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Direct observation of crystallization in silicon by in situ high-resolution electron microscopy
ultramicroscopy
UltramicroscopyNorth.Holland51 (1993)41-45
Direct observation of crystallization in silicon by in situ high-resolution electron microscopy R. Sinclair
a,
j. Morgiel
a,b,
m.s. Kirtikar
a,
I.-W. Wu c and A. Chiang c
a Department of Materials Science and Engineering, Stanford University, Stanford, CA 94305, USA b Institute for Metals Research, Polish Academy of Science, Krakow, Poland c Xerox Corporation, Palo Alto Research Center, Palo Alto, CA 94304, USA Received 13 November 1992
We have studied the nucleation and growth of crystalline silicon by in situ high-resolution electron microscopy. Amorphous silicon thin films, deposited onto oxidized silicon wafers by low-pressure chemical vapor deposition, were heated in the microscope in the nominal temperature range of 700-775°C. Many sub-critical crystal embryos exist at the a-Si/SiO 2 interface, very few of which develop into viable nuclei. One nucleation event was recorded successfully at lattice resolution, allowing an estimate of the critical radius as 2.5+ 1.0 nm and the a-Si/c-Si interfacial energy as 600+200 mJ m -2. The crystal growth was followed for extended periods of time. It appeared to be characterized by sporadic growth bursts rather than a continuous advance of the crystal-amorphous phase interface. We specifically sought evidence for a ledge growth mechanism but did not find such behavior.
I. Introduction One of the most difficult occurrences to investigate directly is the nucleation event of a solidstate phase transformation. This arises because of the statistical nature and the minute scale (e.g. 1-10 nm) on which nucleation takes place. However, as first reported by Sinclair and Parker [1], it is now possible to study reactions in solids at the atomic level, at controlled elevated temperatures, by in situ, high-resolution electron microscopy (HREM) [2]. Accordingly we have attempted, and succeeded in, recording such behavior, using as an example the crystallization of amorphous silicon thin films. Silicon, being a simple element which can be easily fabricated in many forms, provides an excellent means of carrying out basic scientific experiments of this type. In addition, processes such as crystal nucleation and growth are extremely important during semiconductor-based, integrated circuit manufacture, hence such work can have technological significance. Recently Reiche and Hopfe have also
shown that in situ hot-stage TEM techniques can be used effectively to study crystallization in amorphous silicon [3]. In our work, we have focussed on an atomic-level study of the ongoing interfacial processes.
2. Experimental procedure and results The system we chose for this project consists of amorphous silicon thin films, formed by lowpressure chemical vapor deposition techniques onto thermally oxidized single-crystal silicon wafers. A schematic cross-section view is shown in fig. 1, and details of the material preparation are described in full elsewhere [4]. Its technical application lies in the development of polycrystalline silicon thin-film transistors, and so the achievement of large (e.g. 1-10/~m), defect-free crystals from the amorphous state is the practical goal. Preliminary work using cross-section electron microscopy showed that crystal nucleation initiates during the silicon deposition, resulting in
0304-3991/93/$06.00 © 1993 - Elsevier Science Publishers B.V. All rights reserved
42
R. Sinclair et al. / Direct obserL,ation of crystallization in silicon by in situ HREM
a small grain size. Subsequent Si2+s ion-implantation renders the film completely amorphous again, as verified by H R E M observation, and a further annealing treatment at about 600°C brings about the desired structure and properties [5]. It is this latter stage which we investigated by in situ H R E M , employing the heating holder, specimen preparation and 300 kV transmission electron microscope methods reviewed earlier [2]. The temperatures quoted here are nominal, being measured by a thermocouple attached to the heating holder. It is thought that the actual temperature in the field of view is about 100°C lower, because of a problem with thermal gradients. However, observations, and their kinetic rate, at a particular nominal temperature were reproducible from one specimen to another. The crystallization process is illustrated by the in situ annealing experiment shown in fig. 2. It is clear that a nucleus is formed at the amorphous silicon (a-Si)/silicon oxide (SiO 2) interface, that it grows to consume the amorphous matrix in its immediate vicinity and that a high density of microtwins (as confirmed by electron diffraction and H R E M ) and stacking faults is created. The sequence is qualitatively similar to that obtained from a series of specimens annealed in a bulk furnace and subsequently sectioned for examination, fulfilling one of the criteria we have proposed to ensure that the in situ experiment is representative of bulk behavior [2]. It should be noted that surface nucleation does not occur, and this is further confirmed from a separate bulk experiment: when SiO z is deposited on top of the
Fig. 1. Schematic cross-section view of the material used for the present experiments. Crystal nucleation occurs predominantly at the a-Si/SiO 2 interface.
Fig. 2. Sequence of images showing the growth of a silicon crystallite in amorphous silicon in a ten-minute interval at a nominal temperature of 750°C.
amorphous silicon, nucleation takes place at both upper and lower a-Si/SiO 2 interfaces. The nucleation event is clearly heterogeneous, and we can take advantage of this observation to focus in on the preferred nucleation sites, viz. the a-Si/SiO 2 interface. Although both phases have typical random H R E M image appearance, with little difference in contrast level, the interface can be located fairly easily as the SiO 2 image is quite dynamic at the elevated temperatures of interest here (e.g. 600-800°C nominal), whereas that of a-Si is static. Presumably this arises from random atomic motion in the SiO2, which might be slightly viscous at these temperatures, and is a phenomenon which is sometimes noticeable under pronounced beam heating during conventional H R E M imaging (e.g. ref. [6]). During the time taken to ramp-up to the observation temperature and to stabilize the image (a few minutes) it is very noticeable that many small crystallites have formed, approximately 1-5 nm in size and random in orientation, at the a-Si/SiO 2 interface ( H R E M imaging prior to heating shows definitively that no crystallites are present originally). The vast majority does not grow, even upon prolonged observation (e.g. up to thirty minutes at "750°C"), and occasionally they are consumed by
R. Sinclair et al. / Direct observation of crystallization in silicon by in situ HREM
the lateral expansion of crystals which have nucleated elsewhere along the interface. Such crystallites must be regarded as "embryos", the size of which must be lower than the critical nucleus radius, r*, at which the free energy of the system is lowered upon growth (i.e., when the bulk free energy change increases more rapidly than the interfacial energy). It is somewhat surprising that so many embryos are visible (e.g. one every few tens of nanometers along the interface), and probably there are more in orientations unfavorable for imaging. Perhaps this is the first direct observation of embryos prior to nucleation and so it will be interesting to see if equivalent behavior is encountered in other material transformations in the future. In our work to date, we have only recorded once on video-tape the progression of an embryo becoming a viable nucleus, Its original size was approximately 2.5 nm and after starting to grow it increased to about 5 nm after 10 s and completely filled the field-of-view (6.5 nm square) after about 20 s. This nucleus was internally twinned and it is very noticeable that the growth rate is significantly faster (about t w o - t h r e e times) in a direction parallel to the twins (fig. 3). This has been deduced before from conventional electron microscopy observations (e.g. refs. [7,8]) but is seen here directly at the atomic level. We can therefore place the value of r * for this embryo as certainly less than 5 nm and probably in the range 2 - 4 rim. The "contact" angle of the heterogeneous nucleus at the a-Si/SiO 2 interface is approximately 90 °, which is consistent with similar c-Si/SiO 2 and a-Si/SiO 2 interracial energies. Using nucleation theory [9] and the suggested free energy change for silicon crystallization [10], we can estimate the a - c Si interracial energy for the growth front as about 600 mJ m -2 (see ref. [11] for more details). This is within a factor of two of that estimated by Spaepen [310 mJ m-Z], from physical modeling studies of the interface structure [12] and values of 430 and 480 mJ m -2 given by Roorda and Sinke [13] and Tu et al. [14], respectively. Clearly further experiments will allow us to refine the experimental value to a more acceptable scientific level. It is not clear at this stage how trace impurities associated with the
43
CVD process or the ion-implantation influences the critical nucleus size, but given the consistency of the results maybe it is a small effect. The dominant (fast) growth direction is in the (112) twinning direction. Occasionally a twin is formed in an alternative orientation: this generally impedes the growth rate locally and the surrounding material grows around such "obstacles" (fig. 4). The original twins are also replicated laterally (i.e. sideways along the SiO 2 interface), as can be seen also from the low-magnification views in fig. 2. This, of course, promotes the growth rate of this particular crystallite, but since there is a finite twin boundary energy for silicon, the twin formation must be offset by some other factor (perhaps a reduction in the transformation strain energy arising from the amorphous-crystalline volume change, which is about a 3 - 6 % decrease for ion-implanted silicon [15]).
Fig. 3. A crystal embryo at the a-Si/SiO 2 interface spontaneously grows into a viable nucleus within less than a ten-second interval at "750°C '', as seen by in situ HREM. The SiO 2 interface is in the extreme lower right-hand corner in the lower figure.
44
R. Sinclair et al. / Direct obsen,ation o f crystallization in silicon by in situ H R E M
Fig. 4. Sequence of hot-stage H R E M micrographs showing the growth of a crystalline silicon grain, consuming the amorphous silicon matrix. A twinned crystal in a different orientation (circled) locally impedes the growth and the advancing crystal proceeds around it. T h e white line, indicating the twinning and the circle are in the same location in each photograph. The time sequence is approximately one minute at a nominal temperature of 750°C.
Once a nucleus is formed, it is distinctly easier experimentally to follow the advance of the growth front. There seems little question that the film crystallization is dominated by growth of the heterogeneous nuclei, with no homogeneous nucleation occurring under our experimental conditions. The advancing interface is often serrated, achieving {111} facets at each twin as shown in fig. 5. The interface can remain static for several seconds and then move forward rapidly for a short period, in the process encompassing hundreds or thousands of atoms locally into the crystal. Similar "bursts" of reaction have been reported by us in other interfacial reactions [2,16], perhaps most notably in the solid-phase epitaxial (SPE) regrowth of silicon [1] which is essentially equivalent to the present silicon growth. However, no reaction reversals were encountered in the present work.
The advance of the serrated interface tends to occur first, with outward growth in other crystallographic orientations (e.g. (100)) following. This
Fig. 5. In situ H R E M view of the serrated c-Si/a-Si interface, during a momentary arrest of the advancing interface at ,,750oc ,,.
R. Sinclair et al. / Direct observation of crystallization in silicon by in situ HREM
contrasts with the situation for S P E for which (100) growth is the most rapid [15]. It is t h o u g h t that this might arise from nucleation of interface growth ledges at the re-entrant twin interfaces [7], and so we carefully examined our recordings to capture this process. W e were not successful. Even in experiments at lower t e m p e r a t u r e s at which the growth rate is extremely slow, no distinct ledge nucleation was f o u n d nor a distinct ledge m e c h a n i s m for the growth. Rather, growth appears to p r o c e e d without the necessity of extensive ledges. Likewise when highly crystallographic (e.g. {111}) facets were observed to advance, no clearcut ledge m e c h a n i s m was seen. It could be said that our inspection has not b e e n sufficiently extensive to isolate the ledge nucleation event. However, in our experience, when a ledge growth m e c h a n i s m is present, it is especially noticeable (see ref. [2] for several examples). W e therefore suggest that while the re-entrant twin interfaces might assist the attachment of silicon atoms to the crystal, no extensive {111} plane ledge m e c h a n i s m is necessary to advance the crystalline front. This proposal is consistent with our prior observations of S P E growth, for which a ledge m e c h a n i s m was also not observed.
3. Conclusions W e believe that the significance of this work lies in the successful d e m o n s t r a t i o n that crystal nucleation and growth can be studied directly at the atomic level, by in situ H R E M . T h e system chosen (amorphous-crystalline silicon) is relatively simple, yet the process we have studied has immense practical importance. Some of the observations are possibly unique at the present time, particularly the discovery of a large n u m b e r of non-viable crystal embryos and the recordings of nucleus growth. Likewise the advancing crystal interface has b e e n followed over p r o l o n g e d periods of time. F u r t h e r work is necessary to obtain reliable values for the critical nucleus size, to determine the importance of internal twinning to the nucleation and growth, and to establish (as we indicate) that an extensive ledge m e c h a n i s m is not important to crystal growth in this system.
45
However, it is clear that our a p p r o a c h can be employed to study the initial stages of a phase transformation or reaction and that it has the capacity to reveal many, possibly unexpected, p h e n o m e n a . By this means we can test, directly and rigorously, models of nucleation and growth in solids.
Acknowledgements This work has b e e n supported by the National Science F o u n d a t i o n ( G r a n t no. D M R 8902232) and by Xerox P A R C . O n e of us (J.M.) wishes to thank the Fulbright F o u n d a t i o n for a fellowship, during the tenure of which m u c h of the experimental work was carried out.
References [1] R. Sinclair and M.A. Parker, Nature 322 (1986) 531. [2] R. Sinclair, T. Yamashita, M.A. Parker, K.B. Kim, K. Holloway and A.F. Schwartzman, Acta Cryst. A 44 (1988) 965. [3] M. Reiche and S. Hopfe, Ultramicroscopy 33 (1990) 41. [4] I.-W. Wu, A. Chiang, M. Fuse, L. Ovecoglu and T.Y. Huang, J. Appl. Phys. 65 (1989) 4036. [5] I.-W. Wu, A.G. Lewis, T.Y. Huang and A. Chiang, IEEE Elect. Dev. Lett. 10 (1989)123. [6] S. lijima, in: Proc. 47th Annual EMSA Meeting, Ed. G.W. Bailey (San Francisco Press, San Francisco, 1989) p. 126. [7] R. Drosd and J. Washburn, J. Appl. Phys. 53 (1982) 397. [8] A. Nakamura, F. Emoto, E. Fujii, Y. Uemoto, A. Yamamoto, K. Senda and O. Kano, Jpn. J. Appl. Phys. 27 (1988) L2408. [9] J.W. Christian, in: Theory of Transformations in Metals and Alloys, Part 1 (Pergamon, Oxford, 1975) p. 418. [10] E.P. Donovan, F. Spaepen, D. Turnbull, J.M. Poate and D.C. Jacobson, Appl. Phys. Lett. 42 (1983) 698. [11] A.S. Kirtikar, J. Morgiel, R. Sinclair, I.-W. Wu and A. Chiang, Proc. Mater. Res. Soc. 202 (1991) 627. [12] F. Spaepen, in: Amorphous Materials: Modeling of Structure and Properties, Ed. V. Vitek (The Metallurgical Society of AIME, New York, 1983) p. 265. [13] S. Roorda and W.C. Sinke, Appl. Surf. Sci. 36 (1989) 588. [14] K.N. Tu, J.W. Mayer and L.C. Feldman, in: Electronic Thin Film Science for Electrical Engineers and Materials Scientists (Macmillan, New York, 1992) p. 267. [15] M.A. Parker, PhD Thesis, Stanford University (1988). [16] R. Sinclair, M.A. Parker and K.B. Kim, Ultramicroscopy 23 (1987) 383.
UltramicroscopyNorth.Holland51 (1993)41-45
Direct observation of crystallization in silicon by in situ high-resolution electron microscopy R. Sinclair
a,
j. Morgiel
a,b,
m.s. Kirtikar
a,
I.-W. Wu c and A. Chiang c
a Department of Materials Science and Engineering, Stanford University, Stanford, CA 94305, USA b Institute for Metals Research, Polish Academy of Science, Krakow, Poland c Xerox Corporation, Palo Alto Research Center, Palo Alto, CA 94304, USA Received 13 November 1992
We have studied the nucleation and growth of crystalline silicon by in situ high-resolution electron microscopy. Amorphous silicon thin films, deposited onto oxidized silicon wafers by low-pressure chemical vapor deposition, were heated in the microscope in the nominal temperature range of 700-775°C. Many sub-critical crystal embryos exist at the a-Si/SiO 2 interface, very few of which develop into viable nuclei. One nucleation event was recorded successfully at lattice resolution, allowing an estimate of the critical radius as 2.5+ 1.0 nm and the a-Si/c-Si interfacial energy as 600+200 mJ m -2. The crystal growth was followed for extended periods of time. It appeared to be characterized by sporadic growth bursts rather than a continuous advance of the crystal-amorphous phase interface. We specifically sought evidence for a ledge growth mechanism but did not find such behavior.
I. Introduction One of the most difficult occurrences to investigate directly is the nucleation event of a solidstate phase transformation. This arises because of the statistical nature and the minute scale (e.g. 1-10 nm) on which nucleation takes place. However, as first reported by Sinclair and Parker [1], it is now possible to study reactions in solids at the atomic level, at controlled elevated temperatures, by in situ, high-resolution electron microscopy (HREM) [2]. Accordingly we have attempted, and succeeded in, recording such behavior, using as an example the crystallization of amorphous silicon thin films. Silicon, being a simple element which can be easily fabricated in many forms, provides an excellent means of carrying out basic scientific experiments of this type. In addition, processes such as crystal nucleation and growth are extremely important during semiconductor-based, integrated circuit manufacture, hence such work can have technological significance. Recently Reiche and Hopfe have also
shown that in situ hot-stage TEM techniques can be used effectively to study crystallization in amorphous silicon [3]. In our work, we have focussed on an atomic-level study of the ongoing interfacial processes.
2. Experimental procedure and results The system we chose for this project consists of amorphous silicon thin films, formed by lowpressure chemical vapor deposition techniques onto thermally oxidized single-crystal silicon wafers. A schematic cross-section view is shown in fig. 1, and details of the material preparation are described in full elsewhere [4]. Its technical application lies in the development of polycrystalline silicon thin-film transistors, and so the achievement of large (e.g. 1-10/~m), defect-free crystals from the amorphous state is the practical goal. Preliminary work using cross-section electron microscopy showed that crystal nucleation initiates during the silicon deposition, resulting in
0304-3991/93/$06.00 © 1993 - Elsevier Science Publishers B.V. All rights reserved
42
R. Sinclair et al. / Direct obserL,ation of crystallization in silicon by in situ HREM
a small grain size. Subsequent Si2+s ion-implantation renders the film completely amorphous again, as verified by H R E M observation, and a further annealing treatment at about 600°C brings about the desired structure and properties [5]. It is this latter stage which we investigated by in situ H R E M , employing the heating holder, specimen preparation and 300 kV transmission electron microscope methods reviewed earlier [2]. The temperatures quoted here are nominal, being measured by a thermocouple attached to the heating holder. It is thought that the actual temperature in the field of view is about 100°C lower, because of a problem with thermal gradients. However, observations, and their kinetic rate, at a particular nominal temperature were reproducible from one specimen to another. The crystallization process is illustrated by the in situ annealing experiment shown in fig. 2. It is clear that a nucleus is formed at the amorphous silicon (a-Si)/silicon oxide (SiO 2) interface, that it grows to consume the amorphous matrix in its immediate vicinity and that a high density of microtwins (as confirmed by electron diffraction and H R E M ) and stacking faults is created. The sequence is qualitatively similar to that obtained from a series of specimens annealed in a bulk furnace and subsequently sectioned for examination, fulfilling one of the criteria we have proposed to ensure that the in situ experiment is representative of bulk behavior [2]. It should be noted that surface nucleation does not occur, and this is further confirmed from a separate bulk experiment: when SiO z is deposited on top of the
Fig. 1. Schematic cross-section view of the material used for the present experiments. Crystal nucleation occurs predominantly at the a-Si/SiO 2 interface.
Fig. 2. Sequence of images showing the growth of a silicon crystallite in amorphous silicon in a ten-minute interval at a nominal temperature of 750°C.
amorphous silicon, nucleation takes place at both upper and lower a-Si/SiO 2 interfaces. The nucleation event is clearly heterogeneous, and we can take advantage of this observation to focus in on the preferred nucleation sites, viz. the a-Si/SiO 2 interface. Although both phases have typical random H R E M image appearance, with little difference in contrast level, the interface can be located fairly easily as the SiO 2 image is quite dynamic at the elevated temperatures of interest here (e.g. 600-800°C nominal), whereas that of a-Si is static. Presumably this arises from random atomic motion in the SiO2, which might be slightly viscous at these temperatures, and is a phenomenon which is sometimes noticeable under pronounced beam heating during conventional H R E M imaging (e.g. ref. [6]). During the time taken to ramp-up to the observation temperature and to stabilize the image (a few minutes) it is very noticeable that many small crystallites have formed, approximately 1-5 nm in size and random in orientation, at the a-Si/SiO 2 interface ( H R E M imaging prior to heating shows definitively that no crystallites are present originally). The vast majority does not grow, even upon prolonged observation (e.g. up to thirty minutes at "750°C"), and occasionally they are consumed by
R. Sinclair et al. / Direct observation of crystallization in silicon by in situ HREM
the lateral expansion of crystals which have nucleated elsewhere along the interface. Such crystallites must be regarded as "embryos", the size of which must be lower than the critical nucleus radius, r*, at which the free energy of the system is lowered upon growth (i.e., when the bulk free energy change increases more rapidly than the interfacial energy). It is somewhat surprising that so many embryos are visible (e.g. one every few tens of nanometers along the interface), and probably there are more in orientations unfavorable for imaging. Perhaps this is the first direct observation of embryos prior to nucleation and so it will be interesting to see if equivalent behavior is encountered in other material transformations in the future. In our work to date, we have only recorded once on video-tape the progression of an embryo becoming a viable nucleus, Its original size was approximately 2.5 nm and after starting to grow it increased to about 5 nm after 10 s and completely filled the field-of-view (6.5 nm square) after about 20 s. This nucleus was internally twinned and it is very noticeable that the growth rate is significantly faster (about t w o - t h r e e times) in a direction parallel to the twins (fig. 3). This has been deduced before from conventional electron microscopy observations (e.g. refs. [7,8]) but is seen here directly at the atomic level. We can therefore place the value of r * for this embryo as certainly less than 5 nm and probably in the range 2 - 4 rim. The "contact" angle of the heterogeneous nucleus at the a-Si/SiO 2 interface is approximately 90 °, which is consistent with similar c-Si/SiO 2 and a-Si/SiO 2 interracial energies. Using nucleation theory [9] and the suggested free energy change for silicon crystallization [10], we can estimate the a - c Si interracial energy for the growth front as about 600 mJ m -2 (see ref. [11] for more details). This is within a factor of two of that estimated by Spaepen [310 mJ m-Z], from physical modeling studies of the interface structure [12] and values of 430 and 480 mJ m -2 given by Roorda and Sinke [13] and Tu et al. [14], respectively. Clearly further experiments will allow us to refine the experimental value to a more acceptable scientific level. It is not clear at this stage how trace impurities associated with the
43
CVD process or the ion-implantation influences the critical nucleus size, but given the consistency of the results maybe it is a small effect. The dominant (fast) growth direction is in the (112) twinning direction. Occasionally a twin is formed in an alternative orientation: this generally impedes the growth rate locally and the surrounding material grows around such "obstacles" (fig. 4). The original twins are also replicated laterally (i.e. sideways along the SiO 2 interface), as can be seen also from the low-magnification views in fig. 2. This, of course, promotes the growth rate of this particular crystallite, but since there is a finite twin boundary energy for silicon, the twin formation must be offset by some other factor (perhaps a reduction in the transformation strain energy arising from the amorphous-crystalline volume change, which is about a 3 - 6 % decrease for ion-implanted silicon [15]).
Fig. 3. A crystal embryo at the a-Si/SiO 2 interface spontaneously grows into a viable nucleus within less than a ten-second interval at "750°C '', as seen by in situ HREM. The SiO 2 interface is in the extreme lower right-hand corner in the lower figure.
44
R. Sinclair et al. / Direct obsen,ation o f crystallization in silicon by in situ H R E M
Fig. 4. Sequence of hot-stage H R E M micrographs showing the growth of a crystalline silicon grain, consuming the amorphous silicon matrix. A twinned crystal in a different orientation (circled) locally impedes the growth and the advancing crystal proceeds around it. T h e white line, indicating the twinning and the circle are in the same location in each photograph. The time sequence is approximately one minute at a nominal temperature of 750°C.
Once a nucleus is formed, it is distinctly easier experimentally to follow the advance of the growth front. There seems little question that the film crystallization is dominated by growth of the heterogeneous nuclei, with no homogeneous nucleation occurring under our experimental conditions. The advancing interface is often serrated, achieving {111} facets at each twin as shown in fig. 5. The interface can remain static for several seconds and then move forward rapidly for a short period, in the process encompassing hundreds or thousands of atoms locally into the crystal. Similar "bursts" of reaction have been reported by us in other interfacial reactions [2,16], perhaps most notably in the solid-phase epitaxial (SPE) regrowth of silicon [1] which is essentially equivalent to the present silicon growth. However, no reaction reversals were encountered in the present work.
The advance of the serrated interface tends to occur first, with outward growth in other crystallographic orientations (e.g. (100)) following. This
Fig. 5. In situ H R E M view of the serrated c-Si/a-Si interface, during a momentary arrest of the advancing interface at ,,750oc ,,.
R. Sinclair et al. / Direct observation of crystallization in silicon by in situ HREM
contrasts with the situation for S P E for which (100) growth is the most rapid [15]. It is t h o u g h t that this might arise from nucleation of interface growth ledges at the re-entrant twin interfaces [7], and so we carefully examined our recordings to capture this process. W e were not successful. Even in experiments at lower t e m p e r a t u r e s at which the growth rate is extremely slow, no distinct ledge nucleation was f o u n d nor a distinct ledge m e c h a n i s m for the growth. Rather, growth appears to p r o c e e d without the necessity of extensive ledges. Likewise when highly crystallographic (e.g. {111}) facets were observed to advance, no clearcut ledge m e c h a n i s m was seen. It could be said that our inspection has not b e e n sufficiently extensive to isolate the ledge nucleation event. However, in our experience, when a ledge growth m e c h a n i s m is present, it is especially noticeable (see ref. [2] for several examples). W e therefore suggest that while the re-entrant twin interfaces might assist the attachment of silicon atoms to the crystal, no extensive {111} plane ledge m e c h a n i s m is necessary to advance the crystalline front. This proposal is consistent with our prior observations of S P E growth, for which a ledge m e c h a n i s m was also not observed.
3. Conclusions W e believe that the significance of this work lies in the successful d e m o n s t r a t i o n that crystal nucleation and growth can be studied directly at the atomic level, by in situ H R E M . T h e system chosen (amorphous-crystalline silicon) is relatively simple, yet the process we have studied has immense practical importance. Some of the observations are possibly unique at the present time, particularly the discovery of a large n u m b e r of non-viable crystal embryos and the recordings of nucleus growth. Likewise the advancing crystal interface has b e e n followed over p r o l o n g e d periods of time. F u r t h e r work is necessary to obtain reliable values for the critical nucleus size, to determine the importance of internal twinning to the nucleation and growth, and to establish (as we indicate) that an extensive ledge m e c h a n i s m is not important to crystal growth in this system.
45
However, it is clear that our a p p r o a c h can be employed to study the initial stages of a phase transformation or reaction and that it has the capacity to reveal many, possibly unexpected, p h e n o m e n a . By this means we can test, directly and rigorously, models of nucleation and growth in solids.
Acknowledgements This work has b e e n supported by the National Science F o u n d a t i o n ( G r a n t no. D M R 8902232) and by Xerox P A R C . O n e of us (J.M.) wishes to thank the Fulbright F o u n d a t i o n for a fellowship, during the tenure of which m u c h of the experimental work was carried out.
References [1] R. Sinclair and M.A. Parker, Nature 322 (1986) 531. [2] R. Sinclair, T. Yamashita, M.A. Parker, K.B. Kim, K. Holloway and A.F. Schwartzman, Acta Cryst. A 44 (1988) 965. [3] M. Reiche and S. Hopfe, Ultramicroscopy 33 (1990) 41. [4] I.-W. Wu, A. Chiang, M. Fuse, L. Ovecoglu and T.Y. Huang, J. Appl. Phys. 65 (1989) 4036. [5] I.-W. Wu, A.G. Lewis, T.Y. Huang and A. Chiang, IEEE Elect. Dev. Lett. 10 (1989)123. [6] S. lijima, in: Proc. 47th Annual EMSA Meeting, Ed. G.W. Bailey (San Francisco Press, San Francisco, 1989) p. 126. [7] R. Drosd and J. Washburn, J. Appl. Phys. 53 (1982) 397. [8] A. Nakamura, F. Emoto, E. Fujii, Y. Uemoto, A. Yamamoto, K. Senda and O. Kano, Jpn. J. Appl. Phys. 27 (1988) L2408. [9] J.W. Christian, in: Theory of Transformations in Metals and Alloys, Part 1 (Pergamon, Oxford, 1975) p. 418. [10] E.P. Donovan, F. Spaepen, D. Turnbull, J.M. Poate and D.C. Jacobson, Appl. Phys. Lett. 42 (1983) 698. [11] A.S. Kirtikar, J. Morgiel, R. Sinclair, I.-W. Wu and A. Chiang, Proc. Mater. Res. Soc. 202 (1991) 627. [12] F. Spaepen, in: Amorphous Materials: Modeling of Structure and Properties, Ed. V. Vitek (The Metallurgical Society of AIME, New York, 1983) p. 265. [13] S. Roorda and W.C. Sinke, Appl. Surf. Sci. 36 (1989) 588. [14] K.N. Tu, J.W. Mayer and L.C. Feldman, in: Electronic Thin Film Science for Electrical Engineers and Materials Scientists (Macmillan, New York, 1992) p. 267. [15] M.A. Parker, PhD Thesis, Stanford University (1988). [16] R. Sinclair, M.A. Parker and K.B. Kim, Ultramicroscopy 23 (1987) 383.