# Three-Dimensional Silicon-Germanium Nanostructures for CMOS Compatible Light Emitters and Optical Interconnects

Authors: L. Tsybeskov; E.-K. Lee; H.-Y. Chang; B. V.  Kamenev; D. J. Lockwood; J.-M. Baribeau; T. I. Kamins
Journal: Advances in Optical Technologies
 (2008)

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## Abstract

Three-dimensional SiGe nanostructures grown on Si (SiGe/Si) using molecular beam epitaxy or low-pressure chemical vapor deposition exhibit photoluminescence and electroluminescence in the important spectral range of 1.3–1.6μm. At a high level of photoexcitation or carrier injection, thermal quenching of the luminescence intensity is suppressed and the previously confirmed type-II energy band alignment at Si/SiGe cluster heterointerfaces no longer controls radiative carrier recombination. Instead, a recently proposed dynamic type-I energy band alignment is found to be responsible for the strong decrease in carrier radiative lifetime and further increase in the luminescence quantum efficiency.


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## Body

## 1. Introduction

Optical
interconnects in the form of fiber optics have been used for many years in
different long-distance communication applications [1, 2]. With the
microprocessor clock speed approaching 10 Gbps, optical interconnects are now
being considered for board-to-board and on-chip technology as an alternative to
metal wires with their unavoidable RC delay, significant signal degradation, problems
with power dissipation, and electromagnetic interference [2–5]. Two major
avenues toward optical interconnects on a chip comprise the hybrid approach
with III-V optoelectronic components densely packaged into complementary
metal-oxide semiconductor (CMOS) architecture [6–9] and the all-group-IV (e.g.,
Si, SiGe, SiGeC, etc.) approach with the all major components, for example,
light emitters, modulators, waveguides, and photodetectors, monolithically
integrated into the CMOS environment [10]. There have been great efforts over
the past several decades to obtain technologically viable and efficient light emission
from group-IV materials. In the visible spectral region, the main emphasis has
been on porous silicon [11–13] and other Si nanostructured systems such as
silicon/silicon dioxide superlattices [14–17] and silicon nanoprecipitates in
silicon dioxide [18, 19]. In the near infrared spectral region, materials and systems
such as erbium in silicon [10, 20], silicon/germanium quantum wells [21, 22], and,
more recently, iron disilicide [23] offer potentially useful routes. However, no
approach has so far been applied commercially. The reasons for this are the
lack of a genuine or perceived compatibility with conventional CMOS technology,
the long carrier radiative lifetime in Si-based nanostructures, and, especially
in the case of near-infrared emitters, the high thermal quenching leading to
extremely poor room-temperature luminescence efficiencies [10].In
the 1990s, an interesting form of semiconductor nanostructure, namely, the
three-dimensional (3D) self-assembled system produced by the Stranski-Krastanov
or cluster-layer, growth mode in lattice mismatched materials, has been demonstrated
[24, 25]. In the case of 3D Si/SiGe nanostructures (NSs), the growth of mostly
dislocation-free samples has successfully been achieved using both molecular
beam epitaxy (MBE) and chemical vapor deposition (CVD) [26–31]. It has been
shown that the nonplanar geometry is mainly responsible for the significant
increase of the critical layer thickness in 3D Si/SiGe NSs grown on Si [28, 29].
It has also been found that compared to planar Si/SiGe quantum wells (QWs), the
photoluminescence (PL) and electroluminescence (EL) quantum efficiency in 3D
Si/SiGe NSs is higher, especially at temperatures T>50 K [32–35]. Despite
many successful demonstrations of PL and EL in the spectral range of 1.3–1.6μm, which is important
for optical fiber communications, the proposed further development of 3D
Si/SiGe NS-based light emitters was discouraged by several studies confirming a
type-II energy band alignment at Si/SiGe heterointerfaces [36–40], where the
spatial separation of electrons (located in Si) and holes (localized in SiGe)
was thought to make carrier radiative recombination very inefficient. Later, it
was also shown that 3D Si/SiGe NSs exhibit an extremely long (of the order of
10-2 seconds) luminescence lifetime [41], which is of the order of a
million times longer than those found in
III-V semiconductors. In addition, single crystal Si and Ge are indirect band
gap semiconductors and it has been declared that since carrier radiative
recombination is an indirect process in these materials, in both real and
reciprocal space, this process should be extremely inefficient. Thus, according
to this analysis, 3D SiGe NSs cannot be used to achieve efficient and
commercially valuable light emitting devices.In
this review paper, we show that despite the fact that bulk Si and Ge are indeed
indirect band gap semiconductors and that the Si/SiGe heterointerface most
likely exhibits type-II energy band alignment, it is still possible to obtain
conditions favorable for efficient carrier radiative recombination. We
demonstrate that the recent revised understanding of basic physics in such
systems has already helped to achieve nearly constant luminescence intensity at
temperatures4K<T<250K, and that the radiative carrier
recombination lifetime can successfully be reduced from 10-2 second
to 10-7 seconds, which is only ~10 times longer compared to those
found in direct band gap III-V semiconductors.


## 2. Growth and Structural Properties of Si/SiGe Three-Dimensional Nanostructures

The
standard approach to the fabrication of a 3D Si/SiGe NS, which for most
experiments discussed here is a multilayer Si/SiGe cluster system, is based on
the sequential physical sputtering of Si and Ge (SiGe) in MBE or the thermal
decomposition of SiH4 and GeH4 in CVD at a temperature in
the range T=550–650°C. At that temperature, both a high Ge solid
solubility in Si as well as strain induced SiGe interdiffusion due to the ~4%
lattice mismatch between Si and Ge are important. The MBE growth provides better
control over the average SiGe cluster composition, although, because of
interdiffusion during growth, the composition is not uniform within the cluster
volume [42–45]. Figure 1 shows typical transmission electron micrographs (TEM)
together with a summary of analytical TEM studies confirming both the 3D
geometry of the SiGe NSs and the complex atomic composition. In general, SiGe
cluster growth commences with the spontaneous development of a Si1-xGex planar, <1 nm thick, wetting layer where x varies, mainly due to uncontrolled SiGe
interdiffusion. With the further influx of Ge and Si, the growth mode then switches from two dimensional (a layer) to
3D (a cluster), which helps release some of the lattice-mismatch induced strain
[44]. The fully grown 3–10 nm high and initially nearly pyramidal-shaped SiGe
clusters have a Ge-rich (~50% depending on the Ge flux) core, although the
exact final cluster shape and composition strongly depends on the fabrication
conditions [45]. If the SiGe cluster is covered by a Si cap, the initial
pyramid-like cluster top is smoothed out (Figure 1(a)), and the capping Si layer
is locally strained, mostly near the top of the SiGe cluster. In multilayer
Si/SiGe cluster samples, this strain field propagates in the growth direction,
and it induces vertical SiGe cluster self-ordering (Figure 1(a)). Detailed
structural analysis also indicates that the Si in the valleys between SiGe
clusters is slightly compressed [30, 44]. To summarize, buried SiGe clusters
with the highest Ge composition of near 50% in the middle of the clusters are
surrounded by Si, which is tensile strained above each cluster and compressed
laterally between clusters to maintain a low overall strain. Each SiGe cluster
consists of Si1-xGex crystalline alloys with x increasing
toward the cluster center [43–45]. Thus, despite being fully crystalline
alloys, 3D Si/SiGe NSs with a high (~50%) Ge atomic concentration exhibit
significant embedded strain and compositional disorder [46]. These conclusions
regarding the structural properties of 3D Si/SiGe NSs are critically important
in understanding their optical characteristics and light-emitting properties.(a) Transmission of electron micrograph
showing a cross-section of an MBE grown, vertically self-aligned Si/SiGe
cluster (or 3D) multilayer structure. (b) Schematic showing typical Si and Ge
atomic concentrations within an SiGe cluster measured along a horizontal axis running
through the middle of the cluster (see the dashed line in Figure1(a)).
(a)(b)


## 3. Optical Properties of Si/SiGe Three-Dimensional Nanostructures

In
a system with strong selection rule relaxation, quasidirect carrier
recombination is possible, and it should provide a higher PL quantum efficiency
compared to that in an indirect band gap semiconductor, for example, single
crystal Si and Ge. In fact, many PL measurements in SiGe bulk alloys and
Si/SiGe NSs reveal a significantly enhanced intensity ratio between no-phonon
(NP) luminescence and phonon-assisted luminescence, as shown in Figure2(a). On
the other hand, in undoped bulk crystalline Si (c–Si), the NP PL
line intensity is negligible compared
to the transverse optical (TO) and transverse acoustic (TA) phonon-assisted radiative
transitions (see Figure 2(b)). Thus, systematic studies of the PL spectra in 3D
Si/Si1-xGex NSs with control over the average Ge atomic
concentration x provide very important information regarding changes in the
carrier recombination mechanism (e.g., selection rule relaxation, conduction
and valence band alignment, etc.) as x increases from 0 (c–Si) to ~55%.Comparison between
low temperature (T=4 K) PL spectra in (a) an SiGe wetting layer with ~20% Ge
concentration and (b) undoped crystalline Si (c–Si). Note the dramatically
different intensity ratios between the NP and TO phonon PL lines of each
sample.
(a)(b)Figure3(a) compares PL spectra in MBE samples with x=0.096, 0.16, and 0.53. In
addition to PL related to c–Si, we observe for x=0.096 relatively narrow PL
bands at 1.05 and 1.11 eV (Figure 3(a)) attributed to NP and TO phonon PL
bands in Si-rich SiGe alloys [33–35, 47, 48]. Note that these PL bands are in
the vicinity of c–Si luminescence, and that the intensities of these two sets
of PL bands are comparable. Therefore, we conclude that a small amount of Ge
(<10%) slightly reduces the SiGe band gap and mainly relaxes selection
rules, increasing the ratio of NP/TO PL band intensities. The observed
broadening of these two PL bands, compared to c–Si related PL, is, apparently,
due to the composition disorder resulting from the introduction of substitutional
Ge atoms (~10%) into the c–Si matrix.Figure 3
(a) Low
temperature (T=4 K) PL spectra in (a) MBE grown
Si/Si1-xGex 3D NSs with the indicated average Ge atomic
concentration x and (d) in Si/SiGe 3D NSs grown by CVD. The characteristic
phonon-assisted transitions are indicated. PL spectra in CVD samples show a
strong spectral shift toward higher photon energy when the excitation intensity
is increased by approximately 5 times from trace (i) to trace (ii) (note the
linear intensity scale).Increasing
the average Ge composition within Si1-xGex clusters up to
x=0.16 results in significant changes in the PL spectrum (Figure 3(b)). An intense PL band peaked at 0.95 eV
has appeared, showing an effective SiGe band gap reduction of ~150 meV
compared to c–Si. This broad and featureless PL band with a full width at half
maximum (FWHM) of ~70 meV indicates a much stronger compositional
disorder compared to 3D Si/SiGe NC samples with x=0.096.The
PL spectrum in the samples with an average Ge composition close to 53% (Figure3(c))
depicts a broad feature with a major PL peak centered at a photon energy of
0.75 eV, and this peak energy is close to the band gap of crystalline Ge
(c–Ge) at 4 K [49]. A second PL peak is found at ~0.85 eV. Such samples
have the highest PL quantum efficiency as compared to other samples with lower
x values. Both PL bands are quite broad, most likely due to compositional
disorder, which is in an agreement with the Raman scattering measurements
(discussed below).In
contrast to MBE grown samples, CVD growth of 3D Si/SiGe NSs does not provide precise
control over the Ge atomic composition, and, most likely, it produces more
Si/SiGe interdiffusion at heterointerfaces [42–44, 50]. This is well reflected
in PL spectra, where no fine structure has been found (see Figure 3(d)). However,
the broad and asymmetric PL peak is well fitted by two Gaussian bands often
identified as the NP and TO phonon lines [46, 47] that are separated by ~48 meV,
which is close to the energy of characteristic SiGe phonons. Thus, there is at
least a qualitative similarity between PL spectra in MBE and CVD grown 3D
Si/SiGe NSs.Figure4 summarizes the PL intensity as a function of excitation intensity in MBE samples
having different average Ge atomic concentrations. The same linear dependence
(on a log-log plot) for PL associated with c–Si and a sublinear (close to
square root) dependence for PL associated with a Ge-rich SiGe cluster core has
been found in nearly all 3D Si/SiGe NSs grown by both CVD and MBE [35, 51–53].The PL intensity
as a function of excitation intensity in MBE grown Si/Si1-xGex 3D NS samples with the indicated average Ge concentration x. The PL bands
associated with the cluster core, wetting layer (WL), and characteristic TO
phonons are indicated.
(a)(b)(c)Studies
of the PL temperature dependence show that at low temperatures the PL intensity
is nearly temperature independent. At higher temperatures the PL intensity
drops exponentially, and the activation energies of PL thermal quenching are
shown in Figure5. There is a clear correlation between the Ge composition in Si1-xGex 3D nanostructures and the PL intensity temperature dependence. For samples with
a low (x=0.096) Ge composition, the activation energy Ea is
~10 meV, and the PL has almost vanished by 40 K. This result could
be explained by the thermal dissociation of a nearly free exciton: a
localization mechanism associated with SiGe stochiometric fluctuations has been
proposed in
[53–55]. Figure 4(a) shows that the PL intensity as a
function of the excitation intensity is linear (again, on a log-log scale), and
that is consistent with the assumption of nearly free exciton PL.The PL intensity
temperature dependence in MBE grown Si/Si1-xGex 3D NS
samples with the indicated average Ge concentration x measured for different PL
bands. Figure 5(c) compares the PL intensity temperature dependencies measured
for PL bands with peaks at 0.75 eV (SiGe core) and 0.85 eV.
(a)(b)(c)In
samples with a higher (x=0.16) Ge composition, the activation energy of PL
thermal quenching is increased and the PL is observable up to ~100 K (Figure
5(b)). The broad PL band observed at 0.95 eV for the Si/Si1-xGex sample of x=0.16 exhibits an activation energy of ~25 meV, which
combined with a sublinear dependence in PL intensity as a function of
excitation intensity (Figure 4(b)) is in contrast with the x=0.096 sample. This
observation suggests that the nonequilibrium carriers are spatially localized
and that Auger recombination contributes to the overall recombination mechanism
even at very low (~100 mW/cm2) excitation intensity [54, 55].
Most likely, in 3D Si/Si1-xGex NSs with x=0.16, the
carriers (possibly holes) are localized within 3D SiGe quasiwells.In
3D Si/Si1-xGex nanostructures with x=0.53,
the PL spectrum contains two bands peaked at 0.85 eV and 0.75 eV
(Figure 3(c)). There are no characteristic phonons in the Si/SiGe system having
an energy of ~100 meV, and it is thus reasonable to assume that the
observed PL bands are associated with carrier recombination within two
different regions of the 3D SiGe nanostructures. The PL band peaked at
0.85 eV has almost the same PL quenching activation energy (~20 meV) as in the sample with x=0.16, while the PL
intensity as function of excitation intensity is linear over a wide range of
excitation intensities. The second PL band peaked at 0.75 eV has an
activation energy of ~60 meV (Figure 5(c)) and is nearly temperature
independent up to 100 K. Because of its high quantum efficiency, it can be
monitored almost up to room temperature. This data suggests that 3D Si/Si1-xGex nanostructures with x=0.53 contain coupled subsystems with different (lower
and higher) Ge concentrations. It is quite possible that a spatial localization
of electron-hole pairs within 3D regions of higher Ge concentration and thus
having a lower band gap could be responsible for the observed sublinear
excitation dependence of the PL band at 0.75 eV (Figure 4(c)). The remaining 3D regions with a lower Ge
concentration (e.g., having a higher band gap) have a lower carrier
concentration and the PL band (peaked at 0.85 eV) exhibits a liner excitation
dependence.A
continuous shift of the PL band from ~1 to 0.75 eV has been found previously in
SiGe alloys with increasing Ge concentration [56]. Instead, in these 3D Si/Si1-xGex samples with a Ge concentration higher than 50%, a simultaneous threshold-like
appearance of two clearly resolved PL peaks at 0.85 and 0.75 eV is observed.
This suggests that Ge segregation might take place as x increases up to ~0.5.
Since the PL peak at 0.75 eV essentially matches the value of the band gap
in pure c–Ge, we propose that such a segregation results in a Ge-rich core within
an SiGe shell forming the 3D Si/Si1-xGex NSs embedded
within a pure Si matrix.The
PL thermal quenching observed in 3D Si/Si1-xGex samples grown
by MBE could be associated with different mechanisms. As mentioned earlier, the
activation energies of the PL thermal quenching in samples with x=0.096 and x=0.16 are close to the observed exciton binding energy in SiGe alloys and
Si/Si1-xGex nanostructures [53–57]. In samples with x=0.53,
a greater activation energy could be attributed to carrier diffusion from a 3D
potential well. This assumption is justified by the expected type-II band
alignment in SiGe nanostructures with a deep (>100 meV) potential well for
holes and a relatively small potential barrier for electrons [38, 40, 41]. In
this model, phonon-assisted carrier tunneling can produce the observed ~60 meV
activation energy for thermal quenching of the PL intensity. The PL intensity behavior
as a function of temperature in CVD grown samples is more complex and will be
discussed later.In
3D Si/SiGe NSs, the PL properties strongly correlate with the sample structural
properties, and Raman scattering under excitation conditions similar to PL
measurements has been shown to be a very informative characterization technique
[35]. In fact, in SiGe materials and systems, Raman spectroscopy is a unique
characterization technique due to the multimodal nature of Raman scattering
from optical phonons in SiGe, which reveals all three major vibrational modes known,
respectively, as the Si–Si vibration at ~500 cm-1, the Si–Ge
vibration at ~400 cm-1, and the Ge–Ge vibration at ~300 cm-1 (see Figure 6, e.g.). In addition, Raman phonon spectroscopy is a very sensitive
probe of global and local embedded strain and of compositional and structural
disorder. Figure 6 compares Raman spectra for MBE samples of different Ge composition.
In Si1-xGex clusters with x=0.096, an intense optical
phonon Raman signal associated with Si–Si vibrations is observed at ~520 cm-1 together with a weaker feature at ~300 cm-1 related to second-order
scattering from Si acoustic phonons [58]. No significant Raman peaks related to
Ge–Ge vibrations at ~290 cm-1 and Si–Ge vibrations at ~420 cm-1 [59] and to an amorphous Si phase at ~480 cm-1 [60] were found in
this or the x=0.16 sample (see Figure 6, noting that the intensity axis has a log
scale).Raman spectra measured
using 514 nm excitation in MBE grown Si/Si1-xGex 3D NS
samples with the indicated average Ge concentration x.
(a)(b)(c)A
further increase in Ge composition tox=0.53 in 3D Si1-xGex NSs produces strong Raman signals at ~290 cm-1 (related to Ge–Ge
bonds) and at ~420 cm-1 (related to Si–Ge bonds), as shown in Figure 6(c).
Using a short wavelength photoexitation (458 nm) and multilayer Si/SiGe 3D
samples, the Raman signal from the Si substrate is minimized (see Figure 7). In
addition to the major vibrational modes indicated in Figure 7(a), a doublet near
520 cm-1 related to compressed (~522 cm-1) and strained
(~507 cm-1) Si is observed (see Figure 7(b)). The spectrum also
exhibits weak and broad (background) Raman peaks in the vicinity of 480–490 cm-1 and 250 cm-1 related to disordered phases of Si and Ge, respectively.
Thus, despite the absence of purely amorphous Si and Ge phases in these 3D
nanostructured fully-crystalline alloys, compositional disorder and
strain-induced lattice distortion combine to produce similar effects in the Raman
spectrum.Raman spectra measured
using 458 nm excitation in MBE grown Si/Si1-xGex 3D NS
samples with x approaching 50%. (a) A full-range spectrum comparing the major
SiGe Raman modes with the positions of the Raman peaks seen in a-Si and a-Ge.
(b) Raman spectrum on a linear intensity scale showing Si–Si vibrational modes
including the vibration associated with strained Si at ~507 cm-1.
(a)(b)


## 4. Polarized Raman Scattering in Si/SiGe Three-Dimensional Nanostructures

In
addition to conventional inelastic light scattering spectroscopy, polarized
Raman scattering provides information on Raman scattering intensity as a
function of polarization angle. The result is presented employing polar plots and
has been shown to be an informative tool for the analysis of local embedded
strain in 3D Si/SiGe NSs with larger dome-shaped SiGe clusters grown on Si
(100) substrates [61]. In these measurements, the incident light is usually polarized
in the plane of the incident and scattered light, and polarization of the
scattered light is analyzed using a thin-film polarizer that can be rotated
through 360°.Figure8 shows angular Raman polarization diagrams for three different alloy vibrational
modes (Si–Si at 520 cm-1,
Si–Ge at 415 cm-1,
and Ge–Ge at 298 cm-1)
for multilayer dome/island samples, as well as the Raman polarization dependence
for the Si–Si vibration at
520 cm-1 measured in 〈100〉 oriented single-crystal Si as a
reference. The observed angular dependencies in the polar plots of the Ge–Ge and Si–Ge Raman mode
intensities are nearly identical to that in a 〈100〉 Si single crystal.
However, a quite different behavior is observed in the Raman polarization
dependence for the Si–Si vibration at ~519 cm-1.Figure 8
Polarization Raman
polar diagrams in CVD grown SiGe 3D NS samples measured (a) for Si–Si, Si–Ge,
and Ge–Ge vibrational modes using 458 nm excitation and (b) for the Si–Si
vibrational mode using 458 and 514 nm excitation. The results obtained are
compared with that of c–Si. (c) Raman spectra on a linear intensity scale in
the vicinity of Si–Si vibrational modes in CVD-grown SiGe 3D NSs measured using
two different polarization angles compared with that of c–Si.These
different results for the Si–Si mode are emphasized in Figure8(b), where the
Raman spectrum collected at 70° exhibits a major Raman peak at ~520 cm-1 slightly shifted toward higher wave numbers compared to the data collected at
340°. If this shift is associated with a built-in local strain, the strain can
be estimated to be ~0.1 GPa [62, 63]. Higher resolution Raman measurements show
that the major Raman peak at ~520 cm-1 is not only just slightly
shifted but also broader compared to the spectrum collected from single crystal
Si. In addition, we observe a second, much broader peak at ~504 cm-1. The angular polarization
dependence of the Raman signal at 504 cm-1 also deviates from the
〈100〉 oriented single-crystal Si reference, although this deviation is
less significant compared to the Raman peak at ~520 cm-1. The strong
localization of strain observed in the Si matrix is consistent with our understanding
of the nature of Stranski-Krastanov growth,
where vertical self-ordering is produced by strain propagation through the
separating Si layers [64].


## 5. Photoluminescence Properties of Si/SiGe
Three-Dimensional Nanostructures as a Function of Excitation Intensity

It
has been known for some time that the PL spectra in 3D Si/SiGe NSs, which is similar
to that in III-V quantum wells with type II energy band alignment [65], exhibit
a blue shift as the excitation intensity increases [55, 66]. In our studies, this
effect is found in both MBE and CVD grown samples. Figure 9(a) shows PL spectra
in a CVD grown sample measured under excitation intensities varied from 0.1 to
100 W/cm2. At the lowest excitation intensity used (0.1 W/cm2), the PL
peaks at ~0.8 eV. With increasing excitation intensity, a continuous almost
parallel PL blue shift of 30–40 meV per decade of excitation intensity increase
is observed. At an excitation intensity of 100 W/cm2, the PL peak
reaches ~0.92 eV. Under photoexcitation of 1–10 kW/cm2, the low
energy part of the PL spectrum does not shift further, while the high energy
part continues shifting toward higher energy.PL spectra in CVD
grown Si/SiGe 3D NSs (a) measured atT=4 K under different excitation
intensities from Imin (0.1 W/cm2) to Imax (100 W/cm2) and (b) under a fixed excitation intensity of 5 W/cm2 at the indicated temperatures (the PL spectra have been shifted vertically for
clarity).
(a)(b)Figure9(b) compares PL spectra measured with a fixed excitation intensity (~5 W/cm2)
for different temperatures in the range 8–210 K. These
measurements clearly show that with increasing temperature, the PL peak
associated with SiGe clusters at 0.85–0.9 eV shifts
toward lower photon energies. This “red” PL spectral shift is, most likely,
associated with the SiGe band gap decrease as the sample temperature increases.
Thus, it is opposite to the “blue” PL spectrum shift observed with an
excitation intensity increase, and thus this blue shift cannot simply be
explained by sample heating due to the intense incident laser beam. There was
some sample heating in the measurements, but it was not significant except at
the higher laser powers. From the observed broadening of the PL spectral features
associated with c–Si, the sample temperature is estimated to increase from 4 to
60 K under the highest excitation intensity used.Despite
the clear shift toward higher photon energy under increasing excitation
intensity, the entire set of PL spectra recorded at photon energies <0.95 eV are well-fitted by two Gaussian peaks (shown in
Figure10(a) for the PL spectrum (1) obtained with 0.1 W/cm2 excitation). Figure 10(b) summarizes the excitation intensity dependence of the
energies of these two PL peak and their widths. For all PL spectra obtained at
different excitation intensities, the two PL peaks at photon energies of ≤0.95 eV are found to be always separated by ~43 meV with a constant FWHM of ~78 and
~47 meV, respectively. These two PL peaks are attributed to the NP transition
and TO phonon replica in the SiGe clusters, respectively, and the separation
energy of ~43 meV is close to the energy of the TO Si–Ge phonon
[30, 32–35, 47, 52, 53, 54, 55]. The constant energy separation of the two PL peaks is additional evidence
that, at the excitation intensities used, the observed PL blue shift is not due
to thermal broadening caused by sample heating with the laser beam, and that
the integrated PL intensities measured at different excitation intensities can
be directly compared.(a) Normalized PL
spectra in CVD grown Si/SiGe 3D NSs showing the PL spectral shift to higher
photon energy under increasing excitation intensity. Each spectrum can be
fitted with two (NP and TO) Gaussian spectral bands, as shown, for example, by
the dashed lines under trace (1). (b) Summary of PL spectra changes as a
function of excitation intensity.
(a)(b)Figure11 shows a modified
(note the double logarithmic scale) Arrhenius plot of the normalized integrated
PL intensity of a 3D Si/SiGe multilayer sample grown by CVD and measured at
different excitation intensities. The normalized PL intensity temperature
dependencies are fitted by a standard equation: (1)IPL(T)=1[1+C1⋅exp(−E1/kT)+C2⋅exp(−E2/kT)], (see, e.g., [41]) with two thermal
quenching activation energies E1
and E2. Here T is the
temperature, k is Boltzmann’s constant, and C1 and C2 are
scaling coefficients. In all measurements for all samples, the PL thermal
quenching activation energy E1≈15 meV and E1
is independent of excitation intensity. In contrast, the activation energy E2
depends significantly on
the excitation intensity: the PL temperature dependence shows a step-like
behavior, and E2
increases
dramatically from ~120 to 340 meV as the excitation intensity increases from
0.1 to 10 W/cm2.Figure 11
Typical
integrated PL intensity for CVD grown samples as a function of the reciprocal
temperature measured under different excitation intensities, as indicated. The
activation energies obtained from fits (shown by the different lines) to the
data (shown by the different points) are also given.We start our discussion by
noting that the observed excitation-independent PL thermal quenching activation
energy of ~15 meV is close to the exciton binding energy in SiGe alloys and
Si/SiGe superlattices [3, 14]. Thus, we conclude that one of the mechanisms of
PL thermal quenching is the thermal dissociation of excitons. The activation
energy of ~15 meV can therefore be associated with exciton localization in specific
regions of the clusters associated with variations of the SiGe composition [1, 4]. Hence, the nonuniform SiGe cluster composition and, perhaps, variations in
SiGe cluster size and shape could be responsible for the observed relatively
broad PL spectra.It has been proposed
previously that 3D Si/SiGe nanostructures can be modeled using type-II energy
band alignment, where, depending on the SiGe cluster size, valence band energy
quantization is possible [41, 67]. Using the same model, we focus next on
nonradiative carrier recombination and the different mechanisms of
electron-hole separation. Electron transport in 3D Si/SiGe nanostructures is
limited by a small (≤ 10–15 meV)
conduction band energy barrier and SiGe compositional disorder [41]. Thus, the
PL thermal quenching activation energy of ~15 meV could also be associated with
electron diffusion in Si/SiGe 3D NSs.In contrast, hole diffusion
in 3D Si/SiGe multilayer NSs with a high Ge content is controlled by large
(>100 meV) valence band energy barriers at Si/SiGe heterointerfaces [30, 40, 52, 68]. In this system, we consider two major mechanisms of hole
transport: (i) hole tunneling and (ii) hole thermionic emission. Hole tunneling
in 3D Si/SiGe NSs with thin (5–7 nm) Si
separating layers and nearly perfect SiGe cluster vertical self-alignment could
be very efficient. These nanostructures are usually grown by MBE and exhibit a
PL thermal quenching activation energy of ~60 meV [41]. The same PL thermal
quenching activation energy is found for the lowest excitation intensity in our
CVD-grown samples with 7.5 nm thick Si separating layers. We suggest that in 3D
Si/SiGe multilayer NSs with thin Si layers at low excitation intensity, the electron-hole
separation and nonradiative carrier recombination are mainly controlled by hole
tunneling between SiGe clusters. Due to significant variations in the SiGe
cluster size, shape, and chemical composition, the process of hole tunneling
could be assisted by phonon emission and/or absorption [69]. Therefore, the observed
PL thermal quenching activation energy is close to the Si TO phonon energy
[14]. In 3D Si/SiGe multilayer samples with 20 nm thick Si layers, where SiGe
cluster vertical self-alignment is practically absent [42], the probability of
hole tunneling is reduced, and hole thermionic emission over the Si/SiGe
heterointerface barrier is playing a bigger role. Thus, in these samples the PL
thermal quenching activation energy is expected to be greater, as has been
found in our experiments (see Figure 11).In this simple model, efficient
hole tunneling between adjacent SiGe nanoclusters requires not only reasonably
low and thin energy barriers but also a low carrier concentration (i.e., a
large enough number of empty adjacent SiGe clusters). By increasing the photoexcitation
intensity (i.e., the number of photogenerated carriers), hole tunneling can
effectively be suppressed since fewer empty adjacent SiGe clusters can be
found. At high excitation intensity, assuming (i) a negligible value of the
conduction band offset compared to that in the valence band and (ii) a nearly
pure Ge composition in the SiGe cluster core, the maximum anticipated PL
thermal quenching activation energy should beE2≤EgSi−EgGe≤400 meV. This value sets the upper limit in the activation energy of PL
intensity thermal quenching in 3D SiGe multilayer NSs, and it is close to the
activation energy of E2≈340 meV that has been found under the highest
excitation intensity (see Figure 11).


## 6. Photoluminescence Dynamics in Si/SiGe Three-Dimensional Nanostructures

The
PL dynamics, that is, the PL intensity decay under pulsed laser excitation, is
an important technique for studying the carrier recombination mechanism. The PL
decays discussed here were obtained using excitation by the second harmonic of
a Q-switched YAG : Nd laser with photon energyℏω=2.33 eV, pulse duration τ≈6 nanoseconds, and repetition rate
v=10 Hz. The laser energy density on the sample was of the order of
1 mJ/cm2. The PL signal in the 0.77–0.9 eV
spectral region was recorded with an InGaAs photomultiplier tube (PMT) and
stored in a digital LeCroy oscilloscope. The overall time resolution of the
entire system was better than 20 nanoseconds.Figure12 shows the normalized PL decays collected from a CVD grown sample at 4 K. The
initial PL decay is fast—close to the
resolution of our detection system (<20 nanoseconds). The longer-lived PL
shows a strong dependence on the detection photon energy, as summarized in Figure
13: the PL lifetime at photon energies below 0.8 eV is found to be ~20 microseconds, and it
drastically decreases to ~200 nanoseconds for the PL component measured at
0.89 eV.Figure 12
Typical low
temperature (T=4 K) PL dynamics for CVD grown samples measured at the
indicated photon energies using a short (~6 nanoseconds) excitation pulse.Figure 13
Summary of PL
lifetimes for CVD grown samples at different photon energies.Such
a dramatic (100 times) decrease in the PL lifetime provides key information
about the carrier recombination mechanism and this will be discussed later.


## 7. Electroluminescence in Si/SiGe Three-Dimensional Nanostructures

Electroluminescence
(EL) in 3D Si/SiGe NSs was found almost simultaneously with the first investigations
of the PL [22, 30]. In many of the Si-based nanostructures with promising PL
properties (e.g., Si nanocrystals embedded in silicon oxide, Er+ in
silicon-rich silicon oxide, etc.), EL is difficult to obtain due to poor
carrier transport [11, 15, 20]. In contrast, vertical carrier transport in
Si/SiGe multilayers and 3D Si/SiGe NSs is very efficient, and a simple device
where Si/SiGe multilayers are embedded into a p-i-n diode or a similar structure
can easily be fabricated [22, 33]. Figure 14 proves this statement by showing EL
spectra for a CVD grown sample measured under a relatively low value of forward
bias of 3–6 V applied to a Schottky-barrier
type structure. The measured EL spectrum is broad with an asymmetric spectral shape,
which, similar to the PL spectra, can be well fitted by two Gaussian bands
separated by ~45 meV. This separation energy is close to an SiGe characteristic
phonon energy, proving that the EL mechanism is nearly identical to the PL one,
that is, it is due to radiative electron-hole recombination in 3D Si/SiGe
layered NSs. On increasing the applied voltage, we observe (similar to that in
PL spectra under increasing photoexcitation intensity) a noticeable EL spectral
shift toward greater photon energies, that is, a “blue shift” (see Figure 14).
The integrated EL intensity is nearly a linear function of the applied voltage
(Figure 15). The EL intensity as a function of temperature is also similar to that
found in PL, and the EL thermal quenching activation energy is ~130 meV (Figure 16(b)).
Interestingly, in the same sample, the device current as a function of
temperature depicts nearly an exact anticorrelation with the EL intensity and
exhibits an activation energy of ~140 meV (Figure 16(a)).Figure 14
EL spectra of a CVD
grown sample atT=30 K detected using pulsed electrical excitation at the
indicated voltage amplitudes.Figure 15
Integrated EL
intensity over the 0.75–0.95 eV spectral
region of a CVD grown sample as a function of the pulsed voltage amplitude
measured at the indicated temperatures. The dashed straight line is a guide to
the eye.(a) Current and
(b) EL intensity as a function of the reciprocal temperature for a CVD grown sample.
The activation energies obtained from fits to the data (dashed lines) are also
shown.
(a)(b)


## 8. Mechanism of Carrier Recombination and
Light Emission in Si/SiGe Three-Dimensional Nanostructures

In
our discussion on carrier recombination in 3D Si/SiGe NSs, we focus on MBE and
CVD grown samples with an average Ge atomic composition close to 50%. Several
experimental results, including the PL spectral distribution extending well
below the band gap of pure Ge [33–35] and the extremely long carrier radiative
lifetime of ~10 milliseconds [41], as well as the ~30 meV per decade PL
spectral shift toward higher photon energies as the excitation intensity
increases [70], point out strong similarities between the PL in 3D Si/SiGe NSs
and the PL in III-V quantum wells with type-II energy band alignment [65, 66]. Generally,
a type-II energy band alignment at the heterointerface is a strong disadvantage
for light emitting devices due to a weak overlap between spatially separated electron
and hole wave functions. In reality, however, the critical limitation in the
efficiency of light-emitting structures is rather the presence of competing
nonradiative recombination channels for excess electrons and holes. The most
important nonradiative mechanism is carrier recombination via defects,
especially heterointerface structural defects such as propagating dislocations
and dislocation complexes. The 3D Si/SiGe NSs investigated here, grown by both
MBE and CVD processes, show an almost undetectable density of dislocations [24, 30]. Thus, in these well grown 3D Si/SiGe NSs, we can neglect nonradiative
carrier recombination via structural defects, and this explains the
experimentally observed high quantum efficiency of PL with photon energy <0.9 eV, which is associated with SiGe clusters, at low excitation intensities.It
has been suggested that SiGe Stranski-Krastanov (S-K) clusters with a small (3–5 nm) height and greater
than 10 : 1 base-to-height aspect ratio can be modeled as nanostructures with a type-II
energy band alignment and possibly with SiGe cluster valence-band energy
quantization in the direction of growth [39–41, 53]. Strained Si and Si-rich
SiGe alloy regions near the base of the clusters (also called SiGe wetting
layers) also need to be considered [42]. In our samples, the PL associated with
the c–Si separating layers exhibits unusual doublet-like structures (Figure 3), most
likely due to the built-in strain. Including the effect of strain, the observed
PL bands at 0.916 and 0.972 eV indicate a composition of the Si1-xGex transition region, which is presumably located near the bottom of the Ge/Si
pyramid-like clusters, to be close to x≈ 0.2 [30]. This conclusion is supported by
recent direct analytical TEM measurements [42–45].It
has also been proposed that the broad PL band with a peak energy of ~0.8–0.9 eV is
due to the recombination of carriers localized in the Ge-richest areas of the
clusters, which is close to the center of “pancake” shaped SiGe clusters [43, 44].
We suggest that at the lowest excitation intensity, the PL arises from
electron-hole recombination between holes localized in the Ge-richest regions
of the cluster and electrons localized in the strained SiGe alloy region near
the cluster base. This immediately explains the extension of the observed PL
spectrum below the pure Ge band gap energy.On
increasing the excitation intensity, we find that the high-photon-energy edge
of the Ge-rich cluster PL eventually overlaps with PL originating from the Si1-xGex alloy region with x≈0.2. Theoretical calculations predict that
strained Si/Si1-xGex two-dimensional NSs with x≈0.2 might have type-I energy band alignment [36].
The SiGe clusters embedded in the Si matrix can induce local strain mainly in the
SiGe alloy region near the bottom of the Ge cluster. We attribute the observed
decrease of the PL lifetime (presumably radiative lifetime) detected between
0.8 and 0.9 eV in our experiments to an increasing contribution from fast
radiative transitions involving locally strained SiGe regions of lower Ge
concetration at the bottom of SiGe clusters.Figure17 summarizes in schematic form the proposed model of radiative carrier
recombination in 3D Si/SiGe NSs. It illustrates that, at low excitation
intensity, the observed PL at photon energies ℏω<0.8 eV is due to slow recombination between electrons localized within a
strained Si-rich Si1-xGex region with x≈0.2 and holes occupying the lowest energy
states in the Ge-rich cluster core. An increase in the excitation power in PL
experiments, as well as the current density in EL experiments, usually leads to
changes in the emission spectra; for example, the PL and EL peaks blue shift, that
is, they shift toward shorter wavelength (see Figures 9, 10, and 14). At the same
time, the PL intensity as a function of excitation intensity in MBE grown 3D
Si/SiGe NSs displays a sublinear dependence on a logarithmic scale (Figure 4). It
has been pointed out that possible transformations of hole energy spectra due
to quantization and/or strain might dramatically increase the rate of
nonradiative Auger recombination, by more than 100 times compared to that in
bulk Si and Ge [71]. However, in a quantum well with smoothed (e.g., diffused)
heterointerfaces, Auger processes are expected to be less efficient [35, 41].
We find strong evidence that in CVD grown 3D Si/SiGe NSs the reduction of the Auger
rate, most likely due to diffused Si/SiGe interfaces, does take place. Thus, on
increasing the pumping power in PL experiments on CVD grown samples, we observe
a relatively small deviation from a linear function in the PL intensity
dependence on excitation (see Figure 9). In addition, we find that the previously
mentioned PL spectral blue shift, with the PL peak shifting from a photon
energy of 0.78 to 0.88 eV, correlates with a strong (~100 times) decrease in the
carrier radiative lifetime (Figures 9, 10, 12, and 13). The hole energy barrier between the strained
Si0.8Ge0.2 transition region and Si is ≥ 200 meV, in good agreement
with a larger value of the observed PL thermal quenching activation energy,
while the smaller value of 15 meV is closer to the exciton binding energy in
SiGe alloys, as has been previously discussed [35, 41, 72].(a) Schematic
representation of the atomic structure of an SiGe nanocluster on a wetting layer
with a core Ge concentration close to 50% embedded into an Si matrix and (b)
the corresponding energy band diagram with two possible radiative transitions
indicated.
(a)(b)Figure17 also proposes an energy band diagram, which takes into account the previously
described complex compositional structure of 3D Si/SiGe heterointerfaces. It
shows a modified type-II energy band alignment at the SiGe cluster core with a
compositional transition toward the Si/SiGe heterointerface and a nearly type-I
alignment at the cluster base, mainly due to the strain in the SiGe wetting
layer. A continuous change in the Ge atomic concentration is reflected by a
gradually increasing energy band gap from the cluster center toward the SiGe
wetting layer, where the Ge atomic concentration is estimated to be ~20%. In
this band diagram, two types of radiative transitions are shown (Figure 17(b)).
First, there is a slow recombination between electrons localized mainly in the
SiGe wetting layer and holes localized within SiGe clusters, mostly near the SiGe
cluster core. Second, a faster and more efficient radiative recombination is
achieved between electrons and holes leaking from the SiGe cluster core toward
the cluster base and SiGe wetting layer. This latter recombination mechanism,
which becomes dominant under a high photoexcitation intensity when the “slow”
recombination channel for spatially separated electrons and holes is saturated,
has been called a “dynamic type-I” energy band alignment [41, 70]. It explains
very well both the PL spectral blue shift under increasing excitation intensity
and the dramatic (~100 times) decrease in carrier radiative lifetime measured
at photon energies from 0.77 to 0.89 eV. This explanation is also consistent
with the previously discussed PL intensity temperature dependence, which shows a
different PL thermal quenching activation energy at different excitation
intensities (see also [72]).


## 9. Conclusion

In
conclusion, we present comprehensive experimental studies on Raman scattering
and light-emitting (PL and EL) properties of 3D Si/SiGe NSs. We show that these
nanostructures emit light at the technologically important 1.3–1.6μm wavelength region. The
highest PL and EL quantum efficiency is found in SiGe clusters with an ~50% Ge
composition near the cluster core. The highest luminescence quantum efficiency
is observed at low excitation intensity. However, the PL quantum efficiency decreases
as the excitation intensity increases, most likely due to competition with a
faster nonradiative Auger recombination. Using time resolved PL measurements,
we found that the suspected type-II energy band alignment at the Si/SiGe
cluster heterointerface is responsible for the observed long carrier radiative
lifetime (10-4–10-3 second). We also found that within
the broad PL band, the part of the PL spectra associated with higher photon
energies exhibits an ~100 times faster radiative transition, which is explained
by a proposed “dynamic type-I” energy band alignment, for example, radiative
recombination between holes occupying excited states in SiGe clusters and
electrons mostly localized close to the Si1-xGex wetting
layer with x≈0.2.
These experimental observations suggest that a commercially useful SiGe light-emitting
device can be fabricated and integrated into the traditional CMOS environment.



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*Source: 218032-2008-06-12.xml*