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Transcript
Chapter 1
Introduction
1.1 Motivation
The rapid growth market and need of wireless network communication are vast and visible.
SiGe HBTs have more advantages to replace GaAs or InP material systems. Low cost and
high compatibility with Si CMOS are the major attractiveness. The SiGe bipolar/BiCMOS
technology provides high speed/bandwidth, high breakdown and low noise characteristics as a
figure of merits in analog and power circuit applications. The record fT of SiGeC HBT is 350
GHz, which had been created by IBM in 2002 [1]. In general, the typical transient speeds of
SiGe/SiGeC HBTs made by other worldwide companies or institutes are typically 150 GHz ~
200 GHz. For improving this index, the shrinkage of vertical layouts is indispensable to
reduce the base transient time furthermore.
The scaling of transistors to ever smaller size has created an increasing demand for
controlling over dopant profiles in silicon devices. As device dimensions have shrunk, control
of the dopant diffusion length becomes ever increasingly critical in defining the minimum size
of a device. Past generations of dopant profile engineering have relied on shrinking the
thermal budgets of the device fabrication process to reduce the thermal broadening of profiles
enough to remain within the limits of the design parameters for junction depth and
gate-lengths. Smaller sizes have, therefore, traditionally meant lower temperatures, however,
processing steps like ion implantation, which create point defects like the silicon
1
self-interstitial that promote diffusion of common dopants like boron and phosphorus, limit
the success of this approach. Ion implantation is known to produce enough excess silicon
self-interstitials to enhance boron and phosphorus diffusivities thousands of times above their
intrinsic diffusivities, making these effects the dominant source of diffusion at low
temperature.
One novel approach to control boron diffusion has been the incorporation of carbon in SiGe,
which dramatically reduce the boron diffusivity and has been shown to suppress the otherwise
enhanced diffusivities resulting from processing steps like ion implantation. For this reason,
the incorporation of carbon has been used to significantly improve the thermal stability of
important device structures, e.g. SiGe heterojunction bipolar transistors (HBTs). The
substitutional C reacts directly with the silicon self-interstitials in a simple one-to-one reaction
(“kick-out”), depleting the local interstitial concentration and thereby reducing the local boron
diffusivities. In this thesis, we utilized UHV/CVD (Ultra High Vacuum Chemical Vapor
Deposition) to grow the SiGeC films using C2H4 as C sources.
For increasing the transient speed of bipolar devices, except the methodology used in
reducing the base transient time, other efforts were put at the collector design of SiGe HBTs.
To reduce the parasitic delay caused by the capacitive and resistive components of junction at
the B-C interface, the considerations including the lateral and vertical layout, SIC (Selective
Implantation Collector) and SOI (Silicon On Insulator) are of important concerns.
1.2 Outline
In chapter 2, the effect of C incorporations into the SiGe alloy is investigated. We had
grown thin (several tens of nms) SiGe0.2C with different C concentrations using C2H4 as C
precursors and UHV/CVD. Material, electrical and optical characteristics are characterized by
2
double crystal x-ray diffraction (XRD), atomic force microscope (AFM), transmission
electron microscopy (TEM), secondary ion mass spectrometry (SIMS), photoluminescence
(PL), and capacitance-voltage measurement (CV). A heterojunction bipolar transistor (HBT)
made from this SiGeC layer is also studied in the following chapter.
In chapter 3, the SiGeC HBT was designed in purpose of purchasing the highest transient
speed for optical communications. We focus on the methodology of C incorporations, and the
design of ultra thin bases. The position and site of C incorporated into the stack SiGe base
regions are closely relevant to DC and RF performances of bipolar devices. We had
characterized the SiGe/SiGeC HBTs by electrical measurements and made a comparison
between Gummel plots, noise figures and cut-off frequencies.
In chapter 4, a novel design of SiGe HBTs fabricated on SOI can provide a higher fmax and
breakdown voltage (BVceo) than conventional bipolar devices. In this innovative idea, it had
been demonstrated by IBM’s announcements at 2002 VLSI and BCTM conference that Si
collector above this SOI has an optimal thickness of 0.15 um. We found that the thickness of
buried oxide embedded in SOI is also effective to affect DC and RF characteristics of the
SiGe HBTs. Through program simulations of 2D and 3D schematic structures by the
commercial simulator, important parameters including the electric field distribution, depletion
region, electrical currents, breakdown voltage and fmax can be extracted to study.
In chapter 5, we comment the influences on the SiGe alloy with the incorporations of C and
summarize the performance of SiGeC HBTs compared with SiGe HBTs. We found that a
good deposition and incorporation of substitutional C in the SiGe alloy is the key step for the
fabrication of high-performance SiGeC HBTs. To achieve this requirement, the choice of C
precursors and the growth parameters involving growth temperature and pressure are crucial.
3
References
[1] J. -S. Rieh, B. Jagannathan, H. Chen, K. T. Schonenberg, D. Angell, A. Chinthakindi, J.
Florkey, F. Golan, D. Greenberg, S. -J. Jeng, M. Khater, F. Pagette, C. Schnabel, P. Smith, A.
Stricker, K. Vaed, R. Volant, D. Ahlgren, G. Freeman, K. Stein, and S. Subbanna, Tech. Dig.
Int. Electron Devices Meet. 2002.
4
Chapter 2
C in SiGe alloy
2.1 Introduction
In this chapter, the effect of C imposed on B diffusion in the SiGe/Si alloy will be
examined by using C2H4 as C precursors and UHV/CVD. Because the lattice constant of
carbon crystallized in the diamond structure is considerably smaller than that of silicon (aC is
0.354 nm), carbon incorporation in SiGe to form SiGeC is also of great interest for its
potential to compensate the compressive stain in SiGe layers grown commensurate to Si and
therefore relax the critical thickness constraint on SiGe layers. This chapter includes a
discussion of the role of silicon interstitials in solid-state diffusions of B in silicon, a
quantitative examination of the influence of substitutional C incorporations on the B
diffusivity in the SiGe alloy, and a observation of thermal behavior of substitutional C after
heat anneal. Atomic interactions of B, Ge, and intentional C dopant are observed in SIMS
profiles. By XRD analysis, the incorporations of C lead somewhat intrinsic changes in the
strain originally induced by SiGe layers epitaxial grown on Si buffer substrates. On the other
hand, AFM indicates surface morphology differences as C incorporated into the SiGe alloy.
For optical characteristics of Si/SiGeC/Si single quantum wells (SQWs), PL intensity of
SQWs monotonically shrinks as the increasing C2H4 flow rate. Electrical characteristics of the
MOS (metal-oxide-semiconductor) structures with 8.8 nm thermal oxide grown upon bulk
Si0.8-yGe0.2Cy (y = 0 or 0.0014) were probed using the CV technique. Normalized inversion
5
capacitances (Cinversion/Coxide), conducted with several operation frequencies (50 Hz ~ 0.2
MHz), may be related to the generations of trap centers accompanied with C incorporations.
These highlight technological skills of controlling B diffusions in the HBTs by using this
“C-doping” technique, and motivate the later chapter investigating more physical
characteristics of complete single-poly SiGeC HBT devices.
2.2 B and Ge in SiGeC alloy
We investigate the microscopic model of B diffusion in Si in this section. Solid-state
diffusivity of B is known to increase during the ion implantation process. The implantation
enhanced diffusion (TED) of B is attributed to the interactions between the substitutional B
and the excess interstitials injected during dopant implantations for reducing resistances. In
typical TED experiments, it’s found that the B diffusivity is relatively enhanced thousands of
times above its intrinsic diffusivity due to the point defects created in high energy/dosage
implantation processes [1-2].
In a simple (n-)Si/(p+)SiGe/(n-)Si HBT, B serves as major p-type dopants in the SiGe
layer. As B diffuses from the p+ SiGe base into the n-Si emitter and collector (figure 2.1), the
parasitic barrier formed in the conduction bands, which impedes the flow of electrons from
emitter to collector (figure 2.2) [3].
6
SiGe
As-grown
SiGe
Annealed
Boron inside
SiGe layer
p+
n-
Activation
anneal
Concentration
Concentration
n-
Boron outside
SiGe layer
n-
Depth
p+
n-
Depth
(a)
(b)
Figure 2.1 Schematic diagrams of (a) the as-grown dopant profiles of a SiGe HBT structure
and (b) the profile after anneal. B is shown to diffuse faster than the surrounding SiGe layer in
P doped Si regions.
The parasitic barrier arising from B outdiffusions is strongly dependent on the B
concentrations that diffuse into the Si, and a small amount of B outdiffusion (Ld ~ 10 A0) can
already cause large parasitic barriers evident in HBT’s (figure 2.2).
Figure 2.2 Schematic conduction band diagram of a (n-)Si/(p+)SiGeC/(n-)Si HBT as-grown
7
and after anneal showing the creation of a parasitic conduction band barrier as a result of B
diffusion from the base into the n-type Si emitter and collector regions [3].
The possible mechanism that B diffusivity can be suppressed by C incorporations is
schematically shown in figure 2.3. The random thermal B diffusion is mediated by Si
interstitials. B is known to dissolve onto the substitutional lattice sites of silicon at the
concentration below its intrinsic solid solubility in Si, and designated as B in the
substitutional states, Bs. Bs is relatively strongly bound in this sate and considered as
immobile, therefore the diffusivity of Bs is for all intents and purposes zero [4]. The mobile B
defect state is formed when as silicon point defects, i.e. silicon vacancy or interstitial, are
captured by the substitutional B atom. It has been assumed, then, that the observed B
diffusion is mediated by the mobile B defect state, Bi, that has non-zero diffusivity. This
chemical reaction leading to the negatively charged defect is expressed as: Bs  I  Bi ,
where Bs- is the ionized substitutional B atom and I is the silicon interstitial [5].
B
B
I
I
I reacts with
substitutional
Boron
Si Interstitial (I)
Boron interstitial (Si and B share site)
is highly mobile
Figure 2.3 Schematic representation of an interstitial silicon atom near a substitutional B
atom in a silicon lattice and the resulting mobile B interstitial (BI) formed by the binding of
the interstitial silicon atom to a B [6].
8
2.2.1 Test experiment using SiGeC
All SiGeC samples in this thesis were grown on 100-mm (100) Si wafers by UHV/CVD.
The SiGeC layers were epitaxially grown using C2H4 as C source. Czochralski silicon (100)
substrates were cleaned for 20 min in H2SO4/H2O2 followed by an etch in a dilute HF solution.
Then the wafers were loaded into the reactor via a loadlock. The growth pressure and
temperature were 1 mTorr and 550 0C, respectively. Low growth temperature is preferred to
have high substitutional C incorporation [7-8]. The gas flows of GeH4 (5% in He), B2H6 (5%
in He), and SiH4 (in He) are 36, 50, and 70 sccm, respectively. This yields a SiGe layer with
the Ge fraction of 0.2. The various C concentrations in the SiGeC layers were obtained by
using different C2H4 (2 % in He) flow rates.
To identify the mechanism of C interacting with Ge and B dopants in the SiGe layers, a
Si/SiGe(C)/Si single quantum well (SQW) was fabricated. This schematic structure is
exhibited in figure 2.4.
Si cap 10 nm
Si 40 nm
Si0.75-yGe0.25Cy 20 nm, [B = 6E19 cm-3;
y = 0, 0.1 %, 0.5 %]
Si 40 nm
Si0.75Ge0.25 20 nm, [B = 6E19 cm-3]
Si 40 nm
Si 20 nm [B = 6E19 cm-3]
Si seed layer 40 nm
Si substrate
9
Figure 2.4 A SQW structure includes a SiGe layer strained by seed Si, and a SiGeC alloy
layer intended for investigating the effect of C incorporations. All the Ge concentration
profiles are box-like [9].
From figure 2.4, three splits (sample A, B and C) were fabricated with different C
concentrations in Table 2.1.
Splits no
A
B
C
Seed Si layer thickness
40 nm
40 nm
40 nm
Si0.75Ge0.25 thickness
20 nm
20 nm
20 nm
Si0.75-yGe0.25Cy, y =
0%
0.1 %
0.5 %
Si0.75-yGe0.25Cy thickness
20 nm
20 nm
20 nm
Si cap layer thickness
10 nm
10 nm
10 nm
Table 2.1 Three Si/SiGe(C)/Si SQWs with different C concentrations (y = 0 %, 0.1 %, or 0.5
%) [10].
SIMS
The SIMS measurement clearly indicates the C concentration incorporated in the SiGe
layer ranging from 1E18 ~ 2E20 cm-3. The lowest detection limit of C in this SIMS profile is
at the concentration of 5E17 cm-3. An intrinsic residue C can’t be avoidable and appear at the
interfaces during each epitaxial layer growth.
10
0.5% C
no C
0.1 % C
0.5 % C
20
-3
C concentration (cm )
10
0.1% C
19
10
No C
18
10
0
500
1000
1500
2000
2500
0
Depth (A )
Figure 2.5 Differences in the C concentrations of SiGe(C) peaks among three SQWs (sample
A, B and C) [11].
The distributions of B dopant concentrations are exhibited in figure 2.6. When C
concentration increases up to 0.5 % (sample C), it shows the B concentration slightly
dropping to 5E19 cm-3. That may indicate a “kick-out” of a small fraction of substitutional B
by substitutional C incorporations if C in the B-doped SiGe layers at the concentration greater
than 0.5 %. Oppositely, the other two B peaks between samples A (no C) and B (C = 0.1 %)
doesn’t exhibit obvious differences, which may imply a little of substitutional C within total C
concentration of 0.1 % when C2H4 is utilized as C sources.
11
21
-3
B concentration (cm )
10
no C
0.1 % C
0.5 % C
no C
0.5% C
20
10
0.1% C
19
10
18
10
17
10
0
500
1000
1500
2000
2500
0
Depth (A )
Figure 2.6 Difference in B concentrations among three SQWs (sample A, B and C). B
concentration degrades about 30 % in the Si0.745Ge0.25C0.005 HBT with comparisons to the
SiGe HBT [12].
Figure 2.7 describes the behavior of B outdiffusion after thermal processes, which is a main
topic of our project and research. B profile becomes severely widen after thermal anneal at
1000 0C for 10 seconds in N2. Heat treatments and transient enhanced diffusion (TED) caused
by implantation damage can broaden the B profile into the adjoining Si regions [13]. As stated
in previous sections, B diffusion in Si is sensitive to the thermal budget and increases in high
temperature anneal. One of the key targets in the SiGe HBT technology is to retain a shallow
as-grown B profile within the SiGe base layer and a nominal junction profile of p-n dopants
after post-epitaxial processing.
12
as-grown
0
as-grown
-3
B concentration (cm )
after anneal at 1000 C
for 10 seconds in N2
after anneal
1E20
1E19
1E18
0
500
1000
1500
2000
2500
0
Depth (A )
Figure 2.7 B concentrations of the sample A (no C incorporation) before and after thermal
anneal at 1000 0C for 10 seconds in N2. An apparent broaden B profile appears after anneal
than as-grown [14].
The phenomenon that C incorporated into the SiGe layer will replace a fraction of Ge
bonded at substitutional sites is also observed in figure 2.8. After subtracting that of the
sample with no C (sample A), the mole fraction of Ge degrades almost 0.05 % for the sample
with 0.5 % C incorporation (sample C), which is right equal to the intended-doped C
concentrations. Nevertheless, the Ge concentration of SiGe layers in other samples (sample A
and B) remained unchanged, which clearly exhibits a competition of occupy of substitutional
sites
between
the
doped-carbon
atoms
13
and
germanium
atoms
incorporations.
0.1 % C
no C
0.25
Ge concentration (%)
0.5 % C
0.20
0.15
no C
0.1 % C
0.5 % C
0.10
0.05
0.00
0
500
1000
1500
2000
2500
0
Depth (A )
Figure 2.8 Differences in Ge concentration between the SiGeC layers of samples A, B and C
(C = 0 %, 0.1 % and 0.5 %, respectively). No significant change of Ge concentrations in the
SiGe layers [15].
To realize the influence of C incorporations in the SiGe alloy more, a multiple-layer
SiGe(C) structure was grown as after the 40 nm undoped Si buffer growth, six SiGeC layers
with C2H4 flow of 0, 25, 50, 65, 75, 100 sccm were grown with 40 nm spacers between each
layer. Each Si/SiGe(C) layer thickness was measured by fitting the XRD and calibrated by
cross-sectional TEM. The SiGe(C) layer was also doped with B and there is 4 nm Si cap on
the top. The schematic diagram of this MQW is shown in figure 2.9.
14
2.2.2 Si/SiGe(C)/Si MQWs with different C concentrations
Si cap-40nm
SiGe0.2C 40nm [B=1E19 cm -3; C flow rate=100 sccm]
Si spacer-40nm
SiGe0.2C 40nm [B=1E19 cm-3; C flow rate=25 sccm]
Si spacer-40nm
SiGe0.2 40nm [B=1E19 cm-3]
Si spacer-40nm
Epitaxy Si crystal -40nm
Si buffer layer-40nm
p-type Si substrate
Figure 2.9 A Si/SiGe(C)/Si MQW structure consisted of six SiGeC layers with C2H4 flow rate
of 0, 25, 50, 65, 75, 100 sccm [16].
The SIMS depth profiles of this sample before and after anneal at 1000 C for 10 seconds
are shown in figure 2.10. Measured C concentrations range from 11019 to 31020 cm-3
correspondent to the flow of 0 ~ 100 sccm. Accurate controls of epitaxial process can be
revealed from precise concentrations of multiple Ge peaks (0.2  0.6 %). It indicates that the
Ge fraction isn’t substantially changing until the C2H4 flow up to 100 sccm. Obstruction of Ge
incorporations is observed to degrade 5% fraction (Ge ~ 0.15) because of supersaturated
15
substitutional C in the SiGe alloy, which is consistent with x-ray diffraction results that the
low diffraction peak intensity of the SiGe peak is observed for the sample E (figure 2.11). The
result is conflicted with the observations by Vinh et al. [17] and Mocuta et al. [18] that the C
incorporations from SiCH6 does not influence the Ge concentrations.
C (as grown)
C (after anneal)
Ge (as grown)
0.21
-3
2.8x10
0.18
20
2.4x10
20
0.15
20
0.12
20
0.09
19
0.06
19
0.03
2.0x10
1.6x10
1.2x10
8.0x10
4.0x10
100
200
300
400
Ge mole fraction
C & B concentration (cm )
20
B (as grown)
B (after anneal)
0.00
500
Depth from surface (nm)
Figure 2.10 SIMS depth profiles of a Si/SiGe(C)/Si multiple quantum well before and after
thermal anneal at 1000 oC for 10 seconds in N2 [19].
X-ray diffraction is an extremely sensitive method for determining the lattice constant of
crystals. Because the lattice constants of Si and Si1-xGex shrink due to the incorporation of
smaller C atoms on substitutional sites in the crystal, the totoal amount of substitutional C
may be deduced by the lattice constant of the SiGeC alloy layer by using a Ge to C strain
16
compensation ratio n = 10 [20]. We investigate the strain relaxation by C incorporation into
the SiGe layers by using the same structures of MQWs shown in figure 2.9. As a small
amount of C2H4 was progressively added to the reaction chamber, there is a shift in the (400)
x-ray diffraction peak of the resulting 40 nm SiGeC layers from that of a control SiGe sample
(figure 2.11), due to the decrease of lattice constant by substitutional carbon incorporation. A
largest shift (Δθ = 100 arcsec) is observed at the C2H4 flow of 75 sccm, obtained from XRD
simulation with the assumption that the Ge content of SiGeC is the same as the SiGe control
sample, corresponding to 0.09  0.01 % C. As the C2H4 flow increases from 10 to 75 sccm,
the substitutional C content increases with the increasing C2H4 flow, and the growth rate
decreases from 2.2 nm/min to 1.9 nm/min. At C2H4 flow rate of 100 sccm, little increment of
the substitutional C fraction is observed as compared to 75 sccm C2H4 flow, and the resulting
SiGeC layer is of poor crystal quality, indicated by the broadened and weak diffraction with
X-ray intensity (cps)
the absence of thickness fringes in the XRD spectrum (figure 2.11).
(400) Si
10
6
10
5
10
4
10
3
(B) C 2 H 4 10 sccm
10
2
(A) no C 2 H 4
10
1
(E) C 2 H 4 100 sccm
(D) C 2 H 4 75 sccm
(C) C 2 H 4 50 sccm
0
10
-3000
-2500
-2000
-1500
-1000
arc-second
17
-500
0
500
Figure 2.11 Double crystal x-ray rocking curves of the SiGeC alloys with Ge fraction of ~ 0.2.
Additions of C2H4 into the SiGe alloy exhibits good substitiional C incorporations by the
symptom of peak sequentially shift toward Si (400) peak [21].
Estimated total and substitutional C (CS) concentrations as a function of the C2H4 flow are
represented in figure 2.12. The behavior of Cs concentrations with the C2H4 flow is somewhat
different from those exhibited by other groups using SiCH6 growth techniques [22-23]. Under
the conditions of the same growth temperature (550 C), the reported researches indicated the
higher saturation concentrations of Cs (> 0.5 %) can be achieved in the Si1-x-yGexCy layers by
the incorporation of SiCH6 than those (~ 0.1 %) in using C2H4.
0.6
C fraction (y%)
0.5
total C
0.4
0.3
0.2
subst. C
0.1
0.0
20
40
60
80
C2H4 flow (sccm)
18
100
Figure 2.12 Estimated total and substitutional C concentrations as a function of the C2H4
flow.
Figure 2.13 represents the C peak concentrations and the full widths at half maximum
(FWHM) as a function of the C2H4 flow. Comparisons of FWHM between the carbon profiles
with or without thermal treatment provide one path to qualitatively check if the interstitial C
(CI) exists. The same FWHM observed after thermal annealing at high C2H4 flow (> 75 sccm)
indicated that the formation of CS-CI clusters or silicon-carbide precipitation [24], especially
pronounced apparent at 100 sccm. This represents a large part of CI generations
accompanying with high vapor flow of C2H4. CS was able to suppress outdiffusion of B via the
substitutional-interstitial mechanism [25-26]: CS  I  CI
The substitutional carbon
reduction can be enhanced by interstitial silicon (I) injection [27]. Coupled diffusions of B
and point defects are restricted because the Si self-interstitials reacted with immobile CS.
19
20
1000
-3
800
as grown
after anneal
as grown
after anneal
20
3x10
600
20
2x10
400
FWHM (nm)
C concentration (cm )
4x10
20
1x10
200
20
40
60
80
100
C2H4 flow (sccm)
Figure 2.13 C peak concentrations and FWHM from figure 2.10 as a function of the C2H4
flow before and after thermal anneal at 1000 oC for 10 seconds.
As indicating in figure 2.14, the outdiffusion of B after anneal at 1000 C for 10 seconds
will be apparently significant at low flow (< 20 sccm) of C2H4 reactant, representing that the
small amount of Cs is insufficient to reduce the diffusion of B. If the flow is elevated over 60
sccm, however, the increasing CS yielded the unaffected FWHM of B profiles in spite of high
temperature anneal.
20
500
5.00E+019
as-grown
after anneal
-3
400
4.00E+019
300
3.50E+019
as-grown
after anneal
3.00E+019
FWHM (nm)
B concentration (cm )
4.50E+019
200
2.50E+019
2.00E+019
20
30
40
50
60
70
80
90
100
100
110
C2H4 flow (sccm)
Figure 2.14 B peak concentrations and FWHM from figure 2.10 as a function of the C2H4 flow
before and after thermal anneal at 1000 oC for 10 seconds.
The C concentrations adopted from the SIMS (figure 2.10) and XRD [28-29] (figure
2.11) measurements were extracted for comparisons. When the higher flow of C2H4 was used,
the curve is getting more flat with a smaller slope. The about ratio of Csubstitutional/Ctotoal 0.49 at
the C2H4 flow of 75 sccm was less than 0.69 for 50 sccm, implying that the larger portion of
CI generations in ternary SiGeC system at the high C2H4 flow.
21
2.3 Thermal stability
A loss of substitutional C in the SiGe layers is apparently observed after thermal process. A
test structure of Si/SiGeC/Si SQW with C concentrations of 0.1 % (sample B in table 2.1) is
used in the thermal stability test. The incorporated C peak of the SiGeC diffuses significantly
after anneal at 1000 0C for 10 seconds in an insert ambient. As mentioned in the previous
paper [3], the root reason for the outdiffusion of C with heat treatment can be attributed to the
reaction of Cs with Si interstitials and formation of interstitial C. Except this intended doped
C, the other two C peaks coming from contaminations also exhibit severe diffusion after
thermal anneal.
as-grown
20
0
after anneal at 1000 C
for 10 seconds in N2
-3
C concentration (cm )
10
10
as-grown
19
after anneal
10
18
0
500
1000
1500
2000
2500
Depth (nm)
Figure 2.15 C profiles of the Si/SiGeC/Si SQW with C concentration of 0.1 % (sample B in
table 2.1) before and after anneal at 1000 0C for 10 seconds in N2 ambient [30].
22
Figure 2.16 (a) and (b) reveal the thermal effect on the strain of SiGe alloy. These samples
constitute Si0.75Ge0.25 (20 nm) (sample A) or Si0.75-yGe0.25Cy (y = 0.5 %) (sample C) layers
grown on Si substrate, as previously shown in figure 2.4. In figure 2.16 (a), a high
temperature (~1000 0C) anneal provides activation energy to relax the lattice strain. A large
shift in the SiGe peak indicates the strain relaxation and dislocation formations. In figure 2.16
(b), due to substitutional C reducing strain tension of the SiGe alloy, a good thermal stability
of this SiGeC SQW is shown in the XRD profile. There is no apparent shift of SiGe peak of
the Si0.75-yGe0.25Cy (y = 0.5 %) layer before and after anneal. It’s beneficial to engineers that
substitutional C will not improve the thermal stability of SiGe heterojunction structures after
thermal processing, but also prevent B outdiffusion in the base of SiGeC HBT, providing the
useful knowledge for devices design.
X-ray intensity (CPS)
No C
10
6
10
5
10
4
10
3
10
2
10
1
as-grown
0
after anneal at 1000 C for 10 seconds
as-grown
after anneal
0
10
-2500
-2000
-1500
-1000
-500
0
500
arc-second
Figure 2.16 (a) XRD analysis of the sample with no C incorporation before and after anneal
at 1000 0C for 10 seconds in N2 [31].
23
0.5 % C
6
10
as grown
X-ray intensity (CPS)
0
after anneal at 1000 C for 10 seconds
5
10
after 1000oC
10 sec anneal
4
10
as-grown
3
10
2
10
1
10
0
10
-3000
-2500
-2000
-1500
-1000
-500
0
500
arc-second
Figure 2.16 (b) XRD analysis of sample with 0.5 % C incorporation before and after anneal
at 10000C for 10 seconds in N2 [32].
The thermal behaviors of SiGeC alloy are probed using X-ray analysis in figure 2.17 (a) &
(b). These sample structures are a muli-MQWs of Si/SiGe(C)/Si layers, which consisted of
same ingredients as shown in figure 2.4. As C is added into this SiGe MQWs in figure 2.17
(a), more the peak shifts of SiGe layer are observed, representing an increasing substitutional
C content and partial strain relaxation due to smaller lattice constants own by C compared to
Si or Ge. After subjected to thermal anneal, same as the above conclusion, substitutional C
will diffuse via Si interstitial and become interstitial C. A loss of substitutional C will help in
remaining the original strain caused by Ge effects in SiGe alloy. So no significant peak shifts
are observed among the samples of different C concentrations after high temperature anneal in
24
figure 2.17 (b).
as-grown
5
X-ray intensity (CPS)
10
4
10
0%C
0.1 % C
0.5 % C
0.5% C
0.1% C
No C
3
10
2
10
1
10
0
10
-3000
-2500
-2000
-1500
-1000
-500
0
500
arc-second
Figure 2.17 (a) XRD of as-grown SiGeC layers for C = 0, 0.1% or 0.5 %.
X -ra y in te n s ity (C P S )
0
a ft e r a n n e a l a t 1 0 0 0 C 1 0 s e c
10
7
10
6
10
5
10
4
10
3
10
2
10
1
10
0 % C
0 .1 % C
0 .5 % C
0.5% C
0.1% C
No C
0
-3 0 0 0
-2 5 0 0
-2 0 0 0
-1 5 0 0
-1 0 0 0
a rc -se c o n d
25
-5 0 0
0
500
Figure 2.17 (b) XRD of SiGeC layers for 0, 0.1% or 0.5 % after anneal at 1000 0C for 10
seconds in N2.
2.4 Defect formation
AFM
In figure 2.19 (a), AFM is used to investigate the surface status of Si/SiGeC/Si SQW in
figure 2.4. The results of roughness on the surface Si cap of Sample A (no C incorporation)
are similar to that of pure Si. Roughness systematically increases as C concentration increase
(as shown in Fig. 2.19 (b) & (c)).
Figure 2.19 (a) AFM measurement of Sample A in table 2.1 (with no C incorporation),
roughness = 0.2 nm [33].
26
Figure 2.19 (b) AFM measurement of Sample B in table 2.1 (with C concentration of 0.1 %),
roughness = 0.27 nm [34].
27
Figure 2.19 (c) AFM measurement of Sample C in table 2.1 (with C concentration of 0.5 %),
roughness = 1.66 nm [35].
The relationship of roughness with the mole fraction of C incorporations is described in
figure 2.20. This roughness enhancement due to C effect maybe be explained as following:
when C incorporated into the SiGe alloy, C of substitutional and interstitial sites formed at
same time, depending on the growth condition and temperature. Interstitial C will make
irregular alignments of constituent atoms in the lattice, and hence induce surface roughness.
1.8
1.6
Si0.75-yGe0.25Cy (before anneal)
1.4
Rrms(nm)
1.2
1.0
0.8
0.6
0.4
0.2
0.0
0.0
0.1
0.2
0.3
0.4
0.5
0.6
y (%)
Figure 2.20 Roughness as a function of C concentration with the samples in table 2.1. As C
increasingly incorporated into the SiGe alloy, the roughness becomes significantly larger.
PL
Three Si/Si1-x-yGexCy/Si samples (with no C2H4 as a control sample, C2H4 flow of 10 sccm
and 50 sccm) were grown for PL measurements. The fractions of Ge (x ~ 0.2) and total C
28
contents (y = 0, 0.02 % or 0.05 %) in the Si1-x-yGexCy layer were determined by the SIMS
profiles. Those Si0.8-yGe0.2Cy layers were capped with about 10 nm of silicon grown at 550 C
using a mixture of 15 sccm siliane and He gas. For increasing the PL intensity, thick Si cap is
intended to use [36]. Each thickness of these quantum wells was identified by fitting the XRD
spectra. Neither samples showed evidence of SiC precipitates, confirmed by the cross section
TEM pictures. PL studies were performed with a 488nm Ar+ laser beam at the power density
of about 0.2 Wcm-2. The circular spot size was a diameter of 0.5 um. The temperature of the
samples during measurements was cooled and held constant at 15k. PL was taken by standard
lock-in technique using 1m grating monochromator and a Ge photodetector cooled with
liquid-nitrogen [37]. Spectral variations of PL for the SiGeC SQWs as a function of the C2H4
flow were presented in figure 2.21. Those SiGeC SQWs had qualitatively similar PL spectra.
The overlay of the transverse optical (TO) phonon replicas results in the peak on the low
energy side of the NP line. Except the Si peak at 1.10 ev, the strongest peak is the no-phonon
(NP) transition due to the lattice disorders (interface roughness and alloy fluctuations), which
relaxes the momentum-conservation requirement [20] and the peak at 1.04 ev of SiGeC PL
spectra is the two phonon replica of bound excitation emission from Si [37]. Significant
attenuations of PL intensity were observed at both NP and TO peaks as the C2H4 flow
increases. The integration of PL intensity with subtraction of background noise, shown in the
inset of figure 2.21, shrunk monotonically with the increasing flow rates. Expected energy
shifts were not observed in the spectra, since the concentrations of CS were not enough to
significantly change the bandgap of SiGeC alloy [38]. It exhibited that the severe decreasing
density of PL peak at the C2H4 flow of 50 sccm was caused by the carrier scatterings from
extra CI atoms.
29
0.002
SiGeC 15K PL
Integration of PL Intensity (au)
PL Intensity (a.u.)
0.0013
0.001
0.0012
0.0011
0.0010
0.0009
0.0008
0.0007
NP
0.0006
0
10
20
30
40
50
C2H4 flow rate (sccm)
A (0 sccm)
B (10 sccm)
C (50 sccm)
TO
0.000
0.9
1.0
1.1
Energy (ev)
Figure 2.21 15 K photoluminescence spectra of Si/Si0.8-yGe0.2Cy/Si quantum wells at the
power density of 0.2 Wcm-2 with the C2H4 flow of 0, 10, and 50 sccm. The inset shows that
the integration of PL intensity shrinks with higher flow of C2H4 [39].
CV
To obtain the response of electric carriers to external exerted bias, the MOS diodes
fabricated with the Si/Si0.8-yGe0.2Cy/Si substrates had Al gate electrodes on the front and rear
sides with various circular areas. The SQWs substrate composed of undoped Si thin buffer
(thickness 10 nm), p-doped (B: 1101 9 cm-3) Si0.8-yGe0.2Cy with y = 0 or 0.2 % (thickness 90
nm), and undoped Si cap (thickness 9 nm). After cleaning the substrate in diluted HF, the
oxidation was performed by furnace oxidization process at 900 C for 30 min to form 8.8 nm
thermal oxides. The normalized capacitance (C/Coxide) characteristics were measured at room
temperature using a capacitance meter in a light-tight box with bias sweep from -3V to 4V
30
and operation frequency from 50 Hz to 0.2 MHz. In order to estimate the trap concentrations
in the substrate [40], the normalized capacitances (Cinversion/Coxide) biased in strong inversion
regions (bias > 1 V) as a function of operation frequency were shown in figure 2.22. No
matter in the SiGe or SiGeC samples, normalized capacitance inversely increases when the
operation frequencies decrease. Within the same operation frequency, the SiGe samples of C
incorporation had larger Cinversion/Coxide than SiGe, indicating extra-induced minority carriers
response with exerted bias. They can be generally attributed to the generation-recombination
of trap centers generated during incorporation of C, which afford the medium for carriers to
response with outer voltages.
1.0n
Si0.8Ge0.2 HBT
Si0.798Ge0.2C0.002 HBT
Capactiance (F)
800.0p
600.0p
Si0.8Ge0.2 HBT
Si0.798Ge0.2C0.002 HBT
400.0p
200.0p
0.0
-5
-4
-3
-2
-1
0
1
2
3
4
Voltage (V)
Figure 2.22 Normalized capacitance (Cinversion/Coxide) of MOS diodes fabricated on the
Si/Si0.8-yGe0.2Cy/Si substrates with y = 0 or 0.2 % measured under several operation frequencies
(from 50 Hz to 0.2 MHz) and exerted biases (from -4 V to 3 V) [41].
31
2.4 Summary
The microscopic mechanisms of solid-state B diffusion in silicon were examined and B
diffusivity was shown to be majorly dependent on the silicon interstitial concentration.
Because current transistor dimensions require extreme control over the B profiles, common
fabrication steps that produce extra silicon interstitials in the silicon interstitials in the silicon
lattice, eq, ion implantation and oxidation, meet significant challenges to fabricating high
performance transistors.
To the overall view, the introduction of substitutional C to the SiGe base is shown to reduce
the effective intrinsic B diffusivity by XRD/SIMS measurements and sinks almost all excess
interstitials that reach SiGeC layer for this process. Nevertheless, the creation of interstitial C
or other complex accompanied with the incorporations of substitutional C renders negative
effects of forming defects and inducing roughness. The optimization conditions for growing
SiGeC epitaxial films commensurate on Si substrate are crucial to good-quality devices.
Acknowledgement
The author wants to acknowledge for the help of Prof. Y. F. Chen and I. J. Wu (Physics
department in NTU) in PL measurements, T.-C. Chen (EE department in NTU) in CV and PL
measurements, L. S. Chiang and Y. M. hsu (both in ERSO) for electrical and noise
measurements and grateful to thank Y. D. Zeng and P. S. chen (both in ERSO) for all XRD
measurements, and SiGe/SiGeC epitaxy growth. This work was in finial supported by ERSO /
ITRI, Taiwan, ROC, under contract No: A321XS6700 for publishing this article.
32
References
[1] L. D. Lanzerotti, J. C. Sturm, Appl. Phys. Lett., vol. 70, pp. 23, 1997.
[2] R. Scjpz, U. Goesele, J. Y. Huh, and T. Y. Tan, Appl. Phys. Lett., vol. 72, pp. 2, 1998.
[3] E. J. Prinz, P. M. Garone, P. V. Schwartq, X. Xiao, and J. C. Sturm, IEEE Elec.
Dev. Lett., vol. 12, pp. 661, 1991.
[4] R. D. Harris, J.L. Newton, and G.D. Watkins, Phys. Rev. B, vol. 12, pp. 1094, 1987.
[5] A. Ural, P. B. Griffin, and J. D. Plummer, J. Appl. Phys., vol. 85, pp. 6440, 1999.
[6] Malcolm S. Carroll, Ph.D. Thesis, Princeton, JUNE, 2001.
[7] C. W. Liu, A. St. Amour, and J. C. Sturm, J. Appl. Phys., 80, 3043, 1996.
[8] A. C. Mocuta and D. W. Greve, J. Vac. Sci..Technol., A 17, 1239, 1999.
[9] Process was executed in ERSO, ITRI, Hsinchu.
[10] Process was executed in ERSO, ITRI, Hsinchu.
[11] Process was executed in ERSO, ITRI, Hsinchu.
[12] Process was executed in ERSO, ITRI, Hsinchu.
[13] H.-J. Gossmann, T.E. Haynes, P. A. Stolk, D. C. Jacobson, G. H. Gilmer, J. M. Poate, H.
S. Luftman, T. K. Mogi, and M. O.Thompson, Appl. Phys. Lett, vol. 71, pp. 1988, 1998.
[14] Process was executed in ERSO, ITRI, Hsinchu.
[15] Process was executed in ERSO, ITRI, Hsinchu.
[16] Process was executed in ERSO, ITRI, Hsinchu.
[17] Vinh LeThanh, C. Calmes, Y. Zheng, and D. Bouchier, Appl. Phys. Lett. vol. 80, pp. 43,
2002.
[18] A. C. Mocuta and D. W. Greve, J. Vac. Sci..Technol. A 17, pp. 1239, 1999.
[19] Process was executed in ERSO, ITRI, Hsinchu.
33
[20] C. W. Liu, A. St. Amour, and J. C. Sturm J. Appl. Phys. 80, pp. 3043, 1996.
[21] Process was executed in ERSO, ITRI, Hsinchu.
[22] A. C. Mocuta and D. W. Greve, J. Appl. Phy. vol. 85, pp. 1240, 1999.
[23] A. C. Mocuta and D. W. Greve, J. Vac. Sci..Technol. A 17, pp. 1239, 1999.
[24] J. W. Strane, H. J. Stein, S. R. Lee, S. T. Picraux, J. k. Watanabe, and J. W. Mayer, J. Appl.
Phys. vol 76, pp. 3656, 1994.
[25] R. Scholz, U. Gosele, J.-Y. Hsu, and T. Y. Tan, Appl. Phys. Lett. vol. 72, pp. 200, 1998.
[26] R. F. Lever, J. M. Bonar, and W. Willoughby, J. Appl. Phy. vol. 83, pp. 1988, 1998.
[27] C.W. Liu, Y. D. Tseng, and Y. S. Huang, Appl. Phys. Lett. vol. 75, pp. 2271, 1999.
[28] J. L. Regolini, F. Gisbert, G. Dolino, and P. Boucaud, Mater. Lett. vol. 18, pp. 57, 1993.
[29] B. D Cullity, Elements of X-ray Diffraction, 2nd ed. (Addision-Wesley, Reading, MA,
1978), P. 102.
[30] Process was executed in ERSO, ITRI, Hsinchu.
[31] Process was executed in ERSO, ITRI, Hsinchu.
[32] Process was executed in ERSO, ITRI, Hsinchu.
[33] Process was executed in ERSO, ITRI, Hsinchu.
[34] Process was executed in ERSO, ITRI, Hsinchu.
[35] Process was executed in ERSO, ITRI, Hsinchu.
[36]C. W. Liu, A. St. Amour, and J. C. Sturm, J. Appl. Phys. vol. 80, pp. 3043, 1996.
[37] J. C Sturm, A. St. Amour, Y. Lacroix, and M. L. W. Thewalt, Appl. Phys. Lett., vol. 64,
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[38] C. Penn, S. Zerlauth, J. Stangl, G. auer, G. Bdrunthaler, and F. Schaffler, J. Vac. Sci.
Technol. B 16, pp. 1713, 1998.
[39] Process was executed by T.-C. Chen .in NTU.
[40] Nicollian, E. H. MOS (Metal Oxide Semiconductor) physics and technology (Wiley,
34
1982).
[41] Process was executed by T.-C. Chen .in NTU.
35
Chapter 3
Characteristics of SiGeC HBTs
1. Introduction
SiGe bipolar/BiCMOS technology has become an important technology with many new
and exciting product applications. There is a growing demand today for high performance
devices with improved current gain (beta), higher cut-off frequency (fT) and a larger Early
voltage value (VA). All these key parameters of figures of merit or improved with the
incorporation of Ge into the Si base of a transistor, giving birth to a SiGe heterojunction
bipolar transistor (HBT). Despite many advantages, these HBTs suffer from many vital
difficulties, such as the problem of a lower value of critical thickness due to 4.2 % lattice
mismatch between Si and Ge, the barrier in process integration and B outdiffusion from the
base across the collector-base and emitter-base junction [1]. The lower value of critical
thickness causes worse film stability [2-3] and the outdiffusion is responsible for the
degradation of device performance. These problems can be mitigated by incorporating small
amounts of C in the base of a SiGe HBT [4].
“C-doped” was commonly considered effective to maintain the sharp B profiles by
eliminating transient enhanced diffusion (TED) [5-7]. Adding C into the SiGe material
systems virtually extends the design freedom for SiGe-based heterostructures [8-9]. The
technique of C-doped had became prevailing in modern electronic devices like novel Silicon
Ge heterojunction bipolar transistors (SiGe HBTs) [10-11], p-metal-oxide-semiconductor
36
field-effect transistors (PMOSFET) [12], and photodetectors [13]. With the aid of C, the HBTs
can have the trendous transmission speed (fT = 350 GHz) for the wireless communication
demand. The key role of this technology relies on the quality of SiGeC epitaxial films [14].
The apparatus and recipes of growing SiGeC epitaxial films are crucial key to fabricate high
performance nano-size devices.
The bulk solid solubility of C is higher in neither Si nor Ge. No stable G-C phases are
known above this solid solubility limit of Ge. In Si, SiC is the only thermally stable phase.
Above it, C has been identified in various meta-stable states including substitutional C,
interstitial carbon (Ci), and several carbon interstitial complexes X-Ci including Bs-Ci, and
Cs-Ci. These different states fo carbon in Si are all undesirable, except for substitional C,
because of their effects on the electrical properties of the material. It’s difficult to alloy an
appreciable amount of C on substitutional sites in either Si or SiGe at thermal equilibrium.
Figure 3.1 shows the considerations in designing high-speed SiGe HBTs. It is well known, for
the applications in high cut-off frequency circuits, reducing the parasitic delay and
majority/minority carrier transient time is critical to the performance of HBTs. The
out-diffusion of base dopant (B) makes the increase of quasi-neutral base width and base
resistance, in turn inducing the parasitic barrier of the conduction band at B-C junction and
degrading fT and fmax. “C-doped” technology is both effective in keeping steep base profile
and minimizing base width, so beneficial to improve the fT over 200 GHz. Figure 3.2 shows
the schematic concepts of C incorporations in the SiGe HBTs.
37
Boron porfile
TED / OED
degradation
base
wilder base
smaller Nb
limit fT
RB increase
C doped
Solution
Enhanced
performance !
Figure 3.1 Methodology of base design in remaining steep B profiles and high base
concentrations.
38
Figure 3.2 Schematic diagrams of C incorporations in the SiGe HBTs.
3.2 Fabrication and characterization
To investigate minority carrier transport behavior, SiGeC HBTs were fabricated and
compared to SiGe HBTs. The nominal contents (Ge: 0.2; C: 0.002) and concentration (B:
21019 cm-3) in the base region of the SiGeC HBTs were confirmed by SIMS spectrum. Upon
finishing the pedestal collector, the patterned wafers were loaded into UHV/CVD chamber to
finish epitaxial base growth after 10: 1 HF-last ex-situ wet chemical treatment. The
metallurgical base width was 15 nm and the spacers at either side were 30nm / 20nm to
provide buffer spaces for B outdiffusion. Electrical properties and flicker noise fluctuations
were measured thru metal contacts connected to poly emitter, extrinsic base and collector
sinker, respectively.
39
3.2.1. SIMS
The dopant distributions in the SiGe/SiGeC HBTs can be identified in the following SIMS
profiles, figure 3.3 and 3.4. Cs+ ion beam was used to probe about 2 um depth using a low
energy of 1 kev and a incident angle of 600. It should be noticed that Cs ion beam is not
preferred to analyze B dopant distributions due to its negative electricity making chemical
segregation effects, rendering undesired tails of B profiles. To get ”real” B dopant profiles,
oxygen ion beam is desirable to detect B. Interfacial C and O peaks can also be apparently
observed in this spectrum due to the unavoidable contaminations and negative oxides
existence. Compared to figure 3.3, the intended incorporated C peak is observed at the spacer
side of the base regions, and the concentration is approximately equal to 1E20 cm-3 in figure
3.4. All the other dopants (B, P, As, and Ge) distributions are also shown in SIMS profiles
with the detection limits of 1E16 cm-3.
SiGe HBT SIMS profile
Ge
As
C
O
1E20
0.20
0.15
0.10
1E19
P
1E18
0.05
B
Sb
1E17
0.3
0.4
0.5
Depth (um)
40
0.00
0.6
Ge mole fraction
Dopant conc
1E21
As
B
Ge
P
O
Sb
Ge
Figure 3.3 SIMS profile of the SiGe HBT. The dopants include B, P, As, and Ge [15].
SiGeC HBT SIMS profile
As
As
B
P
Sb
C
O
Ge
Ge
1E20
C
0.14
0.12
0.10
0.08
1E19
P
1E18
0.06
0.04
Ge mole fraction
Dopant conc.
1E21
O
Sb
0.02
1E17
0.3
0.4
0.5
0.00
0.6
Depth
Figure 3.4 SIMS profile of the SiGeC HBT. The concentration of intended incorporated C
peak is observed at 1E20 cm-3, equal to 0.2 % mole fraction [16].
Correspondent to our device configurations, the metallurgic base width ranges from 0.32
um to 0.38 um in figure 3.4. P dopants distribute from 0 to 0.32 um, which are sum of the
emitter poly of thickness 3000 nm and single crystal emitter of thickness 200 nm. Si epitaxial
collector ranges from 0.38 um to 0.70 um in depth. The thickness of 0.32 um is consistent
with nominal 0.3 um and the in-situ doped P concentration of 1E16 cm-3 is almost uniform
across the collector region. The distribution of Sb species, used to implant the Si substrate for
forming a low resistive contact with collector sinker, ranges from 0.70 um to 1.2 um. Figure
3.5 shows the comparison of B profiles between the SiGe and SiGeC HBTs. It’s exhibited that
the C incorporation of SiGe alloy is able to suppress the outdiffusion of B in the base region.
41
Nevertheless, it will diffuse more about 30 nm appealing in the SiGe HBT than in the SiGeC
HBT. Oppositely, the result that C will induce more diffusion of As dopant into the base
region is also observed in figure 3.5. This phenomenon is same with previous reported papers
because of As dopant diffusion exclusively by a vacancy mechanism [17].
As no C
B no C
As with C
B with C
-3
As & B concentration (cm )
1E21
1E20
1E19
1E18
1E17
0.30
0.35
0.40
0.45
0.50
Depth (um)
Figure 3.5 Comparisons of the As and B dopants profiles in the SiGe and SiGeC HBTs from
figure 3.3 and 3.4. C can suppress the diffusion of B, while induce the diffusion of As [18].
3.2.2 XRD
The DCD X-ray spectrum of SiGe and SiGeC control wafers is presented in figure 3.6. The
SiGe peak of the control wafer of the SiGeC HBT indicates apparently a angle shift compared
to that in the SiGe HBT. These results are consistent with those in last chapter. The strain
originally induced by a lattice constant mismatch of 4 % between Si and Ge atom will relax
42
by the incorporations of substitutional C. A shift (Δθ of 60 arcsec) is observed, assuming a
compensation of ratio of 1% C to 10.8 % Ge, this corresponds to 0.06 % of substitutinal C
incorporations.
10000
X-ray intensity
1000
SiGeC
HBT
SiGe HBT
SiGeC HBT
SiGe HBT
100
10
1
-3000
-2500
-2000
-1500
-1000
-500
0
500
arc-second
Figure 3.6 XRD spectra of the Si0.8Ge0.2 and Si0.798Ge0.2C0.002 HBTs. A angle shift is observed
at (400) SiGe peak [19].
3.2.3 DC measurements
The importance of limiting B diffusion during thermal processing is highlighted by the
electrical characteristics of the SiGe HBTs, which are sensitive to B diffusion lengths as small
as 1-2 nm. Gummel plots and common emitter characteristics of the processed HBT’s are
exhibited in figure 3.7 and 3.8. Collector currents (IC) and base currents (IB) were measured
under forward base-emitter voltage (VBE) sweep (from 0 to 1 V) and fixed reverse
43
collector-base voltage (VCB) 2 V. Similar IC and significant different IB curves are observed in
the Gummel plot, as shown in figure 3.7. Compared to the Si0.8Ge0.2 HBT (n = 1.1), the ideal
factor of IB curve in the Si0.798Ge0.2C0.002 device is estimated to n = 2.3 at VBE below 0.75 V.
Despite Ic is almost of no differences, but the magnitude of IB of the Si0.798Ge0.2C0.002 HBT
increases 2 orders at VBE = 0.4 V in comparison with that of Si0.8Ge0.2 HBT, hence eroding the
current gain (IC / IB). While referred with other reported results of SiGeC HBTs with SiCH6 as
precursors, the base currents were independent of the additions of C [20-21]. As the result, the
increase of IB in our samples may be probably due to the interface defects between Si and
SiGeC layers, which act as recombination centers to deplete back-injected holes, which
Ib,Ic (A)
renders negative affects on the base currents.
10
-2
10
-3
10
-4
10
-5
10
-6
10
-7
10
-8
10
-9
10
-10
10
-11
10
-12
10
-13
10
-14
Ib of SiGe HBT
Ic of SiGe HBT
Ib of SiGeC HBT
Ic of SiGeC HBT
0.2
0.4
0.6
0.8
1.0
VBE(V)
Figure 3.7 Current voltage diagrams of SiGe and SiGeC HBTs. Higher non-ideal base current
at the low applied voltage in the SiGeC HBT is exhibited than in the SiGe HBT [22].
44
A improved collector current of the SiGeC HBT is observed in figure 3.8. Under a fixed
base-emitter forward bais, the better current-driving capacity is exhibited in the SiGeC HBT
than SiGe HBT. In the bipolar devices, the collector current is porptional to Vbe, Nieb (the
intrinsic number in the base)
800n
Vbe=0.65V
700n
600n
IC(A)
500n
400n
300n
SiGe HBT
SiGeC HBT
200n
100n
0
0.0
0.5
1.0
VCE(V)
1.5
2.0
Figure 3.8 Output characteristics of SiGe and SiGeC HBTs. The Ic was measured at fixed
Vbe = 0.65 V [23].
The variety of base currents in the current voltage diagrams of the SiGe and SiGeC HBTs is
calculated over average 15 dies on one wafer and exhibited in Weibull plot. Collector currents
are almost the same of magnitude, but a large deviation of base currents on the SiGeC HBTs
is observed in figure 3.9.
45
3
2
ln(-ln(1-F))
1
0
-1
no carbon
with carbon
-2
-3
0.0
1.0x10
-7
2.0x10
-7
-7
3.0x10
-7
4.0x10
-7
5.0x10
6.0x10
-7
Ib (A)
Figure 3.9 Weiblll plots of base currents from average 15 dies on the SiGe and SiGeC HBTs.
3.3.4 Noise figure
Figure 3.10 shows the noise figure characteristics of the Si0.8Ge0.2 and Si0.789Ge0.2C0.002
HBTs. Each area of test pads in both devices is 0.7 * 20 μm². Both traces achieve minimum
noise figure (NFmin) at about 0.8 GHz frequency, then ascend as the exerted frequency
increases. General speaking, the thermal noise is related with fT and current gain by the
following formula: NFmin  1 
1

 2 g m rb
2
 f 
   [24]. More defects generated with C
  fT 
1
incorporations making larger surface recombinations if more interstitial C act as the
recombination traps. That increases base currents and decreases the current gain (Ic/Ib) in turn.
A lower current gain is dominant in explaining slightly larger NFmin in the SiGeC HBT
46
compared to the SiGe HBT.
4.0
Noise Figure (dB)
3.5
3.0
2.5
with carbon
2.0
1.5
no carbon
1.0
0
1
2
3
Frequency (Gz)
Figure 3.10 Characteristics of noise figure as a function of the operated frequency. Larger
NFmin appears in the SiGeC HBT compared to that in SiGe HBT [25].
Figure 3.10 indicates the dependence of NFmin with the collector current. These
common-emitter characteristics are measured at VCE = 1.5 V and exerted frequency = 1.2
GHz. As expected, when the bias increases, the NFmin of SiGeC HBT becomes larger than
that of SiGe HBT. This issue may be attributed to the effect of forming conjugated complex
from high mobile interstitial C atoms.
47
2
SiGe HBT(0.7X20_6m )@Vc=1.5V&Freq=1.2GHz
2.4
NFmin (dB)
2.0
with Carbon
1.6
no Carbon
1.2
2
4
6
8
10
12
Collector current IC (mA)
Figure 3.11 Characteristics of noise figure minimum (NFmin) as a function of the collector
currents. Larger NFmin appears in the SiGeC HBT compared to that in SiGe HBT [26].
3.3.5 Flicker noise
The Flicker noise of SiGe HBTs is dominated by defects states existence in the base by the
expression: SIB = K × IB2 / f (K: the factor related to the defect density; f: the operation
frequency). The SIB value exhibits the noises generated during the transportation of holes
from the base to the emitter. The results of flicker noise measurements of SiGe and SiGeC
HBTs are shown in figure 3.12. Least SIB of the device without C incorporation was observed,
indicating the lower defects levels in the SiGe alloy. However, as C incorporated into the
stacked base region, the defects induced by C of conjugated complexes will cause larger
noises. Figure 3.12 indicates that the noises of Si0.798Ge0.2C0.002 HBT increase 8X than those
48
of S0.8Ge0.2 HBT when measuring at fixed IB of 100 uA and operation frequency of 10 Hz, but
no larger noises are observed in the reported papers using SiCH6 as precursors. Nevertheless,
the conclusion of noise measurements is consistent with those of above noise figure
characteristics and Gummel plots.
Si0.798Ge0.2C0.002 base
1E-18
2
SIB(A /Hz)
1E-17
1E-19
Si0.8Ge0.2 base
1E-20
10
100
1k
Frequency(Hz)
Figure 3.12 Flicker noises of Si0.8-yGe0.2Cy HBTs with y = 0 or 0.2 % measured at the
frequency from 1 Hz to 10 KHz and exerted base current of 100 uA [27].
3.3.6 RF characteristics
The incorporations of C were demonstrated to have no negative influence on the
performance of SiGe HBT [28]. C was proven to provide reduced B diffusivity about 1/8 [29]
or enhanced thermal stability of SiGe epitaxial bases about 100 0C. In our experiments, the fT
peak value in the SiGeC HBT appears at 74 GHz larger than 72 GHz in the SiGe HBT, as
49
shown in figure 3.13. The fT is related with parasitic delay, base and collector transient
1
1
Cbc  Cbe    b   e  Cbc rb  rc  [24]. Although no

2fT g m
times by the formula:
significant improvement within C incorporations using C2H4 as C sources, the transient speed
seemed to be not affected by the defects formed by interstitial C in the SiGe base. A little
larger fT peak with the SiGeC HBT can be attributed to that the metragullar base width of
SiGeC HBT is shorter about 30 nm than that of SiGe HBT due to less outdiffusion of B.
Fmax can be related with fT by the relationship: f max 
fT
. Capactiances of
8Cbc Rb
Base-Collector and fT are major concerns in achieving high fmax. Smaller Cbc in the SiGeC
HBT will be apparent than that in SiGe HBT due to less B outdiffusion into the collector
region.
fT,fmax(GHz)
0
Jc(A/cm2)
-3
5x10
80
75
70
65
60
55
50
45
40
35
30
25
20
15
10
5
0
-2
1x10
2x10
-2
-2
2x10
3x10
-2
3x10
-2
-2
4x10
-2
4x10
-2
5x10
fT SiGe HBT
fmax
fT SiGeC HBT
fmax
0
5m
10m 15m 20m 25m 30m 35m 40m 45m
Ic(A)
Figure 3.13 RF comparisons between the SiGeC and SiGe HBTs. Higher fT and fmax in the
50
SiGeC HBT can be observed than in SiGe HBT [30].
3.4 Summary
A comparison of performances between the SiGe and SiGeC HBT is investigated in this
chapter. According to previous statements and experiments, we conclude that the
incorporation of C into the SiGe alloy would be difficult due to the little fraction of
substitutional C contents by using C2H4 as C sources. From the DC and RF characteristics
measured through the SiGeC HBTs, the larger non-ideal component of base currents indicates
defects formation influencing the carrier transport, as consistent with the results of noise
measurements. By suppressing the B outdiffusion in the base region, the SiGeC HBT has a
slight larger peak value of fT and fmax traces, but poor yield and high non-uniformity make it
a larger challenge in commercial fabrications.
Acknowledgement
The author wants to acknowledge for the help of L. S. Chiang and Y. M. hsu (both in ERSO)
for electrical and noise measurements and grateful to thank Y. D. Zeng and P. S. chen (both in
ERSO) for all XRD measurements, and SiGe/SiGeC epitaxy growth. This work was in finial
supported by ERSO / ITRI, Taiwan, ROC, under contract No: A321XS6700 for publishing
this article.
51
References
[1] L. D. Lanzerotti, J. C. Sturn, E. Stach, R, Hull, T. Buyukimanli. And C. Magee, Appl. Phys.
Lett. 70, 3125, 1997.
[2] J. D. Cressler, J. H., Comfort, E. F. Crabbe, G. L. Patton, J. M. Stork, J. Y-C. Sun and B. S.
Meyerson, Trans. Electron Device, 40, 525, 1993.
[3] D. L. Harame, J. H., Comfort, E. F. Crabbe, J. Y-C. Sun, B. S. Meyerson and T. Tice, Trans.
Electron Device, 42, 442, 1995.
[4] H. J. Osten, M. Kun, G. Lippert and P. Zaumseil, Thin Solid Films, 294, 93, 1997.
[5] C.W. Liu, Y. D. Tseng, and Y. S. Huang, Appl. Phys. Lett. 75, 2271 (1999).
[6] A. Gruhle, H. Kibbel, and U. Konig, Appl. Phys. Lett. 75, 1311 (1999).
[7] H. Rucker, B. Heinemann, W. Ropke, R. Kruger, G. Lipert, and H. J. Osten, Appl. Phys.
Lett. 73, 1682 (1998).
[8] L.D. Lanzerotti, A. St. Amour, C. W. Liu, J. C. Sturm, J. K. Watanabe, and N.D.
Theodore, IEEE Electron Device Lett. 17, 334 (1996).
[9] J. Mi, P. Warren, P. Letourneau, M. Judelewicz, M. Gailhanou, M. Dutoit, C. Dubois, and
J. C. Dupuy, Appl. Phys. Lett. 67, 259 ~1995.
[10]I. M. Anteney, G. Lippert, P. Ashburn, H. J. Osten, B. Heinemann, G. J. Parker, and D.
Knoll, IEEE Electron Device Lett. 20, 116 (1999).
[11] S´ebastien Jouan, IEEE Trans. Electron Devices 7, 46 (1999)
[12] S. Im, J. Washburm, R. Gronsky, N. W. Cheung, Y. M. Yu, and J. W. Ager, Appl. Phys.
Lett. 63, 2682 (1663)
[13] F. Y. Huang and K. L. Wang, Appl. Phys. Lett.. 69, 2330, 2002.
[14] C. W. Liu, A. St. Amour, and J. C. Sturm J. Appl. Phys. 80, 3043 (1996).
52
[15] Process was executed in ERSO, ITRI, Hsinchu.
[16] Process was executed in ERSO, ITRI, Hsinchu.
[17] H. Rucher, B. Heniemann, D. Bolze, D. Knoll, D. Kruger, and P. Zaumseil, IEDM, 345,
1999.
[18] Process was executed in ERSO, ITRI, Hsinchu.
[19] Process was executed in ERSO, ITRI, Hsinchu.
[20] D. Knoll, B. Heinemann, H.J. Osten, K. E. Ehwald, B. Tillack, P. Schley, R. Barth, M.
Mathes, Kwang Soo Park, Young Kim, and W. Winkler, Tech. Dig. Int. Electron Devices Meet.
703 (1998).
[21] A. C. Mocuta and D. W. Greve, J. Appl. Phy. vol. 85, pp. 1240, 1999.
[22] Process was executed in ERSO, ITRI, Hsinchu.
[23] Process was executed in ERSO, ITRI, Hsinchu.
[24] J. D. Cressler and G. Niu, “Silicon-Germanium Heterojunction Bipolar Transistors”,
ARTECH HOUSE, INC, 2002
[25] Process was executed in ERSO, ITRI, Hsinchu.
[26] Process was executed in ERSO, ITRI, Hsinchu.
[27] Process was executed in ERSO, ITRI, Hsinchu.
[28]H. J. Osten, D. Knoll, B.Heinemann, H. Rucker, and B. Tillack, Proceeding of BCTM, PP.
109, 1999.
[29] Malcolm S. Carroll, Ph.D. Thesis, Princeton, JUNE, 2001.
[30] Process was executed in ERSO, ITRI, Hsinchu.
53
Chapter 4
SiGeC HBTs on SOI
4.1 Introduction
SiGe BiCMOS using SOI technology has been demonstrated for microwave and
high-speed digital applications [1-5]. The devices fabricated on SOI can improve the RF and
reduce cross-talk for mixed signal circuits [6]. This schematic diagram of the device adapted
from IBM research group is exhibited in figure 4.1. Typically, SOI serves as thin collector and
buried oxide work as excellent isolations from substrates.
Figure 4.1 Schematic diagram of SiGe HBT fabricated on SOI substrate [7].
54
The thin buried oxide thickness (< 0.2μm) used in SOI devices can accommodate the
BiCMOS process and are mostly used in high-speed and high power applications. IBM group
[6-7] has demonstrated that a SiGe HBTs on SOI substrates with a fully depleted collector can
offer higher base-collector breakdown voltage, higher Early voltage and better BVCEO-fT
tradeoff through experiment and simulation. In this work, the device parameters of SOI such
as the buried oxide thickness and lateral distance between the collector and the reach-through
regions (Lcol) are optimized.
4.2 Structures and Simulation method
The schematic structure of a Si1-xGex HBT on SOI substrate is shown in figure 4.2.
Figure 4.2 Schematic of a SOI SiGe HBT. Lcol: the distance from the n+ reach-through to the edge of
the intrinsic base. Tox: buried oxide thickness.
The bulk device for comparison has collector of 0.3 μm thick and a heavy n+ doped buried
layer. The buried oxide thickness is varied from 0.06 μm to 0.2 μm. The distance form the n+
reach-through to the edge of the intrinsic base (Lcol) is varied from 0.1 μm to 0.2 μm for SOI
devices and is fixed at 0.15 μm for bulk SiGe HBT. For both type of devices, the base widths
55
are ~75 nm, the emitter sizes is 0.14 μm×10 μm. Doping profile and Ge mole fraction follows
the SIMS data of ERSO baseline SiGe HBT process. Emitter doping is ~ 10 20cm-3, base
doping is ~1018cm-3, and collector doping is 1017cm-3. For simplification, only single-emitter
structure is adopted in the device, and no metal depositions are considered. The substrate is
grounded in our simulation. Both DC and AC characteristics are studied using commercial
device simulator from ISE (DESSIS) [8].
For the SiGe material parameters, we used a linear energy band gap reduction with Ge [9].
The band gap narrowing due to heavy doping is also considered. At the Si/SiGe interface the
conduction band discontinuity ΔEC was taken as zero, and the valence band discontinuity is
equal to the band gap reduction ΔEV = ΔEg [10]. The effective density of states for the
conduction and valence bands, NC and NV, respectively, were calculated using Ref. 11. For
low doping levels, the electron’s minority carrier lifetime in the base was taken as 5 μs, and
decreases with increased doping, similar to silicon. The electron mobility and its doping
dependence on high doping levels and Ge content was taken as that of silicon using
Klaassen’s model [12]. For the dependence of electron saturation velocity on Ge, the results
of Ref. 13 are used. The dependence of hole mobility on the Ge concentration was taken as a
quadratic function with the same dependence on doping as in silicon [14].
4.3 Buried oxide thickness effect
DC Characteristics
A comparison of the output characteristics is shown figure 4.3. The larger Early voltage V A
are observed in SOI device due to lower base-collector capacitance CBC. The formula for
Early voltage VA is as [15]
VA 
qD nB ( WB ) n 2 ieB ( WB ) WB
N B (x)
dx

2
C BC
0 D nB ( x ) n ieB ( x )
The silicon collector on buried oxide is fully depleted, and the depletion region is extended
laterally. Thus, the total effective depletion width at B-C junction will increase, i. e. CBC
decreases. The CBC value in the equivalent small signal circuit is extracted from simulated S
parameters.
56
40
-7
IB=10 A/m
Ic(A/m)
30
Bulk
20
SOI
10
-8
IB=10 A/m
0
0.0
0.2
0.4
0.6
0.8
1.0
VCE (V)
Figure 4.3 Output characteristics of bulk and SOI SiGe HBTs. The base current varies form
0.01 μA/μm to 0.1 μA/μm. Early voltage of SOI device is larger for than that of bulk
Figure 4.4 shows Early voltage and CBC as functions of the buried oxide (BOX) thickness.
CBC keeps relatively the constant as the buried oxide thickness is thicker than 0.08 μm, but is
50f
500
40f
400
30f
300
20f
200
Early Voltage (V)
CdBC (fF/m)
smaller than that of bulk devices.
10f
100
0
0
0.08
0.1
0.15
Oxide thickness (m)
0.2
Figure 4.4 Early voltage and CBC v.s oxide thickness. CBC is almost not change as oxide thickness
increases. VA decreases slightly as oxide thickness increases.
57
Figure 4.5 shows BVCEO and maximum electric field at B-C junction as function of BOX
thickness. The base open breakdown voltage (BVCEO) increases as BOX thickness increases.
The electric field at B-C junction decreases as BOX thickness increases, due to the extra
voltage drop in the BOX. The impact ionization rate is a strong function of electric field. The
lower electric field of the SOI devices indicates the higher breakdown voltage (BVCEO)
4.10
5
4.05
5
BVCEO (V)
1x10
4.00
5
1x10
3.95
5
1x10
Electric Field (V/cm)
1x10
3.90
5
0.10
0.15
0.20
1x10
Oxide thickness T ox (m)
Figure 4.5 BVCEO and maximum electric field at B-C junction vs oxide thickness. BVCEO
increases as oxide thickness increases. Peak electric field decreases as oxide thickness
increases.
AC Characteristics
The AC characteristics are obtained from small signal analyses. The unilateral power gain
and current gain are calculated using S parameters. The maximum oscillation frequency and
cut-off frequency are extracted at unitary power gain and current gain, respectively. Figure 4.6
shows cutoff frequency fT and CBC are as functions of BOX thickness. As BOX thickness
increases, fT starts to decrease rapidly and then decreases slightly. This is due to CBC almost
remaining the same value as BOX thickness is larger than 0.08 μm. This behavior is
confirmed by SPICE simulation. BOX thickness affects the collector-substrate capacitance
CCS. Effective CCS is a series of Cox (BOX capacitance) and CCS_SOI (BOX –substrate
capacitance). It keeps the same in all BOX thickness devices and Cox decreases as BOX
58
thickness increases. SPICE simulation shows that variation of CCS do not affect cutoff
frequency seriously.
Figure 4.6 Cutoff frequency and CBC v.s oxide thickness. Cutoff frequency and CBC both decrease as
oxide thickness increases.
Figure 4.7 shows maximum oscillation frequency fmax and CCS are as functions of BOX
thickness. As BOX thickness increases, fmax starts to decrease due to the increase of CBC and
then increases. The capacitance Ccs starts to increase due to the CCS_SOI and then decreases as
BOX thickness increases. SPICE simulation demonstrated that variation of CCS affects fmax.
The variation of Cox is from variation of CCS_SOI due to change of BOX thickness.
59
1.4f
265
1.3f
fmax(GHz)
260
1.2f
255
1.1f
250
Ccs (fF/m)
270
1.0f
245
900.0a
240
0
0.06
0.08
0.1
0.15
Oxide thickness (m)
0.2
Figure 4.7 fmax and Ccs vs oxide thickness. fmax increases as oxide thickness increases while
Ccs decreases as oxide thickness increases.
4.4 Lateral scaling
The BOX thickness and the distance between the intrinsic base edge and the reach-through
region (Lcol) are optimized. Given BOX thickness of 0.15 μm, and collector doping of
1017cm-3, the both fT and fmax decreases as Lcol decreases as shown in figure 4.8. The increase
of the lateral depletion region as Lcol increases should be responsible for this.
60
300
fT/fmax(GHz)
250
fT
f max
200
150
100
50
0
0.10
0.15
Lcol (m)
0.20
Figure 4.8 fT and fmax vs Lcol. Both fT and fmax decrease with oxide thickness increasing.
However, the increase of Lcol can decreases the peak electric field (see figure 4.9) in the
devices, and thus increases the BVCEO, as shown in figure 4.10. The tradeoff between fT/fmax
and BVCEO due to the Lcol exists [6,7]. If fT*BVCEO product is used as a figure of merit, the
maximum value is obtained at Lcol = 0.15μm. Given Lcol length of 0.15μm, and collector
doping of 1017cm-3, the fT decreases as BOX thickness increases while fmax shows the opposite
behavior. However, the increase of BOX thickness (Tox) can decreases the peak electric field
in the devices, and thus increases the BVCEO, as shown in figure 4.10. The tradeoff between
fT/fmax and BVCEO due to the BOX thickness exists. If fT*BVCEO product is used as a figure of
merit, the maximum value is obtained at Tox = 0.1μm for fixed Lcol.
61
Figure 4.9 Lateral electric field distribution with three different lengths of Lcol. Buried oxide
thickness is fixed as 0.15μm
4.25
BVCEO (V)
4.00
3.75
3.50
0.10
0.15
0.20
Lcol (m)
Figure 4.10 Base open breakdown voltage (BVCEO ) as function of Lcol. BVCEO increases as
Lcol increases.
62
4.5 Summary
In summary, as compared to bulk SiGe HBTs, the SiGe HBTs on SOI substrates have a
higher Early voltage, maximum oscillation frequency, and breakdown voltage. Buried oxide
(BOX) thickness effect and the lateral distance between collector and reach-through region of
SiGe HBT on SOI substrates are investigated. A SiGe HBTs on SOI substrates with larger
buried oxide thickness can achieve a higher BVCEO and fmax with a comparable fT and VA. The
maximum value is obtained at Lcol = 0.15μm for fixed BOX thickness and Tox = 0.1μm for
fixed Lcol if fT*BVCEO product is used as a figure of merit.
References
[1] K. Washio, et. Al, IEDM Tech. Dig., pp. 741-744, 2000.
[2] L. Larson, IEDM Tech. Dig., pp.737-740, 2000.
[3] T. Hiramoto et.al, IEDM Tech, Dig., p39, 1992.
[4] E. Leobandung et al., IEDM Tech. Dig., p. 679, 1999.
[5] J. W. Sleight et al., IEDM Tech. Dig., p.245, 2001.
[6] J. Cai, et al., Symp. VLSI Tech. Dig., p. 172, 2002.
[7] From IBM, EE times special issues, 2002.
[8] Q. Ouyang et al., BCTM, pp. 28-31, 2002
[9] DESSIS User’s Manual, ISE, 2002.
[10] Z. Matutinovic-Krystelj et al. IEEE Trans. Electron Dev., Vol. 43, p. 457, 1996.
[11] J. S. Yuan et al., SiGe, GaAs and InP HBTs. New York, Wiley, p. 6, 1999.
[12] A. Neugroschel et al., IEEE Trans. Electron Dev, Vol. 47, p. 187, 2000.
[13] D. B. Klaassen, Solid-State Electron, Vol. 35, p. 953, 1992.
[14] Bufler et al., IEEE Electron Dev, Lett., Vol. 18, p. 264, 1997.
[15] T. K.Carns et al., IEEE Trans. Electron Dev., Vol. 41, p. 1273, 1994.
[16] Y. Taur et al., Fundamentals of Modern VLSI Devices, New York,Cambridge, p. 317,
1998.
63
Chapter 5
Discussions and Further Works
Discussions
For the wireless communications and portable applications, the SiGe technology had been
widely used in RF IC applications due to its low cost and high compatibility with now
prevailing Si industry. For the further demands of wireless LAN, hardware modules and parts
are nevertheless disposable so that markets grow more and more rapidly in the visible future.
We have successfully fabricated the SiGeC HBTs using C2H4 as C precursors and essential
electrical characteristics and material properties of the SiGe/SiGeC alloy are understood
through instrument analysis and measurements.
Typical performance of our SiGeC HBT compared to the SiGe HBTs is summaried as the
following table.
64
SiGeC
HBT
SiGe
HBT
Collector current(A)
larger
smaller
Base current(A)
larger
smaller
Current [email protected]=0.65V
quite
quite
Noise Figure
minimum(dB)@frequency=1.2G
Hz,Vc = 1.5 V
larger
smaller
Flicker noise(A2/Hz)@Ib=100
uA,frequency=100Hz
larger(10X)
smaller
fTpeak(GHz)
larger(5%)
smaller
fmaxpeak(GHz)
larger(35%)
smaller
Early Voltage(V)
larger(>2X)
smaller
Performance characteristics
Future works
One of the key problems in NPN SiGe technology is to retain the narrow as-grown boron
profile with in the SiGe base layer during post-epitaxial processing. Heat treatments and
transient enhanced diffusion (TED) caused by annealing implantation damage can broaden the
profile into the adjoining Si regions. This can cause parasitic barrier (conduction band
barriers). A way to reduce TED effects is to use a “carbon-doping” technology. Another issue
is about base thickness. For high speed purpose, we want to minimize it to avoid
strain-induced defect formation.
65
The negative effects of adding C into SiGe using C2H4 precursors were investigated. The
results of XRD and SIMS measurements explicitly show that a further increase of the C2H4
flow leads to an increase of the carbon incorporation into interstitial sites, which translates
into a rapid deterioration of the Si1-x-yGexCy layer. For growth conditions similar to our
experiments (550 C growth temperature in UHVCVD, growth rate of 18 A/min and total
carbon concentration of 0.2 % from SiCH6), almost 90% of the carbon is substitutional. Our
data for alloys of similar carbon concentration, we observe that about 30 ~ 40 % of the carbon
is substitutional. The narrow process window for using C2H4 to satisfy the demand of both
reducing CI and suppressing B outdiffusion is restricted at total C fraction < 0.2 % in
Si1-x-yGexCy layers in our system. If the precise control of epitaxy process is executed, the
C2H4 precursors are still utilized to avoid inducing disorders and defects of epitaxial SiGeC
films and beneficial to modern electronic devices. We are interested in fabricating SiGeC
HBTs using other candidates of C precursors, like methyl-siliane.
66