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Growth and Characterization of ZnO Nano- and Microstuctures 319

Figure 9.28 shows the time dependence of the current of a ZnO nanowire with two ohmic

contacts under constant bias (red curve) measured at the wafer–prober in comparison with

that of a PLD thin ﬁ lm (black solid curve). The samples were exposed several days to ambient

light prior to the experiments and at t  10 s the wafer–prober was completely darkened. Unlike

the thin ﬁ lm, the nanowire exhibits almost time-independent conductivity. For the thin ﬁ lm the

inverse effect (switching on the light) can also be observed (dotted black line). ZnO is known to

exhibit persistent n-type photoconductivity [118] and a slow photoresponse [122] . Typically, the

photosensitivity factor [120]

S 

σσ

photo dark

dark

(9.13)

is in the order of 10–100 and in some cases even values above 10

have been reported [120] with

exponential decay constants in the 10 s range. This effect is commonly attributed to oxygen vacan-

cies [118] but the exact mechanism is still under debate since computational studies [17, 123] have

cast doubt on the categorization of oxygen vacancies as shallow donors. Already in the 1950s, oxy-

gen adsorption at the surface was reported to play an important role for time-dependent photocon-

ductivity of ZnO [124] . In this model, adsorbed oxygen molecules can bind one electron each, the

surface is therefore charged negatively, a depletion layer forms and the conductivity is reduced. Upon

illumination the depletion layer vanishes. Also the free carrier concentration as well as the degree

of compensation have a strong inﬂ uence [125] . Due to the larger surface-to-volume ratio and

good crystallinity, surface effects dominate the photoconduction transient in nanowires [122] . For

instance, Li et al . have shown that the transient photocurrent decay time depends strongly on the

ambient atmosphere. This is also the reason why ZnO nanowires exhibit good gas sensing properties.

9.6 Conclusion

In summary, ZnO nanowires have been grown by high-pressure PLD on sapphire substrates cov-

ered with gold colloidal particles as nucleation sites. A detailed study of the growth evolution

and of the nanowire size and length distribution was presented. It was found that the aspect ratio

varies linearly with the deposition time.

We investigated the resonant optical behaviour of ZnO micro- and nanocrystals, giving detailed

insight into optical whispering gallery modes within the micro and nanometer regime. In partic-

ular, we analysed hexagonal WGMs in nanocavities for mode numbers in the range from N  1

to N  6. The energy shift and broadening of the modes for cavity diameter decreasing to zero is

well described without free parameter by Eq. 9.1 from a plane wave model. Furthermore, we have

found RT stimulated emission in a ZnO microcrystal grown by carbothermal evaporation.

Moreover, rectifying as well as ohmic contacts attached to single microwires were character-

ized. We ﬁ nd that the conductivity is strongly temperature dependent in the range from 200 to

300 K. From this, carrier activation energies were determined by taking the sample geometry

into account. The values are between 18 and 40 meV while the estimated carrier concentration

is around 10

 3

at room temperature.

Acknowledgements

This work was supported by Deutsche Forschungsgemeinschaft within FOR 522 (Project Gr

1011/12-1), by the European Network of Excellence SANDIE (NoE 500101) and by the EU

STREP NANDOS (FP6-016924).

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Miniband-related 1.4–1.8 μ m Luminescence of Ge/Si

Quantum Dot Superlattices

V. G. Talalaev ,

1,2,3

G. E. Cirlin

1,4

A. A. Tonkikh ,

N. D. Zakharov ,

P. Werner

U. Gösele,

J.W. Tomm,

and T. Elsaesser

Max-Planck-Institut für Mikrostrukturphysik, Weinberg 2, 06120 Halle/Saale, Germany;

Max-Born-Institut für Nichtlineare Optik und Kurzzeitspektroskopie,

Max-Born-Strasse 2A, 12489 Berlin, Germany;

V.A. Fock Institute of Physics, St Petersburg State University, Ulyanovskaya 1,

198504 Petrodvorets, St Petersburg, Russia;

Ioffe Physico-Technical Institute RAS, 194021 Polytekhnicheskaya 26, St Petersburg, Russia

10.1 Introduction

The development of efﬁ cient silicon-based light-emitting devices is a challenging task in modern

semiconductor physics and optoelectronics. Due to the indirect nature of the band gap, bulk

silicon has a poor luminescence. One of the promising approaches to increase the luminescence

efﬁ ciency from Si-based materials is to apply a concept of nanostructures based on Ge inclusions

into an Si matrix. The development of epitaxial methods enabled the growth of ultra-thin Ge/Si

layer superlattices. Structures that are based on the Brillouin folding concept can result in quasi-

direct optical transitions near 1.55 μ m wavelength [1–4] . However, due to the small localiza-

tion potential the photoluminescence (PL) from Ge/Si layer superlattices is observed only at low

temperatures [3–6] . The concept of a quantum-cascade Ge/Si-based laser has also been demon-

strated. However, it emits only in the mid-infrared region and shows electroluminescence at low

temperatures [7] .

Several attempts have been made to obtain stronger localization and to achieve a room temper-

ature luminescence at 1.55 μ m (0.8 eV) in Ge quantum dots (QDs) embedded into an Si matrix.

Although arrays of self-assembled Ge islands grown by the Stranski–Krastanov mode on silicon

have been studied quite intensively, just a few research teams have reported about PL [8–9] and

electroluminescence (EL) data [11–14] . Probably, due to the low emission intensity, the internal

quantum efﬁ ciency (QE) measured at room temperature for a Ge/Si light-emitting diode (LED)

at 1.42 μ m has been reported only in two studies with the following results: 5  1 0

 4

% [12]

and 0.015% [14] . In one case [14] an external QE of 3.4  1 0

 4

% was measured for EL close to

1.5 μ m. It should be mentioned that all publications mostly concentrate on QDs, i.e. on the hole

subsystem. It is well known [15] that Ge/Si heterostructures have a type-II band alignment, in

which both carriers have opposite locations in relation to the heterointerface: electrons in Si and

holes in Ge. Due to the small overlap of electron and hole wavefunctions (not more than 15%)

the oscillator strength of indirect excitons in a Ge/Si heterostructure is quite low [16] . This is one

of the key factors for the weak near-infrared luminescence of Ge/Si heterostructures.

CHAPTER 10

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Miniband-related 1.4–1.8 μ m Luminescence of Ge/Si Quantum Dot Superlattices 325

In this work a novel approach is presented, which overcomes the disadvantages mentioned

above and allows efﬁ cient luminescence at room temperature from Ge/Si heterostructures to be

achieved. Special regimes of molecular beam epitaxy (MBE) with Sb-doping enabled the growth

of highly strained dislocation-free Ge/Si quantum dot multilayer structures. Our samples consist

of stacked QD arrays, in which the columns of Ge QDs are well correlated vertically. We charac-

terize such structures as Ge/Si QD superlattices (QDSLs), because there exists a conﬁ nement of

holes in the Ge QDs, as well as a conﬁ nement for electrons in the Si spacer layers. The electron

states in the Si layers, which behave as real quantum wells (QWs), are characterized by activation

energies in the range of 45 to 85 meV, i.e. are stable up to room temperature. The optimization

of QD columns have provided conditions, which are favourable for vertical electron tunnelling

and for the formation of a conduction miniband. The different investigations presented in the fol-

lowing give strong indications that a conversion to quasi-direct excitons occurs. This conversion

enables an external QE of 0.04% for 1.55 μ m EL maximum at room temperature.

10.2 Experimental techniques

The Ge/Si structures are grown on Si(001) 5-inch substrates using the MBE set-up Riber SIVA

45. Samples consist of an undoped Si buffer layer with 100 nm thickness, a Ge/Si multilayer

structure and a 50 nm Si cap layer. The growth temperature is 600°C. The growth rates of Si

and Ge are 1.0 and 0.2 Ås

 1

, respectively. As the nominal thickness of the single Ge layers values

between 0.7 nm and 0.9 nm are chosen. The MBE growth of the heterostructures is monitored

in situ by a reﬂ ection high-energy electron diffraction system (RHEED).

Due to the difference in the lattice parameters between the bulk of Ge and of Si (4%) the

initial stage of the Ge deposition is accompanied by the formation of nano-size islands known in

the literature as Stranski–Krastanov QDs. The formation of Ge nano-islands in all layers of each

sample is evident from RHEED by observing the spot patterns. These measurements document

the transition from the 2D growth mode of layers to the 3D growth of islands (QDs). The effec-

tive thickness of an Si spacer is varied in the range 5.5–9.5 nm. A doping of the Si spacer layer

by antimony is performed in the following way: a growth interruption after Ge QD formation is

used. During the growth interruption the Ge QD array is exposed to the Sb ﬂ ux. A Sb deposition

rate was 2  1 0

 4

n m s

 1

. The time of the exposition is tuned between 0 and 40 s. After exposi-

tion the ﬁ rst part of the Si spacer layer (2 nm) is also doped by Sb, whereas the remaining part of

the spacer layer is kept undoped. Sb concentration proﬁ ling in QDSLs is carried out by secondary

ion mass spectroscopy (SIMS) using a Cameca set-up. In order to optimize the structure for get-

ting the most effective luminescence at 1.55 μ m the Ge and Si thickness as well as the time of Sb

exposition have to be optimized [17] .

Two types of boron-doped Si substrates are used: p -type ( ρ  5 Ω cm) for PL investigations

and p



-type ( ρ  0.015 Ω cm) for EL measurements and LED fabrication. The capping layer in

PL structures is undoped, in LED structures it is doped with Sb. SIMS data for optimized QDSLs

show that due to diffusion of Sb atoms to the growth surface, both structures have an increased

Sb concentration in the cap layer: 2  1 0

c m

 3

for PL and 2  1 0

c m

 3

for LED structures

( Fig. 10.1c ). Thus we deal with QDSLs embedded into a p – n -junction (PL structures) and into a



– n



-junction (LED structures). Ohmic contact to the cap layer is formed using Al/Au deposited

by magnetron sputtering. Indium is used to form the backside metal contact. The LED is glued

onto a copper heat-sink and the top contacts are bonded with gold wires using standard thermo-

compression bonding.

The structure of the grown samples is investigated by transmission electron microscopy (TEM)

and selected area electron diffractometry using JEM 4010 and CM 20 microscopes operating

at acceleration voltages 400 kV and 200 kV, respectively. The Ge concentration proﬁ le is deter-

mined by applying an image analysing technique of bright-ﬁ eld cross-section TEM micrographs

recorded on a slow scan charge-coupled-device (CCD) camera to keep the linearity between elec-

tron intensity and contrast.

Raman spectra are recorded in backscattering geometry at ambient temperature. The Raman

scattered light is collected by a microscope and subsequently analysed by a Dilor X-Y triple

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326 Handbook of Self Assembled Semiconductor Nanostructures for Novel Devices in Photonics and Electronics

spectrometer equipped with a liquid nitrogen-cooled Si CCD camera. The 488 nm line of an Ar



laser is used for excitation with a typical power below 1 mW. The spectra are taken for incident

and scattering light polarized both parallel or perpendicular to each other. The excitation laser

spot has a diameter of 1 μ m. The spectral resolution is 0.5 cm

 1

Steady-state (cw) PL measurements are carried out in a standard lock-in conﬁ guration. The PL

spectra are excited by the 488 nm line of an Ar



laser. The laser spot on the sample has a diameter

of 1.5 mm. For studying the PL intensity versus excitation power density the laser beam is focused

down to a 100 μ m diameter and adjusted by neutral ﬁ lters. For low-temperature PL measurements

the samples are cooled in a continuous-ﬂ ow He cryostat. The PL signal is collected by mirror

optics and dispersed by a single 0.5 m grating monochromator coupled with a liquid nitrogen-

cooled Ge detector (Edinburgh Instruments) having the photoelectric threshold at 1.7 μ m

(0.73 eV). In order to register luminescence at longer wavelengths the monochromator exit is

equipped with a liquid nitrogen-cooled InGaAs photodiode G7754-01 (Hamamatsu) having a

cut-off wavelength of 2.4 μ m (0.5 eV). PL spectra are always normalized to the spectral photo-

detector sensitivity. EL spectra are obtained for in-plane and edge-emitting geometries in the con-

stant current mode using the same set-up. The diode chip is mounted on a water-cooled pedestal

in order to prevent EL degradation by heating of the active zone during electrical excitation.

EL spectra, current–voltage characteristics and photo-voltage saturation measurements are

performed using a Keithley 2400 source measure unit. All temperature dependent PL measure-

ments are carried out in the short-circuit regime of the p – n -junction.

Time-resolved PL (TRPL) measurements are performed in an He-closed-cycle cryostat at 15 K.

The excitation wavelength is 395 nm, the pulse width is 1 ps, and the repetition rate is 1 kHz.

These excitation pulses are generated by a system consisting of a Spectra Physics Tsunami seed

laser (mode locked Ti–sapphire, 82 MHz, 100 fs) followed by a Spectra Physics Spitﬁ re regenera-

tive ampliﬁ er (Ti–sapphire). An optical BBO crystal is used for the second harmonic generation.

The laser beam is focused onto the sample down to a 100 μ m diameter. The resulting PL is dis-

persed by a 0.25 m grating monochromator and detected by a Judson Technologies J16D-M204

Ge detector having a cut-off wavelength of 1.6 μ m. The signal is analysed by a digital oscilloscope

Agilent Inﬁ nium 54833A. The temporal resolution of the total system is limited to ~10 ns by the

response time of the detector–preampliﬁ er (Femto HCA-100) combination. Neutral density ﬁ lters

are used for adjusting the excitation power density.

In order to determine the external efﬁ ciency of EL from Ge/Si QDSLs at 1.55 μ m, we developed

an absolute calibration method based on the steady-state PL set-up. An integrating 115 mm

diameter sphere is used. The diode is mounted on the entrance port. The exit port has a detach-

able opal–glass diffuser and is always positioned in the monochromator focal plane. Before these

measurements a calibration of photodetector spectral sensitivity is performed using the black-

body simulator. Subsequently, a power normalization of the photodetector is carried out using

the 1.55 μ m line from a calibrated laser diode (0.9 mW).

10.3 Experimental data and interpretation

10.3.1 Structural properties of Ge/Si QDSL

Cross-section TEM images of undoped and Sb-doped multilayer structures are compared in Fig.

10.1a and b . The undoped structure ( Fig. 10.1a ) shows smeared heterointerfaces and a relatively

ﬂ at upper surface. The lateral size of the Ge islands increases with the layer number from bottom

to the top. The Sb-doped structure ( Fig. 10.1b ) is characterized by sharp heterointerfaces and a

highly corrugated upper surface. The lateral size of the Ge clusters is nearly the same for the top

and bottom layers. No misﬁ t dislocations are observed in the samples.

A plan-view TEM image of the Sb-doped 20-layer Ge/Si structure is shown in Fig. 10.2 . The Ge

clusters look like squares with the edges oriented along  100  crystallographic directions and are

assembled into a kind of square lattice. The average base size of the squares is 60  6 0 n m

. The

array surface density is 1.1  1 0

c m

 2

. The array is characterized by high uniformity of the Ge

cluster sizes. It should be mentioned that every surface cluster is related to a single Ge/Si column.

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Miniband-related 1.4–1.8 μ m Luminescence of Ge/Si Quantum Dot Superlattices 327

A cross-section TEM image of a Ge/Si multilayer column is shown in detail in Fig. 10.3 (left).

Lens-shaped vertically correlated Ge clusters separated by Si spacers are seen. The column is

characterized by a high uniformity of the lateral sizes (55 nm  5%). It was established that the

height of the clusters in the column has a  10% deviation from the average value ( B  4.5 nm).

We will follow the traditional name of these clusters as quantum dots taking into account that the

quantum conﬁ nement of the carriers is valid in the growth direction while in the in-plane direc-

tion only the carrier localization takes place. The Si spacer thickness in the column ( W ) has a still

smaller deviation (  5%).

(c)(b)(a)

50 nm 50 nm

Si cap

Ge/Si QDSL

Si buffer

Sb concentration (cm

3

)

100

z (nm)

Figure 10.1 TEM cross-section images of Ge/Si heterostructures with 20 Ge layers: (a) undoped; (b) Sb doped.

The dark regions correspond to Ge layers. (c) SIMS data of the Sb concentration proﬁ le in QDSL similar to (b)

embedded into a p



– n



-junction. z -growth direction.

Figure 10.2 [001] TEM plan-view image of the Ge/Si QDSL. The insert shows the electron diffraction pattern

taken along [001].

z (nm)

0 0.5

Ge content

Si buffer

Figure 10.3 Cross-section TEM image of the Ge/Si QDSL. The average height of QD in column B  4.5 nm (left).

Right: Ge content proﬁ le along the growth direction z for the same sample.

The Ge content distribution across the QDSL in the growth direction is demonstrated in Fig.

10.3 (right). Ge/Si interfaces are rather sharp. Very little intermixing has taken place and the Ge

content in the QDs is x  0.8 (  10%), and in the Si spacer layers it is x  0.1 (  10%). It should

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328 Handbook of Self Assembled Semiconductor Nanostructures for Novel Devices in Photonics and Electronics

be pointed out that all TEM images show a thin Ge wetting layer (WL), which is typical for the

QD arrays grown by the Stranski–Krastanov mode. TEM data show that the Sb-doped structures

under investigation are mostly defect free.

Figure 10.4 shows Raman scattering spectra for the Sb-doped Ge/Si QDSL with B  3 . 6 n m

measured in two different conﬁ gurations. The quasi-periodical structure in the 20–200 cm

 1

spectral range, see Fig. 10.4 (bottom), is attributed to folded longitudinal acoustic (FLA) pho-

nons, which have already been reported for Ge/Si QDSLs in [18] . According to the selection rules

[19] , our geometry, namely z ( x  y  ) ž , is not Raman active for FLA phonons in ﬂ at Ge/Si layers.

Their presence in layers with QDs, however, is explained by symmetry lowering [20] . The high

structural quality of the studied QDSLs, in particular the presence of sharp Ge/Si interfaces, is

conﬁ rmed by the observation of 20 FLA modes. Their intensity distribution appears to be non-

monotonous. The observed beats at 7–9 meV and 14 meV are tentatively explained by electron–

phonon interaction in the QDSL.

z(xy)z

–

FLA

z(x’y’)z

–

100 150 200

100 200 300 400 500

Si–Si

Ge–Si

(Si–Si)

Ioc

Ge–Ge

Raman shift for z(xy)z

–

(cm

1

)

Raman shift for z(x’y’)z (cm

1

)

Intensity for z(x’y’)z (a.u.)

Intensity for z(xy)z

–

(a.u.)

Figure 10.4 Raman spectra of a Ge/Si QDSL measured in two backscattering conﬁ gurations. The terms z , x , y , x 

and y  refer to the [001], [100], [010], [11

–

0] and [110] directions, respectively.

Raman spectra in the 200–500 cm

 1

spectral range, see Fig. 10.4 (top), are recorded in z ( xy ) ž

geometry. According to the selection rules for this geometry all features are assigned to longitu-

dinal optical (LO) phonon modes: Ge–Ge, Ge–Si, (Si–Si)

loc

, and Si–Si. The frequency of the Ge–Ge

mode (297.5 cm

 1

) is below the typical Ge bulk value (301 cm

 1

) [21] . This shift of the Ge–Ge

LO mode in the QDSL is likely to be caused by phonon conﬁ nement [22] . It should be noted that

in those structures with the thicker Ge QDs, e.g. B  5.8 nm, the Ge–Ge mode frequency agrees

well with the bulk value.

The Ge-Si interface mode at 415 cm

 1

has an amplitude comparable to that of the Ge–Ge LO–

mode. The additional features (Si–Si)

loc

at 435, 450 and 465 cm

 1

are attributed to local vibra-

tions (Si–Si)

loc

under the inﬂ uence of Ge atoms in their vicinity [23] . These features reﬂ ect the

presence of a near-range order at the interfaces between the Ge QDs and the adjoining Si layers.

Thus, the Raman spectra indicate an ordered surface of the Ge–Si interfaces in Sb-doped QDSLs.

This does not contradict the observed sharpness, since this order is still on atomic-scale dimen-

sion. Raman spectra of undoped Ge–Si QDSL look quite different: the FLA and Si–Si local modes

are absent, and the Ge–Si interface mode is always weaker than the Ge–Ge mode.

10.3.2 Luminescence properties and initial electronic structure

The inﬂ uence of Sb doping on the PL spectrum of Ge/Si QDSLs at room temperature is demon-

strated in Fig. 10.5 . Undoped structures had always a poor PL – in the low-energy part of the

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