG(r**) and r**(T) are reported. required for overcoming large.
the obstacles is very
Finally, Fig. 12 shows that the free enthalpy of activation that we have calculated is of the same order as that found in Si [6,7], GaAs [lo] and InSb [IO]. Furthermore it seems that for a given stress the free enthalpy of activation of III-V compounds increases with their bandgap, as show in [24]. 5. CONCLUSION: MECHANISMS
THE DEFORMATION OF UNDOPED InP
5.1. At low temperature
The deformation is thermally activated, but the activation volume is too high to reveal a Peierls mechanism; it is transition zone between the high and low temperature mechanisms of deformation. We think that in such a temperature range, where 2~ is no longer negligible compared with z **, the athermal stress helps the nucleation and migration of double kinks. There may also be an increased contribution from inverse jumps. Acknowledgements-The experimental part of the study was performed in the Laboratoire d’optique Electronique du CNRS, in Toulouse, and we are grateful to Dr. B. Jouffrey for supporting the project. The thin foils for electron microscopy were prepared by J. Crestou.
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(i) The stress is thermally activated and the activa-
tion volume is small; this means that the microscopic mechanism controlling the deformation is localized like the nucleation and migration of double kinks on a Peierls potential. (ii) Transmission electron microscopy shows that the screw dislocations control the deformation: thus, the parameters that have been determined in this study are those of the motion of screw dislocations. The activation enthalpy for the motion of the screw
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3-
aSi 15)
2~~~__,:_z~~~~~_.
;I -.
l-
0
-.-_
..*-*-+
1
t
1
20 T**
( MPa
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.-
10
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I
I
30
40
1
Fig. 12. Comparison of the AG(r **) curves in Si, GaAs, InSb and InP.
C4, 469 (1983).
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