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OCR for page 15
2
State of the Art of Wide Bandgap Materials
This chapter surveys the state of the art for the three
major wide bandgap materials for high-temperature
semiconductor devices: silicon carbide (SiC), the nitrides,
and diamond. This chapter is not a comprehensive
examination of all the properties of the different materials,
but does examine closely those properties related to high-
temperature operation. The intrinsic properties of the wide
bandgap materials versus those of the more common
silicon and gallium arsenide (GaAs) materials are
compared in Table 2-1.~ Although silicon and GaAs are
not considered in this chanter because of their expected
-, devices and
interconnects of these materials are discussed in
Appendices A and B. respectively.
limited high-temperature applicability
SILICON CARBIDE
Materials Description and Properties
Of the wide bandgap materials, SiC is by far the most
developed. The earliest reported recognition of the silicon-
carbon (Si-C) bond is by Berzelius in 1824. SiC has been
produced in the United States since 1891 when Eugene G.
Acheson (1893) of Monongahela City, Pennsylvania,
melted a mass of carbon and aluminum silicate by passing
a current through a carbon rod immersed in the mixture.
A variety of vapor-transport furnaces have been used in
this century to grow boules of single-crystal SiC. In
addition, high-purity homo-epitaxial single-crystal films of
SiC have been grown in both horizontal and vertical
chemical vapor deposition (CVD) reactors.
~ Table 2-1 was developed for comparative purposes using the data that
was available during the course of this study. This table should not be
considered a definitive tabulation of the properties of these materials,
.. ~ . ~ . ~ ~
since new, more accurate data are constantly being accumulated for most
of these materials.
15
Moisson reported in 1904 and 1905 that hexagonal
crystals of SiC were present in meteoritic specimens from
Canyon Diablo, Arizona. Naturally occurring SiC was
viewed as exclusively of extraterrestrial origin until 1957.
However, SiC has recently been discovered in alluvial
sands and in Kimberlite breccia in a number of locations
on the earth.
SiC forms in a variety of crystal structures, termed
polytypes, of which over 175 have been described in the
literature (Verma and Krishna, 1966; Pandy and Krishna,
19831. Only simple polytypes are of interest for SiC
devices. Their basic crystallographic stacking sequences
and most common notations are illustrated in Table 2-2
(Verma and Krishna, 19661. The optical properties of SiC
do not differ very much from polytype to polytype
(Figure 2-1~.
To better understand SiC, a brief discussion of
electronic band structure is warranted. Band-structure
calculations for SiC have been made for the past 30 years,
but theorists have concentrated on the zincblende 3C-SiC
polytype and the wurtzite 2H-SiC structure since the other
polytypes are much more complicated due to their much
larger unit cells. The accuracy of such calculations has
recently been considerably improved and currently there
is a sizable effort to work on the band structures of 4H-,
6H-, and l5R-SiC. Early band-structure calculations of
3C and 2H are shown in Figures 2-2 and 2-3 to provide
a qualitative "feel of the neighborhood" where the maxima
in the valence band and the minima in the conduction
band are likely to be located. Since both 3C-SiC and 2H-
SiC are ircdirect-gap semiconductors, it is reasonable to
assume that all polytypes are indirect-gap semiconductors.
Indeed, experiment has verified that in addition to 3C-
and 2H-SiC, 4H-, 6H-, 8H-, 15R-, 21R-, 27R-, and 33R-
SiC are also indirect-gap semiconductors. Figure 2-4
summarizes the experimentally observed excitor bandgaps
OCR for page 16
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OCR for page 17
State of the Art of Wide Bar~gap Materials
TABLE 2-2 Notations for Selected SiC Polytypes
Ramsdell
Notation
3C
2H
Stacking
Sequence
...ABC...
...AB...
Zhdanov
Notation
11
4H . ABAC 22
6H . ABCBAC 33
15R
. ABCBABCACBCABAC (32)3
and their temperature variation. Experiment has also given
an estimate of the binding energy (27 meV) of the excitor
in 3C-SiC. Assuming that this value will not be very
different in the other polytypes, the actual bandgap, KG,
can be estimated by adding 27 meV to the known value of
ant
no
JO
o
1n-2
1 no
10-4
- (a)
Face
plane
- (2110)
o-1 _
_
_ _
_
_
i 1111"1 ' ' '11' 1''1 1 1 11""1 1 1 1]
~| E
no ~ 7! it/
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v k
I! o/
) Plasma /
J Frequency /
1 ~
~ two ~ ED
10-5
10-l 10°
, , . I, , . . .
1o1
Wavelength (pLm)
FIGURE 2-1 Average values of Me optical constants of SiC from We
vacuum ultraviolet to Me middle infrared. NOTE: n0 = index of
refraction; Lo = extinction coefficient.
17
the excitor bandgap, EGX. Estimates of room-temperature
values of both EG and EGX are given in Figure 2-4. The
thermal conductivity is shown in Figure 2-5.
The electrical properties in the various polytypes can
be very different because the actual conduction-band
minima in the various polytypes will not be in exactly the
same positions in the Brillouin Zone. In addition, there is
the extra complication of having a different number of
nonequivalent sites in different polytypes as a consequence
of different size unit cells. This is illustrated for the donor
nitrogen in Table 2-3. SiC may be doped e-type with
nitrogen up to at least 10~9 cm~3. The acceptors aluminum
and boron can be used to dope SiC p-type to at least
5 x 10~8 cm~3. Nitrogen is difficult to keep out of the
growth process, and at present unintentional
concentrations of nitrogen in the range of 10~4 cm~3 are
found in the best epitaxial films. This is sufficiently low
not to interfere with current device fabrication.
Deep electronic states due to scandium (Tairov et al.,
1974), titanium (Patrick and Choyke, 1974), and
vanadium (Mater et al., 1992) have been studied in some
detail in various polytypes of SiC. Other deep states,
termed Do and D,,, due to implantation or radiation
damage have also been widely studied. Many other
impurity defect complexes have been observed during
annealing of irradiated samples from 0-2000 °C.
Methods of Fabrication
Bulk Growth
The commercial potential of SiC semiconductor
technology has been enhanced by recent significant
progress in the growth of large single-crystal SiC boules.
OCR for page 18
Materials for High-Temperature Semiconductor Devices
20
,2 1
10
8 r15
6
4 r,
_ 2
a) 0 F15
>`
Al
UJ
-10
-20
X3
~1
x1 ~ r,s
-~
~ 1
1 ~
I; K X ~r A L
/~L
r
k
A1 /
FIGURE 2-2 Calculated band structure of 3C-SiC. SOURCE: Based on
Hemstreet and Fong (1974).
For many years, the lack of suitable SiC crystal-growth
processes inhibited the commercialization of this
promising semiconductor material. There are two
properties of SiC that make the growth of bulk single
crystals more difficult than that of silicon. First, it does
not melt under any reasonably attainable pressure; rather
it sublimes at temperatures above 1800 °C. Thus,
conventional growth-from-melt techniques (e.g., as for
silicon or GaAs) cannot be used for SiC crystal growth.
Second, different polytypes with different electronic
characteristics can grow under apparently identical
conditions (Knippenberg, 19631. A completely satisfactory
model for the formation of the various polytypes does not
exist. Despite these difficulties, major progress has
recently been made in SiC boule growth. The diameter of
commercially grown, single-crystal boules is typically 30
mm, and prototype boule diameters have exceeded 50
mm.
Currently, there is interest in at least five of the SiC
polytypes: 3C-SiC, 2H-SiC, 4H-SiC, 6H-SiC, and 15R-
SiC. Boules of 4H, 6H, and 15R have been grown, and
wafers from 4H and 6H boules are commercially
available. No s~gn~ficant-s~zed boules of 3C have been
18
reported. To date, 2H has only been grown in the form of
63 small, millimeter-sized needles.
' There are several key review papers that discuss the
growth of bulk SiC single crystals (Knippenberg, 1963;
Tairov and Tsvetkov, 1983; Powell and Matus, 19891.
This section summarizes some of the early work and
L3 describes recent developments for which information is
publicly available. Much of the current technology is
considered to be proprietary and has not been published.
Although growth-from-solution techniques have been
L3 tried, the most successful growth techniques are based on
the sublimation of SiC.
Background. Prior to the mid-1950s, small
, hexagonally shaped SiC platelet crystals were available
through the industrial Acheson process for making
abrasive material (Knippenberg, 1963~. In 1955, Lely
developed a laboratory sublimation process for growing
crystals that were much purer (Lely, 19551. In the Lely
process, a hollow cavity was formed inside a charge of
polycrystalline SiC. The charge was heated to about 2500
°C in a graphite tube furnace at which point the SiC
sublimed and condensed on slightly cooler parts of the
cavity. Growth took place on a thin, porous graphite
cylinder that formed the wall of the cavity. Nucleation
was uncontrolled and the resulting crystals were randomly
sized, hexagonally shaped or-SiC platelets. These platelets
often exhibited a layered structure of various (x polytypes.
The predominant polytype (generally more than 75
6
3
2
1
o
. ~
-M-am t
_A56 1 r
\
.. R L U M
At:
W:
M1
r
< nit
/ Band ~
it/ \
~ 1/
~ r ~
/
H3
As.6
Or\
,
1 K2
.3~:
S
k
FIGURE 2-3 Calculated band structure of 2H-SiC. SOURCE: Based
on Hemstreet and Fong (1974).
OCR for page 19
State of the Art of Wide Bandgap Matenals
~[it
3.3063.300 3.327
---em SiC
+~~~~~"~ 2H
'4H "
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OH
`;33R
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a'
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x
(4.2K) (RT) (8T)
EGX
(2H) 3.33C
(4H) 3.265
(OH) 3.023
(33R) 3.003
(15R) 2.986
(21 R) 2.853
c,
>`
~ (24R) 2.73 ~
o
c'
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(3C) 2.39O
i5R -
~ ~ ,-_ ' 2~1 R
8H`"
'24R~ ``
3C 2.30
"` 2.20
I I ,
0 200 400 600 800
Temperature (K) _
.3.235 3.262
3.20
3.10
3.00~2.995 3.022
·2.972 2.999
.2.957 2.984
2.90
28o~.82 2.827
·277
2.70 -2.71 2.74
2.60
2.40
.2.360 2.387
Energy Gap - EG(eV) = EGx(eV) + (BE)X
Exciton Binding Energy = ~ B E) x ~ 0.027 ev
FIGURE 2-4 Summary of Me experimentally observed excitor bandgaps
and Weir temperature variation for Me different SiC polylypes.
percent) was 6H, followed by 4H and 15R. Although
much was learned about SiC from investigations of these
crystals over the next 30 years, the process was not
suitable for commercial development of SiC.
In the 1970s, Tairov and Tsvetkov (1978, 1981)
developed a modification of the Lely process (now
commonly called the modified sublimation process, or the
modified Lely process) in which growth occurred on a
seed crystal. Although some research groups have been
somewhat slow in adopting this process, it is now being
developed in many labs in Russia, Germany, Japan, and
the United States.
The basic elements of the modified sublimation
process are shown in Figure 2-6, which is a schematic
diagram of the configuration used by Westinghouse.
Nucleation takes place on a SiC seed crystal located at
one end of a cylindrical cavity. A temperature gradient is
established within the cavity such that the polycrystalline
SiC is at approximately 2400 °C and the seed crystal is at
approximately 2200 °C. At these temperatures and at
reduced pressures (argon at 200 Pa), SiC sublimes from
the source SiC and condenses on the seed crystal. Growth
rates of a few millimeters per hour can be achieved.
Current Status. Cree Research Incorporated of
Durham, North Carolina, is the only commercial source
in the world of SiC wafers produced from boules. Cree is
currently selling 30-mm-diameter wafers of both 4H- and
6H-SiC. Other companies and institutions, known to be
producing SiC boules for internal consumption, include
Westinghouse, ATM, Siemens, Sanyo, Nippon Steel,
280 Kyoto University, and Kyoto Institute of Technology.
Both Cree and Westinghouse have demonstrated boules
(and wafers) of up to approximately 50 mm in diameter.
Despite the fact that SiC is extremely hard (between
sapphire and diamond in hardness), techniques for cutting
and polishing wafers are currently in use. However, the
capability is far short of that for silicon. As a result, the
polished surface of commercial SiC wafers contains many
scratches and defects. Some defects introduced into the
wafer by cutting and polishing can be removed by suitable
pregrowth (i.e., prior to epitaxy) etching processes
(Powell et al., 1991~.
Currently, SiC boules (and the commercially available
wafers) do contain defects and impurities. One of the
most significant defects is a distribution of tubular voids,
called micropipes, in the order of a micrometer in
diameter (Koga et al., 19921. The micropipes are oriented
with respect to their long axis and are approximately
parallel to the crystal c-axis; density is typically several
hundred per square centimeter. In addition, wafers contain
line defects (dislocations) intersecting the surface with a
density of 104 to 105 cm~2. The most common background
impurities are nitrogen, aluminum, boron, and metals that
can act as deep-level traps.
It has been shown that the micropipes can cause
premature reverse breakdown in pen junctions (Neudeck
and Powell, 19941. Evidence shows that microplasmas
form in the micropipe at reverse voltages of several
hundred volts. We current micropipe density limits the
area of high-voltage devices to about 3 mm2; hence, this
defect must be significantly reduced before high-power
devices are practical. Several theories have been proposed
19
OCR for page 20
1o2
_ 1o1
y
5
-
Is
E
a,
1
10-2
1/
~ _
. _
Materials for High-Temperature Semiconductor Devices
fi:
l
I 1 1 1 ~1 1 1 1 1 ~
10° 1o1 1o2 103
Temperature (K)
FIGURE 2-5 Thermal conductivity of two single crystals of SiC.
SOURCE: Adapted from Slack (1964).
to explain the formation of micropipes. One theory is
based on the presence of contaminant particulates during
nucleation and boule growth (Yang, 19931. Another
theory is based on the presence of super-screw
dislocations (Wang et al., 1993~. In this latter theory,
hollow cores would form to relieve stress caused by screw
dislocations.
Progress is being made in reducing the density of
micropipes. In a recent paper, growth of boules in the
(1010) direction significantly reduced the formation of
micropipes (Takahashi et al., 19941. However, the
dislocation density is very high in these crystals.
Researchers at Cree have reported (J.W. Palmour,
personal communication, 1994) that the density of
micropipes has been reduced significantly in the last year.
It should be noted that the research team directed by
Professor Yu Vodakov at the Ioffe Institute in St.
Petersburg, Russia, have produced small single polytype
SiC boules (1.5 cm diameter and 7 mm thick) that are
claimed to have no micropipes (Y. Vodakov, personal
communication, 1994~.
Another impediment to wide use of SiC technology is
the cost of wafers. At present, there is only one
commercial supplier of wafers in the world. The current
price per 30-mm-diameter wafer is more than $1,000.
This high price can be expected to drop considerably
during 1995 as other sources enter the market. The
primary reason for this price being lower than GaAs is
that both silicon and carbon are 100 times cheaper than
gallium.
Epitaxial Growth
Semiconductor-quality a-SiC epitaxial films can now
be grown routinely on or-SiC wafers by CVD. In addition,
in situ CVD doping processes can produce both e-type
and p-type epitaxial films with net carrier concentrations
from the 10~4 cm~3 range to greater than 10~9 cm~3. This
technology, which has largely been developed in the last
few years, has allowed the development of SiC devices
with record-setting performance.
Background. The growth of epitaxial SiC films has
many similarities with the growth of epitaxial silicon;
however, it has only been recently that the differences in
growth processes have been appreciated. While
conventional semiconductors are grown at approximately
two-thirds of their melting temperatures, these
temperatures are not practical with wide bandgap
materials. For this reason, the substrate temperature
cannot be used to assure that all components of the
activation energy have been exceeded. In addition, only
one crystal structure can be produced in silicon, whereas
many crystal structures are possible in SiC. Thus, the
polytypic structure of the film must be controlled during
formation. The factors that control SiC structure are the
crystal orientation and perfection of the substrate. The
presence of defects and contamination can also
significantly affect the resulting structure. In this report,
the term "homo-epitaxy" is used for growth in which the
film and substrate are the same polytype, and "hetero-
epitaxy" is used when the SiC polytype of the film is
different from the substrate. With respect to doping, the
incorporation of dopants is dependent on the ratio of the
silicon and carbon sources during the growth process and
also on the crystal orientation.
20
OCR for page 21
State of the Art of Wide Bandgap Materials
TABLE 2-3 Exciton Binding, Nitrogen Ionization, and Valley-Orbit Splitting Energies and Effective Mass for SiC Polyn,rpes
Exciton Nitrogen Ionization Energy Valley-Orbit Effective
Binding to 4D Splitting Mass
ED (meV) EvO(meV)
EBX
SiC L: {) (PL) | (Haynes) (JR) (2EL) (Hall) | (JR) (ERS) |
ml, mll
(Cyclotron
. resonance)
3C 10 57 53.6 20-47 - 8.37 0.247,
0.677
7 47 52.1 45 7.6 0.42,
20 96 91.8 - 100 - 0.29
6H 16 81 81.0 85 12.6 13.0
31 136 137.6 125 60.3 0.42,
32 140 142.4 62.3 2.0
15R 7 47 49.3
9 54 59.6 53
19 91 99
20 96 -
Techniques used to produce epitaxial SiC films
include CVD (Davis et al., 1991), the "sublimation
sandwich" process, and liquid-phase epitaxy (Ivanov and
Chelnokov, 1992~. Homo-epitaxial growth of cx-SiC on a-
SiC substrates has been achieved by all three techniques.
The lack of 3C-SiC substrates has led to a variety of
hetero-epitaxial processes to produce 3C-SiC epitaxial
films. The 3C-SiC polytype has been grown on silicon,
TiC, and o`-SiC substrates. These processes are examined
in the following sections.
Cal) of ~x-SiC Epitaxial Films. For both ax- and 3C-
SiC, the CVD process is the current method of choice
because. of the three techniques. it yields better films at
- 7
Al ~ ~ _
the lowest temperature. It is also adaptable to commercial
production.
A typical SiC CVD growth chamber, shown in Figure
2-7, is similar to chambers used for silicon (Powell et al.,
19871. The quartz chamber is water-cooled because
growth temperatures are generally higher than those used
for silicon epitaxy. The substrates are heated by an
inductively heated SiC-coated graphite susceptor.
Hydrogen is used as a carrier for various process gases.
Prior to growth, the substrates are frequently subjected to
21
an etch with hydrochloric acid (HC1) to reduce defects
and contamination. Silane (SiH4) and propane (COHN) can
be used as sources of silicon and carbon during growth.
Important system parameters for growth include the
growth temperature, flow rates of the various gases, and
the silicon/carbon ratio in the gas. Important substrate
parameters include the orientation and polarity of the SiC
substrate. Typical growth rates are in the 1- to 5- ~mlh
range. In situ doping is achieved by adding nitrogen or
phosphorous for e-type and aluminum
(trimethyl aluminum , TM A) or boron (diborane) fo r p -type
material. Particular growth and doping processes are
discussed.
SiC Epitaxy in the Claris Direction. An important
discovery in SiC epitaxy was that the crystalline
orientation of the growth surface is an important growth
parameter. In the past, much of the growth was carried
out on the "as-grown" (0001) surface (the basal plane) of
Lely crystals; that is, growth was in the c-axis direction.
The (0001) SiC crystals with polished surfaces have
vicinal (0001) orientations, that is, the growth surface
may be tilted slightly "off-axis" with respect to the (0001)
crystallographic plane. The size of this tilt angle can have
OCR for page 22
Materials for High-Temperature Semiconductor Devices
temperatures because deposited atoms cannot migrate to
the steps on large terraces. Also, mobility of deposited
atoms is reduced at these lower temperatures and
deposited atoms do not reach the steps.
Work at the NASA Lewis Center demonstrated that
homo-epitaxy of 6H-SiC on vicinal (0001) 6H-SiC can be
achieved at 1450 °C with tilt angles as low as 0.1°
(Powell et al., 1991~. As a consequence of this result, it
was proposed that the cause of the 3C-SiC nucleation was
due to defects and contamination on the growth surface.
By a suitable pregrowth etching process, the defects and
contamination were reduced or eliminated. In effect, there
is a competition between defects and surface steps. At
sufficiently large tilt angles (high step density), homo-
epitaxy will occur even in the presence of defects. At low
tilt angles (low step density), any defects that are present
will dominate and act as nucleation sites for 3C-SiC.
Thus, growth must occur at atomic steps if homo-epitaxy
of 6H-SiC is to be achieved. In addition, suitable
pregrowth etches can be effective in reducing or
eliminating defects caused by cutting and polishing the
SiC substrate.
Homo-epitaxial SiC films on vicinal (0001) SiC
substrates have been obtained with the 4H-, 6H-, and
15R-SiC polytypes. These films exhibit a variety of
surface features that include hillocks and depressions.
Structural defects that occur include the micropipes and
dislocations that propagate from the substrate into the film
(Powell et al., 19941. Although excellent devices have
been fabricated using these films, much work remains to
improve the surface morphology and to reduce the defect
density.
The electrical quality achievable in SiC epitaxial CVD
films was significantly improved recently by the
development of the "site-competition epitaxy" process by
Larkin et al. (1993) at the NASA Lewis Center. In this
process, the incorporation of nitrogen and aluminum into
a SiC epilayer grown on a silicon-face vicinal (0001)
plane is controlled by setting the silicon/carbon ratio in
the precursor gases to appropriate values. The nitrogen
donor atoms that reside on carbon sites in the SiC crystal
lattice compete with carbon atoms during growth.
Increasing the carbon concentration (i.e., decreasing the
silicon/carbon ratio) decreases the nitrogen incorporation
in the epilayers. On the other hand, the aluminum
acceptor atoms that reside on silicon lattice sites compete
with silicon atoms during growth. Increasing the silicon
~ ~ ~ ~[~]
~ ~ 1 ~1~:
3
3
~ Seed
I)
3 Crystal
~ Growth
=> Cavity
3 sic Charge
Crucible
-Thermal
Insulation
-; Quartz
Tubes
1 ~
Water
Cooling
FIGURE 2-6 Schematic showing Me basic elements of He modified
sublimation process. SOURCE: Hobgood (1993). Courtesy of
Westinghouse, Inc.
a dramatic effect on the structure of an epitaxial film. In
subsequent discussions in this report, SiC substrates
having tilt angles of about 3° are referred to as being
"off-axis" and substrates with tilt angles of less than 0.5°
are "on-axis." The polarity (i.e., silicon face or carbon
face) of the substrate is also an important parameter.
In sublimation sandwich growth, it was found that
homo-epitaxy of the various polytypes was enhanced if the
growth surface of the substrate was polished off-axis by
a few degrees from the (0001) basal plane (Tairov and
Tsvetkov, 19831. The research team of Matsunami at
Kyoto University discovered that the CVD growth
temperature required for producing good-quality 6H-SiC
epilayers on 6H-SiC substrates could be reduced from
about 1750 °C to about 1450 °C if the growth surface
was off-axis by a few degrees from the (0001) plane
(Matsunami et al., 1989~. They called this growth "step-
controlled" epitaxy because growth occurs at steps on the
off-axis surface. The stepped surface automatically
provides the stacking sequence of the substrate polytype.
Hence, homo-epitaxy takes place. 3C-SiC was found to
grow at small tilt angles (e.g., less than 1.5°) or at low
22
OCR for page 23
. ~ ,'~¢,,, ~ s ',.' ~ '.
a ~ B ~ i ~ 9 ~ ~ S ~ ~ ~ S . ~ D 8 ~ ~ ~ D - D - ~ n
~ A ' · ~ PHI ~ ~
s D ~ a 7 R ^3 o ~ ~ D ~ ~D ~ ~ 7 S _ ~ 7 S
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it.
OCR for page 24
Materials for High-Temperature Semiconductor Devices
To eliminate the problem of the large lattice
mismatch, titanium carbide (TiCX) with a lattice match
within 1 percent was investigated (Parsons, 1987~.
Somewhat improved growth of 3C-SiC films was
reported, but great difficulties in producing defect-free,
single-crystal TiCX has hindered its use as a substrate for
SiC growth.
In a previous section, it was pointed out that 3C-SiC
generally grows on vicinal (0001) ~x-SiC with small tilt
angles if there is contamination or defects on the growth
surface. Unfortunately, 3C-SiC films grown in this
manner typically have a defect known as double-
positioning boundaries. This defect arises because there
are two possible orientations of the 3C-SiC film that can
nucleate on an a-SiC substrate; these two orientations are
rotated 180° about the c-axis with respect to each other.
When nuclei with both orientations occur on the substrate,
the intersection of domains with different orientations are
not coherent and they form double-positioning boundaries
that are electrically and chemically active.
Recent work at Kyoto University has shown that the
density of double-positioning boundaries in 3C-SiC films
grown on vicinal (0001) 15R-SiC is less than that found
in 3C-SiC films grown on 6H-SiC (Chien et al., 1994~.
Chien and colleagues presented a model that predicts 3C-
SiC films that are tens of micrometers thick and grown on
(0001) 15R-SiC should be free of double-positioning
boundaries. Unfortunately, the stacking-fault density
appears to be very high in these 3C-SiC films.
Another approach investigated at the NASA Lewis
Center is to limit the epitaxial growth areas to small
mesas on vicinal (0001) 6H-SiC substrates and then limit
the nucleation of 3C-SiC to the highest atomic planes on
the mesa (Powell et al., 1991~. With nucleation limited to
a very small region on each mesa, 3C-SiC films will grow
laterally and will subsequently cover the mesa with a
double-positioning boundary-free 3C-SiC film. This
approach has been successful obtaining double-positioning
boundary-free 3C-SiC films on 1 mm2 mesas. These films
also have a lower stacking-fault density than previously
reported 3C-SiC films grown on SiC substrates.
Combining this technique with the site-competition epitaxy
process for doping SiC epitaxial films, p-njunction diodes
with reverse breakdown voltages exceeding 300 V were
fabricated (Neudeck et al., 19931. This breakdown voltage
is four times that of any previously reported 3C-SiC
diode.
24
Other Epita~cial Processes. The sublimation sandwich
process (Ivanov and Chelnokov, 1992) is similar to the
modified sublimation process. In the sublimation sandwich
process, the substrate is placed near a solid SiC source
that is sublimed at temperatures greater than 1800 °C.
The resulting vapor condenses on the substrate that is held
at a slightly lower temperature. The high temperature
required by this process is its main disadvantage.
In the liquid-phase epitaxy technique (Ivanov and
Chelnokov, 1992), the substrate is placed in liquid silicon
that is saturated with carbon at a temperature in the range
of 1500-1700 °C. If the temperature is lowered, SiC is
deposited from the supersaturated silicon solution onto the
substrate. In one version of this process, the liquid silicon
solvent is suspended by an electromagnetic field; this
"containerless" approach avoids contamination of the
solvent by a crucible. The higher temperature required
and the difficulty of control are disadvantages of this
approach.
Summary. Excellent epitaxial films of or-SiC polytypes
can now be grown on ~x-SiC substrates. Both e-type and
p-type films with net carrier concentrations from 10~4 cm:3
to greater than 10~9 cm~3 can be routinely achieved. The
growth of large-area epilayers that are free of micropipes
will only be possible when micropipe-free substrates are
available. In the future, it will probably be desirable to
reduce the growth temperature from the present 1450 °C;
this may be beneficial for some device fabrication
processes.
NITRIDE MATERIALS
There are four major nitride semiconductors and
several minor ones. The four major nitride
semiconductors are indium nitride (InN), gallium nitride
(GaN), aluminum nitride (A1N), and boron nitride (BN).
For high-power electronics applications, there is yet
another nitride (iron nitride) that, although not a
semiconductor, warrants attention. These materials are
composed of cations from Group III of the periodic table
and a nitrogen anion from Group V. They are often
referred to as III-N materials. A1N, GaAlN, and GaN
have been studied for some time, but due to the lack of
good single crystals, the electronic, optical, and physical
properties of single-crystal nitrides are not extensively
OCR for page 25
State of the Art of Wide Barcdgap Materials
known. Interest in the nitride materials has dramatically
increased with the recent introduction of bright blue light-
emitting diodes (LEDs) by Nichia, the successful growth
of better samples, and the accumulation of more precise
data (Strife and Morkoc, 1992; Choyke and Linkov, 1993;
Lin et al., 1994; Morkoc et al., 1994~.
Properties
The most intriguing aspect of the large bandgap
nitrides (i.e., A1N, GaN, and InN) is the fact that they
form a continuous alloy system with room-temperature
direct bandgaps varying from 6.2 eV for A1N to 3.45 for
GaN, to 1.9 for InN. In addition, there is a small lattice
mismatch ~ < 1 percent) between (wurtzite, 2H) A1N and
2H-SiC, and between cubic EN (cBN) and diamond. The
band structure of the hexagonal and cubic modifications of
A1N and GaN are given in Figures 2-8 and 2-9
(Lambrecht and Segall, 19921.
Boron nitride is the least understood of the nitrides.
Most work is directed towards the synthesis and
characterization of cBN as it is believed to exhibit an
indirect bandgap in excess of 6.4 eV. The relative
dielectric constant of cBN is 6.5 and its hardness is 4,500
kg/mm2 compared with 3,980 for SiC and 10,400 for
diamond (Davis, 1992~. Its thermal conductivity is
believed to be 1,300 W/m °C, or more than twice that of
SiC and about 60 percent that of diamond. Young's
modulus is 5.2 MPa compared with 4.0 MPa for SiC. The
Poisson ratio for cBN is 0.2 or equal to that of both
diamond and SiC. The thermal expansion coefficient of
cBN is 3.7 x 104/°C, which is the same as that of SiC but
greater than the 2.3 x 10~ of diamond. Unfortunately,
cBN has also been the most difficult to synthesize.
Aluminum nitride exhibits a direct bandgap of 6.2 eV
in its hexagonal form. Like diamond, A1N exhibits
negative electron affinity (Benjamin et al., 1994~. While
cubic A1N has recently been synthesized as a thin film on
cubic (3C) SiC, its bandgap has not been ascertained but
is believed to be somewhat less than 6.2 eV and is most
probably indirect. The thermal conductivity of
polycrystalline A1N is 3.0 W/cm °C at room
temperature-over twice that of silicon and 60 percent
that of SiC. Its relative dielectric constant is 10.0, or 85
percent of that of SiC. The best crystallinity reported to
date using X-ray 8/2-8 diffraction data is 90 arc-seconds
for films grown hetero-epitaxially on sapphire. While
25
there are reports in the literature of both n- and p-type
A1N having been synthesized, these reports are not recent
(Chu et al., 1967; Rutz, 1976~. Aluminum in A1N has an
affinity for oxygen and oxygen appears at a deep level in
A1N. Oxygen is typically found in AlN in concentrations
of 102° cm~3, rendering it extremely difficult to obtain AlN
with either p- or e-type conductivity.
There is currently a considerable amount of work
underway addressing the A1N doping issue. Alloys of AlN
and SiC have recently been made. These may not be true
alloys, however, as there is no measurable interdiffusion
at temperatures up to 1900 °C. Nevertheless, absorption-
band edge measurements on this alloy appear to track the
mole-fraction composition. The low mass of nitrogen
engenders A1N with a high optical phonon energy; for this
reason, the charge-carrier velocitY could be very hiah and
approach that of diamond.
·r ~
The lattice parameters for A1N are a = 3.112 A and
c = 4.982 A (293 K). The A1N linear thermal expansion
coefficient is a' = 5.27 x 10~ K-~; T = 20-800 °C; and
a 11 = 4.15 x 10~ K-~. The thermal conductivity of A1N is
k = 2 W/cm °C at room temperature. The density of
A1N is d = 3.244 g/cm3. Phonons for A1N are in the
frequency range of 895 cm~~ to 303 cm~~. The dielectric
constants for A1N are e(0) = 9.14 and c(~) = 4.84
(300 K).
Gallium nitride exhibits a direct bandgap of 3.5 eV in
its hexagonal form and apparently slightly less in its cubic
form. Its lattice constant is 3.189 A, or about 4 percent
greater than that of SiC. Its dielectric constant has been
measured at 8.9 to 9.5, or just less than that of SiC. It is
thought to have an effective electron mass of 0.20, but
this figure should not yet be taken as definitive. Along the
a-axis, the coefficient of thermal expansion is
5.5 x 10~ K-~. The thermal conductivity at room
temperature of GaN is 1.3 W/cm °C, nearly equal to the
1.45 W/cm °C of silicon and about three times higher
than GaAs. After two decades of research, both p-type
and e-type GaN have now been produced in hetero-
epitaxial thin films. In the best of this material (e.g., with
X-ray 8/2-8 diffraction data exhibiting a full width at half
maximum of about 27-28 arc-seconds; Plano et al., 1994),
acceptors freeze out at about 205 K and holes exhibit a
mobility of 450 cm2/V s. Electron mobilities of 1,200
cm2/V s at room temperature have been observed. Both of
these values exceed those of SiC but not of diamond.
With better crystallinity, these mobilities may perhaps be
OCR for page 26
Materials for High-Temperature Semiconductor Devices
a)
a)
Is
, ~
-
_/
1 ,3\
1 ,3
1 ,3_
2,4~
-
_/
1 ,3
_
_
_
1 ,3
_
.
-
_
L M K
(Wurtzite)
~ conduction
J\ Bands
2 \
6
~ ~-
//Val~ W.
and ~ .
3\ 3 ~
r
k
you
~: ~s6
M X r A
4 ~Z
l ~
3
(Zincblende)
I A\ I Conduction I /
L M K r M X
k
FIGURE 2-8 Band structure of hexagonal and cubic modifications of A1N. SOURCE: Based on Lambrecht and Segall (1992).
expected to further improve. The dielectric strength is
believed to be about equal to that of SiC and its computed
peak electron velocity exceeds 2 x 107 cm/s.
A large number of luminescent features have been
reported between 1.65 and 3.5 eV in GaN. These have
been attributed to a variety of impurities and defects.
However, there is a great deal of controversy in the
literature as to the various interpretations. A number of
articles have reviewed the literature of luminescence and
absorption lines in A1N and GaN (Strife and Morkoc,
1992; Choyke and Linkov, 19931. Although there is
currently great commercial potential for GaN optical
devices, the beneficial impurities and defects of this
26
r A
material for luminescent features are detrimental for high-
temperature device operation.
Indium nitride exhibits a direct bandgap of 1.9 eV
and an indirect bandgap only slightly higher. Its thermal
conductivity and most other properties have not yet been
definitively ascertained. Its lattice constant of 3.5 A
considerably exceeds that of SiC, A1N, and GaN. Double
heterostructures of GaN/InGaN/GaN currently exhibit the
brightest purple, blue, and blue-green LED s ever made.
The blue and blue-green devices are commercially
available from Japan and exhibit an operating efficiency
of about 2.7 percent.
OCR for page 27
State of the Art of Wide Bandgap Materials
1 ,3\
_
1 ,3
~ _
>O .1:,
_ ~ .
L1J 2,4~ .
~ .
1 ,3
-1 ,3
(Wurtzite)
~:
my_ at.
~
.; ~
1 1
1 1
. _
3 I_
_ _
~ _
1 _ _
3 3
/`Conduction ~
\ Band ~
2 \:
b ~/ 16~
N3/ /Valence \
,~ Band <\_
_
; - ~
3
'4=
l ~ ~
1
K r M X r
A
k
FIGURE 2-9 Band structure of hexagonal and cubic modifications of
The nitride materials, like diamond, are very difficult
to etch with liquid etchants. This is an active area of
research, however, and phosphoric acid and sodium
hydroxide have recently been found to work on both A1N
and GaN. Also like diamond, the nitride semiconductors
can be left exposed to the atmosphere at high humidity for
months at a time without becoming oxidized or otherwise
having their surface properties changed. Unlike GaAs, the
nitrides do not exhibit self-depleting surface states. For
this reasons devices employing pen junctions do not
exhibit high surface recombination velocities, which
should lead to a long laser-operating life and to extremely
27
,3
5,6
,3
1 ,3
1 ,3
(Zincblende)
154
<7CK
Conduction ~
Band /`
_ _
1 SIX_
Lop
':
. ~
. ,
. ~
Lo
Vale rice
/ Band \
L1
~ /
~ =_
L1
r,
M F< r
k
./
M X
GaN. SOURCE: Based on Lambrecht and Segall (1992).
V
tar
it,
X1~
1 K\
1 \
X/~`
L1
. _ . .
.=
1 -
l
1
r A
long charge-storage capability (e.g., millennia) and
extremely low-leakage devices suitable for applications
such as nonvolatile memories.
Crystal Growth
Very little work has been done in attempting to
synthesize boules of the nitrides. Japanese researchers
have synthesized small boules of cBN via high-
temperature, high-pressure processes and have even made
light-emitting pen junctions of the material. The material
was so contaminated by impurities, however, that the
absorption band edge was not very sharp and its bandgap
OCR for page 28
Materials for High-Temperature Semiconductor Devices
was difficult to ascertain. The pen junctions emitted light
in both the visible and in the ultraviolet. Most of the cBN
research in the United States has been directed at thin-film
synthesis on diamond. It is exceedingly difficult to obtain
films thicker than 20 nm that are not polycrystalline or
fractured. The only high-temperature, high-pressure
attempts to synthesize boules of GaN have been in Poland
(Perlin, 1993) however, the boules were very small.
growth has been attempted with A1N, but
these efforts nave generally not been successful to date.
Some boule
i. ^~ . ~
DIAMOND
Materials Description and Properties
Diamond has been admired as a jewel since antiquity
and has been studied for a very long time. In fact, Sir
Isaac Newton made measurements of the index of
refraction of diamond some time around 1665. Large
single crystals are found in nature, and synthetic diamonds
have been made in high-temperature, high-pressure anvil
machines for about 40 years. However, from a
semiconductor standpoint, only limited impurity and
defect control has been possible to date. Low-temperature
and low-pressure polycrystalline film growth has been
actively pursued in the last 10 years, but no high-quality,
single-crystal films have yet been obtained. The major Synthesis
properties of crystalline diamond are well understood, and
two excellent books chronicle the development of artifact
diamond and tabulate its known properties (Davis, 1992;
Spear and Dismukes, 1994~.
Diamond is an indirect-gap semiconductor, with the
lowest minima of the conduction band being located along
the delta axes (k = 0.76~1,0,011. The valence band
maximum has a structure that is common to all Group IV
semiconductors. There are three bands that are degenerate
at the It point when spin is neglected. A band calculation
by Chelikowsky and Louie (1984) is shown in Figure
2-10. The indirect energy gap at room temperature is 5.5
eV, and between 135 K and 300 K the variation of the
bandgap with temperature is given by -5 >< 10-5 eV/K.
The excitor binding energy, Ex, is about 80 med.
ma = 0.36 me, and ms0 = 0.15 ma. Hall mobilities have
been obtained for n- and p-type diamond as follows:
He = 2,200 cm2/V s(RT),and~h = 1,600 cm2/V s(RT).
Diamond is famous for its excellent thermal
conductivity. In the last few years single isotope diamond
has been produced, and it has a higher thermal
conductivity than natural diamond (i.e., 32 W/cm.K;
Anthony, 19941. Ordinary isotopic ratio diamond has a
thermal conductivity as shown in Figure 2-11. The
dielectric constant of diamond measured at 300 K is 5.5.
The lattice parameter a is 3.56683 A at 298 K. Linear
thermal expansion coefficients for various temperatures
are
-lxlo-6 K-1
-3x10-6 K-1
~4 x 10-6 K~1
-5 x 10-6 K-1
(300 K),
(600 K),
(900 K), and
(1200 K).
The density of diamond is 3.51525 g/cm3, as calculated
from the lattice constant. The second-order elastic
constants for diamond are On = 10.764 x 1O12 dyne/cm2
(296 K), C12 = 1.252 x 10~2 dyne/cm2 (296 K), and
c" = 5.774 x 1012 dyne/cm2 (296 K).
Methods of Synthesis and Characterization
Aside from the high-pressure, high-temperature
synthesized boules of diamond, virtually all diamond films
are grown by plasma-assisted methods in the presence of
an abundance of atomic hydrogen. The feedstock typically
consists of 99 percent H2 and 1 percent CH4. There are
many variations of this basic method. While conventional
(lower bandgap) materials are typically synthesized at
substrate temperatures approaching two-thirds of the
melting temperature, this is not possible with diamond and
some of the higher bandgap materials. The typical
substrate temperature of 900 °C cannot be used to assure
that feedstock species alighting on the growth surface are
fully "activated." The activation-energy threshold (Ea.) is
generally composed of three parts: (1) energy sufficient to
dissociate the molecule of radical, (2) energy sufficient to
chemisorb rather than physisorb the feedstock species, and
(3) energy sufficient to ensure that the adsorbed species
Effective masses have been measured and calculated
in high-quality bulk diamond crystals. The electron
effective masses are m' = 0.36 ma, and ml = I.4 ma.
The effective masses of the holes are mh = 1.08 mO, arrive at a proper lattice site. In diamond-film synthesis,
28
OCR for page 29
State of the Art of Wide Bandgap Materials
_ I
_ _ f
l _ /Valer ce Band ~
>
-
o,
LL
10
5
o
-10
-20
Kin
k
FIGURE 2-10 Band-structure calculation of diamond.
SOURCE: Based on Berman and Martinez (1976).
/ Conduction
Band
r'S ~ ~
r25
it\ C/~
X U,K
r
this is typically accomplished by providing kinetic energy
to the feedstock species and to the atomic hydrogen. This
kinetic energy is typically imparted via a plasma. These
plasmas are generated by direct current (including arc
jets), radio frequency fields, microwave fields, or by
combustion (e.g., oxy-acetylene torch). Diamond-film
synthesis has an additional complication unknown to the
synthesis of any other semiconductor. The lowest energy
form of its surface is that of an sp2-configured or bond that
is graphitic in character. When the diamond surface is so
constructed, only graphite can be grown on it. To prevent
this unwanted surface construction, the surface is
terminated by hydrogen, but this hydrogen must be
removed and quickly replaced by carbon to grow the
diamond. Since hydrogen bonds to the diamond surface
with an energy of 104 kcal/mol versus the 88 kcal/mol of
the carbon-carbon bond, it is not easy to remove.
Removal requires energetic hydrogen in a "sea" of
hydrogen atoms and (typically) methyl radicals or
acetylene. For every 104 hydrogen atoms removed from
the diamond surface, only one Is replaced by a
carbonaceous radical; the remainder are replaced by
another hydrogen. The growth process is thus slow and
relatively inefficient, although DC arc jets and combustion
jets have grown diamond at rates exceeding 100 micron/in.
29
The cost and safety aspects of methane and gaseous
hydrogen storage are in some cases circumvented by an
alcohol and water plasma process. Several halogen-based
techniques are also currently under investigation to reduce
the synthesis cost. The target for diamond is to grow
large areas of single crystals. This has not yet happened.
Single crystalline diamond has not yet been grown on any
substrate except natural diamond and cBN. Cubic BN is
much less plentiful than natural diamond and available
only in very small sizes. Attempts to synthesize diamond
on all other foreign substrates has resulted in films
characterized by numerous low-angle grain boundaries
and much worse. Large area arrays of seeded natural
diamond on silicon has resulted in mosaic diamond films.
Characterization
The preferred method of characterization of the
diamond films is by Raman spectroscopy. An unstrained
diamond film exhibits a Raman signal at 1,332 cud. The
full width at half maximum of this Raman signal is an
1o2 _
_
y 10
E
5
._
10°
o
C:
Ct
BE
10-1
10-2
., --~
1 1 1 11 1 1 1 11 1 1 1 11
10° 101 1o2 103
Temperature (K)
FIGURE 2-11 Thermal conductivity of two Type IIa diamonds.
SOURCE: Based on Berman and Martinez (1976).
OCR for page 30
Materials for High-Temperature Semiconductor Devices
indication of the quality of the film. The best of type IIa
natural diamond (diamond with no active optical centers)
typically exhibits a full width at half maximum of 2.2 cm~i
or slightly less. The best of artifact homo-epitaxial
diamond films was grown by microwave plasma-enhanced
CVD and was characterized by a Raman signal of 1.7
cod. The full width at half maximum Raman signature
has been correlated with the thermal conductivity of
diamond. In polycrystalline diamond films, a full width at
half maximum Raman signature of 3.2 cm~~ generally
ensures that the thermal conductivity (in the direction of
growth) exceeds 17 W/cm K and the electrical resistivity
exceeds 106 Q cm. Diamond films can be grown with
intrinsic resistivities of 10~° Q cm and larger, but at
slower growth rates (e.g., ~ 1 ,um/h).
30
Diamond Processing
Diamond is not etched by boiling acids or bases.
There are two preferred methods of processing diamond.
The first is the use of kinetic energy beams of oxygen or
oxygen-containing molecules or radicals. The second
method is by electrolytic etching. The electrolytic etching
is generally limited to removal of defect-ridden or
otherwise conducting regions of diamond. Boron is the
only universally recognized acceptor impurity that can be
controllably introduced into diamond. Until recently, only
a small portion of the boron in diamond was electrically
activated. It can now be introduced and nearly 100
percent electrically activated by a series of implantation
processes to concentrations of 1 x 10~9/cm3. By a similar
procedure the same investigator stated that he has
activated phosphorous in diamond at 80 MeV (Prinz,
19941.
Representative terms from entire chapter:
tilt angles
