Metastable Solvent Epitaxy of SiC,
the Other Diamond Synthetics 7
liquid
Ni
diamond
graphite
ba
temperature [K]
atomic fraction of carbon
1800
1750
1700
1650
Ni
0 0.1 0.2 0.996 0.998 1.0
graphite
graphite + diamond
1728 K
diamond
1740 K
liquid + diamond
liquid Ni(C)
α
α
+
liquid
1667 K (Ni + diamond)
1661 K (Ni + graphite)
liquid +
graphite
Fig. 6. Schematic drawings of a diamond synthesis and b enlarged sections of the Ni-C phase

Ostwald ripening is the reaction in a solid solutions, which means the life time of metastable
phase may be longer at lower temperatures and l ower d iffusivity. Does such a m etastable
phase directly melt in liquid?
The typical example of the behavior of the coexistence of stable and metastable phases is
observed in the Fe-C system. As mentioned before, the double phase diagram of Fe-C and
Fe-Fe
3
C systems is well studied and established. The m elting behavior of metastable Fe
3
C
phase has investigated in detail by Okada et al. (1981). They measured the differential thermal
analysis (DTA) curves for the white, gray and mixture cast irons at the eutectic temperature
and composition region. Fig. 8 shows the summarized results of DTA curves as well as the
schematic double phase diagram. The endothermic temperatures shift due to the kinetic
reason of the measuring apparatus, but the corrected temperatures show the stable and
metastable eutectic temperatures o f 1426K and 1416K, respectively. The specimens of gray
cast iron contains stable phases of graphite and fcc-Fe(autenite), where they all melt only
59
Metastable Solvent Epitaxy of SiC, the Other Diamond Synthetics
8 Will-be-set-by-IN-TECH
Concentration
Free energy G
x
α
L
x
α
S
x
β

3
C phase occurs. For the specimens of mixture cast iron, the reactions are complicated but
the melting and s olidifying occur simultaneously. These experimental results i ndicate that the
metastable phase is so stable that can melt directly.
Temperature
Exothermic Endothermic
gray cast iron
white cast iron
mixture iron
ΔT
a
b
Fig. 8. a DTA curves of cast irons(Okada et al., 1981) and b the double phase diagram of
equiblibrium Fe-Graphite and metastable Fe-Fe
3
Csystems.
3.5 Speculated mechanism
From the experimental result shown in Fig. 4, 4H-SiC is expected to be more stable than
3C-SiC. The Si-C system should show a double phase diagram, as schematically shown in
Fig. 9a. The corresponding free-energy vs. concentration diagram is also illustrated in Fig. 9b.
60
Silicon Carbide – Materials, Processing and Applications in Electronic Devices
Metastable Solvent Epitaxy of SiC,
the Other Diamond Synthetics 9
The solubility limit of each phase i s determined by the tangent common to the free-energy
curves of the coexisting phases. At the temperature indicated by the dotted line, the solubility
limit of metastable 3C-SiC is x
3C
l
, which corresponds to the dashed line of the liquidus in

functions in Chase (1998) suggest that α(hexagonal) phase is less stable up to 2000K, and they
concluded unlikely the transformation to β(cubic) phase at abo ut 2300K.
The most widely adapted phase diagram should be that by Olesinski & Abbaschian (1996)
as shown in Fig. 1, where the β(cubic) phase is more stable than the α(hexagonal) phase
at any temperatures below the periodic temperature of the decomposition of SiC, 2545
◦
C.
Although the evaluators of Olesinski & Abbaschian (1996) mentioned nothing on the types of
α(hexagonal) phase, the same authors reported the co-existence o f polytypes of α phases, 6H,
15R, and 4H(Olesinski & Abbaschian, 1984). Furthermore, it also mentioned on the report
of Verma & Krishna (1966), the existence of α stability above 2000
◦
C. On the other hand,
Fromm & Gebhardt (1976) reported the different type of phase diagram as shown in Fig. 11,
wherethephasetransitionfromβ to α phases occurs at around 2000
◦
C.
Solubilities of carbon in liquid silicon measured by Hall (1958), Scace & Slack (1959), Dash
(1958), Dolloff (1960), Nozaki et al. (1970), Oden & McCune (1987), Suhara et al. (1989),
Kleykamp & Schumacher (1993), Iguchi & Narushima ( 1993), O ttem (1993), and Yanabe et al.
(1997) are summarized as in Figs. 12. Tw o reported phase diagrams as shown i n Fig. 1 and
Fig. 11 are based on the data given by Dolloff (1960). Dolloff (1960)’s data, however, are
distinctively different from the others, where the solubility limits are larger than the others.
61
Metastable Solvent Epitaxy of SiC, the Other Diamond Synthetics
10 Will-be-set-by-IN-TECH
temperature
liquid Si
SiC
a

4H
Si
S
+ SiC
3C
Fig. 9. Schematic drawings of a the p redicted Si-C double phase diagram, b related
free-energy vs. concentration diagram and c carbon concentration p r ofile in the liquid Si
solvent between the 3C-SiC source and the 4H-SiC fine particles. The metastable eutectic
temperature of the reaction liquid Si
→ Si
S
+ SiC
3C
is lower than the stable eutectic
temperature of the reaction liquid Si
→ Si
S
+ SiC
4H
,whereSi
S
denotes solid Si. The chemical
potentials of C, μ
c
, are given by the intersections of the co mmon tangents with the
pure-carbon line in b, and are spatially different in the liquid Si solvent contacting with
3C-SiC and 4H-SiC in c. The configuration of c is re lated to that of the panel on the left-hand
side in Fig. 2b.
62
Silicon Carbide – Materials, Processing and Applications in Electronic Devices

The reported phase stability between α(hexagonal) and β(cubic) might be 6H and 3C. If this
assumption is true, 4H stability has not been shown experimentally. Furthermore, the distinct
difference of solubility limits indicates that the coexistence of stable and metastable phases.
4.2 Difficulty of the equilibrium state
Although the co nflicts of phase diagrams shown above remain, there have been many
attempts of experimental and theoretical researches on the kinetic p rocess of crystal growth
of SiC polytypes. Famous stability diagrams of SiC polytypes proposed by Knippenberg
(1963) and Inomata e t al. (1968) show that the crystalized phases are controlled both by the
temperature and growth rate of the operations.
The l imit to slow growth rate of kinetic processes or results of long period holding should be
equal to static results. But it is very difficult to dissolve w hole amount of SiC crystals due to
small solubility limit of SiC in liquid Si. If there remain seeds of metastable phases during the
previous pr ocesses, it is difficult to re move all of the m. T he metastable phase also grows due
to the co-exsitence of less stable phases as shown in Ostwald ripening, or from super-saturated
liquid Si. Furthermore, the required high temperature and inert environment make the static
conditions very difficult.
Inomata et al. (1969) performed careful experimental observations on the relationship
between the polytypes of SiC and the supersaturation of the solution at 1800
◦
Cwiththe
solution method, and have shown the following results;
1. β-SiC crystallizes from highly supersaturated solution. The crystals obtained at the
condition of low supersaturation, however, consist of mainly α-SiC such as 4H, 15R and
6H.
2. Relative amount of 4H increased markedly with decreasing the supersaturation.
3. From the results stated above, it is concluded that 4H is the most stable structure at 1800
◦
C
among the basic polytypes of SiC, 3C, 4H, 15R and 6H.
Those results indicate that the difference between 4H and 6H is crucial for determining the

0.006
4H
6H
3C
2H
Temperature [K]
Free energy difference [eV/SiC-pair]
0
1,000
2,000 3,000
Temperature [K]
Fig. 13. First principles calculations of temperature dependency of free energy difference of
6H, 3C, and 2H against 4H Si C(Nishitani et al., 2009). Finite temperature effects are included
through the vibrational free energy calculated by Phonon codes(Medea-phonon, n.d.;
Parlinski et al., 1997).
These first principles calculations were carried out using the Vienna Ab initio Simulation
Package (VASP) code(Kresse & Furthmüller, 1996a;b; Kresse & Hafner, 1993; 1994). The
interaction between the ions and valence electrons was described by a projector
augmented-wave (PAW) method(Kresse & Joubert, 1999). A plane-wave basis set with a cutoff
of 400 eV was used. The exchange-correlation functional was described by the generalized
gradient approximation (GGA) of the Perdew-Wang91 form(Perdew & Wang, 1992). Phonon
calculation was performed by a commercial pre-processor of Medea-phonon(Medea-phonon,
n.d.) with t he direct method developed by Parlinski et al. ( 1997). The volumes and/or c/a
ratios were fitted to the most stable point at each temperature.
Fig. 13 shows the te mperature dependencies of free energy of 6H, 3C, and 2H SiC polytypes
measured from 4H SiC. 4H SiC is most stable at low temperatures, but 6H Si C is most stable
at higher temperatures. 3C SiC is less stable against 4H or 6H SiC except at very high
temperature region. Those results are consistent with the other speculations but the precisions
of the calculations are not enough. Although the more precise calculations will alter the results
of hierarchy of polytypes, their result pointed out the possibility of the phase transition in the

Giardini, A. A. & Tydings, J. E. ( 1962). Diamond synthesis: Observations on the mechanism
of formation, American Mineralogist 47: 1393–1421.
Hall, R. N. (1958). Electrical contacts to silicon carbide, J. Appl. Phys. 29(6): 914–917.
Hansen, M. & Anderko, K. (1958). Constitution of binary alloys, 2nd Ed., McGraw-Hill, New
York, pp. 353–365.
Hofmann, D. H. & Müller, M. H. (1999). Prospects of the use of liquid phase techniques for
the growth of bulk silicon carbide crystals, Mater. Sci. and Eng. B 61-62: 29–39.
Hurle, D. T. J., Mullin, J. B. & Pike, E. R. (1967). Thin alloy zone crystallisation, J. Mater. Sci.
2: 46–62.
Iguchi, Y. & Narushima, T. (1993). 1st. Int. Conf. on Processing Materiasl for Properties,The
Minerals, Metals & Materials Society, pp. 437–440.
Inomata, Y., Inona, A., Mitomo, M. & Sudzuki, H. (1968). Relation between growth
temperature and the structure of SiC crystals grown by sublimation method,
Yogyo-Kyokai-Shi 76(9): 313–319.
Inomata, Y., Inoue, Z., Mitomo, M. & Tanaka, H. (1969). Polytypes of SiC crystals grown from
molten silicon, Yogyo-Kyokai-Shi 77(3): 83–88.
Ishihara, K. N., Nishitani, S. R., Miyake, H. & Shingu, P. H. (1984-5). Rapid solidification
and the metastable phase diagrams of the f e-c, co-c and ni-c systems, Int. J. Rapid
Solidification 1: 51–58.
Izhevskyi, V. A., Genova, L. A., B ressiani, J. C. & Bressiani, A. H. A. (2000). Review article:
silicon carbide. structure, properties and processing, Cerâmica 46: 4 – 13.
Kleykamp, H. & Schumacher, G. (1993). The constitution of the silicon-carbon system, Ber.
Bunsenges. Phy. Chem. 97(6): 799–805.
Knippenberg, W. F. (1963). Growth phenomena i n silicon carbide, Philips Research Reports
18: 161–274.
Kresse, G. & Furthmüller, J. (1996a). Efficiency of ab-initio total energy calculations for metals
and semiconductors using a plane-wave bas is set, Comput.Mat.Sci.6: 15–50.
Kresse, G. & Furthmüller, J. (1996b). Efficient iterative schemes for ab initio total-energy
calculations using a plane-wave bas is set, Phys. Rev. B 54(16): 11169–11186.
66

Olesinski, R. W. & Abbaschian, G. J. (1996). Binary Alloy Phase Diagrams, 2nd ed.,ASM
International, Materials Park, Ohio, pp. 882–3.
Ottem, L. (1993). SINTEF Report STF34 F93027.
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, Phys.Rev.Lett.78(21): 4063 – 4066.
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Metastable Solvent Epitaxy of SiC, the Other Diamond Synthetics
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68

was (1.7−2)×10
17
см
-3
in the layer (n)
6H-SiC and acceptors N
a
−N
d
~ 3×10
18
см
-3
in the layer (p) 3C-SiC. In the spectrum of the
electroluminescence of diodes revealed two bands with maxima hν
max
≈ 2.9 eV (430 nm) and
2.3 eV (540 nm), close the band gaps of 6H-and 3C-SiC. Currently, using the methods of
vacuum sublimation (Savkina et al., 2000), molecular beam epitaxy (Fissel et al., 1996), the
epitaxial and heteropolytype layers based on the cubic 3C-SiC and two hexagonal 6H-SiC,
4H-SiC on substrates of SiC, are grown. By chemical vapor deposition (CVD) (Nishino et al.,
2002) are grown heteroepitaxial layers of 3C-SiC on substrates of Si. At the temperatures
below 1200°C there are conditions for the growth of both poly- and nanocrystalline SiC with
different degrees of crystallinity and structure of the cubic polytype 3C-SiC. Such conditions
were realized in the magnetron sputtering (Kerdiles et al., 2000; Sun et al., 1998), laser
ablation (Spillman et al., 2000) and plasma deposition (Liao et al., 2005), plasma-enhanced
chemical vapor deposition (George et al., 2002; Pajagopalan et al., 2005), molecular beam
epitaxy (Fissel et al., 2000). At temperatures below 1500°C in the direct deposition of carbon
and silicon ions with an energy of ~100 eV, the growth of nanocrystalline films with a
consistent set of the polytypes 3C, 21R, 27R, 51R, 6H is possible (Semenov et al., 2008, 2009,

ions (Е = 200
кэВ, D = ~10
17
см
-2
(Borders et al., 1971) or Е = 40 кэВ, D > 10
17
ион/см
2
(Baranova et al.,
1971)) into a silicon substrate. From the position of the IR absorption band (700–725 cm
-1
),
reducing its half-width after annealing and displacement in the region at 800 cm
-1
,
corresponding to transverse optical phonons of SiC, it was found that the formation of
crystalline SiC phase occurs in the temperature range near 850ºC (Borders et al., 1971) and
900°C (Baranova et al., 1971). Silicon carbide was identified using transverse optical phonon
spectra in most of the above work on ion implantation, as well as in (Gerasimenko et al.,
1974, Wong et al., 1998, Akimchenko et al., 1977a, 1980; Chen et al., 1999, Kimura et al.,
1981). The detection of longitudinal optical vibrations of lattice atoms (LO phonons) and
their changes during film annealing give additional information on the crystallization
processes (Akimchenko et al., 1977b, 1979).
Difficulties associated with the problem of synthesis of a crystalline silicon carbide
prevent the wide use of SiC in microelectronics. The Si-C mixture, after implantation of
large doses of carbon, is assumed to be amorphous (Lindner , 2003; Liangdeng et al., 2008;
Kimura et al., 1981). Carbon atom diffusion in the implanted layer is restricted by the
strong Si-C bonds (Liangdeng et al., 2008). A negative influence of stable C- and C−Si-
clusters (Yan et al., 2000, Chen et al., 2003; Kimura et al., 1982, 1984; Durupt et al., 1980,

Implantation of carbon ions at temperatures of silicon substrate well above room
temperature by in-situ annealing helps to form SiC crystalline layers immediately during
implantation or after annealing at lower temperatures (Durupt et al., 1980; Frangis et al.,
1997; Edelman et al., 1976; Simon et al., 1996; Preckwinkel et al., 1996). High temperature of
the substrate can be achieved also by using beams with high current density of carbon ions
(Reeson et al., 1990; Alexandrov et al., 1986). Treatment of carbon implanted silicon layers
by power ion (Liangdeng et al., 2008; Bayazitov et al., 2003), electron (Theodossiu et al.,
1999) or laser (Bayazitov et al., 2003a, 2003b) beams like thermal annealing also leads to the
formation of a polycrystalline β-SiC layer.
In this paper, the composition and structure of homogeneous SiC
1.4
, SiC
0.95
, SiC
0.7
, SiC
0.4
,
SiC
0.12
and SiC
0.03
layers, received by multiple high-dose implantation of carbon ions with
energies of 40, 20, 10, 5 and 3 keV are investigated. The influence of decay of carbon- and
carbon-silicon clusters during thermal annealing or hydrogen glow discharge plasma
processing on the formation of tetrahedral Si–C-bonds and crystallization processes in
silicon layers with high and low concentrations of carbon, is studied.
2. Experimental
Single-crystal (100) silicon wafers of sizes 7×12×0.4 mm
3

with respect to the normal to the sample
surface were measured. The composition of the layers was examined by Auger electron
spectroscopy. The parameters were as follows: incident electron beam of diameter 1 μm,
energy 10 keV, angle of incidence 45°, diameter of scanning region 300 μm, vacuum 1.33
×10
-8
Pa, angle of Ar
+
beam incidence 45°.

Silicon Carbide – Materials, Processing and Applications in Electronic Devices

72
The glow discharge hydrogen plasma was generated at a pressure of 6.5 Pa with a capacitive
coupled radio frequency (r.f.) power (27.12 MHz) of about 12.5 W. The temperature of
processing did not exceed 100°С and it was measured by thermocouple. The processing time
was 5 min.
3. Results and discussion
3.1 Depth profiles of multiple-energy implanted carbon in Si
To produce a rectangular profile of the distribution of carbon atoms in silicon five
energies and doses have been chosen in such a way (Table 1) that the concentration ratio
of C and Si atoms to a depth (up to ~120 nm) was equal to the values of N
C
/N
Si
= 1.0 , 0.8,
0.5, 0.3, 0.1 and 0.03. The calculated profiles N
C
(Burenkov) are the sums of the
distributions (Figs. 1 and 2), constructed for the chosen values of energies (E) and doses

D
Nx
RR
π
−
=−
ΔΔ
, (1)
where х is the distance
from the surface.
Figs. 1 and 2 also show the experimental profiles N
C
(20°С), N
C
(1250°С) and N
О
(1250°С),
obtained by Auger electron spectroscopy, which show the concentration ratio of carbon and
oxygen atoms (N
C
/N
Si
and N
О
/N
Si
) over the sample depth after implantation and annealing
at 1250°C for 30 min in an argon atmosphere with low oxygen content. Fig. 1 shows that the
average concentrations of carbon and oxygen were: a)N
C

/N
Si
= 0.12.
Thus, the average carbon concentration over the depth exceeded the corresponding
calculated values (N
C
/N
Si
= 0.1; 0.3; 0.5; 0.8 and 1.0) and led to the formation of layers
SiC
0.12
, SiC
0.4
, SiC
0.7
, SiC
0.95
and SiC
1.4
(Table 2). The average value of the ratio of N
О
/N
Si

exceeded the stoichiometric value for SiO
2
(N
О
/N
Si

23 2 2
C
SiC O SiO CO
−°
+ ⎯⎯⎯⎯⎯→+.

The Formation of Silicon Carbide in the SiC
x
Layers
(x = 0.03–1.4) Formed by Multiple Implantation of C Ions in Si

73 0.0
0.5
1.0
1.5
2.0
2.5
3.0
0 40 80 120 160 200
х, nm
N
C
/ N
S i

С
(5 keV)
N
C
(20°C)
c) 0.0
0.5
1.0
1.5
2.0
2.5
0 40 80 120 160 200
х, nm
N
C
/ N
S i
( N
O
/ N
S i
)
N
C
(Gibbons)
N
C

1.0
1.5
2.0
2.5
0 40 80 120 160 200
х, nm
N
C
/ N
S i
( N
O
/ N
S i
)
N
C
(Gibbons)
N
C
(40 keV)
N
C
(20 keV)
N
C
(10 keV)
N
C
(3 keV )

C
(Burenkov) and N
C
(Gibbons) are the profiles calculated according to
Burenkov et al. (1985) and Gibbons et al. (1975), respectively, where N
C
(Gibbons)

= N
C
(40
keV) + N
C
(20 keV) + N
C
(10 keV) + N
C
(5 keV) + N
C
(3 keV). N
С
(20
0
С), N
С
(1250
0
С) and
N
О

N
C
(40 keV)
N
C
(20 keV)
N
C
(10 keV)
N
С
(5 кэВ)
N
C
(3 keV)

0.0
0.1
0.2
0.3
0.4
0.5
0.6
0 40 80 120 160 200
х
, nm
N
C
/N
Si

O
(20°C)
а)

Fig. 2.
12
С distribution profiles in Si produced by ion implantation (see table 1). (a) SiC
0.4
; (b)
SiC
0.12
. N
C
(Burenkov) and N
C
(Gibbons) are the profiles calculated according to Burenkov et
al. (1985) and Gibbons et al. (1975), respectively, where N
C
(Gibbons)

= N
C
(40 keV) + N
C
(20
keV) + N
C
(10 keV) + N
C
(5 keV) + N

17
c
m
-2
4.48 1.54 0.792 0.264 0.184
D
(
SiC
0.5
)
, 10
17
c
m
-2
2.80 0.96 0.495 0.165 0.115
D
(
SiC
0.3
)
, 10
17
c
m
-2
1.68 0.576 0.297 0.099 0.069
D
(
SiC

(Е).
nm
46.0 28.3 16.9 10.2 7.2
N
C
(Gibbons) profile
(Gibbons et al., 1975)
R
p
(Е).
nm
93.0 47.0 24.0 12.3 7.5
ΔR
p
(Е).
nm
34.0 21.0 13.0 7.00 4.3
Table 1. Values of energy, E, dose, D, projected range, Rp(E), and straggling, ΔRp(E), for
12
C
+

ions in Si, used for constructing a rectangular distribution profile
The Formation of Silicon Carbide in the SiC
x
Layers
(x = 0.03–1.4) Formed by Multiple Implantation of C Ions in Si

75
In (Mandal et al., 2001) reported 100% conversion of SiC into SiO


Depth range, nm 40–110 20–120 25–120 10–120 20–120
SiC
х
(Burenkov) SiC
1.0
SiC
0.8
SiC
0.5
SiC
0.3
SiC
0.1

SiC
х
(Gibbons) SiC
1.38
SiC
1.06
SiC
0.67
SiC
0.40
SiC
0.13

SiC
х

formed in the surface layer of single crystal Si wafers using multiple ion implantation, after
annealing at 1200°C for 30 minutes were investigated. Figure 3 schematically shows a
section of the objects. The investigated area can be divided into three sections: section 1
consists of a layer SiC
x
; section 2 includes a [transition layer "Si-SiC
x
" + layer SiC
x
]; section 3
is a three-layer structure [layer-Si + transition layer "Si-SiCx "+ layer SiCx].

5
SiC
x
SiC
x
SiSi 1
22
33
1
4

Fig. 3. Schematic cross section of the sample under study: (
1) SiC
x
regions, (2) regions of an
intermediate Si-SiC
x
layer, (3) double-diffraction regions, (4) a through hole, and (5) the area

С for 30 min: rings – SiC, point reflections – Si, bright regions
– SiC
1.4
layer, dark regions – c-Si. a) SiC
1.4
region; b) SiC
1.4
regions + intermediate Si– SiC
1.4

layer + с-Si: (
1) SiC
1.4
regions, (2) regions of an intermediate Si– SiC
1.4
layer, (3) double-
diffraction regions. a) b) c)
1
4
3 2
1

Fig. 5. Electron diffraction patterns and microstructure (×50000) of the SiC
0.95
layer after
annealing at temperature 1200
0

1
2
3

Fig. 6. Electron diffraction patterns and microstructure (×50000) of the SiC
0.7
layer after
annealing at temperature 1200
0
С for 30 min: rings – SiC, point reflections – Si, bright regions
– SiC
0.7
layer, dark regions – c-Si. a) SiC
0.7
region; b) SiC
0.7
regions + intermediate Si–SiC
0.7

layer + с-Si, c) SiC crystallites in dark-field image regime: (
1) SiC
0.7
regions, (2) regions of an
intermediate Si–SiC
0.7
layer, (3) double-diffraction regions
The value of D⋅d of device, calculated from the point patterns of the silicon lattice, was equal
D⋅d = (3.53±0.01) nm⋅mm. The diameters D of the most intensive rings of SiC
1.4
pattern were

−10 nm (Akimchenko et al., 1977, 1980,
Calcagno et al., 1996). It should be noted that by X-ray diffraction the average crystallite size
was determined. In the layer SiC
0.7
due to the high concentration of carbon atoms one can
expect a large number of stable clusters that hinder the crystallization process. In this case a
significant prevalence of nanocrystals of a few nanometers sizes can be expected that are
making a major contribution to the value of the average grain size. Small nanocrystals
should give a reflection of the size in hundredths and tenths of a millimeter on the pattern, it
is difficult to distinguish them and they are observed as a bright diffuse background
between the larger crystallites (Fig. 6c).

Silicon Carbide – Materials, Processing and Applications in Electronic Devices

78
Transition layer with a lower concentration of carbon between the Si substrate and SiC
0.7
is
not uniform. It can be assumed that excess silicon atoms that are between the large SiC
grains, in the process of recrystallization are combined with the substrate, forming a
sawtooth SiC–Si structure (Fig. 6b).
3.3 Investigation of the structure by X-ray diffraction
After annealing of the SiC
1.4
layer in vacuum 10
-4
Pa at temperatures 1200 and 1400°С in the X-
ray diffraction patterns the intensive lines of polycrystalline silicon carbide are observed (Fig.
7a, b). There are also weak lines of polycrystalline silicon, apparently, from the transition layer
“Si–SiC

b)
0,2
0,4
0,6
0,8
1,0
1,2
10 30 50 70 90 110
2
θ
, de
g
ree
Si
(
111
)
Si (220
)
SiС (111
)
SiС (200
)
SiС (220
)
SiС (311
)
0,6
0,8
1,0

(c) after annealing at 1400°C and processing by glow discharge hydrogen plasma for 5 min.
An absence of SiC crystallites and the corresponding X-ray lines at high annealing
temperatures 900–1100°C indicates a low ability of atoms to diffusion in the layer SiC
1.4
. This
may be caused by high concentrations of stable double and triple bonds like Si=C, Si≡C,
C=C, C≡С, Si=Si, Si≡Si and strong clusters that prevent the diffusion of atoms in the layer
and are decomposed at temperatures of 1200°C and above. An attempt was made to use
processing by the hydrogen glow discharge plasma for the modification of the structural
properties of the SiC
х
films. For this purpose, the SiC
1.4
layer annealed at 1400ºC was treated
by the hydrogen glow discharge plasma (27.12 MHz, 12.5 W, 6.5 Pa, 100°С, 5 min).
After treatment by H-plasma is not taken place the complete destruction of the SiC crystallites
(Fig. 7c). However, the intensity of X-ray lines of SiC decreased in comparison with the
corresponding lines on the diffraction pattern before treatment (Fig. 7b). This demonstrates the
The Formation of Silicon Carbide in the SiC
x
Layers
(x = 0.03–1.4) Formed by Multiple Implantation of C Ions in Si

79
destructive effect of plasma on the structure of the β-SiC crystallites, although the capacity of
plasma was only 12,5 W, and SiC on the hardness scale, according to Knoop holds a high place
(SiC, 2480) after diamond (C, 7000), boron carbide (B4C , 2750), and aluminum boride (AlB,
2500). In addition, the processing by H-plasma for 5 minutes led to complete disappearance of
the lines Si(111) and Si(220), reflected by the crystallites of silicon in the transition layer “film-
substrate" ("SiC-Si"). It follows that the effect of hydrogen plasma propagates on the entire
d)
c)
0
0,1
0,2
0,3
0,4
0,5
0,6
10 30 50 70
2
Η
, degree
I
, ar
b
.un
.
Si (111)
Si (220)
SiС (111)
SiС (200)
SiС (220)
SiС (311)
b)
а)
0,6
0,7

at 1100ºC (Fig. 9a) indicates on intensive process of crystallization of β-SiC due to lower
content of stable clusters in comparison with the layers SiC
1.4
, SiC
0.95
and SiC
0.7
. High
amplitude of β-SiC peaks after annealing at 1250
о
С confirms this assertion. In diffraction
pattern of SiC
0.12
layer after implantation a broad diffuse line of amorphous silicon Si(111)
(θ = 14.3º) is observed (Fig.10a). After annealing at 800ºC a narrowing of this line, at 900ºC -
a sharp narrowing of the line Si(111) and the appearance of Si(220) and Si(311) lines of
polycrystalline Si phase, as well as two weak lines of β-SiC, at 1250ºC - a decrease of the
integrated intensity of Si lines, are observed.

0,4
0,6
0,8
1,0
1,2
1,4
1,6
10 20 30 40
θ
,
de

ree
I, arb.un
.
Si (111
)
SiC (111
)
Si (220
)
SiС (220
)
Si (311
)
а)

Fig. 9. X-ray diffraction patterns of the SiC
0.4
layer annealed for 30 min at (a) 1100ºС and (b)
1250°C.
The implantation of carbon causes the formation of weakly ordered set of randomly
oriented Si regions of size ~1.5 nm. Temperature dependence of the average crystallite size
Si (Fig. 11a, curve 1) and SiC (curve 2) in plane (111) shows that at low temperatures the
curve 1 is characterized by slow growth in the size of weakly ordered Si nanocrystals. Their
transformation into well ordered Si crystallites at 800ºC is taken place. Average size of Si
crystallites is 47 nm after annealing at 1250ºC. The formation of β-SiC crystallites prevent a
complete recrystallization of layer at this temperature.
The integrated intensity curve of Si(111) line, which is proportional to the phase volume,
has three sections in the temperature ranges of 20– 800ºC, 800−900ºC and 900−1250ºC (Fig.
11b). In the range of 20-800ºC I
int

, de
g
ree
Si (111)
Si (220)
Si (311)
Si
(
422
)
Si
(
511
)
Si
(
331
)
Si (531)
Si
(
matrix
)
SiC (111)
SiC (311)
SiC (220)
SiC (422)
SiC (331)
SiC (400)
0,0

I, arb. units
Si (111)
Si (matrix)
a) b) c)
θ, degree
θ, degree
θ, degree

Fig. 10. X-ray diffraction patterns of the SiC
0.12
layer (a) before and after annealing for 30 min
at (b) 1100ºС and (c) 1250°C.

0,00
0,04
0,08
0,12
0,16
20 300 600 900 1200
Temperature,
o
С
I
int
, arb.un.
2
3
1
b)
0

and phase volume, as well as the average grain size as a function of annealing temperature
(a) and scheme of Si and SiC crystallite formation in this layer (b). The diagram is based on
the curves in Fig. 11a, b for SiC
0.12
. Regions "amorphous Si-C mixture” and "c-Si" are formed
adhering to the relation: "amorphous Si-C mixture” + "c-Si" = 100% − (“poly-SiC” + “poly-
Si"+"amorphous -Si"). Fig. 12. Crystallization of the SiC
0.12
layer: (a) the phase volumes at various annealing
temperatures and (b) the formation of crystalline Si and SiC crystallites in the temperature
range 900−1000°C.
In the range of 800−900°C the increase in both size crystallites Si (from 2.2 to 4.7 nm) and the
volume of polycrystalline phase of Si (Fig. 11a and 12) is observed. This is due to the
formation of SiC crystallites at 900°C in the regions of carbon accumulation and the joining
of excess silicon atoms to Si crystallites. So, k is increased: k
900
≈ 82%. This leads to a decrease
in Si−C-mixture volume (Fig. 12). After annealing at 1000ºC and 1250ºC, an increase in the
phase volume and crystallite sizes of β-SiC as well as a decrease in k due to recrystallization
of regions near the Si substrate, are taken place: k
1000
≈ 73% and k
1250
≈ 46%. Si crystallite
sizes increased almost 10 times (up to 47 nm). Probably, there is destruction of defective
silicon crystallites and the uniting of their atoms into crystallites with a perfect structure or
with the substrate. As can be seen from the diagram, after annealing at 1200°C in the layer

0,75
10 30 50 70 90
2
θ
, degree
I, arb.units
Si (111)
.
Si (220)
.
SiС (111)
.
SiС (200)
.
SiС (220)
.
SiС (311)
.

Fig. 13. X-ray diffraction pattern of the SiC
0.12
layer after annealing at 1400°C for 30 min.
The X-ray diffraction results are in accordance with the data of Auger electron spectroscopy,
whereby the concentration of carbon atoms in a layer is N
C
/N
Si
= 0.12/1 (Fig. 2). Then the
maximum possible ratio of atoms, forming part of SiC and Si, will be: N
SiC

disappearance after annealing at 1250ºC due to recrystallization of the layer.
For comparison, in Table 3 the grain sizes of silicon and silicon carbide in plane (111) in
layers Si(111) and β-SiС(111) are given. Dimensions of weakly ordered regions in SiC
0.03

layer, contributing to the intensity of the Si(111) line, exceed the same values for SiC
0.12
. This
is associated with a lower concentration of carbon and larger volume of Si regions with
extremely low concentrations of carbon. At temperatures of 1200 and 1250ºC a decrease in
the average crystallite size up to 10 nm is observed, which is associated with the process of
recrystallization near the substrate.