- Email: [email protected]
Electron-beam enhanced growth of CVD-deposited silicon on alumina
Journal of Crystal Growth 59 (1982) 485-- 498 North-Holland Publishing Company
485
ELECTRON-BEAM ENHANCED GROWTH OF CVD-DEPOSITED SILICON ON ALUMINA K. HEINEMANN and T. OSAKA
*
Stanford/NASA Joint Institute for Surface and Microstructure Research, Department of Materials Science and Engineering. Stanford University, Stanford, California 94305, USA Received 30 December l98l~ manuscript received in final form 5 May 1982
Silicon was deposited onto polycrystalline alumina and single crystal sapphire substrates by low-pressure chemical vapor decomposition (CVD) of silane. The experiments were conducted in a custom UHV specimen/reaction chamber of a transmission electron microscope and recorded in-situ using an image intensifier—TV chain. Sapphire substrate areas were obtained by in-Situ electron beam flash-heating of electron-transparent alumina films prepared by anodic oxidation of aluminum foil. Silicon was observed to grow epitaxially on those sapphire substrate areas that were impinged by the electron beam during the CVD process. while in all other recrystallized substrate areas no silicon was deposited. Silicon also deposited the original amorphous/polycrystalline alumina areas that had not been heated by or exposed to the electron beam. On the other hand, no deposit was noted in alumina areas heated by the electron beam to over 1000°C but below the sapphire recrystallization temperature. We speculate that, once the substrate surface is clean, no silicon will deposit, independent of the state of crystallinity; only where the electron beam locally stimulates the decomposition of silane, deposition of silicon on the alumina substrate is observed.
I. Introduction Compared to the extensive information and cxperience available on the growth of silicon on various substrates by chemical vapor decomposition (CVD) of silane, only little has been reported about the nucleation and early growth stages of this process, and about the effect of the substrate surface orientation and cleanliness, as well as a concomitant electron beam, on the nucleation and growth characteristic, We have adapted the in-situ transmission dcctron microscopy (TEM) technique [l—3] to investigate the nucleation and early growth of silicon on sapphire substrates by low-pressure CVD. Abrahams et al. [4,5] studied the early growth of silicon on (0112)-oriented sapphire surfaces by TEM, preparing electron-transparent films by ion milling after conventional reactor growth of the silicon deposit to average thicknesses ranging from 3 to 23 nm. They obtained information about tOn leave from Waseda University, Tokyo, Japan.
0022-0248/82/0000--0000/$02.75
©
epitaxial relationships and twinning in the island deposits, but their results do not allow conclusions with regard to the influence on the nucleation and growth behavior of the substrate surface orientation, the substrate cleanliness, or a concomitant electron beam. Most earlier work has been done in CVD systems operating near atmospheric pressure. Our approach necessitates working at pressures in the 1 to l0~ Torr range. Reactors working at reduced pressures have recently been introduced [6,7]. Typical reported pressures ranged from 100 to 102 Torr [8,9]. 2. Experimental A cross-sectional schematic drawing of the microscope with reaction chamber is shown in fig. 1. We have modified a Siemens Elmiskop 101 highresolution TEM by inserting a bakeable, metalsealed chamber above the objective lens. The chamber is differentially pumped by a He cryopump. The vibrations caused by the expander
1982 North-Holland
486
K Ileinenion,i, 11 Osaka
ELECI
(Irowth i~I~,1 /).JE’p(EEItESI
si/li
~oi on (J/ilfrfriJnil
/
HON GUN
VA~Er~~/~
/O~A
HAMBER
¶f,
QUARTZ
SPECIMEI~
~
4 CRYSTAL
EXIT MEMBRANE
RELLOW~~~~
~r~IN1EHMEDIATE
~
LENS
r -
EXPANDER
PROJECTIVE
L -
J
LENS
~
EM
10
(i
PRESSURE, TORE
Fig. 2. Diff~ientiaI pressure characicristic between chamber and microscope; electron-beam entrance and exit apertures closed (curve A) and open (1 mm diameter openings, curve B).
~
/
B
/
VIEWING SECTION
_________________
II~II __________
~
SCREEN CAMERA 5,14GB INTENSIFIER
LIGHr0PTIcS TV CAMERA - ~-‘
Fig. 1. Schematic cross-section of TEM with VHV chamber for in-situ experimentation,
module of the pump have been successfully filtered out by a stainless steel bellows arrangement be-
and exit apertures being tn place (soltd line) and I . mm DIA aperture openings (dash line). Silane is admitted through a jet located 10 mm above the specimen and directed toward it. The jet arrangement consists of a 10 cm long stainless steel tithe with 0.5 mm inner diameter. During standard silane pyrolysis experiments, a chamber pressure of 2 z< 10 Torr is maintained. (Tonsidcling the pumping speed of the cryopump as conductance-limited by its cross-sectional entrance
area and considering all the geometries involved, one can estimate the silane pressure P
The electron beam enters and exits the UHV chamber via small x, ~‘-adjustable apertures. the lower one being mounted immediately above the objective aperture slider. A chamber pressure in the low io~Torr range is reached without baking within 18 h after pumpdown. Improvement is cxperienced by baking and — at a loss of image quality — by replacing the chamber entrance and
0 at the exit of the jet to be P1 0.03 Torr [101. TIns agrees with the pressure measurement of 5 Torr in the silane supply tube between the needle valve and the jet. In agreetnent with a paper by Madey [11] we then estimate the pressure P~at the site of the specimen to be P., 0.147P11 4.5 X 10 Torr. The specimen holder was designed for hightemperature in-situ CVD experiments. Temperatures well above 1000°Cneeded to he achieved at negligible outgassing rates and without having parts hotter than the specimen itself in the vicinity of the specimen. The holder is essentially a directgrid heater, using a 0.25 mm diameter nichrome wire, spotwelded around the periphery of a 200-
exit apertures with thin film membranes. The differential pressure characteristic between the
mesh Mo-specimen grid, as the main conductor. The specimen temperature was measured in-situ
TEM—column vacuum and the chamber vacuum is shown in fig. 2 for the case of thin film entrance
with an IRCON Radiation Thermometer. calibrated for the open-area-ratio of the grid and
tween expander and pump housing. Some residual vibration influence is noticeable in the electron image only when operating at highest resolution.
However, the cryopump can be shut off, without loss of pumping speed, for short durations such as required for taking high-resolution micrographs.
K. Heinemann, T Osaka
/
Growth of C VD-deposited silicon on alumina
the emissivity of molybdenum. This hot stage was tested up to 1300°C. However, most silane CVD experiments were conducted in the 900—1000°C range. A heating current of 4.5 A is required for 1000°C, but no uncorrectable astigmatism problems were encountered. The temperature equilibrium is obtained almost instantaneously, and thermal drift subsides to a level commensurate with video-image recording within one minute of a temperature change. The alumina substrates were prepared by anodic oxidation at 45 V of 0.025 mm thick aluminum foil and by subsequent dissolution of the unoxidized aluminum in a saturated mercury chloride solution. The resulting approximately 100 nm thick amorphous alumina films were washed by repeated transfer to a dish with distilled water and finally transferred to the grid heater,
487
After establishing the desired substrate temperature with the grid heater, which transformed the entire alumina film into a fine-polycrystalline state, some areas of 10 to 100 ~tm in size were recrystallized to sapphire by electron-beam flash heating [12]. Immediately following the preparation of a few satisfactory sapphire regions, an appropriate field of view usually one containing portions of several grains with different crystallographic surfaces was selected, and the silane was admitted to start the CVD process. —
—
3. Results Virtually immediately upon admitting silane, strong local changes in diffraction contrast occur in the sapphire substrate, such as would be ex-
200nm L A Fig. 3. Typical SOS growth sequence during low-pressure, in-situ CVD: left 6 mm, middle 16 mm, right 26 mm. after admission of silane.
K. l-lcin&’niann. I. Osaka
488
&romcth of (
pected if the excitation error changes due to bending of the substrate. These changes subside after a few seconds. About five minutes later, the first silicon particles can be distinguished in the dcctron image. A typical sequence of micrographs taken during a low-pressure silane CVD experiment is shown in fig. 3. Typical growth rates were 1-3 nm/mm. Neither the lateral growth rate nor the initial island number density seemed to he strongly dependent on the growth temperature between 900 and 1000°C. The island nucleation number densities were always very high. and coalescence started to
~D-de’posited s,luon on aluniinii
occur at a very early stage of growth, as soon as individual islands could be detected in the electron image. However, the number densities and average sizes of the particles varied markedly from grain to grain. Some of these differences disappeared with increasing deposit thickness. Variation of the smiane pressure during the CVD experiments by one order of magnitude also did not affect the observed nucleation and growth behavior. Fig. 4 shows typical electron diffraction patterns of silicon deposits on vartoLts sapphire surfaces and on polycrystalline alumina. Whereas the highest degree of epitaxy was generall~ noted
I
1102
220
111
1120
220 S
1102
111 111
311 Fig. 4. Typical selected—area diffraction patterns of sapphire ssoh v,mrii~u’.c\pIISed surfaces mu 1’.) amid pl~~ i s’.;,mlllne CVD-deposiied ~mlmcumi. (Hexagonal indices refer to sapphire ~uhstr,mteo
.ilutniu.i
I
di
iii
K. Heinemann, T Osaka
/ Growth of CVD-deposited silicon on
alumina
489
on the (1012) planes (“r” planes) of sapphire (fig. 4a), no consistent temperature dependence of the
with silicon, as is evidenced in fig. 5a. However, we never found any evidence of “graphoepitaxial”
degree of epitaxy was observed in the 900—1000°C range.
alignment of the silicon deposit with respect to such distinct topographical features, as is demon-
Polycrystalline alumina areas that had not been
strated by the silicon reflections in the diffraction
subjected to electron-beam heating were heavily
pattern in fig. Sd which is representative of the
decorated with silicon during silane CVD, as is demonstrated in fig. 5. The silicon particle density and area coverage was considerably higher than in the recrystallized sapphire areas. The oxide films
area of fig. 5a and does not exhibit any azimuthal preference. Once silicon had been grown on polycrystalline alumina, it was possible to re-evaporate this de-
often contained unidirectional topographical features (stemming from the fabrication process of the original aluminum foil), which were decorated
posit by localized heating with the electron beam. This evaporation occurred at a temperature below the recrystallization temperature for sapphire (yet
311 iii 220 SILICON
.lpm -
..
Fig. 5. In-situ grown silicon on alumina: (a) decoration of unidirectional surface relief structures in polvcrvstalline substrate region; (b) border-region between polycrystalline alumina and sapphire; (c) magnification of (a); and (d) diffraction pattern of (a).
490
A ll,’incman,i, T ( 15111,0
P
J
I~
ti
( ,os ii: ,,/ I i l).h pmtI J
1i
‘in ‘n i/uniina
si
~
~.
~,,
~.
;~
..4~
I
t
I iii 220 SILICON
311
S
I ig. 0. Localmied re—eNaporatmon
of
silicon from pol\crvstalline alumina film during ~—heam flash-heating Ia) and diffraction patterns
of original area with silicon deposit (b) and of silicon re-evaporated area (c).
above 1000°C at which the primary silicon (TVD growth had taken place). An example is given in fig. 6, showing the typical random-orientation diffraction patterns of silicon and alumina in the
as-deposited areas (fig. 6b). whereas the diffraction pattern taken from an electron-beam heated area (such as the one shown in the center of fig. 6a) evidences only diffraction intensities from alumina and not from silicon (fig. 6c).
The growth of silicon on recrystallized sapphire
areas by low-pressure CVD of silane was found to be strongly enhanced in the presence of the imaging electron beam. Fig. 7a shows the typical result obtained on an area that was exposed to the electron beam (100 keV, approximately 1 X 10
A/cm2) during the CVD reaction. Fig. 7b shows an adjacent recrystallized sapphire area which had not been exposed to the electron beam but had been subjected otherwise to the identical treatment. No silicon deposit was found in this or in
K. lIe,nemann, T. Osaka
/
491
Growth of CVD-deposited silicon on alumina
211m ___
Al 2 0.0112 .~
Si 111
IS—
A1 203-01T2 3 A cm2
Fig. 7. Two electron beam adjacent duringe-beam deposition; recrystallized (b) no c-beam sapphire exposure areas after during low-pressure deposition.silane CVD experiment: (a) area exposed to i0
any recrystallized sapphire area that had not been irradiated by the imaging electron beam during the CVD process. Further evidence of the influence of the electron beam during the growth of silicon by lowpressure CVD of silane is presented in fig. 8. Immediately following the recrystallization to sap-
phire, the electron beam (1 X l0~ A/cm2) was confined to approximately 2 ~tm diameter and maintained in a fixed position within the recrystallized region during the entire subsequent CVD experiment. Fig. 8a presents an overview of the recrystallized area and its surroundings, also demonstrating the silicon re-evaporation zone around
492
A. Ileini’mann. T. Osaka
(,roscth 11/ I. I l)—leplislll’il
the sapphire region of this spectnien that had been Si-deposited in a previous experiment. Fig. 8h depicts a section of fig. 8a at a higher magnification, showing the beam-enhanced silicon growth area. Figs. 8e and 8f show successively increased enlargements of the respective areas outlined in the previous micrographs. Fig. 8h is a dark field image of exactly the area shown in fig. 8f, showing moire fringes between A1 205{006} and Si{220} type planes. As is demonstrated with fig. 8d - a Si{ 111) selected-zone dark field image of the exact area of fig. 8c no silicon image intensities were detected outside the irradiated region, i.e., no silicon deposit was registered within the silicon re-evaporation zone. Fig. 9 illustrates these findings schemati-
—
~—
.5//Icon
in?
iluoiuoi
cally. If I and i~ are the electron beam intensity during microscopy and c-beam recrystallization. respectively, there appear to he three critical dcctron irradiation intensities: for I~> 1r alumina is recrystallized, for I~> / no silicon is deposited unless a (low-intensity) c-beam is present. and for I> l~enhanced growth of silicon is observed.
4. Discussion 4.1. Nucleation and ear/v growth An initially very high nuclei number density. followed by coalescence already during the very early growth is consistent with earlier in-situ
~ff4’W
K. Hm’mnemann, T. Osaka
/
_
I
lOOnm A
~
/a
~‘Sm-111
I
A1
203
-
493
Growth of CVD.depo.c:ieil silicon on alumina
0
_
lOnm
5,.
0006
*
Fig. 8. Electron-beam enhanced growth of SOS: (a) overview of recrystallized sapphire area produced by e-beam flash-heating of polycrystalline alumina film previously deposited with silicon, showing silicon re-evaporation zone in polycrystallmne alumina region; (b), (c), (e), (f) enlargements of respective outlined regions within c-beam exposed growth area; (d) Si{ Ill) selected-zone dark field (SZDF) micrograph of area shown in (c): (g) SAD pattern of area shown in (e); and (h) central dark field image of area shown in (f), using Si 220 plus sapphire 0006 type reflections.
7
‘N,
/‘
/ Ir
DURING
/
nucleation studies of gold on sapphire [13]. Abrahams et al. [4] also reported that the highest partide number densities occurred at their earliest 2partide number and 20 nm diameter), growth stage density (approximately 2 X mean 10~ cm~ and they suggested that coalescence had already occurred at that stage. Our experiments have in-
‘\ ‘\ ~,
(.‘BEAM T~LA~Ho~oA;~NG E-BEAM IRRADIAT
I
‘
dicated that the detection limit of the particles was le
-
5,
,~
~
SI
Fig. 9. Schematic cross-section of condition shown in fig. 8a, POLYCRYSTALLINE ALUMINA
SAPPH IRE (ZONE
~
EXPOSED TO
25O~C(
POLYCRYSTALLIN E ALUMINA
indicating approximate Gaussian distributions of c-beam inten. . . sitmes during recrystallization and observation (dash Curves) and . .
.
. .
critical beam intensities (dotted lines): ‘r for recrystallization, 1~for removal of contaminants, and J~for enhanced growth.
494
K. Heinemann, T. Osaka
Growth of CVD.deposits’i/ si/non on alumina
j
z 12
~
-
40
101
‘:1010
-
.60 -
—
— -5
________ 10 15 — 20 GROWTH
25
Fig. 11. GROWTH F: Mean particle TIME, MIN sapphire substrate grainsdiameter shown inversus fig. 3.growth time for three
TIME, MIN
Fig. 10. Silicon particle number density versus growth time for the three sapphire substrate grains shown in fig. 3 (top right, lower right, and left grains),
particle can be. neglected, i.e. when the particles . . A
are small [141,fig. 11 indicates a dependence d where k is approximately 2. This high value of k may actually suggest that the primary tendency during early growth is not that the coalescing particles sinter into some sort of hemispherically
governed by their number density rather than their size. If one assumes conservatively a detectability limit of average particle-to-particle distances of 2
shaped particle, but predominantly remain flat in that they form “rafts” of the original small partides. This would, in turn, be consistent with
nm and 10H atoms per cm2 in a monolayer, the particle number density versus deposition time characteristic must have peaked at density values above 2. This stipulation is confirmed in cm fig. 10 10t3 which is representative of our experiments at 1000°Csubstrate temperature (fig. 3) and represents a quantitative evaluation for the earliest growth range that we could measure with a rcasonable degree of confidence. (Extrapolation of these results to t~0 would lead to the 10’s cm 2 range.) As is apparent by the negative slope of the curves, this range is already dominated by coalescence. The development of the mean particle size
recent results by Carey [15] who suggests that the planar defects in SOS films originate during the coalescence of silicon islands in the early growth phase of the films. If complete sintering to a new single crystal particle does not occur, then the formation of twins and other planar defects is a likely alternate mechanism of accomodation. However, various other nucleation and growth results have shown that three-dimensional particle growth does not necessarily imply that the new particle is a single crystal (see. e.g.. refs. [16—181).
(fig. 11) is also indicative of coalescence being a domineering mechanism in the early particle growth phase. Whereas a dependence of the mean particle diameter d on the growth time I of d -~ 17 is expected for “regular” three-dimensional partide growth, when the adatom diffusivity on the substrate surface is high, when the capture rate of adatoms is proportional to the circumference of the particle, and when the contribution of the
experiments of gold on amorphized graphite substrates [19]. However, in these experiments the substrate temperature was much lower, and even then complete sintering usually occurred some time later, thus allowing interpretation of the results as a two-step process of coalescence, leading via “raft” formation as precursor state to complete sintering. An attempt to explain the underlying reason for the present results should be consistent with our
vapor beam directly impinging on the growing
specific observations that:
Coalescence without immediate sintering into
particles of essentially three-dimensional shape was found earlier in in-situ nucleation and growth
K. Heinemann, T Osaka
/
Growth of CVD-deposited silicon on alumina
495
(i) A value of k 2 was found not only during the very early stages but up to raft diameters well above 100 nm. (ii) Coalescence into rafts involved particle mo-
upper right, and left grains shown in fig. 3, respectively). Whereby we were not able to determine the exact substrate surface orientation for this exampie, it is clear that in the depicted growth range the
bility (that appeared very similar to the type of mobility seen in earlier experiments [19,13,20]),
particles on the upper right grain (dash curves in figs. 10 and 11) grew in the depicted time interval
(iii) Relatively large areas of uniform diffraclion contrast exist within individual rafts at a later stage of growth, such as shown in fig. 3c, indicating that some re-formation of the rafts into single crystals or less faulted particles does occur, hut the resulting particles essentially reproduce the flatness and substrate area coverage of the original rafts. Whether the low self-diffusivity of silicon (even at the high substrate temperature used) or the relatively low ratio of free surface energies of overgrowth and substrate (~w 1 for SOS, >> 1 for example for Au/NaCI [21,22] which is well documented to immediately sinter into three-dimensional particles upon coalescence [18,23,24]), is the primary cause for the raft formation is difficult to determine at this time. The latter cause would be most far-reaching, in that it would essentially imply that layer-type growth is favored in the initial stages of SOS epitaxy. Our observations of appreciable diffraction contrast changes immediately following silane admittance into the in-situ reaction chamber, i.e. at a much earlier stage of growth than depicted in fig. 3a, well before appearance of the silicon islands, would be consistent with such a model of initial layer growth. Initial layer coverage of silicon could exert strains on the substrate film, causing bending and concomitant diffraction contrast changes. It should be noted, however, that our present experimental evidence is not strong enough to warrant more than a speculation of
markedly faster than the particles on the other two grains, indicating that, all other parameters and
layer-type growth in the initial stages of growth of silicon on sapphire. More work would need to be
done, possibly requiring such elaborate techniques as in-situ micro-area Auger spectroscopy to clarify this important point. Irrespective of these speculations concerning the initial growth of silicon nuclei on c-beam recrystallized sapphire during low-pressure CVD, fig. 3 and, more clearly, figs. 10 and 11 exhibit differences in the decoration results obtained on three different substrate grains (the solid, dash, and dash-dot curves in figs. 10 and 11 refer to the lower right,
conditions being the same, the substrate surface
structure does have a marked influence on the nucleation and growth characteristic of silicon on sapphire. The observation that neither the silane pressure nor the substrate temperature markedly affected the growth rate or the nucleation number density in our experiments is generally unusual in the epitaxial growth of islands and thin films, where one would expect a decrease of the particle number density and the growth rate, as well as usually an improvement in epitaxy, with a decrease in supersaturation. A departure from this rule has been reported earlier for SOS in that for yet not fully understood reasons an optimum substrate temperature (of 960°C,uncalibrated for the emissivity of sapphire as is customary in the semiconductor industry) exists, beyond which the degree of epitaxy decreases again and the number of planar defects increases [15]. The range of substrate temperatures covered in our studies, if the emissivity calibration is considered, falls just below the 960°C mark, and most of our results should, therefore, fall within the range of supersaturations where “normal” growth behavior could be expected. Since the very high initial particle number densities observed in our studies cannot be explained on the basis of high supersaturation conditions, a —
—
conclusion which is emphasized by the fact that
Abrahams et al. obtained quite comparable partide density results at by three orders of magnitude higher silane pressures, and since substrate surface contamination can be excluded as primary reason in our experiments (on re-crystallized sapphire) due to the low chamber pressure (108 Torr) and high substrate temperatures maintained after recrystallization, we conclude that the imaging electron beam has been of primary influence. This conclusion is certainly confirmed in our oh-
490
A. Ils’,ncnia,in.
7T Osaka
Grsniih of I. I ‘1) ilepss,siIci/ .5//us/n in s1/ioiiiiisi
servations illustrated with figs. 7 and 8. Our results (fig. 4) essentially confirm that the degree of epitaxv is highest on the “r” plane (fig. 4a) which is used virtually exclusively tn SOS technology. However, twinning and faulting are also most pronounced in this orientation (most of the diffraction intensities in figs. 4h and 4c are due to random particle orientation and to contrihutions from neighboring not fully recrystallized alumina regions). This seems to be consistent with the high lattice mismatch in these two orientations (figs. 4b and 4c), e.g., for Si{ Ill } H A 203[ 0001 }. where the lattice mismatch is 24%. At such htgh values of mismatch the strain to accomodate the geometrical misfit would he so high that the degree of direct lattice match ceases to play an important
role, and effects caused by the mismatch, such as planar defects and microtwinning. would become less prominent, 4.2. Electron-beam enhanced nucleation and growth
The in-situ growth of silicon was strongly enhanced during electron beam irradiation. It is tinlikely that the striking magnitude of this effect, evidenced in fig. 7 for (lie case of two neighboring recrystallized areas that had been subject to idetitical conditions except for the beam exposure during the CVD process, can he explained as due to defects introduced in the substrate by the imaging electron beam. Since the electron irradiation intensity was of the order of 10 ~ A cm 2 impingement on the substrate surface occurred at a rate of a few electrons per atom site per second (6 ‘>< 1015 cm 2 ~ I) or some 1018 electrons during a typical experiment of some 3 mm duration. Radiation damage tests done on sapphire indicate that during several hours of 100 keV electron irradiation at 10 A cm 2 of 100 nm thick sapphire films no visible damage occurs [25].This computes to some 1022 electrons cm 2 This assumption is very conservative also in the light of Howitt and Mitchell’s recent results [26] who obtained damage in sapphire upon irradiation of several minutes duration at 4 orders of magnitude higher electron current densities and at energies above the threshold for radiation damage, whereas our work was done below this threshold. If we (conservatively) assume
m~displacements per cni2 would that a density of i0 be sufficient to he visible in the electron image (such as by appearance of phase contrast charactermstic of amorphous or amorphized films [27.1 VI). this gives a displacement probability of ‘~ 10 per primary incident electron ativwhere in the film. The probability for this damage to occur near the surface (saY, within the top 10 nih depth from the surface). where its effect could he the creation of a site for preferred nttcleation. is even lower, probably < 10 ~. Thus, during ;m typical experiment, the total impinging electron flux would have the capability to create less than l0’~~ 1(1 — l0~° sites for preferred nucleation per cm2. which is about 3 orders of magnitude less than the numher of preferred sites one should expect if radialion damage in the substrate is argued as the major factor in determining the c-beam enhanced growth reported here. One should also note that in earlier work on in—situ nucleation and growth of metals on recrystallized sapphire surfaces. tio effect of the electron beam was observed during very comparable electron irradiation conditions. These studies were done under standard PVD condittons and involved inert deposit materials such as gold. as
well as Ag, Pd. and Fe [13.14.281. We therefore. with due caution, rule out c-beam induced radiatioti damage as primary cause for the observed c-beam enhanced growth of silicon. One might then he tempted to explain this effect on the basis of an interaction of the electron beam with the S1H 4 flux. such as c-beam stmrntmIated decomposition of silane. This mechanism would essentially postulate that the electron beam intensity is rate limiting during the growth process of the silicon islands. Although the relative insensi-
tivity of our results with changes in silane pressure and in substrate temperature, as well as the apparent fall-off of average particle size toward the edges of the irradiated area (fig. 8c). would be consistent with this explanation. two other ohservations cannot be as readily explained~~. (i) the apparent heavy silicon depositton on the original. not recrystallized and not c-beam heated alumina substrate areas, and (ii) the high nucleation numher densities. The heavy silicon depostt on the original, not recrystallized and not annealed (“asreceived”) alumina surfaces could conceivably he
K. Hei,temann, T. Osaka
/
Growth of C VD-deposited silicon on alumina
497
associated with surface contamination. Considering the circumstance that heating the substrate to 1000°Cfor several hours apparently does not af-
ence of an electron beam on areas free of contaminants re-establishes growth at essentially the same rate that is experienced on contaminated areas
fect the silicon nucleation and growth behavior (i.e., presumably does not remove the contami-
without an electron beam. Further studies are needed to substantiate the role of substrate surface contaminants in the low-
nants), but short-duration heating to a temperature somewhat above 1000°C but below the recrystallization temperature of sapphire (about
pressure CVD process, to quantify the growth-rate limiting influence of the electron beam intensity
I 250°(’) does have that effect, may possibly point
for preferred growth, and to investigate the specu-
to hydrocarbon contaminants or carbon which are known to be removed only during very high-ternperature annealing well above 1000°C[29,30]. The critical electron irradiation intensity I~(fig. 9) would then be the intensity required for removal of contaminants. There being not really any doubt that the as-received alumina substrates are contaminated, the speculated high degree of cleanliness after recrystallization should further he investigated. Furthermore, comparison with similar experiments on bulk sapphire surfaces would be needed to connect with results such as reported by Blanc and Abrahams [5] who claim that their nucleation results are not influenced by contamination.
lated layer growth on clean sapphire substrate areas.
Acknowledgements We thank Dr. H. Poppa of NASA—Ames for stimulating discussions and suggestions during the course of this work, and we appreciate Dr. R. Schindler’s help during the construction and early testing phase of the in-situ chamber. Dr. J. Blanc of RCA—David Sarnoff Research Center in Princeton, NJ, offered constructive criticism and helpful suggestions to this work, especially to the discussion on early growth. We gratefully acknowledge his contribution. Support by NASA (Grant NCC2-ll I) is gratefully acknowledged.
5. Conclusions The nucleation and growth of silicon on sapphire by low-pressure CVD was performed inside a custom, controlled-vacuum specimen/reaction chamber of a high-resolution TEM. The initial nuclei number density is very high, probably > 10 cm2, and coalescence sets in at a very early stage of growth, possibly indicating an initial layer-type growth mechanism of the silicon deposit. Silicon grows highly preferentially on those substrate areas that are exposed to the imaging electron beam. This is believed to be due to electron-beam stimulated decomposition of silane. The observation that silicon also grows on polycrystalline alumina areas that have not been subjected to heat treatment above 1000°C indicates that surface contaminants have a major role in the nucleation and growth of silicon on alumina substrates of various orientations and crystallographic phases. When contaminants are present, silicon grows heavily; without contaminants, growth is inhibited. Pres-
References [I] RD. Moorhead and H. Poppa, in: Proc. 27th EMSA Meeting (Claitor’s, 1969) p. 116. [21 K. Heinemann, D.B Rao and D L. Douglass, Oxidation Metals 9 (1975) 379. 131 H. Poppa and K. 1-leinemann, Optik 56(1980)183. [41 MS. Abrahams, C.J. Buiocchi, R.T. Smith, J.F. Corboy, Jr., J. Blanc and G.W. Gullen, J. AppI. Phys. 47 (1976) 5139. [51J Blanc and M.S Abrahams, J. AppI. Phys. 47 (1976) 5151.
161 171
R.S. RosIer, Solid State Technol. 20 (1977) 63. E. Tanikawa, 0. Takayama and K. Maeda, in: Proc. Conf.
CVD IV, Eds. G. Wakefield and J.M. Blocker, Electrochem. Soc.. 1974. [8) H. Ogirima, H. Saida, M. Suzuki and M. Maki, J. Electrochem. Soc. 124 (1977) 903.
[91L.N.
Aleksandrov, FL. Edelman and V.V. Voskoboinikov,
Thin Solid Films 32 (1976) 241. [101D.J. Santeler, D.H. Holkeboer, Jones and F. Pagano, Vacuum Technology and SpaceD.W. Simulation, Prepared under Contract NASw 680 by AERO VAC Corporation for NASA, 1966 (NAS SP-l05).
495
K. He,nemann, T Olaka /
(,roIi
Ill] I.E. Madey, Surface Sci. 33 (1972) 355. [121 K. Heineinann, R. Anton and H. Poppa, in: Proc. 39th EMSA Meeting (Claitor’s, 1981) p. 158. 113] K. Heinemann. H.K. Kim and H. Poppa. J. Vacuum Sci. Technol. 16 (1979) 622. 114] R. Anton and K. Heinemann. in: Proc. 38th EMSA Meeting (Claitor’s, 1980) p. 402. [15l K. (‘arey. Dept. Mater. Sci. and Eng.. Stanford University, Thesis (1981>. [16] 5. Ino and S. Ogawa, J. Phys. Soc Japan 27 (1967) 1365. [17] K. Heineinann, M.J. Yacaman, C.Y. Yang and H. Poppa. J. Crystal Growth 47 (1979) 177. [181 H. Sato and S. Shinozaki. J. AppI. Phys. 41 ,( 1970) 3165. 1191 K. Heincmann and H. Poppa, Thin Solid Films 33 (1976) 237.
[20) J.J. Metois, K. 1-leinemano and H. Poppa. AppI. Phys. Letters 29 (1976) 134.
th of (.VDolepo.sueij si//ion on a/uniina
[21) S.K. Rhce, J. Am. Ceraoi. Soc. 55 (1972> 31)0. [22] W.S. Verwoerd and F.J. Kok, Surface Sci. 80 (1979) ~j23] K. Yagi, K. Takayanagi, K. Kohayashi and (3. Honjo. J. Crystal Growth 28 (1975) 117. 124] Ml. Birjega. N. Popescu-Pogrion. V. Teodoies~cuand V Topa, Thin Solid Filnis 57 (1979) 263. [251 K. l-leinexnann, R. Anton and M. Avalos-ISoija. unpublished results. [26] DL. Howiti and T.F. Mitchell. Phil. Mag. A44 (1981) 229. [27] K. Heinemann. in: Proc. 28th EMSA Meeting ((‘laOisr’s, 1970) p. 526. 1281 R. Anton, H. Poppa and K. Heinemann. in. Proc. 39th EMSA Meeting (Claitor’s. 1981) p. 218. [29] CC. Chang. J. Vacuum Sci. Technol. 8 (1971) 500. [30] S. Ladas. Dept. (‘hem. Eng., Stanford Universit~. Fhesis (1979).
485
ELECTRON-BEAM ENHANCED GROWTH OF CVD-DEPOSITED SILICON ON ALUMINA K. HEINEMANN and T. OSAKA
*
Stanford/NASA Joint Institute for Surface and Microstructure Research, Department of Materials Science and Engineering. Stanford University, Stanford, California 94305, USA Received 30 December l98l~ manuscript received in final form 5 May 1982
Silicon was deposited onto polycrystalline alumina and single crystal sapphire substrates by low-pressure chemical vapor decomposition (CVD) of silane. The experiments were conducted in a custom UHV specimen/reaction chamber of a transmission electron microscope and recorded in-situ using an image intensifier—TV chain. Sapphire substrate areas were obtained by in-Situ electron beam flash-heating of electron-transparent alumina films prepared by anodic oxidation of aluminum foil. Silicon was observed to grow epitaxially on those sapphire substrate areas that were impinged by the electron beam during the CVD process. while in all other recrystallized substrate areas no silicon was deposited. Silicon also deposited the original amorphous/polycrystalline alumina areas that had not been heated by or exposed to the electron beam. On the other hand, no deposit was noted in alumina areas heated by the electron beam to over 1000°C but below the sapphire recrystallization temperature. We speculate that, once the substrate surface is clean, no silicon will deposit, independent of the state of crystallinity; only where the electron beam locally stimulates the decomposition of silane, deposition of silicon on the alumina substrate is observed.
I. Introduction Compared to the extensive information and cxperience available on the growth of silicon on various substrates by chemical vapor decomposition (CVD) of silane, only little has been reported about the nucleation and early growth stages of this process, and about the effect of the substrate surface orientation and cleanliness, as well as a concomitant electron beam, on the nucleation and growth characteristic, We have adapted the in-situ transmission dcctron microscopy (TEM) technique [l—3] to investigate the nucleation and early growth of silicon on sapphire substrates by low-pressure CVD. Abrahams et al. [4,5] studied the early growth of silicon on (0112)-oriented sapphire surfaces by TEM, preparing electron-transparent films by ion milling after conventional reactor growth of the silicon deposit to average thicknesses ranging from 3 to 23 nm. They obtained information about tOn leave from Waseda University, Tokyo, Japan.
0022-0248/82/0000--0000/$02.75
©
epitaxial relationships and twinning in the island deposits, but their results do not allow conclusions with regard to the influence on the nucleation and growth behavior of the substrate surface orientation, the substrate cleanliness, or a concomitant electron beam. Most earlier work has been done in CVD systems operating near atmospheric pressure. Our approach necessitates working at pressures in the 1 to l0~ Torr range. Reactors working at reduced pressures have recently been introduced [6,7]. Typical reported pressures ranged from 100 to 102 Torr [8,9]. 2. Experimental A cross-sectional schematic drawing of the microscope with reaction chamber is shown in fig. 1. We have modified a Siemens Elmiskop 101 highresolution TEM by inserting a bakeable, metalsealed chamber above the objective lens. The chamber is differentially pumped by a He cryopump. The vibrations caused by the expander
1982 North-Holland
486
K Ileinenion,i, 11 Osaka
ELECI
(Irowth i~I~,1 /).JE’p(EEItESI
si/li
~oi on (J/ilfrfriJnil
/
HON GUN
VA~Er~~/~
/O~A
HAMBER
¶f,
QUARTZ
SPECIMEI~
~
4 CRYSTAL
EXIT MEMBRANE
RELLOW~~~~
~r~IN1EHMEDIATE
~
LENS
r -
EXPANDER
PROJECTIVE
L -
J
LENS
~
EM
10
(i
PRESSURE, TORE
Fig. 2. Diff~ientiaI pressure characicristic between chamber and microscope; electron-beam entrance and exit apertures closed (curve A) and open (1 mm diameter openings, curve B).
~
/
B
/
VIEWING SECTION
_________________
II~II __________
~
SCREEN CAMERA 5,14GB INTENSIFIER
LIGHr0PTIcS TV CAMERA - ~-‘
Fig. 1. Schematic cross-section of TEM with VHV chamber for in-situ experimentation,
module of the pump have been successfully filtered out by a stainless steel bellows arrangement be-
and exit apertures being tn place (soltd line) and I . mm DIA aperture openings (dash line). Silane is admitted through a jet located 10 mm above the specimen and directed toward it. The jet arrangement consists of a 10 cm long stainless steel tithe with 0.5 mm inner diameter. During standard silane pyrolysis experiments, a chamber pressure of 2 z< 10 Torr is maintained. (Tonsidcling the pumping speed of the cryopump as conductance-limited by its cross-sectional entrance
area and considering all the geometries involved, one can estimate the silane pressure P
The electron beam enters and exits the UHV chamber via small x, ~‘-adjustable apertures. the lower one being mounted immediately above the objective aperture slider. A chamber pressure in the low io~Torr range is reached without baking within 18 h after pumpdown. Improvement is cxperienced by baking and — at a loss of image quality — by replacing the chamber entrance and
0 at the exit of the jet to be P1 0.03 Torr [101. TIns agrees with the pressure measurement of 5 Torr in the silane supply tube between the needle valve and the jet. In agreetnent with a paper by Madey [11] we then estimate the pressure P~at the site of the specimen to be P., 0.147P11 4.5 X 10 Torr. The specimen holder was designed for hightemperature in-situ CVD experiments. Temperatures well above 1000°Cneeded to he achieved at negligible outgassing rates and without having parts hotter than the specimen itself in the vicinity of the specimen. The holder is essentially a directgrid heater, using a 0.25 mm diameter nichrome wire, spotwelded around the periphery of a 200-
exit apertures with thin film membranes. The differential pressure characteristic between the
mesh Mo-specimen grid, as the main conductor. The specimen temperature was measured in-situ
TEM—column vacuum and the chamber vacuum is shown in fig. 2 for the case of thin film entrance
with an IRCON Radiation Thermometer. calibrated for the open-area-ratio of the grid and
tween expander and pump housing. Some residual vibration influence is noticeable in the electron image only when operating at highest resolution.
However, the cryopump can be shut off, without loss of pumping speed, for short durations such as required for taking high-resolution micrographs.
K. Heinemann, T Osaka
/
Growth of C VD-deposited silicon on alumina
the emissivity of molybdenum. This hot stage was tested up to 1300°C. However, most silane CVD experiments were conducted in the 900—1000°C range. A heating current of 4.5 A is required for 1000°C, but no uncorrectable astigmatism problems were encountered. The temperature equilibrium is obtained almost instantaneously, and thermal drift subsides to a level commensurate with video-image recording within one minute of a temperature change. The alumina substrates were prepared by anodic oxidation at 45 V of 0.025 mm thick aluminum foil and by subsequent dissolution of the unoxidized aluminum in a saturated mercury chloride solution. The resulting approximately 100 nm thick amorphous alumina films were washed by repeated transfer to a dish with distilled water and finally transferred to the grid heater,
487
After establishing the desired substrate temperature with the grid heater, which transformed the entire alumina film into a fine-polycrystalline state, some areas of 10 to 100 ~tm in size were recrystallized to sapphire by electron-beam flash heating [12]. Immediately following the preparation of a few satisfactory sapphire regions, an appropriate field of view usually one containing portions of several grains with different crystallographic surfaces was selected, and the silane was admitted to start the CVD process. —
—
3. Results Virtually immediately upon admitting silane, strong local changes in diffraction contrast occur in the sapphire substrate, such as would be ex-
200nm L A Fig. 3. Typical SOS growth sequence during low-pressure, in-situ CVD: left 6 mm, middle 16 mm, right 26 mm. after admission of silane.
K. l-lcin&’niann. I. Osaka
488
&romcth of (
pected if the excitation error changes due to bending of the substrate. These changes subside after a few seconds. About five minutes later, the first silicon particles can be distinguished in the dcctron image. A typical sequence of micrographs taken during a low-pressure silane CVD experiment is shown in fig. 3. Typical growth rates were 1-3 nm/mm. Neither the lateral growth rate nor the initial island number density seemed to he strongly dependent on the growth temperature between 900 and 1000°C. The island nucleation number densities were always very high. and coalescence started to
~D-de’posited s,luon on aluniinii
occur at a very early stage of growth, as soon as individual islands could be detected in the electron image. However, the number densities and average sizes of the particles varied markedly from grain to grain. Some of these differences disappeared with increasing deposit thickness. Variation of the smiane pressure during the CVD experiments by one order of magnitude also did not affect the observed nucleation and growth behavior. Fig. 4 shows typical electron diffraction patterns of silicon deposits on vartoLts sapphire surfaces and on polycrystalline alumina. Whereas the highest degree of epitaxy was generall~ noted
I
1102
220
111
1120
220 S
1102
111 111
311 Fig. 4. Typical selected—area diffraction patterns of sapphire ssoh v,mrii~u’.c\pIISed surfaces mu 1’.) amid pl~~ i s’.;,mlllne CVD-deposiied ~mlmcumi. (Hexagonal indices refer to sapphire ~uhstr,mteo
.ilutniu.i
I
di
iii
K. Heinemann, T Osaka
/ Growth of CVD-deposited silicon on
alumina
489
on the (1012) planes (“r” planes) of sapphire (fig. 4a), no consistent temperature dependence of the
with silicon, as is evidenced in fig. 5a. However, we never found any evidence of “graphoepitaxial”
degree of epitaxy was observed in the 900—1000°C range.
alignment of the silicon deposit with respect to such distinct topographical features, as is demon-
Polycrystalline alumina areas that had not been
strated by the silicon reflections in the diffraction
subjected to electron-beam heating were heavily
pattern in fig. Sd which is representative of the
decorated with silicon during silane CVD, as is demonstrated in fig. 5. The silicon particle density and area coverage was considerably higher than in the recrystallized sapphire areas. The oxide films
area of fig. 5a and does not exhibit any azimuthal preference. Once silicon had been grown on polycrystalline alumina, it was possible to re-evaporate this de-
often contained unidirectional topographical features (stemming from the fabrication process of the original aluminum foil), which were decorated
posit by localized heating with the electron beam. This evaporation occurred at a temperature below the recrystallization temperature for sapphire (yet
311 iii 220 SILICON
.lpm -
..
Fig. 5. In-situ grown silicon on alumina: (a) decoration of unidirectional surface relief structures in polvcrvstalline substrate region; (b) border-region between polycrystalline alumina and sapphire; (c) magnification of (a); and (d) diffraction pattern of (a).
490
A ll,’incman,i, T ( 15111,0
P
J
I~
ti
( ,os ii: ,,/ I i l).h pmtI J
1i
‘in ‘n i/uniina
si
~
~.
~,,
~.
;~
..4~
I
t
I iii 220 SILICON
311
S
I ig. 0. Localmied re—eNaporatmon
of
silicon from pol\crvstalline alumina film during ~—heam flash-heating Ia) and diffraction patterns
of original area with silicon deposit (b) and of silicon re-evaporated area (c).
above 1000°C at which the primary silicon (TVD growth had taken place). An example is given in fig. 6, showing the typical random-orientation diffraction patterns of silicon and alumina in the
as-deposited areas (fig. 6b). whereas the diffraction pattern taken from an electron-beam heated area (such as the one shown in the center of fig. 6a) evidences only diffraction intensities from alumina and not from silicon (fig. 6c).
The growth of silicon on recrystallized sapphire
areas by low-pressure CVD of silane was found to be strongly enhanced in the presence of the imaging electron beam. Fig. 7a shows the typical result obtained on an area that was exposed to the electron beam (100 keV, approximately 1 X 10
A/cm2) during the CVD reaction. Fig. 7b shows an adjacent recrystallized sapphire area which had not been exposed to the electron beam but had been subjected otherwise to the identical treatment. No silicon deposit was found in this or in
K. lIe,nemann, T. Osaka
/
491
Growth of CVD-deposited silicon on alumina
211m ___
Al 2 0.0112 .~
Si 111
IS—
A1 203-01T2 3 A cm2
Fig. 7. Two electron beam adjacent duringe-beam deposition; recrystallized (b) no c-beam sapphire exposure areas after during low-pressure deposition.silane CVD experiment: (a) area exposed to i0
any recrystallized sapphire area that had not been irradiated by the imaging electron beam during the CVD process. Further evidence of the influence of the electron beam during the growth of silicon by lowpressure CVD of silane is presented in fig. 8. Immediately following the recrystallization to sap-
phire, the electron beam (1 X l0~ A/cm2) was confined to approximately 2 ~tm diameter and maintained in a fixed position within the recrystallized region during the entire subsequent CVD experiment. Fig. 8a presents an overview of the recrystallized area and its surroundings, also demonstrating the silicon re-evaporation zone around
492
A. Ileini’mann. T. Osaka
(,roscth 11/ I. I l)—leplislll’il
the sapphire region of this spectnien that had been Si-deposited in a previous experiment. Fig. 8h depicts a section of fig. 8a at a higher magnification, showing the beam-enhanced silicon growth area. Figs. 8e and 8f show successively increased enlargements of the respective areas outlined in the previous micrographs. Fig. 8h is a dark field image of exactly the area shown in fig. 8f, showing moire fringes between A1 205{006} and Si{220} type planes. As is demonstrated with fig. 8d - a Si{ 111) selected-zone dark field image of the exact area of fig. 8c no silicon image intensities were detected outside the irradiated region, i.e., no silicon deposit was registered within the silicon re-evaporation zone. Fig. 9 illustrates these findings schemati-
—
~—
.5//Icon
in?
iluoiuoi
cally. If I and i~ are the electron beam intensity during microscopy and c-beam recrystallization. respectively, there appear to he three critical dcctron irradiation intensities: for I~> 1r alumina is recrystallized, for I~> / no silicon is deposited unless a (low-intensity) c-beam is present. and for I> l~enhanced growth of silicon is observed.
4. Discussion 4.1. Nucleation and ear/v growth An initially very high nuclei number density. followed by coalescence already during the very early growth is consistent with earlier in-situ
~ff4’W
K. Hm’mnemann, T. Osaka
/
_
I
lOOnm A
~
/a
~‘Sm-111
I
A1
203
-
493
Growth of CVD.depo.c:ieil silicon on alumina
0
_
lOnm
5,.
0006
*
Fig. 8. Electron-beam enhanced growth of SOS: (a) overview of recrystallized sapphire area produced by e-beam flash-heating of polycrystalline alumina film previously deposited with silicon, showing silicon re-evaporation zone in polycrystallmne alumina region; (b), (c), (e), (f) enlargements of respective outlined regions within c-beam exposed growth area; (d) Si{ Ill) selected-zone dark field (SZDF) micrograph of area shown in (c): (g) SAD pattern of area shown in (e); and (h) central dark field image of area shown in (f), using Si 220 plus sapphire 0006 type reflections.
7
‘N,
/‘
/ Ir
DURING
/
nucleation studies of gold on sapphire [13]. Abrahams et al. [4] also reported that the highest partide number densities occurred at their earliest 2partide number and 20 nm diameter), growth stage density (approximately 2 X mean 10~ cm~ and they suggested that coalescence had already occurred at that stage. Our experiments have in-
‘\ ‘\ ~,
(.‘BEAM T~LA~Ho~oA;~NG E-BEAM IRRADIAT
I
‘
dicated that the detection limit of the particles was le
-
5,
,~
~
SI
Fig. 9. Schematic cross-section of condition shown in fig. 8a, POLYCRYSTALLINE ALUMINA
SAPPH IRE (ZONE
~
EXPOSED TO
25O~C(
POLYCRYSTALLIN E ALUMINA
indicating approximate Gaussian distributions of c-beam inten. . . sitmes during recrystallization and observation (dash Curves) and . .
.
. .
critical beam intensities (dotted lines): ‘r for recrystallization, 1~for removal of contaminants, and J~for enhanced growth.
494
K. Heinemann, T. Osaka
Growth of CVD.deposits’i/ si/non on alumina
j
z 12
~
-
40
101
‘:1010
-
.60 -
—
— -5
________ 10 15 — 20 GROWTH
25
Fig. 11. GROWTH F: Mean particle TIME, MIN sapphire substrate grainsdiameter shown inversus fig. 3.growth time for three
TIME, MIN
Fig. 10. Silicon particle number density versus growth time for the three sapphire substrate grains shown in fig. 3 (top right, lower right, and left grains),
particle can be. neglected, i.e. when the particles . . A
are small [141,fig. 11 indicates a dependence d where k is approximately 2. This high value of k may actually suggest that the primary tendency during early growth is not that the coalescing particles sinter into some sort of hemispherically
governed by their number density rather than their size. If one assumes conservatively a detectability limit of average particle-to-particle distances of 2
shaped particle, but predominantly remain flat in that they form “rafts” of the original small partides. This would, in turn, be consistent with
nm and 10H atoms per cm2 in a monolayer, the particle number density versus deposition time characteristic must have peaked at density values above 2. This stipulation is confirmed in cm fig. 10 10t3 which is representative of our experiments at 1000°Csubstrate temperature (fig. 3) and represents a quantitative evaluation for the earliest growth range that we could measure with a rcasonable degree of confidence. (Extrapolation of these results to t~0 would lead to the 10’s cm 2 range.) As is apparent by the negative slope of the curves, this range is already dominated by coalescence. The development of the mean particle size
recent results by Carey [15] who suggests that the planar defects in SOS films originate during the coalescence of silicon islands in the early growth phase of the films. If complete sintering to a new single crystal particle does not occur, then the formation of twins and other planar defects is a likely alternate mechanism of accomodation. However, various other nucleation and growth results have shown that three-dimensional particle growth does not necessarily imply that the new particle is a single crystal (see. e.g.. refs. [16—181).
(fig. 11) is also indicative of coalescence being a domineering mechanism in the early particle growth phase. Whereas a dependence of the mean particle diameter d on the growth time I of d -~ 17 is expected for “regular” three-dimensional partide growth, when the adatom diffusivity on the substrate surface is high, when the capture rate of adatoms is proportional to the circumference of the particle, and when the contribution of the
experiments of gold on amorphized graphite substrates [19]. However, in these experiments the substrate temperature was much lower, and even then complete sintering usually occurred some time later, thus allowing interpretation of the results as a two-step process of coalescence, leading via “raft” formation as precursor state to complete sintering. An attempt to explain the underlying reason for the present results should be consistent with our
vapor beam directly impinging on the growing
specific observations that:
Coalescence without immediate sintering into
particles of essentially three-dimensional shape was found earlier in in-situ nucleation and growth
K. Heinemann, T Osaka
/
Growth of CVD-deposited silicon on alumina
495
(i) A value of k 2 was found not only during the very early stages but up to raft diameters well above 100 nm. (ii) Coalescence into rafts involved particle mo-
upper right, and left grains shown in fig. 3, respectively). Whereby we were not able to determine the exact substrate surface orientation for this exampie, it is clear that in the depicted growth range the
bility (that appeared very similar to the type of mobility seen in earlier experiments [19,13,20]),
particles on the upper right grain (dash curves in figs. 10 and 11) grew in the depicted time interval
(iii) Relatively large areas of uniform diffraclion contrast exist within individual rafts at a later stage of growth, such as shown in fig. 3c, indicating that some re-formation of the rafts into single crystals or less faulted particles does occur, hut the resulting particles essentially reproduce the flatness and substrate area coverage of the original rafts. Whether the low self-diffusivity of silicon (even at the high substrate temperature used) or the relatively low ratio of free surface energies of overgrowth and substrate (~w 1 for SOS, >> 1 for example for Au/NaCI [21,22] which is well documented to immediately sinter into three-dimensional particles upon coalescence [18,23,24]), is the primary cause for the raft formation is difficult to determine at this time. The latter cause would be most far-reaching, in that it would essentially imply that layer-type growth is favored in the initial stages of SOS epitaxy. Our observations of appreciable diffraction contrast changes immediately following silane admittance into the in-situ reaction chamber, i.e. at a much earlier stage of growth than depicted in fig. 3a, well before appearance of the silicon islands, would be consistent with such a model of initial layer growth. Initial layer coverage of silicon could exert strains on the substrate film, causing bending and concomitant diffraction contrast changes. It should be noted, however, that our present experimental evidence is not strong enough to warrant more than a speculation of
markedly faster than the particles on the other two grains, indicating that, all other parameters and
layer-type growth in the initial stages of growth of silicon on sapphire. More work would need to be
done, possibly requiring such elaborate techniques as in-situ micro-area Auger spectroscopy to clarify this important point. Irrespective of these speculations concerning the initial growth of silicon nuclei on c-beam recrystallized sapphire during low-pressure CVD, fig. 3 and, more clearly, figs. 10 and 11 exhibit differences in the decoration results obtained on three different substrate grains (the solid, dash, and dash-dot curves in figs. 10 and 11 refer to the lower right,
conditions being the same, the substrate surface
structure does have a marked influence on the nucleation and growth characteristic of silicon on sapphire. The observation that neither the silane pressure nor the substrate temperature markedly affected the growth rate or the nucleation number density in our experiments is generally unusual in the epitaxial growth of islands and thin films, where one would expect a decrease of the particle number density and the growth rate, as well as usually an improvement in epitaxy, with a decrease in supersaturation. A departure from this rule has been reported earlier for SOS in that for yet not fully understood reasons an optimum substrate temperature (of 960°C,uncalibrated for the emissivity of sapphire as is customary in the semiconductor industry) exists, beyond which the degree of epitaxy decreases again and the number of planar defects increases [15]. The range of substrate temperatures covered in our studies, if the emissivity calibration is considered, falls just below the 960°C mark, and most of our results should, therefore, fall within the range of supersaturations where “normal” growth behavior could be expected. Since the very high initial particle number densities observed in our studies cannot be explained on the basis of high supersaturation conditions, a —
—
conclusion which is emphasized by the fact that
Abrahams et al. obtained quite comparable partide density results at by three orders of magnitude higher silane pressures, and since substrate surface contamination can be excluded as primary reason in our experiments (on re-crystallized sapphire) due to the low chamber pressure (108 Torr) and high substrate temperatures maintained after recrystallization, we conclude that the imaging electron beam has been of primary influence. This conclusion is certainly confirmed in our oh-
490
A. Ils’,ncnia,in.
7T Osaka
Grsniih of I. I ‘1) ilepss,siIci/ .5//us/n in s1/ioiiiiisi
servations illustrated with figs. 7 and 8. Our results (fig. 4) essentially confirm that the degree of epitaxv is highest on the “r” plane (fig. 4a) which is used virtually exclusively tn SOS technology. However, twinning and faulting are also most pronounced in this orientation (most of the diffraction intensities in figs. 4h and 4c are due to random particle orientation and to contrihutions from neighboring not fully recrystallized alumina regions). This seems to be consistent with the high lattice mismatch in these two orientations (figs. 4b and 4c), e.g., for Si{ Ill } H A 203[ 0001 }. where the lattice mismatch is 24%. At such htgh values of mismatch the strain to accomodate the geometrical misfit would he so high that the degree of direct lattice match ceases to play an important
role, and effects caused by the mismatch, such as planar defects and microtwinning. would become less prominent, 4.2. Electron-beam enhanced nucleation and growth
The in-situ growth of silicon was strongly enhanced during electron beam irradiation. It is tinlikely that the striking magnitude of this effect, evidenced in fig. 7 for (lie case of two neighboring recrystallized areas that had been subject to idetitical conditions except for the beam exposure during the CVD process, can he explained as due to defects introduced in the substrate by the imaging electron beam. Since the electron irradiation intensity was of the order of 10 ~ A cm 2 impingement on the substrate surface occurred at a rate of a few electrons per atom site per second (6 ‘>< 1015 cm 2 ~ I) or some 1018 electrons during a typical experiment of some 3 mm duration. Radiation damage tests done on sapphire indicate that during several hours of 100 keV electron irradiation at 10 A cm 2 of 100 nm thick sapphire films no visible damage occurs [25].This computes to some 1022 electrons cm 2 This assumption is very conservative also in the light of Howitt and Mitchell’s recent results [26] who obtained damage in sapphire upon irradiation of several minutes duration at 4 orders of magnitude higher electron current densities and at energies above the threshold for radiation damage, whereas our work was done below this threshold. If we (conservatively) assume
m~displacements per cni2 would that a density of i0 be sufficient to he visible in the electron image (such as by appearance of phase contrast charactermstic of amorphous or amorphized films [27.1 VI). this gives a displacement probability of ‘~ 10 per primary incident electron ativwhere in the film. The probability for this damage to occur near the surface (saY, within the top 10 nih depth from the surface). where its effect could he the creation of a site for preferred nttcleation. is even lower, probably < 10 ~. Thus, during ;m typical experiment, the total impinging electron flux would have the capability to create less than l0’~~ 1(1 — l0~° sites for preferred nucleation per cm2. which is about 3 orders of magnitude less than the numher of preferred sites one should expect if radialion damage in the substrate is argued as the major factor in determining the c-beam enhanced growth reported here. One should also note that in earlier work on in—situ nucleation and growth of metals on recrystallized sapphire surfaces. tio effect of the electron beam was observed during very comparable electron irradiation conditions. These studies were done under standard PVD condittons and involved inert deposit materials such as gold. as
well as Ag, Pd. and Fe [13.14.281. We therefore. with due caution, rule out c-beam induced radiatioti damage as primary cause for the observed c-beam enhanced growth of silicon. One might then he tempted to explain this effect on the basis of an interaction of the electron beam with the S1H 4 flux. such as c-beam stmrntmIated decomposition of silane. This mechanism would essentially postulate that the electron beam intensity is rate limiting during the growth process of the silicon islands. Although the relative insensi-
tivity of our results with changes in silane pressure and in substrate temperature, as well as the apparent fall-off of average particle size toward the edges of the irradiated area (fig. 8c). would be consistent with this explanation. two other ohservations cannot be as readily explained~~. (i) the apparent heavy silicon depositton on the original. not recrystallized and not c-beam heated alumina substrate areas, and (ii) the high nucleation numher densities. The heavy silicon depostt on the original, not recrystallized and not annealed (“asreceived”) alumina surfaces could conceivably he
K. Hei,temann, T. Osaka
/
Growth of C VD-deposited silicon on alumina
497
associated with surface contamination. Considering the circumstance that heating the substrate to 1000°Cfor several hours apparently does not af-
ence of an electron beam on areas free of contaminants re-establishes growth at essentially the same rate that is experienced on contaminated areas
fect the silicon nucleation and growth behavior (i.e., presumably does not remove the contami-
without an electron beam. Further studies are needed to substantiate the role of substrate surface contaminants in the low-
nants), but short-duration heating to a temperature somewhat above 1000°C but below the recrystallization temperature of sapphire (about
pressure CVD process, to quantify the growth-rate limiting influence of the electron beam intensity
I 250°(’) does have that effect, may possibly point
for preferred growth, and to investigate the specu-
to hydrocarbon contaminants or carbon which are known to be removed only during very high-ternperature annealing well above 1000°C[29,30]. The critical electron irradiation intensity I~(fig. 9) would then be the intensity required for removal of contaminants. There being not really any doubt that the as-received alumina substrates are contaminated, the speculated high degree of cleanliness after recrystallization should further he investigated. Furthermore, comparison with similar experiments on bulk sapphire surfaces would be needed to connect with results such as reported by Blanc and Abrahams [5] who claim that their nucleation results are not influenced by contamination.
lated layer growth on clean sapphire substrate areas.
Acknowledgements We thank Dr. H. Poppa of NASA—Ames for stimulating discussions and suggestions during the course of this work, and we appreciate Dr. R. Schindler’s help during the construction and early testing phase of the in-situ chamber. Dr. J. Blanc of RCA—David Sarnoff Research Center in Princeton, NJ, offered constructive criticism and helpful suggestions to this work, especially to the discussion on early growth. We gratefully acknowledge his contribution. Support by NASA (Grant NCC2-ll I) is gratefully acknowledged.
5. Conclusions The nucleation and growth of silicon on sapphire by low-pressure CVD was performed inside a custom, controlled-vacuum specimen/reaction chamber of a high-resolution TEM. The initial nuclei number density is very high, probably > 10 cm2, and coalescence sets in at a very early stage of growth, possibly indicating an initial layer-type growth mechanism of the silicon deposit. Silicon grows highly preferentially on those substrate areas that are exposed to the imaging electron beam. This is believed to be due to electron-beam stimulated decomposition of silane. The observation that silicon also grows on polycrystalline alumina areas that have not been subjected to heat treatment above 1000°C indicates that surface contaminants have a major role in the nucleation and growth of silicon on alumina substrates of various orientations and crystallographic phases. When contaminants are present, silicon grows heavily; without contaminants, growth is inhibited. Pres-
References [I] RD. Moorhead and H. Poppa, in: Proc. 27th EMSA Meeting (Claitor’s, 1969) p. 116. [21 K. Heinemann, D.B Rao and D L. Douglass, Oxidation Metals 9 (1975) 379. 131 H. Poppa and K. 1-leinemann, Optik 56(1980)183. [41 MS. Abrahams, C.J. Buiocchi, R.T. Smith, J.F. Corboy, Jr., J. Blanc and G.W. Gullen, J. AppI. Phys. 47 (1976) 5139. [51J Blanc and M.S Abrahams, J. AppI. Phys. 47 (1976) 5151.
161 171
R.S. RosIer, Solid State Technol. 20 (1977) 63. E. Tanikawa, 0. Takayama and K. Maeda, in: Proc. Conf.
CVD IV, Eds. G. Wakefield and J.M. Blocker, Electrochem. Soc.. 1974. [8) H. Ogirima, H. Saida, M. Suzuki and M. Maki, J. Electrochem. Soc. 124 (1977) 903.
[91L.N.
Aleksandrov, FL. Edelman and V.V. Voskoboinikov,
Thin Solid Films 32 (1976) 241. [101D.J. Santeler, D.H. Holkeboer, Jones and F. Pagano, Vacuum Technology and SpaceD.W. Simulation, Prepared under Contract NASw 680 by AERO VAC Corporation for NASA, 1966 (NAS SP-l05).
495
K. He,nemann, T Olaka /
(,roIi
Ill] I.E. Madey, Surface Sci. 33 (1972) 355. [121 K. Heineinann, R. Anton and H. Poppa, in: Proc. 39th EMSA Meeting (Claitor’s, 1981) p. 158. 113] K. Heinemann. H.K. Kim and H. Poppa. J. Vacuum Sci. Technol. 16 (1979) 622. 114] R. Anton and K. Heinemann. in: Proc. 38th EMSA Meeting (Claitor’s, 1980) p. 402. [15l K. (‘arey. Dept. Mater. Sci. and Eng.. Stanford University, Thesis (1981>. [16] 5. Ino and S. Ogawa, J. Phys. Soc Japan 27 (1967) 1365. [17] K. Heineinann, M.J. Yacaman, C.Y. Yang and H. Poppa. J. Crystal Growth 47 (1979) 177. [181 H. Sato and S. Shinozaki. J. AppI. Phys. 41 ,( 1970) 3165. 1191 K. Heincmann and H. Poppa, Thin Solid Films 33 (1976) 237.
[20) J.J. Metois, K. 1-leinemano and H. Poppa. AppI. Phys. Letters 29 (1976) 134.
th of (.VDolepo.sueij si//ion on a/uniina
[21) S.K. Rhce, J. Am. Ceraoi. Soc. 55 (1972> 31)0. [22] W.S. Verwoerd and F.J. Kok, Surface Sci. 80 (1979) ~j23] K. Yagi, K. Takayanagi, K. Kohayashi and (3. Honjo. J. Crystal Growth 28 (1975) 117. 124] Ml. Birjega. N. Popescu-Pogrion. V. Teodoies~cuand V Topa, Thin Solid Filnis 57 (1979) 263. [251 K. l-leinexnann, R. Anton and M. Avalos-ISoija. unpublished results. [26] DL. Howiti and T.F. Mitchell. Phil. Mag. A44 (1981) 229. [27] K. Heinemann. in: Proc. 28th EMSA Meeting ((‘laOisr’s, 1970) p. 526. 1281 R. Anton, H. Poppa and K. Heinemann. in. Proc. 39th EMSA Meeting (Claitor’s. 1981) p. 218. [29] CC. Chang. J. Vacuum Sci. Technol. 8 (1971) 500. [30] S. Ladas. Dept. (‘hem. Eng., Stanford Universit~. Fhesis (1979).