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MOVPE growth and characterization
of GaN/InGaN nanowires and
microrods for next generation solid-
state-lighting applications
Von der Fakultät für Elektrotechnik und Informationstechnik
der Rheinisch-Westfälischen Technischen Hochschule Aachen
zur Erlangung des akademischen Grades
eines Doktors der Ingenieurswissenschaften genehmigte Dissertation
vorgelegt von
Diplom-Ingenieur
Bartosz Foltyński
aus Breslau, POLEN
Berichter: apl. Prof. Dr.-Ing. Michael Heuken
Univ. –Prof. Dr. rer. nat. Wilfried Mokwa
Tag der mündlichen Prüfung: 28.06.2016
Diese Dissertation ist auf den Internetseiten der Hochschulbibliothek online verfügbar.
I
Contents
Chapter 1 Introduction ............................................................................................................ 1
Chapter 2 GaN nanowire advantages compared to bulk GaN ............................................ 6
2.1 Crystal structure and basic properties of III-nitride semiconductors.................. 6
2.2 Nanostructures advantages compared to the bulk materials ................................ 9
2.3 GaN NW applications for next generation devices ............................................... 10
Chapter 3 GaN NW growth techniques state-of-the-art by MOVPE ............................... 14
3.1 Advantages and challenges of GaN NWs growth by MOCVD ........................... 14
3.1.1 Challenges and solutions for GaN-on-Si integration .................................... 14
3.1.2 NW morphology as a polarity consequence ................................................... 16
3.2 Au- catalyst induced VLS growth mode of GaN NWs ......................................... 18
3.3 Selective area growth ( SAG) of GaN-based nanocolumn-arrays ....................... 20
3.4 Self-assembled growth of GaN NWs ...................................................................... 22
Chapter 4 Experimental setup and characterization methods .......................................... 26
4.1 Metalorganic vapor phase epitaxy (MOVPE) ....................................................... 26
4.2 European Synchrotron Radiation Facility ............................................................ 28
4.2.1 Nanofluorescence XRF .................................................................................... 29
4.2.2. X-ray absorption near edge structure XANES.............................................. 30
4.3 Complementary characterization methods ........................................................... 31
Chapter 5 Antisurfactant role of SiH4 during vertical growth of GaN NWs ................... 32
Chapter 6 VLS Au-initiated growth of GaN NW on Sapphire substrates ........................ 36
6.1 Experimental procedure for NWs growth and characterization techniques ..... 37
6.2 Atomic composition of coaxial InGaN/GaN quantum wells in NWs .................. 40
6.3 Conclusions of VLS Au-initiated growth of GaN NW on Sapphire substrates . 44
Chapter 7 Selective Area Growth of GaN microrods on Si(111) substrates..................... 45
7.1 Experimental procedure ......................................................................................... 46
7.1.1 Template preparation for GaN-based microcolumn arrays growth ........... 46
7.1.2 The controlled SAG of GaN microrods on Si(111) ........................................ 48
7.2 Optical properties of GaN microcolums determined by photoluminescence .... 52
7.3 Structural properties of GaN microcolumns determined by Raman
Spectroscopy ....................................................................................................................... 53
7.4 Conclusions of Selective Area Growth ................................................................... 56
Chapter 8 Growth of self-assembled GaN NW on Si(111) substrates ............................... 57
8.1 AlN buffer on Si(111) as a basis for GaN NW growth ......................................... 58
8.1.1 Impact of asynchronous introduction of precursors on the AlN polarity ... 58
II
8.2 Investigation of growth parameters on the GaN NW growth and morphology 59
8.2.1 Substrate preparation ...................................................................................... 59
8.2.2 Impact of NH3/TMAl predose before AlN deposition on Si(111) on the NWs
growth 59
8.2.3 Proposed optimization model – nanostructures density as a function of the
key process parameters .................................................................................................. 60
8.2.3.1 Impact of SiNx in-situ masking layer deposition time on the NW density
62
8.2.3.2 Impact of growth temperature on the NW density ................................... 64
8.2.3.3 Impact of silane injection time on the NW density .................................... 66
8.3 Optimized growth conditions for GaN NW growth on Si(111) ........................... 68
8.4 Effect of AlN susceptor coating on NW growth homogeneity ............................. 69
8.5 Conclusions of self-assembled growth of GaN NW on Si(111) substrates.......... 69
Chapter 9 Optical and structural properties of self-organized GaN NW on Si(111)
substrates ................................................................................................................................. 70
9.1 Structural properties and In incorporation in InGaN/GaN nanowires by
µPhotoluminescence ........................................................................................................... 70
9.2 InGaN distribution in GaN/InGaN core/shell heterostructures by nano-scale
Cathodoluminescence mapping ........................................................................................ 73
9.2.1 Nano-scale cathodoluminescence mapping of Sample A .............................. 74
9.2.2 Nano-scale cathodoluminescence mapping of Sample B .............................. 76
9.2.2.1 Non-hexagonal microrod ............................................................................. 76
9.2.2.2 Typical hexagonal microrod ........................................................................ 77
9.2.3 Nano-scale cathodoluminescence mapping of Sample C .............................. 79
9.3 Advanced structural characterization of GaN microrods grown under different
conditions by TEM ............................................................................................................. 82
Chapter 10 Summary and conclusions ................................................................................. 88
References ............................................................................................................................... 93
List of Figures ....................................................................................................................... 101
List of Tables ......................................................................................................................... 106
List of Abbreviations ............................................................................................................ 107
Scientific appendix ............................................................................................................... 108
Acknowledgements ............................................................................................................... 111
1
Chapter 1
Introduction
The long list of semiconductor devices such as transistors, memories, amplifiers,
switches, sensors and more are building blocks for commonly used applications. Nowadays,
microchips, mobiles phones, tablets, solar cells and other high-tech solutions are commercially
available and widely used. They all owe its origin to modern electronics. There are two main
driving forces, which accelerate innovation in the field of semiconductor devices. The first
factor is large volume and high perfection of synthesized semiconductor material. The mature,
well developed Si technology requests an integration of other material systems with silicon
base. Yet, it is still challenging for compound materials like GaN or GaAs, commonly used in
optoelectronics or high-power and high-speed applications. The second factor for innovations
in the semiconductor area is a demand of integration density increase, which also defines
the costs. Currently, lateral structure size already approaches the physical limits. Hence, new
innovative solutions are expected.
The nanowire (NW) structures offer new possibilities to meet these expectations and
demands as well as to overcome the limits of conventional planar devices. Nanowires,
benefiting from their unique morphology, promise a high crystal quality of a material grown on
foreign substrates. The defect-free structure and possibility to release strain are only some of
the main advantages of such structures in comparison to bulk materials.
The nanowires might be utilized as a basis for a high efficiency light source.
The emission wavelength of the light emitting diode (LED) can be tuned by controlling
the alloy composition of GaN with InN or AlN [1]. Therefore, by varying the contents of
the compound elements, a continuous emission spectrum that covers UV and visible ranges
might be achieved (Fig. 1.1). The multi quantum well (MQW) can be deposited between p-n
junction of the single nanowire and therefore the performance of nanoLED can be improved.
2
Figure 1.1: Bandgap of binary InN, GaN, and AlN and their ternary alloys as a function of in-plane lattice
constants (no bowing assumed).
It is particularly motivating to conduct scientific research when a technological benefit
from the scientific results is conceivable. The GaN nanowire-based devices are in focus of
interest to study due to the key technological innovation factor in the field of semiconductors.
The guideline of this thesis is the MOCVD growth of InGaN/GaN nanowire heterostructures
on Si(111) substrates as a building block for a LED application. This study has two prime
interests. First one is a successful growth of vertical GaN nanocolumns on Si(111) substrates
by MOCVD. The second one is an analysis of optical and structural properties of grown
heterostructures and its correlation with the process parameters aiming the optimization of
the material properties. The comprehensive growth investigations are conducted based on three
growth approaches: vapor-liquid-solid (VLS) mode utilizing Au as a catalyst, selective area
growth (SAG) as well as self-organized growth.
3
The thesis manuscript is subdivided into ten chapters (including introduction as a chapter 1):
Chapter 2 will first give a background of nitride material systems: crystal structure and
fundamental properties of III-nitride semiconductors. Afterwards, the main
nanostructures advantages compared to bulk material will be discussed. The rest of the
chapter will resume the state of the art of GaN nanowire-based devices for next
generation applications.
Chapter 3 will deal with the state-of-the-art in GaN nanowire growth by MOCVD.
First, the advantages as well as challenges of GaN MOCVD growth will be discussed.
The challenge for GaN-on-Si integration as well as nanowire morphology as polarity
consequence will be underlined. Then, three main bottom-up growth techniques will be
presented, namely: Au- catalyst initiated vapor-liquid-solid (VLS) growth, selective
area growth (SAG) and self-organized growth. The properties of GaN nanostructures
synthesized by mentioned techniques will be described.
Chapter 4 will address the experimental setup and characterization techniques utilized
during this work. At the beginning the background of MOCVD growth will be
presented. In the second part, the background and principles of several non-standard
synchrotron based measurement techniques will be introduced. The end of the chapter
will summarize the rest of the complementary characterization methods utilized during
the research.
Chapter 5 will give an explanation and background for antisurfactant role of silane
during vertical growth of GaN nanowires by MOCVD.
Chapter 6 will present experimental results and explanations of the Au-initiated vapor-
liquid-solid (VLS) growth of GaN nanowires on sapphire substrates by MOCVD. In the
first section, the experimental procedure with growth condition selection for GaN
nanorods synthesis will be described. Afterwards, the characterization results performed
at European Synchrotron Radiation Facility (ESRF) in frame of the Marie-Curie
Nanowiring project will be presented. The characterization section will deal with atomic
composition of coaxial InGaN/GaN quantum wells in nanowires.
4
Chapter 7 will describe the selective area growth (SAG) of GaN nanowires on silicon
substrates by MOCVD. The template preparation for GaN-based microcolumns arrays
as well as growth conditions will be discussed. The next sections of the chapter will
provide information about optical and structural properties of GaN microrods
characterized by photoluminescence (PL) and Raman spectroscopy. Understanding of
growth and resulting properties will shed a new light on the GaN rod on SI as a building
block for nanoLED.
Chapter 8 will deal with self-organized growth of GaN nanowires on silicon substrates
by MOCVD and present understanding of the developed growth mechanism. First,
the AlN buffer on Si(111) as a basis for GaN rods will be introduced. The reactor
conditioning to ensure reproducible starting point will be proposed. The impact of
asynchronous introduction of precursors on the AlN buffer polarity will be discussed.
The second section of the chapter addresses the detailed investigations on the GaN NW
growth mechanism. The impact of growth parameters on the GaN nanowire growth and
morphology will be studied and discussed. The substrate preparation as well as impact
of the NH3/TMAl predose prior AlN buffer on the nanowire morphology will be
presented. Afterwards, the optimization model based on the nanostructure density as
a function of the key process parameters will be proposed. Precisely, impacts of SiNx
in-situ masking layer deposition time, growth temperature and silane injection time will
be investigated. At the end of the chapter the optimized growth parameters as well as
effect of AlN coating on the GaN nanowire growth homogeneity will be presented.
The results presented in this chapter describe widely and precisely for the first time
the innovative procedure of self-assembled growth of GaN nano and microrods on
Si(111) substrates by MOCVD.
Chapter 9 will address the optical and structural properties of self-organized GaN
nanowires on Si(111) grown by MOCVD. The structural properties and In incorporation
in InGaN/GaN nanorods will be characterized by microphotoluminescence (μPL). In
the second part of the chapter, the InGaN distribution in GaN/InGaN core/shell
heterostructures will be characterized by nanoscale cathodoluminescnce mapping.
The last section of this chapter will provide information about structural properties of
GaN and InGaN microrods grown under different conditions and characterized by TEM.
5
All of these complex characterization techniques lead to understand and describe
the current status of the GaN rod on Si as the building block for the nanoLED.
Chapter 10 will finally summarize the most important results of this thesis and will
propose a few perspectives of this work. The main achievements during the PhD
research (in-line with the manuscript order) are:
o Characterization and understanding of structural and optical properties of
individual GaN NW on sapphire grown by Au-initiated VLS method
o Development of the SAG of GaN microrods on Si(111) substrates by MOCVD
o Characterization and understanding of structural and optical properties of
individual GaN rods on Si(111) grown by SAG method
o Development of novel self-organized growth of GaN NW on Si(111) substrates
o Understanding of self-organized growth mechanism of GaN NW on Si(111)
o Characterization in the nano-scale and understanding of structural and optical
properties of InGaN/GaN NW on Si(111) grown by self-assembled growth
technique
o Understanding and description of the current status of GaN NW on Si growth
and the role of GaN rod on Si as the building block for the nanoLED
o Identification of the current challenges and setting the perspective of the future
research
6
Chapter 2
GaN nanowire advantages compared
to bulk GaN
This chapter provides general theoretical background on the crystal structure and basic
properties of III-nitride semiconductors with the main focus on the GaN. In the second section,
more specifically background information on the GaN nanowire advantages compared to bulk
material is discussed. Due to benefits gained by nanostructures, new applications might be
developed addressing the constant need of improvement in terms of device efficiency,
performance, size and price. The last section of this chapter provides briefly the background
and principles of selected GaN NW-based applications for next generation devices, namely:
nLED (nano Light emitting diode), transistors, solar cells and others.
2.1 Crystal structure and basic properties of III-nitride semiconductors
Gallium Nitride – GaN – as one of the most important semiconductor in the group III-
Nitride material system possesses superior material properties (such as wide and direct energy
bandgap, better thermal and chemical stability, and high electron drift velocity) as compared to
silicon, GaAs and other III-V compound materials. It was strongly investigated since 1970s by
pioneering works by Pankove [2], Akasaki [3], Nakamura [4], and Dingle [5], who showed its
high potential for optoelectronics. A comprehensive summary and history of research activity
was written by Nakamura [6]. Detailed reviews on the GaN literature can be also found [7].
The main material properties of gallium nitride may be divided into two areas:
structure properties and optoelectronic properties.
Similarly to other main materials of III-Nitride group (AlN, InN) GaN crystallizes
preferentially in the so-called wurtzite structures (Fig. 2.1 a)) with the space group P63mc (no.
186), that has a hexagonal unit cell (Fig. 2.1 c)). GaN exists as well in cubic zincblende structure
with the space group F43m (Fig. 2.1 b)). The wurtzite phase is the most thermodynamically
stable configuration for the III-nitrides under standard conditions [8].
7
Figure 2.1: Perspective views along [0 0 0 1] direction of wurtzite and cubic zincblende GaN, a) and b)
respectively [9]. The large circles represent gallium atoms and the small circles nitrogen. c) The hexagonal unit
cell of GaN defined by the lattice parameters: the length of the hexagon’s side (a), (b) and the height (c) of
the hexahedron.
The wurtzite structure consists of alternating biatomic close-packed (0001) planes of Ga
and N pairs stacked in an ABABAB sequence. Atoms in the first and third layers are directly
aligned with each other. The lack of inversion symmetry in the hexagonal cell leads to very
strong polarization effect in group III-Nitride materials – in case of GaN, crystals surfaces have
either a Ga-polarity (designated (0001) or (0001)A) or a N-polarity (designated (0001) or
(0001)B). Thus, different properties like: surface morphology, chemical reactivity and growth
conditions are reported for these two polarity configurations. Figure 2.2 depicts two possible
polarities of wurtzite c-plane GaN [10]. It is worth to mention that GaN layers grown by
MOCVD are usually Ga-face.
Figure 2.2: Atomic arrangements in two possible GaN polarities: Ga-faced and N-faced [10].
8
The main physical properties of III-N nitrides are presented in Table 2.1.
Material Lattice constants
[nm]
Melting point
[ᵒC]
Thermal conductivity at
300K (W cm-1 K-1)
Bandgap Eg
(300K) [eV]
Thermal expansion
coefficient (10-6K-1)
AlN a0=0.311
c0=0.498
3487 2.85 6.20 5.3 ║ c-axis
4.2 ┴ c-axis
GaN a0=0.318
c0=0.518
2791 1.30 3.40 3.2 ║ c-axis
5.6 ┴ c-axis
InN a0=0.354
c0=0.570
2146 0.80 0.70 [11] 3.7 ║ c-axis
5.7 ┴ c-axis
Table 2.1: Main physical properties of III-N nitrides [12].
GaN is an extremely important semiconductor with a wide, direct band gap ( ~3.4 eV).
Compared to Si ( ~1.1 eV) and GaAs ( ~1.42 eV) the GaN band gap is almost 3 and 2.5 times
bigger, respectively. It results in strong light emission in the blue and ultraviolet spectrum
range. Because of its large Eg, high break down field and high saturation drift velocity, it is
a prime candidate for high-temperature, high-voltage and high-power optoelectronic device
application.
The main optoelectronic properties of III-N nitrides are summarized in the Table 2.2.
Material Bandgap Breakdown
Field (cm-1)
Index of
refraction
Dielectric constants Electron
mobility
(cm2V-1s-1) Type Value
[eV]
Static High
frequency
AlN
Wurtzite
politype
Dir
ect
6.2 (300K)
1.2-1.8x106
2.15 (3eV)
9.14 (300K)
4.6 (300K)
300 (300K) [13]
GaN
Wurtzite
politype
3.42 (300K)
3-5x106
(300K)
2.85 (300K, 3.42eV)
10.4
(E║c)
9.5
(E┴c)
5.35
~1400 (300K)
GaN
Zincblende
politype
3.2-3.28 (300K)
~5x106
2.3, 2.9 (at 3eV)
9.7 (300K)
5.3 (300K)
≤1000 (300K)
InN
Wurtzite
politype
0.7 [11]
1x106 [14]
~2.9 [15]
15.3 (300K)
8.4 (300K)
≤3280 (300K) [16]
Table 2.2: Main optoelectronic properties of III-N nitrides [17].
9
2.2 Nanostructures advantages compared to the bulk materials
A nanowire (NW) can be defined as a nanostructure, with a diameter in the order of
nanometers and a length, which is a multiple of its diameter. Alternatively, nanowires are called
nanostructures, with nanoscale diameter or thickness and an aspect ratio of the length to
the diameter, which is higher than 10. Commonly in literature, wires which exhibit a diameter
below 200 nm are called nanowires. Due to the similar morphology and geometry, but bigger
dimensions, rods with a diameter in the microscale are defined as microrods or microwires.
Due to their unique morphology, such structures open new possibilities to overcome some
limitations of current conventional planar devices. The research on the fundamental
understanding as well as technological development of NWs was of increasing interest from
the past 15 years.
The small lateral dimensions help nanowires to effectively relax the strain generated at
the nanostructure-substrate heterointerface without the formation of extended defects, typically
observed in highly mismatched materials. It is possible, because nano-islands can deform
vertically as well as laterally, which is a benefit over the planar layers. Thus, nanorods usually
exhibit high crystal quality. These properties offer an alternative for widely used planar
(InGaN/GaN) LED devices. The efficiency in standard planar LED, suffers from a high rate of
non-radiative recombinations, which originates from crystal defects such as threading
dislocations (TDs). Conventional GaN epitaxy, due to lack of cheap, large wafer, commercially
available bulk GaN substrates, commonly employs foreign substrates with huge lattice
mismatch, like Al2O3, Si or 6H-SiC. The substrate employed determines the crystal orientation,
polarity, polytype, the surface morphology, strain and defects concentration of the GaN film
[9]. In 2D nitride layers, TDs are caused by the large lattice mismatch between commonly used
substrates like silicon or sapphire and GaN layer. In case of NWs structures this issue might be
overcame.
The second advantage of NWs is their large surface-to-volume ratio in comparison to 2D
layers or bulk materials. This unique morphology, utilizing nanorods side walls, promises to
enhance the external quantum efficiency of LEDs by the improved extraction of light.
Moreover, by increased surface-to-volume ratio, the NW-based chemical and biological sensors
are expected to assist higher adsorption rates of the targeted molecules. Therefore their
sensitivity shall be higher than for planar devices.
10
The longitudinal nanostructure shape and small diameter allows lateral confinement of
carriers, which changes the quantum-mechanical effects compared to bulk material. This paves
the way to understand new aspects of spintronics.
The NW geometry enables another interesting property – the wave-guiding [18]. Therefore,
nanorods are considered as an interesting building block for nanoscale photonic devices.
The diameter of the NWs and the density of the NW assembly might be controlled by the use
of catalyst. Additionally, the pre-patterned substrates allow the precise control of the NW
position on the substrate, which makes further processing steps much easier, for example
contacting.
2.3 GaN NW applications for next generation devices
The studies on III-N material system have their origin 45 years ago, when the first
epitaxial GaN layers where grown [19]. Over the years, the Nitrides (GaN, AlN, InN) and their
alloys were widely used in many applications, like blue, green and white LEDs or blue laser
diodes.
Nowadays, due to superior material characteristics and nanostructure properties, GaN
nanowires are in major focus of interests of the scientists conducting their research in the field
of opto- and nanoelectronics. A number of next generation devices, based on the (In)GaN NWs
were already proposed and discussed by several research groups. The theoretical background
as well as some of these applications will be presented in this paragraph.
GaN NWs architecture enables two different configurations for functional LED devices.
First one is based on axial p-n junctions, whereas second one is formed by coaxial core/shell
structures – Fig. 2.3 a) and b) respectively.
Figure 2.3: Schematic draft of axial and core/shell nanowire, a) and b) respectively.
11
The core/shell type of the NWs offers some advantages over typical axial configuration,
known from planar devices. Beside larger light-emitting area, the enhanced carrier injection
through a larger junction area is reported [20].
There are several possible root causes, why high efficient emission in standard LED
devices, especially in the green spectra (“green gap” [21]) is very difficult to obtain. First of all,
the film deposited at low temperatures, causes the formation of defects, which act as trapping
points - SRH (Shockley-Read-Hall) centres and thus weakening the internal quantum efficiency
(IQE) [22]. Secondly, the radiative recombination efficiency is lowered due to quantum-
confined Stark effect (QCSE). The energy band is tilting, which decreases the overlap integral
of electrons and holes by spatial separation [23]. Further mechanisms proposed as a cause of
the efficiency droop are Auger recombination [24] and carrier leakage [25]. The device
architecture, based on nanowire structures, may help to overcome these issues. Besides the 3D
surface, which helps to release the strain, NWs offer the higher radiative recombination rate in
comparison to planar devices. The calculated IQE versus current density for c- and m-plane is
depicted on Fig. 2.4. Please consider that IQE curves were simulated under both optical
excitation and electrical injection.
Figure 2.4: Calculated internal quantum efficiency versus current density for c-plane [a) and b)] and m-plane [c)
and d)] growth, under zero bias [a) and c)] or a 3.5 V applied voltage [b) and d)]. X indicates the In content in
the InGaN alloy. Figure adapted after [26].
12
The pioneer example of the flexible and controllable nanowire based multicolour LED
was demonstrated by Lieber group [27]. The architecture of this n-GaN/InxGa1-xN/i-GaN/p-
AlGaN/p-GaN core/multishell device is depicted in Figure 2.5. By tuning the In concentration
in the InxGa1-xN shell layer during nanowire growth, the wavelength of emitted light can be
systematically adjusted from 367 to 577 nm.
Figure 2.5: Nanowire-based multicolour LED: a) Schematic of the heterostructure cross-section and energy band
line-up. b) Optical microscopy images collected from around the p-contact of nanowire LEDs in forward bias,
showing different colour of emitted light: purple, blue, greenish-blue, green and yellow. c) Normalized electron-
luminance spectra recorded from five representative forward-biased NW LEDs with 1%, 10%, 20%, 25% and
35% of In content in the InGaN alloy (left to right), respectively. Figure adapted after [27].
The high advantages of nanoLED are currently commercially brought by a GLO
company [28]. This developer of nanowire LED addresses research in the area of nanowire
materials, epitaxial growth conditions on a variety of substrates with various device structures
and fabrication processes [29], [30], [31], [32], [33].
Another device benefiting from nanoarchitecture is GaN nanowire-based transistor [34],
[41], [42]. The example of device realization and performance is depicted in Figure 2.6.
Presented dopant-free GaN/AlN/Al0.25Ga0.75N radial NW heterostructure exhibit very good
temperature-dependent transport data. The intrinsic electron mobility of 3100 cm2/VS and
21000 cm2/VS were measured for room-temperature and 5K, respectively. Moreover,
13
investigated heterostructure exhibited scaled transconductance (420 mS/µm) and subthreshold
slope (68 mV/dec) values which showed the high potential of NW as building blocks for more
complex architecture development.
Figure 2.6: GaN/AlN/AlGaN NW-based transistor. a) Left: cross-sectional, high-angle annular dark-field
scanning TEM image of a radial nanowire heterostructure. Scale bar is 50 nm. Right: Band diagram illustrating
the formation of an electron gas (red region) at the core-shell interface. b) Intrinsic electron mobility of
a transistor as a function of temperature (after correction for contact resistance). c) Logarithmic scale Ids-Vg curve
recorded at Vds = 1.5V (channel length 1 µm, 6 nm ZrO2 dielectric). Inset shows the linear scale plot of the same
data. Figure adapted after [35]
Another GaN nanowire-based device examples, like lasers [36], [37], [38] or solar cells
[39], [40] can be found elsewhere. The high commercialization potential of nanowire devices
manifests in high number of patent applications in the field of growth techniques [43], [44],
[45], [46], [47] as well as device fabrication [48], [49], [50].
14
Chapter 3
GaN NW growth techniques state-of-
the-art by MOVPE
In this chapter the main advantages and challenges of GaN NWs growth by MOCVD
will be presented. The challenges and solutions for GaN-on-Si will be discussed. The NW
morphology as a polarity consequence will be underlined. Moreover, up to date the state-of-
the-art of the MOCVD NWs growth will be summarized. In principle, three main bottom-up
growth techniques will be described: the Vapor-Liquid-Solid (VLS) growth mode utilizing Au
nanoseeds as a catalyst; selective area growth (SAG), based on the mask approach and finally
self-assembled in-situ growth mode. The second alternative nano-technological approach based
on subtractive (“top-down”) methods of NWs realization is not discussed, since it is rather
nanofabrication and processing technique (material excess is physically or chemically removed
from the bulk-like or epitaxially structured material) than growth itself, which is the main focus
of this work. The examples of top-down NWs preparation could be found elsewhere [51].
3.1 Advantages and challenges of GaN NWs growth by MOCVD
3.1.1 Challenges and solutions for GaN-on-Si integration
The GaN-on-Si integration is a very desirable aspect for optoelectronics. Silicon
substrates exhibit plenty of advantages over other wafers commonly used for heteroepitaxy
processes, like sapphire or SiC. The main advantage is a cost factor. Si wafers are much cheaper
than other counter candidates. Furthermore, silicon substrates exhibit a scaling up potential,
which will manifest in further cost savings in the future. Albeit integration of GaN on silicon
substrate is strongly desirable, there are consequently several challenges to overcome. First of
all, huge lattice mismatch between GaN and Si could case high dislocation density and as
a consequence might lead to reduction of usable area and decrease of material quality.
The second point is a crystal symmetry break. GaN crystallizes mostly in wurtzite structure
(hexagonal unit cell), whereas Si represents a diamond structure (cubic unit cell).
The difference in thermal expansion coefficient (TEC) is a root cause of high tensile strain,
15
defect formation and even cracking of the grown film. Another point is a high Ga reactivity
with Si. Such a strong chemical aspect may lead to undesirable melt back etching effect if GaN
is deposited directly onto Si surface. Moreover, Si substrate reacts very easily with ammonia
and forms amorphous SiNx, which passivates the surface and prevents GaN growth. An
example of melt back etching of Si by Ga and formation of amorphous SiNx after deposition of
GaN directly on Si can be seen on Fig. 3.1 [52], [53].
Figure 3.1: a) Meltback etching of Si by Ga [52], b) Ga-rich, Si-rich and SiNx formation after GaN deposition
directly on Si substrate [53].
There are several solutions to deal with the challenges of 2D GaN-on-Si integration
described above. The idea is to introduce a layer stack, in which each single layer addresses one
or more issue. First of all, to protect the Si surface, the AlN nucleation is introduced to the layer
stack. Afterwards, thicker AlN and one or more AlGaN layers are grown to systematically
reduce large lattice mismatch and limit the enhanced misfit dislocations. Such a buffer is
sufficient for GaN growth. However, to compensate the stress and terminate the propagation of
dislocations into subsequent layers the strain and defect engineering layers are introduced.
A thin low temperature AlN interlayer in between GaN parts enables building up compressive
stress during the growth process and finally results in unstressed cooled down heterostructure.
The defect reduction might be achieved by depositing thin in-situ SiNx masking layer between
thicker GaN layers [54].
All of the necessary steps make GaN-on-Si integration a sophisticated growth process.
GaN NWs open new possibilities for this approach. The defect termination layers as well as
strain engineering layers are no longer necessary. The only step, which is in common with 2D
approach, is a thin AlN protecting layer. The unique morphology of nano- and microrods
enables growth on preferential substrate of choice – Si, and moreover, offers plenty of further
16
advantages over the bulk (see chapter 2.2). The MOCVD, as a deposition technique, provides
additional benefits. In comparison to other growth mechanism, MOVPE is very fast and cheap,
which makes it the natural choice for a commercial market applications. Taking all this points
into consideration, the combination of MOCVD as a deposition technique, the nanostructure
advantages over the bulk and easier GaN-on-Si integration opens new way for the next
generation applications ready for market commercialization based on GaN NWs.
3.1.2 NW morphology as a polarity consequence
The very interesting property of the NWs structures is their morphology as a polarity
consequence. Different polarity of the seed layer results in different morphology of the rods.
As discussed in the theoretical introduction (see chapter 2.1), III-N materials exhibit two
possible polarities: metal- and N-polar. These two configurations determine large number of
different material properties [55], [56]: impurity [57], [58] and dopant incorporation [59],
surface reactivity [60] or thermal stability [61]. Moreover, the polarity of the subsequently
grown layer is inherited by the layer grown below. Thus, the polarity aspect is crucial in terms
of buffer development for NWs growth. It is found that the N-face nucleation layers lead to flat
tops of the wires. On the contrary, the metal-face layers result in pyramidal tips of the rods
grown on this type of the seeding layer [62]. Due to the hydrogen passivation effect [63],
the crystallographic planes with higher indexes - {11̅01} r-planes family, have smaller growth
rate than c-plane direction and thus the vertical sidewalls are difficult to produce in that case.
Consequently, the metal-polarity may suppress the vertical growth completely. As a result, only
small pyramids with hexagonal base will be found on the substrate.
This very important aspect of polarity of the rods was deeply studied by several groups
[62], [63], [64]. Besides the morphology as a polarity consequence, the mixed polarity issue
was reported [63], [65], [66], [67]. Basically, it means that the NWs are not grown with single
crystallite. Two opposite polarities are present within the heterostructure. The mixture of
polarity leads to inversion domain boundaries (IDB) between two regions: Ga- and N-polar
part, respectively. The structures containing IDBs may strongly reduce the efficiency of
the final LED device, due to the high yellow luminescence [68], [66].
The mixed polarity phenomenon may manifest as a formation of two regions with
slightly different height on the {0001} top facet, as is shown on Figure 3.2 a). A simple KOH
etching experiment is an easy way to verify the polarity of the rods. It is found, that N-polar
17
regions are etched away by KOH solution, whereas metal-polar areas are etch resistant [69].
The result of etching the mixed polar microrods is presented on the Fig. 3.2.
Figure 3.2: a) GaN NW with two different polarity domains – the pattern on the c-plane top facet induces
the existence of both polarities within the structure, b) the same GaN NW after KOH etch; flat region on the c-
plan facet represents Ga-polarity, whereas rough surface is attributed to N-polar regions [65], c) GaN NW after
KOH etch, arrows indicate the remaining Ga-polar regions [68].
Both images represent the GaN microstructures after performed KOH etching
experiment. As can be seen, some areas are etched away – outer side of the {0001} plane and
top part of microwire – Fig. a) and b), respectively. On the left image, one can see the not etched
remaining IDB, which is covering about 1/8 of the top plane of the rod and which is connected
to the m-plane side wall. Contrary, on the right image, the remaining Ga-polar regions formed
the residual base of the rod with a crown like pattern on top.
To eliminate the mixed polarity issue and enhance the single N-polar growth of the GaN
microrods, two steps SAG process was proposed by Waags group [68], [45]. The idea is based
on the SAG with two growth steps, called: “truncated pyramid + column growth approach”.
First, after the N-polar GaN nucleation, the truncated pyramid is formed. The inner part, grown
through the mask opening is N-polar, whereas the outer part of the structure, partially grown
on the mask is Ga-polar. Since the Ga-polar r-plane is a slow growing plane [63], due to the N-
H bonds passivation [70], the vertical GaN columns are grown directly on the N-polar base.
Therefore, after certain time, pure N-polar GaN rods might be grown.
The polarity control of the seed layer is much easier to obtain on sapphire substrates
than on silicon. The nitridation of the Al2O3 surface favours the growth of N-polar GaN crystal
18
[71], [72]. Hence, the initial conditions enhancing the vertical growth might be established. On
the contrary, the buffer polarity control on the silicon substrates is an open issue.
3.2 Au- catalyst induced VLS growth mode of GaN NWs
The first synthesis of GaN nanowire was reported in 1997, when Han et al. used
a reaction that was confined inside a carbon nanotube [73]. However, the vapor-liquid-solid
growth mode has its origin in the pioneer work on the micro-sized silicon whiskers done by
Wagner and Ellis in 1960s [74], [75], [76]. Since then the VLS growth mode has become
the widely used technique. Many metal elements, like Au, Ni, Fe or In [77], [78] were utilized
as a catalyst to synthesize a large number of inorganic materials: GaN [79], [42], [80], [81], Ge
[82], Si [83], [84], [85], GaAS [86], InAs [87], [86] or GaP [88].
The VLS growth concept is based on nanometer-sized metallic particles, which form
the low-temperature eutectic alloy with NW material [89]. The preferential nucleation of
the material is found at the droplet-crystal interface, since the liquid-solid interface acts as
a sink for the arriving precursor molecules and lead to the reactant incorporation [90].
Figure 3.3: Schematic model of the VLS growth of GaN NWs utilizing Au nanoparticles [91].
Figure 3.3 depicts a model of the VLS growth of a GaN NW, which is initiated by a gold
catalyst particle. The starting point of the VLS approach is always a nanoscale metal seed.
The metal catalyst might be deposited onto the substrate as already formed droplets, from
the commercially available dispersible nanomaterials solutions [92] or as a very thin film.
In the second case, if the catalyst is exposed to elevated temperatures, it melts and forms liquid
droplets. The formation of droplet shape is energetically favourable because it reduces
the surface energy. According to the VLS model by Wagner and Ellis, these droplets create
a highly selective area for the deposition of the semiconductor material, which is supplied from
19
the vapor phase. It results in preferred incorporation of precursors into the metal-catalyst droplet
and therefore an eutectic alloy is formed. By further supply of the growth species,
the concentration inside the droplet rises until supersaturation level of the catalyst metal with
the growth elements. The supersaturation results in precipitation of the growth elements at
the solid-liquid interface of the droplet and the NW. The precipitated growth elements are built
into the underlying crystal and the NW grows under the droplet. The additional growth events
might be described as shown in Fig. 3.3:
1- Mass transport through the vapour phase
2- Dissociation reaction on the catalyst particle, either directly from the vapour or by
diffusion from the side facets of the NW to the catalyst particle
3- Diffusion of reactants through the particle
4- Precipitation of the growth element at the liquid-solid interface forming
the semiconductor NW
5- Absorption on the substrate or NW sidewalls
6- Surface diffusion on the substrate or NW
7- Desorption from the substrate or NW sidewalls
8- Film growth on the NW sidewall or substrate [93].
A typical SEM image of GaN NWs grown on sapphire substrate by VLS mode is
presented on Fig. 3.4. The metal catalyst used in this experiment was a very thin Au film (1 nm
of nominal thickness). The impact of the MOCVD process parameters on the NW morphology
as well as structural characterization by TEM might be found in the literature [94].
Figure 3.4: Typical SEM images of GaN NWs grown on the sapphire substrate by the VLS technique. The metal
catalyst used in the growth experiment was a very thin Au film.
Besides sapphire substrates, the successful Au-assisted growth of GaN NWs on the Si
is also discussed [95], [96].
20
3.3 Selective area growth ( SAG) of GaN-based nanocolumn-arrays
The SAG offers the control of position and size of the rods as well as easier emission
wavelength control in comparison to other NW growth techniques. The wavelength emission
depends on the alloy composition of QWs within the nanostructures and the morphology of
the nano- or microcolumns. Since the VLS and self-organized growth techniques offers rather
a broad statistical distribution of these properties, the SAG becomes currently the leading
approach for GaN-based devices for lighting applications.
The SAG approach employs template preparation for subsequent position controlled
growth of microcolums. Most commonly used process steps for structuring the substrates for
obtaining the desirable patterns are: lithography, based on optical and e-beam solutions;
nanoimprint or etching. The choice of process technology implies a set of features and
properties of the template. For example, e-beam lithography enables the possibility to obtain
the very precise pattern in the nano-scale. On the other hand, in comparison to standard optical
lithography, e-beam is extremely expensive, time consuming technique and it is addressed only
for small areas. Yet, the proper substrate preparation is a very important step in the technology
chain.
Hersee et al. [97] proposed the pulsed MOCVD process, which allowed them to grow
vertical GaN NWs on the GaN template prepared on the SiC, sapphire and Si(111) substrates.
They employed a 30 nm SiNx mask deposited by low-pressure chemical vapor deposition
(LPCVD) on the GaN template and afterwards, by interferometric lithography and dry etching
they formed the circular apertures with the average diameter of 220 nm. Their first observation
was a NWs geometry lost due to lateral growth, in case when they applied continuous flux
growth with V/III ratio of 1500. However, once they switched to the pulsed growth mode,
before the NWs emerged from the growth mask, then the vertical growth of GaN nanocolums
with limitation of lateral growth aspect was accessible. The reported vertical growth rate of
the rods achieved the value of 2 µm/h, whereas the ratio of vertical to lateral growth was in
excess of 1000. The NWs, besides their controlled diameter by the diameter of the growth mask
aperture, exhibit single crystal nature. Moreover, the threading dislocations (TDs) were
observed only in the GaN film beneath the growth mask, but not in the GaN NWs.
The mechanism of selective area growth of GaN nanorods by pulsed mode MOVPE was
also studied by Lin et al. [98]. They proposed the kinetic model, which was then supported by
the performed experiment. They reported two factors, which influence the microrod shape by
the pulsed growth mode. First one is the difference in the adsorption/ desorption behavior of
21
Ga adatoms on the top (0001) c-plane and the boundary {11̅00} m-planes. The second factor is
the growth behavior of the semi-polar planes, in particular the {11̅01} plane. The successful
GaN NWs growth requires the suppression of lateral growth while maintaining vertical growth.
The vertical growth is promoted by Ga adatoms, which possess a longer residual time on the top
c-plane in comparison to m-planes, due to the higher sticking coefficient of Ga adatoms on
the c-plane. GaN nanostructure growth is mainly governed by Ga adatom kinetic behaviour,
which varies under different growth conditions. Thus, an appropriate Ga interruption duration
in the pulsed growth mode shall be optimized in terms of GaN nanostructure growth.
The morphology development of GaN nanowires using a pulsed mode MOCVD was
also investigated by Jung et. al. [99]. The diameter, length and orientation of the NWs are
controlled via growth parameters such as growth temperature as well as precursor injection and
interruption durations. The synthesis of GaN nanorods is governed mostly by the kinetic
behaviour of Ga adatoms, since at higher growth temperatures the enhanced surface diffusion
of Ga adatoms was reported. Additionally, the longer TMGa injection duration resulted in
the higher lateral growth rate of GaN nanostructures. The increased TMGa injection time
eventually led to the transformation of nanostructures into a thin film. Opposite, the longer NH3
injection durations lead to the nanostructure shape change from wires to the hexagonal
pyramids due to the overlap with N and Ga adatoms. Finally, the shape of microrods depends
on the initial nucleation step. The morphology on top of the rod depends on whether or not
the complete filling of the mask opening step was applied. The complete filling process results
in the formation of flat top consisting of (0001) c-plane, while the result of the incomplete
filling step is truncated pyramidal shape consisting of (0001) c-plane and residual {11̅01}
facets.
Chen et al. [64] studied the GaN NWs growth on structured sapphire substrates using
a SiNx mask. The patterning was obtained by nanoimprint lithography leading to an array of
circular holes of 400 nm diameter spaced about 1.1 µm. The proposed process includes three
steps: growth at 950 °C, annealing under NH3 and growth at 1040 °C. The growth temperature
smaller than 1000 °C, favors the pyramidal shape of the grown structures whereas the prismatic
shape is observed for higher temperatures applied. On the other hand, the homogeneity of
nucleation selectivity in SAG is dramatically degraded at higher temperatures. The reason for
that is desorption, precursors migration length on the mask and the “sink” effect due to the non-
uniform GaN seed nucleation. Taking these points into account, the proposed process enables
high homogeneity of nucleation and subsequent growth of prismatic GaN nanocolumns on
structured sapphire substrates.
22
Interestingly the Au catalyst might be used to structure the substrate for the subsequent
selective area growth. Song et al. proposed a SAG of GaN NWs on nano-patterned Si(111)
substrates formed by the etching of nano-sized Au droplets [100]. The 10 nm thick Au film on
Si, formed the Au droplets after annealing in 650 °C under H2 ambient. The subsequent
oxidation step to form a SiO2 layer was performed in the furnace at 800 °C for 10 min.
Afterwards, the Au residuals were removed by the etchant and thus, the nano-pattern was
formed in the places, where previously nano-droplets were localized. To initiate the GaN NWs
growth, the predeposition of TMAl was supplied to deposit Al in nano-patterns.
3.4 Self-assembled growth of GaN NWs
Koester at al. [101] extensively studied the self-assembled growth of catalyst-free GaN
wires by MOCVD. Their research is one of the most cited investigation in the field of the GaN
NWs. They proposed the self-organized growth approach for GaN NWs on sapphire substrates.
The idea of GaN nano- and microstructure growth is based on the in-situ deposition of a thin
SiNx layer prior to the nanowire epitaxy. First step is an Al2O3 wafer preparation. After the high
temperature cleaning process under H2 ambient, the sample surface is nitridized to form a thin
AlN layer. Afterwards, a 2 nm thick SiNx interlayer is deposited by simultaneously injecting
45 sccm of SiH4 and 4000 sccm of NH3. Later experiments showed that the SiNx deposition
time is a critical parameter for a subsequent GaN NWs growth. The SiNx layer thickness was
not changing with respect to longer deposition time after reaching the critical thickness of about
2 nm. However the morphology of the GaN NWs was strongly affected by varied SiN
deposition time. The rod growth was performed in two steps: short nucleation under N2 carrier
gas and subsequent vertical growth of the wires with small V/III ratio of around 15 and in
strongly N2-diluted N2/H2 carrier gas mixture. As mentioned before the thickness of the SiNx
layer does not depend on the growth time, yet the roughness and density related to the surface
chemistry varies. It was found that SiNx deposition times shorter than 50 s do not support
the GaN wire growth. Contrary, 100 s of in-situ masking layer resulted in well vertically aligned
rods. The longer deposition times lead to a degradation of the GaN nanostructures and also lead
to nonvertical growth. Authors assumed that since the SiNx layer is chemically inert with
respect to GaN and it exhibits the temperature stability, the GaN nucleation formation might
origin in the weak points of the mask layer with locally changed density/roughness. The studied
GaN nucleation step revealed the formation of small hexagonal seeds of 50-200 nm height and
0.8 nm RMS after 100 s. The differences in the grain size and facet lengths explain the further
23
wide distribution of wires shape and size. Moreover, the lateral growth of the seeds may result
in the formation of dislocations, pits or kinks in the vertical rods. Furthermore, the coalescence
of the nuclei leads to the grain boundaries. Thus, the shorter nucleation time results in lower
density of the seeds and less grain dislocation. Consequently better structural quality of
the vertical GaN NWs might be reached. The supporting role of silane during the vertical
growth of the rods was shown (see Chapter 5). It was found that SiH4 support might be turned
off after reaching a 5 µm height of the rods. Finally, the growth rate might be controlled by
changing the total carrier gas flow. The 8000 sccm in MO- and hydride- lines leads to
the vertical growth rate of about 30 µm/h. The decrease of carrier gas to 2000 sccm resulted in
a higher deposition rate and the vertical growth rate exceeded 145 µm/h.
The mechanism of nucleation and growth of catalyst-free self-organized GaN columns
on sapphire substrates was also investigated by Wang et. al [102]. Two steps growth mode
including GaN nucleation (100 s) and vertical growth of the wires were studied. GaN microrods
were grown using a small V/III ratio of 8.3 and silane injection of 536 nmol/min under 100%
N2 carrier gas. The growth results showed two different kinetic regimes: mass-transport-limited
growth and thermodynamically limited growth. The first regime was found at lower deposition
temperatures between 960 °C to 1020 °C. This observation is attributed to the dominance of
surface and gas-phase diffusion since the density of nucleation seeds decreases slowly as
the temperature increases. At the higher temperature applied, the reaction between partially
decomposed NH3 and TMGa above the susceptor becomes stronger. The density of nucleation
seeds strongly decreases and the grains become larger. It is attributed to the increase of diffusion
coefficient DGaN with increasing temperature. Interestingly, the height of the columns does not
depend on their diameters for the same growth conditions. A growth rate of about 14 and
150 µm/h can be achieved for N2 and H2 ambient, respectively. There was no saturation of
height observed for structures that reached more than 40 µm. The diameter of the microrods
was not uniform. The upper parts of the wires became broader while the bottom parts with
the smaller diameter remained constant. It happens due to Ga adatoms accumulation near or on
the top surface of the rods. The NWs were found to be mostly Ga-polar, only a few investigated
microrods showed a Ga-polar core and a N-polar shell.
The self-catalyzed GaN nanorods grown on sapphire substrates in horizontal MOCVD
reactor were investigated by Tessarek et al. [66]. The hexagonal, vertically aligned
microstructures with diameters from a few µm down to 100 nm, reached heights up to 48 µm.
The calculated density of the NWs was up to 8x107/cm2 and the vertical growth rate was up to
120 µm/h. This relatively high growth rate was reached by introducing high amount of silane
24
(Ga/Si ratio of 1339). On the other hand, with so high growth rate the morphology of structures
suffered. The sidewalls of the rods were not flat. Contrarily, many steps and tilted facets were
observed. A nucleation step including the formation of nanoflakes, which are subsequently
considered as a nucleation sites for NWs growth was proposed. The growth of flakes takes place
at elevated temperatures with higher V/III ratio (373) and under H2 ambient. The flake-like
structures exhibit a diameter between 50-150 nm and densities from 3x108/cm2 to 1x109 /cm2
depending on the deposition time (1 and 4 s, respectively). After successful formation of
the nucleation sites, the nitridation step is performed and subsequently the vertical growth of
the NW takes places. Smaller V/III ratio and reduced temperature was applied in order to grow
microrods. As expected, there is a correlation between the density of the nanorods and density
of flakes. Not every flake is an origin of NW growth and thus, one order of magnitude difference
was calculated between flakes and NWs densities. Interestingly, there was an anti-correlation
reported between size of the flakes and nanorods. The bigger flakes resulted in the formation
of wires with smaller diameter in comparison to those grown on the flakes with smaller size.
Therefore, it was concluded that the diameter of the rods was determined not by the size of the
flakes but by their density. The structural properties of the rods were characterized by TEM. It
was found that the GaN/sapphire interface was full of dislocations, however after reaching
about 500 nm of microrods height, these defects were eliminated. The dislocations are bending
towards the sidewalls and therefore, the NWs are nearly free of defects and strain. On the other
hand, a pair of vertical planar dislocations occurs in the center of the rods. These extended over
the full height of the structure defects are related to the inversion domain boundaries (IDB).
The inner part of the NWs exhibits Ga-polarity, whereas the shell is N-polar. Moreover, on top
of the c-plane rod facet cylindrical structures might be found. The inclined facets indicate Ga-
polarity of such flakes (see chapter 3.1.2). Finally, the evidence of self-catalyzed, Ga-induced
growth was shown. The driving force for self-assembled GaN NWs growth is based on the VLS
growth mode (see chapter 3.2). In this case, the Ga-droplets are the driving force for columnar
growth. The wide distribution in diameter of the structures can be explained by the Ostwald
ripening [103]. The larger droplets shrink at the expense of smaller droplets. This effect can be
suppressed by increasing the total reactor pressure. Then, the mobility of adatoms on the surface
is reduced and therefore only the formation of nanocolums takes place without an additional
randomly shaped microstructure growth.
Salomon et al. [104] tried to transfer the self-organized growth model proposed by
Koester [101] from sapphire to silicon substrates. The GaN NWs were grown on a thin 10-
50 nm AlN layer deposited on n-type Si(111) substrates. The HRTEM images revealed also
25
the spontaneous formation of a 2 nm thick amorphous (or nanoscrystallized) SiNx layer directly
on top of the Si substrate attributed to the high temperature AlN growth on silicon.
The vertically aligned GaN wires exhibit irregular hexagonal cross section and a quite wide
distribution in length and diameter. The calculated density was approximately 106 wires/cm2.
The structural properties of the NWs were investigated by XRD. The Δω rocking curves of
the GaN (0002) and (0004) Bragg peaks reveal the 1.37° of FWHM. This relatively high value
in comparison the GaN wires on sapphire (0.61° [101]) can be attributed to the AlN buffer
quality and to the nucleation on the defects. The vertical GaN NWs were used as a template for
subsequent MQWs deposition with target of In0.18Ga0.82 (1 nm) / GaN (10 nm) structures.
The electroluminescence (EL) spectra exhibit a violet emission centered at 420 nm and
a weaker low-energy contribution at 460 nm. Interestingly, a defect band (usual yellow band
close to 550 nm) was not observed. The two emission contributions for 420 and 460 nm are
explained by the presents of both radial (420 nm) and axial (460 nm) MQWs. This 40 nm shift
of the emission might originate from the variation of In composition or well thickness.
Interestingly, self-organized growth of GaN microrods might be also realized on
graphene films. Chung et al. [105] reported the GaN microstructure realization on the graphene
templates for flexible light emitting diodes.
26
Chapter 4
Experimental setup and
characterization methods
This chapter provides basic information regarding experimental setup and utilized
characterization methods. First section presents a background of the MOVPE growth by
an AIXTRON 3x2” CCS (close coupled showerhead) platform. The second part includes
the background and principles of several synchrotron based techniques such as X-ray
fluorescence (XRF), X-ray absorption near edge structure (XANES) and nanoscale X-ray
diffraction. The end of the chapter describes also some complementary structural and optical
characterizations methods.
4.1 Metalorganic vapor phase epitaxy (MOVPE)
All epilayers investigated in this study were grown in an AIXTRON 3x2” CCS (Close-
Coupled Showerhead) reactor – Fig. 4.1. The deposition can be performed on three times two
inch wafers (3x2”) or one time four inch substrate in the same process .The platform is equipped
with additional in-situ characterization tools: Argus and LayTec EpiTT-Curve. The first one
gives information about surface temperature distribution of the wafers during the epitaxial
growth. Argus is equipped with a line of six diodes. Diode 1 and diode 6 read the temperature
of the wafer from the center to the edge, respectively. The other four diodes read the surface
temperature in between the center and edge. The outcome of the Argus measurement helps to
obtain the homogeneous, uniform temperature profile across the wafer during MOCVD process
and as result helps to control the curvature of the wafer. The second characterization tool –
LayTec EpiTT-Curve collects information about the temperatures (wafer and thermocouple),
curvature of the wafer as well as reflectance of the grown layers, operating with two light
sources with the wavelengths (λ) of 405 nm and 950.4 nm. For the growth of the layers standard
presursors were used. Trimethylalluminium (TMAl), trimethylgallium (TMGa), triethylgallium
(TEGa) and trimethylindium (TMIn) were used as sources of Al, Ga and In, respectively.
27
During multi quantum well (MQW) deposition TEGa were chosen instead of TMGa. By
changing the Ga precursor from TMGa to TEGa, the carbon content is reduced. The reason for
this is mainly due to different behaviour of carbon atoms in the ethyl and methyl radicals.
The TEGa pyrolyzes unimolecularly by b-hydride elimination with the formation of ethylene
[106] without the production of reactive carbon-containing species. Opposite, TMGa pyrolyzes
by producing highly reactive CH3 radicals, which consequently leads to a higher carbon
contamination [107].
Ammonia (NH3) was employed as N source. As carrier gas, H2 or N2 was introduced.
For AlN and GaN deposition H2 was used, whereas for InGaN growth N2 was utilized due to
etching effect of H2 on In containing layers. The n-type doping source was silane (SiH4). Other
growth parameters like growth temperature, total reactor pressure and V/III ratio will be given
and discussed for the corresponding grown structures in the related chapters.
Figure 4.1: Schematic of an AIXTRON 3x2” MOCVD reactor. A: thermocouple, B: tungsten heater,
C: showerhead, D: reactor Lid, E: optical Probe, F: showerhead water cooling, G: double O-ring seal,
H: susceptor, I: quartz liner, J: susceptor support, K: exhaust.
During MOVPE growth three different, temperature dependent growth regimes are
present – Fig. 4.2. At low temperatures, in the kinetically limited regime - A, the precursors are
not completely pyrolyzed. With increasing temperature more precursors decompose which
leads to an increase of the growth rate. In the mass transport limited regime - B, the growth rate
is nearly temperature independent. In this case the growth efficiency, which is the ratio of
the amount of growth performed and the consumed number of moles of the precursor, is almost
28
constant. It increases slightly due to some minor temperature dependent parameters such as
the diffusion coefficient and the viscosity of the gas. This is the typical growth regime for
MOVPE growth. The optimum temperature depends on the material grown and the precursors
used. In the growth regime C the growth efficiency decreases with increasing temperature.
A reason is that the precursors start decomposing at the hot reactor walls (heterogeneous
reactions) or in the gas phase (homogeneous reactions).
Figure 4.2: Graphical visualization of three different growth regimes for MOCVD process: A – kinetic limited
regime, B – mass transport limited regime, C – reduced growth rate due to desorption from surface instead of
incorporation.
4.2 European Synchrotron Radiation Facility
Figure 4.3: Schematic of a synchrotron facility in Grenoble including an injection system, a storage ring and
beamlines. The injections system consists of an electron gun, a linac and a booster. The parts of storage ring are
radio frequency cavities, bending magnets and undulators or wigglers. Figure adapted from [108].
29
4.2.1 Nanofluorescence XRF
Synchrotron X-ray fluorescence (XRF) is a very powerful non-destructive, multi-
elemental and fast characterization technique for qualitative and quantitative elemental analysis
of materials. It provides a high spatial resolution (around 50 x 50 nm2) and low detection limits.
Thus, it allows the study of very small samples, like quantum dots or nanostructures giving
elemental traces analysis below parts per million levels.
Figure 4.4: Draft of the experimental setup for recording XRF and XANES using a synchrotron X-ray nanobeam
at the beamline ID22. Figure adapted from [108].
The XRF measurements were performed at the beamline ID22 (new beamline ID16B)
using an in-vacuum undulator U23 with the setup shown in Fig. 4.4. XRF maps were obtained
by scanning the sample with a piezo stage in the nanobeam with a step size if 25x25 nm2 and
integration time of 500 ms per point. More details about the setup and instrumentation
at the beamline ID22 can be found elsewhere [109], [110].
Based on the classical model of an atom, a positively charged nucleus is surrounded by
negatively charged electrons which are grouped in shells or orbitals. Each shell contains
a certain maximum number of electrons determined by the Pauli Exclusion Principle [111].
Each electron can be described through four quantum numbers (principal quantum number,
azimuthal quantum number, magnetic quantum number, and spin quantum number) that
uniquely define it and specify the shell it may occupy. X-rays irradiating a sample may undergo
either scattering or absorption by atoms of the sample. During absorption, the incident X-rays
with sufficient energy will eject an electron from an inner shell of the atom, leaving a vacancy
in the core shell. That hole-state of the core shell is refilled quickly by an electron from an outer
30
shell. Such transition occurs via two competitive processes: X-ray fluorescence and Auger
effects.
Figure 4.5 depicts the XRF and Auger electron yields for K-shell as a function of atomic
number. As can be seen, Auger transitions (solid curve) are more probable for lighter elements,
whereas the XRF yield becomes dominant for higher ones.
Figure 4.5: XRF and Auger electron yields for K-shell as a function of atomic number, solid and dotted curve,
respectively. Figure adapted from [112].
The allowed transitions for the XRF emission induced by photoelectric effect are
specified by quantum selection rules [113].
4.2.2. X-ray absorption near edge structure XANES
X-ray absorption spectroscopy (XAS) is another characterization technique commonly
used in ESRF: It enables to study a short-range order as well as electronic structure of a broad
range of crystalline or amorphous materials. XAS spectrum can be divided into two parts: X-
ray absorption near edge structure (XANES) and extended X-ray absorption fine structure
(EXAFS). The transition between XANES and EXAFS is roughly estimated, at the wavelength
of the excited electron which is equal to the distance between the absorbing atom and its nearest
neighbours. Commonly, XANES covers the region within ~ 50 eV of the absorption edge,
31
including pre-edge and edge regions, whereas EXAFS is a region from 50 to 1000 eV above
the edge.
The X-ray absorption process is based on the photoelectric effect. A transition between
two quantum states takes place from an initial state with a presence of X-ray and a core electron
to a final state with a presence of a core hole and a photo-electron. This phenomenon results in
a sharp rise in the absorption intensity, called an absorption edge. XAS is the measurement of
the X-ray absorption coefficient µ(E) of a material as a function of energy. In case of X-ray
fluorescence mode utilized in these PhD research experiments, the absorption coefficient µ(E)
can be obtained through the expression:
µ(E) = 0I
I f,
where If is the intensity of the X-ray fluorescence lines and I0 is the intensity of incoming X-
rays.
4.3 Complementary characterization methods
For optical and structural characterization of the nanowires several standard techniques,
like photoluminescence (PL), cathodoluminescence (CL) and Raman spectroscopy were used.
The details regarding set-up parameters will be given in the related chapters. Moreover,
the morphology of the nanostructures was characterized by Carl Zeiss Gemini scanning electron
microscopy (SEM).
32
Chapter 5
Antisurfactant role of SiH4 during
vertical growth of GaN NWs
In this chapter the antisurfactant role of silane during vertical growth of GaN NWs will
be discussed. The rod morphology consequences will be described giving a growth model based
on the VLS growth approach. The supportive role of SiH4 injection is crucial for selective area
growth as well as for self-organized growth of GaN nanowires and therefore this aspect is
underlined as a separate chapter. Moreover, an explanation for an antisurfactant role of silane
was not known at the beginning of this project. The current state-of-the art shown in this chapter
is in line with the experiments performed in house and discussed in the subsequent chapters of
the thesis (i.e. inhomogeneous distribution of In along the NW with the accumulation of In
material in the top part of the rod).
To realize GaN NW-based devices, intentionally doped heterostructures are essential.
However, the doping itself may strongly affect the growth behaviour and modify
the morphology of GaN columns.
In 1998 Haffouz et al. [117], studied the effect of silicon doping on the lateral
overgrowth of GaN pyramid structures grown selectively using a SiNx mask on the sapphire
substrate. They showed experimentally that the vertical growth rate of GaN can be easily
increased by introducing a high Si concentration in the vapor phase. Small SiH4 flows applied
(0.88 nmol/min) resulted in the formation of pyramidal structures delimited with top (0001) c-
plane and {11̅01} r-side walls. Once the silane flow was increased to the higher value
(0.2 µmol/min), the GaN microstructure morphology changed. The pyramidal form
transformed into the columnar structures, delimited by vertical {11̅00} facets. The reason for
this is the high growth rate in the <0001> c-directions.
In past years, the focus on the GaN NWs research increased, and thus the efforts were
put to understand more deeply the role of silane during vertical growth of GaN NWs by
MOCVD. Waag et al. [118] discussed the influences of silane on the growth kinetics for SAG
of GaN microrods grown on patterned SiOx/sapphire templates. Two different growth
33
conditions supplying 16.5 nmol/min and 165 nmol/min of SiH4 were investigated. It was found
that the growth kinetics of GaN microcolums change due to the higher silane flow applied.
The lower SiH4 supply resulted in two growth regimes in terms of vertical and lateral growth.
There was a critical value found for height and diameter of the microrod (about 3.5 and 2.5 µm,
respectively for the apertures of 1.4 µm and pitch of 6 µm). The vertical growth rate reduces
from 16.9 to 7.2 µm/h from stage I to stage II (before and after reaching the critical values).
Similar tendency was found for the lateral growth rate. It was also reduced from 1.9 to
the constant value of 0.2 µm/h, after reaching the critical diameter. Interestingly, for the higher
silane flow applied, the growth kinetics changed, which was manifested by only one growth
regime. There was neither critical height nor diameter, which is a threshold for different lateral
and vertical growth rates. The lateral growth rate was suppressed once the columns reached
the diameter of about 2.2 µm. On the other hand, the deposition rate on the top surface of
the microrods, were strongly enhanced. The vertical growth rate increased to the value of
28.1 µm/h for the whole growth process.
Finally, the antisurfactant role of Si during the growth of GaN nano- and microrods was
comprehensively investigated by Tessarek et al. [119]. The experimental studies led to
the proposal of a growth model based on the vapor-liquid-solid mode. Once again, it was
shown, that silane supply during the GaN microrods synthesis promotes the vertical growth, as
can be seen in Fig. 5.1 (similar comparisons might be found also elsewhere [101]).
Figure 5.1: Comparison of GaN microrods grown on sapphire substrate by self-organized mode. Left image refers
to the process without silane support, whereas right one depicts GaN microrods grown with SiH4 injection [119].
The growth process without silane support results in the formation of irregular GaN
islands. Some of the GaN clusters coalesce and the aspect ratio of such structures is below 1.
Contrary, the growth process with SiH4 injection enables the formation of well aligned GaN
34
microrods, perpendicular to the sample surface. The grown structures have regular hexagonal
shape, smooth side walls and their height is up to 5 µm leading to the aspect ratio of 3.
The RIE etching experiment revealed that the m-plane side walls of these vertical microwires
were less severely etched in comparison to the top c-plane facets and the volume of
the structures. The spatially resolved energy dispersive X-ray spectroscopy showed that there
is an accumulation of Si and N at the m-plane facets of the rods and on the sapphire substrate,
which is an indication of SiN layer formation. The EDX measurements clearly proved that
the top facet and m-planes of the microrods are the regions of increased Si concentration as
compared to the bulk.
Afterwards, the microrods were covered by 3 InGaN/GaN core-shell MQWs stack.
Interestingly, the TEM showed that the intended 3-fold MQW stack was not realized. The first
GaN barrier as well as InGaN well was not visible in the HAADF-STEM image.
The explanation for it is the antisurfactant effect of a formed SiN layer [120]. At elevated
temperatures, the Ga adatoms mobility on the SiN is increased and no deposition takes place
on the m-plane microrods side walls. Similarly, at lower growth temperatures, the In mobility
on the SiN is also high due to low binding energy of In. Consequently, the first GaN barrier and
the first InGaN well is not grown. However, at lower temperatures, the Ga adatoms mobility is
reduced, which allows the coverage of SiN with GaN. Afterwards, further layers might be
deposited. These experimental results allowed proposing a model of NWs growth (Fig. 5.2),
based on the vapor-liquid-solid mode.
Figure 5.2: The schematic model of GaN NWs growth [119].
35
The proposed model is based on the VLS growth mode. Here, GaN NWs are grown by
self-catalyzing role of Ga-droplet. Similarly to the VLS model, Ga droplet acts as a sink for
the Ga atoms in the gas phase and on the surface. On the other hand, the low solubility of Si
and N in liquid Ga causes less pronounced adsorbance of these atoms in the Ga-homoparticle.
Thus, the concentration of Si and N is higher in comparison with Ga on the sidewalls of
the NWs and the surface than in the Ga-droplet. As a consequence, a SiN layer (or heavily Si
doped GaN layer) might be formed on the m-planes of the rods. Moreover the SiNx layer is also
formed on the surface. Additionally, the Ga droplet on top of the NW prevents the formation of
SiN on the c-plane due to low solubility of Si and N, which could suppress the vertical growth.
Ultimately, the Ga droplet is consumed during the final cool down step, because in presence of
ammonia it forms the GaN layer. Finally, the residual Si is forming the SiN layer, which is
covering the top c-plane facet of the rod.
36
Chapter 6
VLS Au-initiated growth of GaN NW
on Sapphire substrates
The vapor-liquid-solid (VLS) growth approach for a synthesis of nanostructures is still
the most commonly used method for a bottom-up NWs fabrication. The theoretical aspect of
VLS growth mode can be found in the chapter 3.2 of this thesis.
In this chapter, the VLS Au-initiated growth of GaN NWs on sapphire substrates will
be analysed and discussed. Moreover, new experimental results based on the synchrotron
characterization techniques will be studied to understand the atomic composition of coaxial
InGaN/GaN quantum wells in NW. The Au catalyst was chosen with respect to the potential
unintentional incorporation of catalyst metal into the GaN nanostructure during the growth. It
is expected that such an unwanted effect is lower for Au elements in comparison to more widely
used catalyst metals like Ni or Fe. The explanation is due to the higher substitution energy
required for Au to enter a Ga or N lattice during growth [121]. Several experiments were
performed to investigate and understand the Au droplet formation on sapphire as well as on
silicon substrates. Interestingly, the growth on Si wafers was not successful due to etching of
Si surface by AuGa eutectic. Thus, in this chapter only sapphire substrates are considered as
a template for Au-initiated VLS growth of GaN nanowires. Detailed investigation of GaN NW
growth on Si substrates is presented in the following Chapter 7 and Chapter 8.
In the first section the experimental procedure for GaN nanorods synthesis by MOCVD
is presented. A thin Au-film is utilized as a catalyst. After the heating up step, the thin Au film
is forming nano-droplets, which are afterwards used as a nucleation seeds for NW growth.
The growth conditions selection as well as their consequences for GaN NWs formation is
described. In the second section the detailed analysis of structural and optical properties of GaN
NW is discussed in order to qualify the NW as building blocks for nanoLED fabrication.
37
6.1 Experimental procedure for NWs growth and characterization techniques
The vapor-liquid-solid (VLS) Au catalyst initiated growth of GaN NWs was realized by
utilizing sapphire substrates coated ex-situ with an Au film of 1 nm nominal thickness. Several
experiments were performed to investigate the Au droplet formation on sapphire substrates.
The Au-film thickness as well as annealing temperature was studied. It was found that 1 nm of
a thin gold film is optimum value as it comes for subsequent nano-sized droplet formation. For
thicker films, the droplets tend to coalesce and as a result they form micro islands. The optimum
annealing temperature was found to be around 1020 °C. It eases the process sequence, since
the subsequent GaN deposition step might be performed at the same temperature. Thus,
an easier and faster final process is achievable. The lower annealing temperatures applied
resulted in not complete Au droplet formation since the Au thin film needs to melt first and then
from a nano-sized droplets.
The schematic of process sequence steps for Au-initiated VLS growth of GaN NWs is
depicted on Fig. 6.1.
Figure 6.1: Process sequence steps involved in the Au-initiated VLS growth of GaN NWs on sapphire substrates
by MOCVD.
38
At the beginning of the growth process, after reaching the growth temperature of around
1020 °C, only TEGa was applied for the first two minutes to allow Ga enrichment of the Au
catalyst. This predeposition step was followed by the simultaneous introduction of both TEGa
and ammonia supply. The low V/III ratio of around 3 promoted the vertical growth of the NWs.
The three pairs of core/shell GaN/InGaN MQWs were deposited at 730 °C under total reactor
pressure of 400 mbar. Finally, the cooling down was carried out under NH3 stabilization to
avoid excess N desorption. All of the growth parameters mentioned above are found to be
optimal for GaN NWs growth in the 3x2” MOCVD reactor. The parametrical studies were done
to investigate the influence of the process parameters on the morphology of the nanorods [94].
The applied V/III ratio very strongly affects the selective NWs growth. It was found that
only small V/III ratio of around 3 enables the vertical growth of non-tapered GaN nanocolums,
as can be seen on Fig. 6.2. The higher V/III ratio applied resulted in the formation of only small
amount of conical shaped nanostructures with Au droplet on top or even led to coalescence of
the nucleation seeds and formation of micro-islands. The explanation of this observation is
the V/III ratio influence on the adatoms surface diffusion lengths. For the lowest V/III ratios of
3, the surface diffusion lengths are longer, which reduces the nucleation on the surface or on
the side facets of the nanostructures and promotes the vertical growth of the NWs.
Figure 6.2: V/III ratio influence on the GaN NW morphology during the Au-initiated growth of GaN NWs.
The growth temperature was 870 °C and the total pressure applied was 100 mbar.
The total working pressure influences the morphology of GaN NW growth.
The diameter of the structure depends on the value of working pressure during the vertical
growth of nanorods. Interestingly, the droplet localized on the top c-plane facet of each
nanostructure is always smaller than dimension of the rod. The dependence of average NW
diameter as function of total working pressure is depicted on Fig. 6.3.
39
Figure 6.3: Average NW diameter (nm) as a function of total working pressure (mbar).
As can be seen from the graph above (Fig. 6.3), the average diameter of NWs increases
with increasing working pressure applied during the growth of the structures. Additionally,
the higher pressure promotes the formation of vertical columns. The big nucleation island
formation and lateral growth is suppressed whereas taller, but wider GaN nanocolumns are
grown. The growth conditions selected for InGaN/GaN MQW deposition was transferred from
the standard recipes dedicated for planar devices: 730 °C, 400 mbar. All of the grown NWs
exhibit the inclined top facets. It originates at the cooling down step, when the TEGa supply is
shut down, but NH3 is still open to stabilize the GaN surface and suppress the thermal
decomposition of GaN film. All of the wires have a droplet on the c-plane top facet, which act
as a reservoir for the vertical growth. The droplet consists not only with Au, but also with Ga.
Therefore, the Ga residuals are consumed during the cool down and they still contribute to
the vertical growth. The schematic draft of this observation is also shown on Fig. 6.1.
The optical, structural and elemental characteristics of single InGaN/GaN NWs grown
by MOCVD are characterized by means of different complementary non-destructive
techniques. In particular, the QWs emission is studied by photoluminescence (PL), the crystal
quality and the strain through Raman scattering, X-ray diffraction (XRD) and X-ray absorption
40
near edge structure (XANES) spectroscopy with hard X-ray nanoprobe and the elemental
distribution thanks to X-ray fluorescence (XRF) and energy dispersive spectroscopy (EDS).
6.2 Atomic composition of coaxial InGaN/GaN quantum wells in NWs
To examine the atomic composition along the NWs, several single NWs were scanned
through the X-ray nano-beam of the beamline ID22NI of ESRF. The GaN NWs were first
separated from the sapphire substrate and transferred onto the small SiNx membrane. Such
a template was afterwards loaded into the synchrotron probe. At the beginning of
the measurement process, the positioning of the sample needed to be performed in order to
localize the NW of choice on the SiNx membrane. Then, the shutter was opened and
characterization process started. The scanned region of interest was used to define distribution
of elements. The XRF spectrum of each pixel was fitted in order to build elemental distribution
maps with PyMca and therefore to analyse and classify the elements within the NW structure.
The XRF intensity maps showing the distribution of In, Ga and Au are shown on Fig. 6.4.
Figure 6.4: XRF intensity maps showing the distribution of Ga, In and Au along the GaN NW – a), b) and c),
respectively.
The Ga distribution is homogeneous along the c-axis of a GaN NW. On the other hand,
the inhomogeneous distribution of In shows that coaxial InGaN QWs form at the lateral surfaces
of the NW, but they are not completely formed at the base of the NW. Similar observations are
already reported and addressed to the antisurfactant role of silane applied during vertical growth
of nanorods (see Chapter 5). However, in presented case the SiH4 was not supplied. Hence,
the origin of inhomogeneous distribution of In and lack of this element at the base of the rod
must have another origin. Most probably it is related to the nature of VLS approach, where
the growth is driven by the nano-droplet localized on top of the structure. There is material
41
accumulation in the upper part of the structure, but not at the bottom of the rod. The root cause
of this effect shall be thus investigated.
Interestingly, Au catalyst was detected only on the top of the NWs, implying no
incorporation of Au along the NW or incorporation below detection limit. This observation is
very important in terms of future full assembled device, i.e. nanoLED. The unfavourable
incorporation of Au elements may strongly decrease the properties and functionality of
the device. However, presented results proof that Au-initiated VLS growth of GaN NWs may
be considered as an alternative for SAG or self-organized growth of GaN nanostructures.
For obtaining data with a higher spatial resolution (around 4 nm), EDS measurements
were performed. On top of the NWs a stronger signal from In was registered, which is attributed
to locally higher In distribution as can be seen at the XRF intensity map above. In Fig. 6.5
a representative In profile along the radius of the NW and the relative fit are presented.
Figure 6.5: EDS In profile perpendicular to the diameter of the NW. The inset shows a HR-TEM of a single
dispersed NW with the magnification of the bottom part of the same NW.
The QWs and the barriers are clearly observed with the regular intervals. According to
the “thin layer” approximation, the intensity of each element is proportional to the excited
volume and to its concentration. From this assumption, it is possible to build a fitting function
that models the emission and to obtain estimation for the radius of the GaN core, the thickness
of the QWs and the barriers as well as the In concentration. It was found that the radius is
42
approximately equal to 80 nm, the thickness of the QWs and the barrier are around 2 nm and
4 nm, respectively and the concentration of In is around 20 %.
To get further insight into the structural properties and information on the residual strain
in the NWs, nano-XRD mapping was performed. By extracting the values of the refraction
angles, it is possible to calculate the lattice parameters at a local level. Moreover, it is possible
to study the trend of the lattice parameter along the NWs and mapping the strain. Two
representative XRD peaks and the evolution of the lattice parameters along c-axis are presented
in Fig. 6.6 a) and b), respectively.
Figure 6.6: a) Representative (210) and (211) XRD peaks along with their best Gaussian fit. The peak position
and the FWHM are indicated for both fits. b) Evolution of the lattice parameters along the c-axis.
43
Only XRD peaks coming from the inner GaN core were found. The recorded reflections
match perfectly with the foreseen values for wurtzite GaN. The analysis of the XRD map for
different points of a single NW shows no changes in the crystal structure. Thus, one can
conclude that the inner core is free of strain. XANES measurements, which offer more details
on the structural microscopic order of structures also reveal no mixture of phases in the NW.
All of the studied NWs exhibit wurtzite crystal structure.
Figure 6.7 presents representative unpolarized Raman spectra of single NW.
The measured Raman modes of strain-free wurtzite GaN are found in the following
configuration: A1(TO), E1(TO), E2h and E1(LO). An additional peak centred at around
701 cm-1 can be attributed to the surface optical mode (SO). Such an observed Raman spectrum
indicates that reflections dominate over the size effects reported for thinner NWs [125]. There
is no clear enhancement of the Raman signal due to presence of the Au catalyst.
Figure 6.7: Representative unpolarized Raman spectra of a single NW taken with 514 nm laser line.
To study the QWs emission the LT-PL measurements were performed. Figure 6.8 shows
the representative LT-PL spectra of a single NW taken at 5 K. The signal was collected at
the top and bottom of the NW – black and red lines, respectively. The observed spectrum is
dominated by band to band transitions of the InGaN/GaN QWs with an energy between 3.19-
3.26 eV. The band centered on around 3.45 eV is attributed to the emission from the GaN. PL
spectra recorded at the top and bottom part of the single NW show an inhomogeneous
distribution of the InGaN and the In concentration along the core-shell structure.
44
Figure 6.8: Representative LT-PL spectra of a single NW taken at 5 K. The signal from top and bottom of the NW
– black and red lines, respectively.
6.3 Conclusions of VLS Au-initiated growth of GaN NW on Sapphire substrates
The study of Au-initiated GaN NW growth led to the realization of the core/shell GaN/InGaN
NW-based heterostructures. The investigated GaN NW exhibit very good crystal quality.
Understanding of Raman spectroscopy and nano-XRD results led to the conclusion that GaN
NW were grown free of strain. Additionally, based on the nano-XRD and XANES, there was
no cubic inclusions within the investigated structures. All of the studied NW exhibited wurtzite
crystal structures.
Presented GaN/InGaN heterostructure must be yet optimized to meet the expectation of
the nanoLED. The Au catalyst was detected only on top of the NW implying no Au
incorporation or incorporation below detection limit. However, the investigated samples
exhibited inhomogeneous In incorporation, which is unfavourable in terms of the nano-LED
functionality. The XRF mapping and LT-PL characterization techniques clearly revealed
locally higher In incorporation on top of the rods whereas the bottom parts of the structures
were In-free. Such inhomogeneous In incorporation leads to broad, non-uniform emission from
the active region of heterostructures and weakens the usability of the device. Ultimately, it was
not successful to achieve the target, which was the GaN rod on Si for LED application.
Therefore, the new growth method has to be developed to ensure the successful GaN-on-Si
integration for NW growth. In the following chapters the results on the selective area growth
(SAG) and self-organized growth of the GaN NW are presented and discussed – chapter 7 and
8, respectively.
45
Chapter 7
Selective Area Growth of GaN
microrods on Si(111) substrates
In this chapter, an investigation and understanding of the growth mechanism of GaN
microrods on SiNx/Si(111) patterned substrates by MOCVD as well as their structural
properties are presented. To avoid fluctuations in density and in the dimension of
nanostructures, which lead to significant dispersion in the optoelectronic properties of the NWs,
ordered arrays are grown by selective area growth (SAG) on pre-structured substrates.
The focus is demonstration of GaN nanostructure growth as a basic building block for
a complete LED.
In the first section the detailed experimental procedure of template preparation for GaN
microrods growth on Si(111) substrate is presented. The process steps for patterning
SiNx/Si(111) templates is proposed. Afterwards, the detailed controlled growth of GaN
microrods on Si(111) by MOCVD is discussed. Different mask parameters like opening shape,
size and spacing are studied in order to understand and optimize the growth process.
The morphology of the structures is described with giving an explanation for pyramidal tips of
the rods. The observation of the facet development and understanding the rod morphology leads
to the structure polarity determination.
The second section consists of theoretical explanation and understanding of optical and
structural properties of GaN microrods determined by photoluminescence and Raman
spectroscopy, respectively. The optical properties of two samples grown without and with silane
support are compared to understand the impact of the SiH4. The evolution of the Raman peak
frequency and FWHM was studied to understand the properties of the crystal structure along
the NW. Therefore, the strain aspect within the rod could be commented.
46
7.1 Experimental procedure
7.1.1 Template preparation for GaN-based microcolumn arrays growth
There are several important template parameters, which influences the growth of
the microcolumns by SAG (see chapter 3.3). The intuitive parameter is the diameter of
the aperture, which defined the size of the rod. However, the spacing between windows plays
a critical role as well. If the distance between openings is smaller than the diffusion length of
the growth species on the mask, than the growing rods compete over the species in
the overlapping areas. On the other hand, if the available substrate surface area per rod is larger
than the diffusion length, then the amount of material, which diffuses to a NW is saturating
[118]. Therefore, to study the growth mechanism of selectively grown microrods on Si(111)
substrates a special template was prepared. We designed a mask consisting of different shapes:
hexagonal, circle and square openings as well as different diameters of the windows, between
1 – 8 µm and separation distances of the openings between 2 – 16 µm. Figure 7.1 depicts
the utilized mask design. The target of the mask design is to investigate and clarify the rod
growth mechanism control for different mask parameters.
Figure 7.1: The utilized mask designed for SAG GaN microrods growth on Si(111) substrates. Red circled regions
showed the investigated mask units. The labels under each mask unit stand for: H – hexagonal, C – circle,
S – square opening; first number refers to window diameter, second number refers to spacing between two
neighboring windows. For example, Hx8x16 stands for hexagonal opening of 8 µm diameter and distances
between openings of 16 µm.
47
The most commonly used approach is to prepare first the buffer for subsequent NWs
growth. In terms of growth on sapphire wafers it is a GaN nucleation layer; in case of Si
substrates it might be AlN, AlN/AlGaN or AlN/AlGaN/GaN. Such a buffer template is
afterwards processed by structuring the desirable patterns. In presented experiments, we
decided to simplify the process and limit the necessary technology steps. Therefore, we decided
to structure directly the Si(111) wafers.
First, silicon wafers were cleaned in the buffered oxide etch (BOE) solution, rinsed by
a deionized water (DI) and dried on a hotplate. Then, a 80 nm SiNx mask layer was deposited
directly on Si(111) substrate by plasma-enhance chemical vapour deposition (PECVD) at
300 °C. The SiNx/Si(111) template was covered by Hexa Methyl Di Silazane (HMDS) adhesive
layer and then by AY 5214 E resist. The soft bake was performed at 90 °C for 2 min. After 15 s
of UV exposure, the samples were last for the N2 outdiffusion for 10 minutes. After a second
bake at 120 °C for 2 minutes, the development step was performed for 60 s. Finally, the resist
was removed by an etching step in CF4/O2 ambient. Figure 7.2 shows the process technology
steps for SAG template preparation.
Figure 7.2: The process steps for SiNx/Si(111) mask preparation for GaN microrods selective area growth.
48
7.1.2 The controlled SAG of GaN microrods on Si(111)
At the beginning of each process, an in-situ desorption step was applied. Si substrates
were deoxidized under H2 ambient at 975 ºC for 10 min.
Due to the strong tendency to form an eutectic system between Ga and Si at high
temperatures [52], the Si surface needs to be protected by an AlN buffer, unlike to the growth
on Al2O3 substrates. Therefore an AlN nucleation was performed to protect the silicon substrate
and also to create a seeding layer for the GaN rods. Nucleation was carried out for 150 s under
100 mbar at 1040 ºC using a V/III ratio of 210.
Subsequent growth of GaN microrods on the AlN-coated SiNx/Si template was carried
out under H2 ambient in two steps called afterwards: filling step and rod growth step. In the first
step, a GaN nucleation was performed with a low growth rate and TMGa flow of 21.3 µmol/min
(V/III ratio of 21) for 1000 s. In the second step of GaN growth, a higher growth rate – TMGa
flow of 78.6 µmol/min (V/III ratio of 851) was used. In order to enhance vertical growth and
the formation of m-plane GaN, SiH4 was introduced as an antisurfactant with the amount of
200 nmol/min. Both steps were carried out at 980 ºC, using a total reactor pressure of 100 mbar
for nucleation and 200 mbar for vertical rod growth.
In terms of conventional 2D MOCVD growth, high V/III ratios favour smooth 2D layer
deposition, whereas low V/III ratios lead to a 3D-like growth mode. Since the openings in
the mask are rather large (> 1 µm), the growth conditions chosen here are more like those
applied for GaN layer growth than for typical NWs. Here, we are rather filling the openings and
growing selective GaN columns than forming spontaneous 3D rods. Thus, a higher V/III ratio
in comparison to self-organized NW growth as described in literature [101] was used.
The increase of total pressure in the second part of the GaN rod growth also changes the growth
process. Higher pressure increases the decomposition rate of ammonia, subsequently increases
the effective V/III ratio and shortens the mean free path. Finally, higher V/III ratio and elevated
pressure are selected in order to achieve a higher deposition rate.
Figure 7.3: The scheme of SAG of GaN microrods on Si(111).
49
Figure 7.4: SEM image of SAG GaN rods grown on a SiNx/Si(111) template. The nominal diameter of the openings
was 8 µm and the distance between the openings was 16 µm. Samples were grown for 30 min, 1 h, 3 h – images:
a), b), c), respectively. (d) Sketch representing the orientation of m-plane GaN sidewalls of the microrod with
respect to Si(111) flat of the substrate. The nominal diameter of the openings was 4 µm and the distance between
the openings was 8 µm. Sample was grown for 1 h.
Figures 7.4 (a)-(c) show the 45 degrees tilt-view SEM images of GaN microrods grown
with silane support for various times of 30 min, 1 h and 3 h, respectively.
The performed time series reveals the growth mechanism of SAG GaN microrods on AlN-
coated SiNx/Si(111) templates. The SEM images captured at different stages of the growth
allow to study and understand the morphology development of GaN microrods. The formation
and evolution of planes can be followed thus revealing the information of the rod polarity. First
step of growth is filling the openings in the mask. GaN nuclei are formed on the thin AlN
nucleation layer. In Fig. 7.4 a), one can distinguish a GaN truncated pyramid and the initiation
of vertical growth. The pyramidal shape of the structures is attributed to the metal polarity of
the seed layer [62]. The subsequently grown layer inherits the polarity of the base layer. Thus,
Al-face polarity of the AlN top surface determines the Ga-polarity of the GaN layer.
Consequently, GaN pyramids are formed. After additional 30 min of growth – Fig. 7.4 b), GaN
50
continues to grow in the c-direction, increasing the height of the GaN m-plane side wall.
Additionally, six {101̅1} plane facets are found as a part of the truncated pyramidal structure.
Figure 7.4 c) shows the matrix of GaN microrods with pyramidal top after 3 h of material
deposition. The planes were identified by crystallographic angle detection. The last image –
Fig. 7.4 d), depicts the sketch representing the m-plane GaN orientation (red line) with respect
to the <110> direction (orange line). The hexagonal mask openings were aligned to the flat of
the silicon substrate, which in case of Si(111) is the <110> orientation. The microrod side wall
is rotated with respect to the hexagonal mask edge by 30°, which suggest that we observe m-
plane {101̅0} GaN. In case of a-plane {112̅0} GaN, side walls of the microstructure would have
been parallel to the <110> direction.
Adapted growth times are necessary for different opening diameters. Too long growth
time results in a loss of selectivity due to lateral growth and parasitic nucleation on the mask.
Neighbouring structures are starting to coalesce. Moreover, the morphology of structures
changes – growth continues in different directions resulting in inhomogeneously shaped
structures. In our case, the optimized time for microrod structures growth was found to be: 1 h
and 3 h for opening diameters of 4 µm and 8 µm, with the spacing of 8 and 16 µm, respectively.
The graph in Fig. 7.5 shows the measured height of the m-plane side facet and distance
of two opposite side facets of the same GaN microrod from Fig. 7.4 as functions of growth
time. One can observe that the structure development is not proportional in time. There is a very
pronounced lateral growth over the mask during nucleation increasing the diameter of
the microrod from 8 µm (nominal size of the openings) to more than 16 µm (after 30 min of
growth). The aspect ratios for a microwire grown for 3 h are: 0.34 and 1.38 considering height
of the m-plane and total height of the rod, respectively.
51
Figure 7.5: Graph showing height of m-plane side facets (black) and distance between two opposite side facets
(blue) of GaN microrods as functions of growth time. The nominal diameter of the opening was 8 µm and
the opening separation was 16 µm.
The structural geometry (pyramidal top of the GaN column) is similar to those grown
by other groups using MOCVD [126], [99], [97] as well as MBE [127]. Typically, self-
organized structures exhibit flat tops (no pyramidal tip) [101], [104]. The structure morphology
might be taken as an indicator of the polarity of the rods. Pyramidal tips are attributed to Ga-
polarity, whereas flat tops suggest N-polarity [62]. The polarity aspect is an important issue and
was studied by several groups [63], [64], [62], [65]. The structural morphology of
the microstructures tip is also attributed to the initial stage of the GaN growth. In our case,
pyramidal tip of the rods consists of six {101̅1} planes. The filling of the mask openings,
performed during the filling step of GaN growth, is critical in terms of top facet formation.
The GaN column shape depends on whether or not a complete filling process was carried out.
An incomplete filling process results in a governing pyramidal shape with a partial (0001) c-
plane. Vertical hexagonal rods with c-plane-dominant shape and partial {101̅1} semipolar
planes on top can only be realized with a complete filling process [99]. The root cause of
the semipolar planes dominance at the initial stage is originating from the slow growth rate of
the {101̅1} planes resulting from the hydrogen passivation effect [63].
A necessary condition to enhance vertical growth and formation of side facets is to
introduce SiH4 as a dopant [101], [117]. Silane stabilizes the m-plane GaN acting as
an antisurfactant. A SiN passivation layer on the sidewalls hinders the lateral growth and
enhances the mobility of the atoms. Therefore, vertical growth mode is promoted [119]. Figure
7.6 a) depicts the microstructures grown without silane support. Here, one can only observe
52
truncated pyramids without the vertical side facets. There are also some nucleation clusters
grown between the mask openings. Once the SiH4 flow is applied during the GaN growth stage,
vertical growth of microstructures occurs – Fig. 7.6 b). This observation clearly shows that SiH4
injection helps to assure desirable vertical rod formation. Moreover, beside good vertical
alignment and height uniformity, samples grown with silane support exhibit very good
selectivity between grown microrods and the SiNx-covered rodless region.
Figure 7.6: 45 degrees tilt view SEM image of SAG GaN rods grown on a SiNx/Si(111) template. The nominal
diameter of the openings was 2 µm and the distance between the openings was 4 µm. Samples were grown for 1 h
without (image a) and with silane support (image b).
7.2 Optical properties of GaN microcolums determined by
photoluminescence
The room temperature PL spectrum of a single GaN microrod measured at its pyramidal
apex is shown in Fig. 7.7. An intense near band edge Gaussian peak (centered at 3.44 eV and
with 125 meV full width at half maximum) is observed which can be attributed to band-to-band
and excitonic recombination. The defect band of GaN is also detected but with three orders of
magnitude weaker intensity than the near band edge emission. The absence of silane during
the growth not only inhibits the vertical growth but also causes a major drop in the number of
counts of the excitonic emission due to the deterioration of the crystal quality, as can be seen in
Fig. 7.7. The peak of the silane-free grown microrods is centered at 3.41 eV, it has a comparable
full width at half maximum (FWHM = 55 meV) and an integrated intensity of a factor of 300
lower than that of the silane-grown microrods. The redshift of the PL peak is consistent with
a lower free exciton contribution in the microrods grown without silane.
53
Figure 7.7: Room temperature PL spectra of single microrods grown with and without silane support. The spectra
correspond to single microrods of 8 µm nominal diameter and 16 µm distance between openings.
7.3 Structural properties of GaN microcolumns determined by Raman
Spectroscopy
The structural characterization of the rods was performed by Raman scattering
spectroscopy using a JY-T6400 Raman spectrometer equipped with a confocal microscope.
The room temperature spectra were measured with 514 nm laser excitations. For Raman
scattering investigations, an excitation area of 1 µm in diameter was achieved with
a microscope objective of times 100 magnification and 0.9 numerical aperture. The sample was
mounted on a XYZ motorized stage with a minimum step of 0.1 µm in the Z-direction.
The evolution of the crystal quality of SAG GaN microrods as a function of the growth
time was studied by Raman scattering spectroscopy. Figure 7.8 shows the Raman spectra
measured for the 3 individual microrods of 8 µm diameter from SAG samples with growth
times of 30 minutes, 1 h, and 3 h, respectively, grown with SiH4 support. All spectra show three
distinctive peaks attributed to E2h, E1(TO) and quasi LO (q-LO) phonons of wurtzite GaN.
An additional intense peak at 521 cm-1 is attributed to the Si substrate. Raman selection rules
54
only allow E2h and A1(LO) modes in backscattering from c-plane-terminated layers [128].
Therefore, the observation of peaks associated with phonons with E1 symmetry is due to
the refraction of the incoming excitation and the scattered light at the {101̅1} plane terminated
microrods. Fig. 7.8 shows this as the pyramidal tip of the microrods forms with increasing
growth time, the E1(TO) Raman peak increases its intensity and the A1(LO) and the E1(LO)
modes mix giving rise to the q-LO peak. We also observe the onset of a broader band between
the LO and the TO frequencies, compatible with the frequency of surface optical (SO) modes
[129]. The SO modes are confined at the surface and are activated by the breakdown of
the translational symmetry. Thus, the relative intensity of these modes can increase towards
the apex of the pyramid as the edges of the pyramid become closer and the surface-to-volume
ratio increases.
500 600 700
q-LO
Eh
2
E1(TO)
In
ten
sity (
arb
. u
nits)
Raman shift (cm-1)
30 min
1 h
3 h
Si
Figure 7.8: Comparison of Raman spectra measured for time series samples. Microrods were grown for 30 min,
1 h and 3 h. Each spectrum corresponds to a single microrod of 8 µm nominal diameter and 16 µm distance
between openings.
The frequency and FWHM of the E2h mode are good indicators for the assessment of
strain fields and the crystal quality of the material, since this mode, being non-polar, is not
sensitive to the presence of free carriers due to unintentional doping. The values obtained for
the microrods grown for 3 h with the nominal diameter of 8 µm reveal nearly strain-free
structures (central frequency of 567 ± 1 cm-1) of crystal quality (FWHM = 4.2 ± 1 cm-1)
comparable to strain-free bulk crystals [128]. Our confocal microscope allowed us to vary
the depth of the laser focus from the microrod apex (at z=0) deeper along its axis (z < 0).
55
The graphs from Fig. 7.9 (a) show the E2h Raman peaks measured for a microrod grown for 3 h
with nominal opening diameter of 8 µm. The intensity of the peak is proportional to the excited
volume and therefore increases towards the microrod half height. The evolution of the peak
frequency and FWHM are depicted in Fig. 7.9 (b) showing only a very small variation along
the full microrod length. Values of strain below -0.1% can be estimated from the shift between
the microrod and bulk Raman frequencies assuming a biaxial strain field (the growth direction
is free of strain) [128]. Thus, we can conclude that the microrods grow nearly free of strain
during most of the stages of the growth.
550 560 570 580 590
0 -5 -10 -15566.5
567.0
567.5
568.0
568.5
4.0
4.5
5.0
5.5
6.0
6.5
7.0
In
ten
sity (
arb
. u
nits)
Raman shift (cm-1)
z=0
z=-3 m
z=-5 m
z=-8 m
z=-11 m
z=-17 m
(a)
Eh
2(c
m-1)
z (m)
(b)
FW
HM
(cm
-1)
Figure 7.9: (a) Raman peak corresponding to the E2h phonon of an 8 µm diameter microrod measured with
the excitation light focused at different depths (z) along the microrod axis. The apex of the pyramidal tip of
the microrod defines z=0 and for increasing depth z<0. (b) Evolution of the E2h frequency and FWHM as
a function of the depth. The frequency of the bulk E2h is plotted as a dashed line [128].
56
7.4 Conclusions of Selective Area Growth
The SAG study of GaN NW led to the understanding of controlled growth of GaN
microrods on Si(111) substrates by MOCVD. The current state-of-the-art describes in detail
the synthesis of GaN NW on sapphire substrates, yet the understanding of GaN NW on Si is
missing. Presented and discussed results shed the new light on the GaN rod on Si as a building
block for nanoLED.
The designed mask allowed to investigate and clarify the rod growth mechanism.
The observation of facet development led to understanding the NW morphology, which is
attributed to two independent factors: GaN polarity and initial stage of the growth. Metal
polarity of the seeding layer results in a pyramidal top of the rods. Furthermore, six {101̅1}
planes, which the pyramid consists of, originate from the initial filling of the mask openings.
The investigated structures exhibit high-quality structural and optical properties. GaN
microrods were grown nearly strain-free with crystal quality comparable to that of bulk crystals.
Such good properties of the microrods prerequisite them to become building blocks for
nanoLED. The future investigation shall focus on the core/shell MQW realization and
optionally on the improvement of the aspect ratio of the structures. The smaller diameter of
the mask openings could be obtained by employing different structuration technique, i.e. e-
beam lithography.
57
Chapter 8
Growth of self-assembled GaN NW
on Si(111) substrates
Self-organized growth of GaN NWs is a bottom-up alternative approach for SAG. It offers
the process simplicity and does not require any additional ex-situ sample treatment. All of
the steps are performed in-situ in the MOCVD reactor chamber. Thus, easier and faster
synthesis of the nanostructures compared to SAG is possible.
In this chapter, a detailed investigation and new understanding of self-organized growth
mechanism of GaN nanostructures on Si(111) substrates by MOCVD is discussed. The first
section concerns the buffer on Si(111) preparation as a basis for GaN NW growth. Afterwards,
the reactor conditioning procedure to ensure reproducible starting point is described. Next,
the experiments regarding AlN polarity and its influence on GaN NW morphology are
investigated. The explanation of buffer polarity importance for subsequent GaN microrod
growth can be also found in chapter 3.1.2 of this thesis. In the second section of this chapter,
the optimization procedure for GaN nano- and microrod growth is proposed. The model for
optimization based on three steps. A density of nucleation, randomly shaped structures as well
as vertically aligned rods is studied as a function of the key process parameters: silane
deposition time, GaN growth temperature and silane injection time during the GaN nanocolumn
growth. The original developed approach allows obtaining extremely high vertical growth rate
of the GaN NWs, up to 300 µm/h. This value is double in comparison to current state-of-the-
art (see chapter 3.4). Finally, the effect of AlN susceptor coating on NW growth homogeneity
is stressed and explained. The strong AlN coating and the gas flow direction across the wafer
might cause the asymmetry in terms of polarity. Thus, the proper reactor conditioning is crucial
to ensure the reproducible starting point for growth experiments.
58
8.1 AlN buffer on Si(111) as a basis for GaN NW growth
8.1.1 Impact of asynchronous introduction of precursors on the AlN polarity
The subsequently grown layer inherits the polarity of the base layer. Thus, AlN buffer
polarity plays an important role in terms of nanowire growth. For instance, Al-face polarity of
the AlN top surface determines the Ga-polarity of the GaN layer. Additionally, the polarity and
structure morphology thought to be linked. In principle, pyramidal tips of the rods are attributed
to a Ga-polarity, whereas flat tops suggest N-polarity [62]. In other words, N-polarity favours
vertical growth of the NWs, on the contrary to metal-polarity, which suppresses vertical growth
and favours pyramidal shape of the grown structures. To avoid formation of the GaN pyramidal
structures, N-polar buffer is necessary.
The silicon surface modification, by introducing ammonia or TMAl before AlN
deposition, may change the polarity of the buffer and subsequently the nanowire growth might
be controlled.
To determine the polarity of the AlN samples, KOH etching experiments were
performed. The epitaxial growth conditions were kept the same for AlN buffer as for GaN
nanowires growth. The predose time before AlN deposition was varied for NH3 and TMAl.
After 300 s of nucleation, 200 nm of high temperature AlN layer was deposited. Samples were
dipped into the 10% KOH solution at 30 ºC for varied time between 1 min and 60 min.
The morphological characterization of AlN surfaces was performed by means of atomic force
microscope (AFM).
The surface morphology of the nanowire samples was characterized by a Carl Zeiss
Gemini scanning electron microscope (SEM).
Figure 8.1 shows the AFM images of the AlN samples after KOH etching experiments.
Figure 8.1: AFM images of the AlN samples after 60 min of etching in 10% KOH at 30 °C. Left image refers to
5 s of TMAl predose, right image to 5 s of NH3 predose.
59
N-polar layers are etched by the KOH solution, whereas metal polar ones are etching
resistant [69]. Comparing the AFM images of investigated AlN buffer, one can distinguish post-
etch rough surface for both predoses applied. The density of small V-pits is much higher for
the sample with NH3 predosing than this observed for its counterpart with TMAl predosing.
There, the bigger V-pits and flat post-etch surface regions are typical features indicating
the dominant metal polarity of the sample. We conclude that both presented AlN buffers very
likely exhibit mix-polarity, but the area of N-polarity is much more dominating in the sample
grown with ammonia preflow. The initial TMAl preflow acts as a surface stabilizer resulting in
more flat substrate. Opposite, initial NH3 preflow strengthens the etching effect resulting in
rougher N-polar substrate.
8.2 Investigation of growth parameters on the GaN NW growth and morphology
8.2.1 Substrate preparation
In order to prepare the substrate prior the growth, a high temperature treatment is
necessary. An in-situ annealing step under hydrogen at elevated temperature ensures
the deoxidation (removal of a thin native oxide layer formed onto the Si(111) surface) and
cleaning process (removal of moisture and organic materials). This step is enough to allow
a straightforward surface preparation before the growth process. The substrate does not have to
be chemically cleaned ex-situ before it is introduced into the reactor.
In our case, we cleaned the Si(111) wafers under H2 ambient for 10 minutes at 940 °C
using 50 mbar of total reactor pressure, injecting 670 nmol/min of silane to stabilize silicon
surface. The subsequent experiments showed no difference between samples grown with and
without SiH4 injection in the cleaning procedure, thus silane injection step can be omitted.
8.2.2 Impact of NH3/TMAl predose before AlN deposition on Si(111) on the NWs growth
To verify the influence of the AlN buffer polarity on the nanowire growth, two GaN
nanowire samples were grown on different AlN buffers. In the first growth we used 5 s of TMAl
predose before AlN deposition and in the second one we change the preflow to 5 s of ammonia.
Afterwards, we deposited SiNx for 250 s, performed GaN nucleation at around 1000 °C and
finally grew GaN nanowires for 480 s constantly supplying 500 nmol/min of SiH4.
The explanation of the importance of SiNx deposition before GaN growth and SiH4 support on
the vertical NWs growth can be found in the Chapter 3.4 and Chapter 5 of this thesis. Figure
60
8.2 shows the 45 degree tilt-view SEM images of nanowiring samples grown on two different
AlN polarity buffers.
Figure 8.2: 45°-tilted SEM view of self-assembled GaN nanowires grown on Si(111) substrate using different
polarity AlN buffers. The preflow applied before AlN deposition was 5 s of TMAl or NH3 for sample a) and b),
respectively.
The impact of the AlN buffer polarity on the GaN nanowire growth is very strong. On
the left image, attributed to metal-polar buffer, no nanowires were found. The only structures
grown on the SiNx/AlN template are not activated GaN nucleation sites and some bigger GaN
pyramidal structures with very high density. This result validates the statement of vertical
growth suppression on the metal-polar buffers. On the contrary, SEM image on the right shows
vertical GaN nanowire formation. Besides vertical and tilted wires, some not activated
nucleation sites are also found, but with much less density in comparison to the left image.
Thus, the polarity of the buffer layer strongly influences the polarity of the GaN nucleation seed
and has a major impact on the rod growth mode and geometry. If the seed is metal-polar, the 3D
rod growth is not activated and the vertical growth is suppressed. Contrarily, if the seed is N-
polar or at least of mixed-polarity, the nanowire growth mode is accessible.
8.2.3 Proposed optimization model – nanostructures density as a function of the key
process parameters
The process window was defined by studying the influence of the key growth
parameters: silane deposition time, GaN growth temperature and silane injection time during
the GaN nanocolumn growth. The impact of these parameters on the morphology of
the nanostructures was determined. We defined several possible shapes of the nanostructures
which we observed during our studies. To assess the impact of the growth conditions we have
61
investigated the density of the various structures as function of the key growth parameters.
Figure 8.3 shows the 45 degree tilt-view SEM image of GaN nanowires with definition of
the studied structures. We considered not activated GaN nucleation seeds (blue), random
structures (orange) as well as three types of the rods: vertical (red), tilted (green) and
multicolumnar (yellow).
Figure 8.3: 45°-tilted SEM view of self-assembled GaN nanowires grown on Si(111) substrate showing
the description of the measured structures used to determine the density: not activated GaN nucleation seeds
(blue), random structures (orange), vertical wires (red), tilted wires (green), and multicolumnar wires (yellow).
For statistics we investigated always three to eight SEM images. The inspection area
was varied depending on the size of the structures and the surface coverage ratio. For
the nucleation sites the approximate investigated area was from 3x102 up to 1x105 µm2, and for
the other types of the structures the range was from 2x104 µm2 and 4 mm2. For each selected
SEM image, described structures were counted. The final density value is an average number
from three to eight inspected images. In the following discussion sections we present the graphs
showing nucleation density (not activated GaN seeds), all three types of the rods densities
(vertical, tilted and multi-columnar) as well as total structure density, which is the sum of the all
rods and the randomly shaped structures density.
We propose to optimize the self-organize growth of GaN NWs on Si(111) substrates
based on three steps. In the first step, we investigate the formation of GaN nucleation to
eliminate the unfavourable not activated nucleation sites. Ideally, all of the GaN pedestals shall
transform into a GaN nanocolumn. Thus, the density of not activated nucleation sites shall be
as small as possible. Afterwards, we study the density of randomly shaped nanostructures
(the difference between total structures density and sum of the rods density). The goal is set to
62
maximize the column-like nanostructure shape. Finally, all three types of rods: vertical, tilted
and multicolumnar are considered. Ideally, all of the wires shall be well vertically aligned to
the surface. Therefore, density of such type of the rods shall be maximized with minimized
density of other types.
In the following discussion and analysis, the results and conclusions regarding
the demonstration of important impact of SiNx deposition time are made by assuming that there
was actually deposition of SiNx by supplying SiH4 during the process step after AlN buffer
deposition. This step was based on an approach proposed by Koester et al. [101]. It is also
known that besides SiNx deposition, silane (with presence of ammonia) may also lead to etching
the surface without necessarily forming SiNx.
8.2.3.1 Impact of SiNx in-situ masking layer deposition time on the NW density
The deposition of SiNx layer is a necessary step in self-organized growth mode approach
[101]. The SiNx forms a not closed, very thin layer, which acts as an in-situ mask, providing
selectivity for GaN nanocolumns formation. GaN grows in the very small openings, leaving
SiNx islands as the spacers between the neighbouring rods.
In our experiment we have investigated the influence of SiNx deposition time on the wire
formation. In-situ masking layer was grown for varied time between 0 and 600 s. The other key
parameters were set as follow: 5 s of NH3 predose before AlN buffer deposition and 480 s of
GaN nanocolumn growth at around 1000 °C using constant SiH4 flow of 400 nmol/min. Figure
8.4 depicts the graphs representing the density of the structures determined by SEM images.
63
Figure 8.4: Density of the structures as a function of SiNx deposition time. Graph a) shows density of total
structures (black), sum of the rods (red) and not activated nucleation sites (blue). Total structure density stays for
the sum of the rods density and density of the randomly shaped structures. Graph b) depicts the density of vertical
(black), titled (red) and multicolumnar (blue) wires.
64
As can be seen on Graph a) from Fig. 8.4, longer SiNx deposition time results in decrease
of the unfavorable not activated nucleation density. For a short deposition times – shorter than
250 s, the nucleation density is extremely high, but afterwards for times larger than 250 s, it
reaches the levels of less than 1.8x103/mm2. A good difference between total structure density
and sum of the rods density is observed for two values of SiNx deposition time: 250 and 500 s.
The Graph b) from Fig. 8.4 gives information regarding types of the rods. Taking two data
points, selected before, into consideration we can comment on the dominant morphology of
the rods. For the longer SiNx deposition time of 500 s, the density of vertical and tilted wires is
decreasing, and the formation of multicolumar structures is enhanced. Since we would like to
have as much vertical wires as possible, we select 250 s as an optimum deposition time for SiNx
in-situ masking layer. We observe good density of vertical nanowires (2.0x102/mm2) with
negligible density of multicolumnar nanostructures.
8.2.3.2 Impact of growth temperature on the NW density
The growth temperature has a very strong influence on the MOCVD processes. In our
experiment we have investigated the influence of different GaN growth temperatures, between
980 and 1060 °C, on the wire formation. The other conditions were kept to the standard values:
5 s of NH3 predose before AlN buffer deposition, 250 s of SiNx in-situ masking layer and 480 s
of GaN nanocolumn growth using constant SiH4 flow of 400 nmol/min. Figure 8.5 depicts
the graphs representing the density of the structures determined by SEM images.
65
Figure 8.5: Density of the structures as a function of GaN growth temperature. Graph a) shows density of total
structures (black), sum of the rods (red) and not activated nucleation sites (blue). Total structure density stays for
the sum of the rods density and density of the randomly shaped structures. Graph b) depicts the density of vertical
(black), titled (red) and multicolumnar (blue) wires.
66
Based on Graph a) from Fig. 8.5, one can observe that the lower the growth temperature
the higher the density of not activated nucleation. It is decreasing by one order of magnitude
from the value of 9x103/mm2 for 980 °C to 5x102/mm2 for 1020 °C. It is also a general trend
that lower GaN growth temperature resulted in higher density of tilted single columns and
multicolumnar structure formation. The diameters of the nanocolumns are not homogeneous.
They increase from the top of the thick nucleation seed to the top of the column. Similar
observation was already reported by Tessarek [119] and Koester [101].
At higher temperature decomposition of NH3 is more efficient and the surface diffusion
length of Ga adatoms is increased. Consequently, larger nucleation seeds with lower density
are formed and lead to larger height and diameter of the columns. The density of randomly
shaped nanostructures is also decreasing towards elevated temperatures. Considering types of
the rods – Graph b) from Fig. 8.5 – it is worth to mention that for the highest temperatures
applied (above 1040 °C) we observed only multicolumnar formation. At the lowest GaN growth
temperature of 980 °C we observed an interesting phenomenon. Contrary to the general trend,
the density of vertical nanowires and titled single columns is smaller in comparison to the higher
growth temperatures applied. This fact might be explained by the change of dominant
nanostructure morphology. At the lowest temperature the density of nucleation sites reaches
the highest values. Moreover, the density of random structures has also bigger value than those
determined for elevated growth temperatures. It means that increasing the deposition
temperature induces the increase of the activated nucleation seeds. These activated nucleation
sites are afterwards transformed into the rod structures. The GaN growth temperature of
1020 °C is the optimum in terms of a nucleation density, a number of randomly shaped
structures as well as a good compromise between vertical and tilted fraction of the nanowires.
8.2.3.3 Impact of silane injection time on the NW density
The vertical growth of GaN nanowires was performed in two steps: first silane was
introduced as an antisurfactant to initiate vertical growth. Afterwards, SiH4 supply was shut
down, but rods growth continued. The series of different times of silane support were studied,
keeping the total growth time of GaN NWs constant and set to 960 s. The other process
conditions were kept to the standard values mentioned before (see Impact of growth
temperature). Figure 8.6 depicts the graphs representing the density of the structures determined
by SEM images.
67
Figure 8.6: Density of the structures as a function of silane injection time ratio (time with SiH4 supply divided by
total GaN NW growth time). Graph a) shows density of total structures (black), sum of the rods (red) and not
activated nucleation sites (blue). Total structure density stays for the sum of the rods density and density of
the randomly shaped structures. Graph b) depicts the density of vertical (black), titled (red) and multicolumnar
(blue) wires.
68
On the Graph a), there are two points – 25% and 50% of silane injection time ratio,
where the density of not activated nucleation sites reaches minimum values. Moreover, for
these two points also the difference between total structure and sum of the rod density is
minimized. Considering types of the rods on the Graph b) from Fig 8.6, one can observe
the local minimum of the density of all types of the wires for 50% time ratio. Thus, the
concluded optimized time of silane injection is 240 s (for 960s of total GaN wire growth time).
8.3 Optimized growth conditions for GaN NW growth on Si(111)
In previous sections we considered four optimization steps: predose before AlN buffer
growth, SiNx deposition time, GaN growth temperature and silane injection time.
The optimized parameters based on understanding and proposed model for self-organized
nanowire growth were following: 5 s of ammonia preflow before AlN buffer, 250 s of SiNx in-
situ masking layer deposition, 1020 °C GaN growth temperature and 240 s of silane injection
time. Figure 8.7 depicts the 45 degree tilt-view SEM image of the sample grown under
optimized conditions.
Figure 8.7: 45°-tilted SEM view of self-assembled GaN nanowires grown on Si(111) substrate using optimized
growth conditions.
Most of the NWs are vertically aligned to the substrates. Due to the suppression of
multicolumnar structure formation, most of the rods are single, free standing columns.
Moreover, the unfavourable not activated nucleation seed density is limited.
69
8.4 Effect of AlN susceptor coating on NW growth homogeneity
In order to ensure the reproducible starting conditions for growth experiments, the high
temperature reactor cleaning was always applied prior to a growth process. Moreover, all
process parameter variations were carried out in mixed order to exclude the influence of
a potential monotonous drift of reactor conditions on the results. Even then, the highly coated
susceptor may strongly affect the uniform and homogeneous deposition of AlN with desired
polarity. Due to the AlN coating and the gas flow direction across the wafer towards the flat,
sample exhibits the asymmetry in terms of the polarity. The centre part of the wafer exhibits
more metal-dominant character in comparison to the anti-flat region. Comparable observation
was reported by Behmenburg [91] for the growth of AlN on sapphire substrates. Similar
behaviour reported for both Si and sapphire substrates validate the assumption that material is
unintentionally decomposed from the susceptor surface and transported onto the wafer during
initial stages of growth causing Al-polar growth.
8.5 Conclusions of self-assembled growth of GaN NW on Si(111) substrates
A novel approach for self-organised growth of GaN NW on Si(111) substrates by MOCVD was
presented and discussed. The silicon surface modification by introducing NH3 before AlN
deposition allowed to develop a functional buffer for GaN NW growth on Si(111). The 5 s
preflow of ammonia resulted in mixed polarity AlN buffer, which enhanced the vertical growth
of microrods. Opposite, the TMAl preflow stabilized silicon surface resulting in rather metal
polar buffer and thus non activated GaN nucleation sited were grown instead of vertical rods.
New understanding of the growth process and transfer of existing knowledge from sapphire to
silicon substrate led to successful growth of the GaN nanostructures on Si(111). In order to
optimize the growth recipe the model was built based on the GaN NW density as a function of
key process parameters. Impact of SiNx in-situ masking layer deposition time, GaN growth
temperature and silane injection time on the GaN NW density was analyzed and discussed.
Based on the density trend observations of the non-activated nucleation sites, tilted and vertical
wires as well as randomly shaped structures the optimized process window for GaN NW growth
on Si(111) by MOCVD was selected. The reference sample grown under optimized conditions
contained of GaN NW vertically aligned to the substrate. The multicolumnar structure
formation as well as the unfavourable not-activated nucleation seeds were suppressed.
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Chapter 9
Optical and structural properties of
self-organized GaN NW on Si(111)
substrates
In this chapter the optical and structural characterization of InGaN/GaN core/shell
nanowires grown on Si(111) substrate is discussed.
The self-organized GaN NWs were grown on Si(111) substrates using AlN buffers and
in-situ SiN masking layers. The growth conditions were optimized to achieve maximum density
of vertical GaN microrods perpendicularly aligned to the substrate (see Chapter 8). GaN NWs
were grown on the SiNx/AlN/Si(111) buffer using two steps vertical growth. First, SiH4 supply
was opened for 240 s and afterwards nanorods continued to grow for subsequent 720 s without
silane support.
The second set of samples includes MQW incorporation. Here, the vertical nanorods
were grown continuously for 960 s with silane support. Once the growth step of GaN NWs was
accomplished, the carrier gas was switched from H2 to N2 in order to deposite InGaN MQW.
The H2 ambient is beneficial for vertical growth support (N2 enhances the lateral growth rate),
but in case of MQW it is destructive due to the etching of In-reach layers by H2. Moreover,
the total reactor pressure was decreased from 800 mbar, which enhanced the vertical NW
growth, to 400 mbar which is a standard value for the MQW deposition in lateral devices.
Afterwards, three different samples consisting of 3 pairs of core/shell InGaN/GaN MQWs were
deposited at 745 °C with varied TMIn flow of 2.6, 5.2 and 10.4 µmol/min.
9.1 Structural properties and In incorporation in InGaN/GaN nanowires by
µPhotoluminescence
A typical SEM image of self-organized GaN NWs grown on Si(111) substrate by
MOCVD is shown in Fig. 9.1. Most of the structures are well organized wires, vertically aligned
to the sample surface. There are also some tilted nanocolumns and parasitic not-activated GaN
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nucleation – small pyramids found in between the NWs, but with much lower density in
comparison to the vertical rods. We believe that those small GaN pyramids are attributed to
the mix-polarity of the AlN buffer (see chapter 3.1.2). The growth of microrods is
inhomogeneous – the diameter and height of the structures varies. The smallest NWs are about
a few µm in height and less than 1 µm in diameter, but the biggest are up to 80 µm in height
and more than 5 µm in diameter. All of the wires, irrespectively to the orientation, exhibit
similar shape. They have a broader hexagonal base, which is a starting point for the vertical
growth. In most cases, the diameter of the rods is increasing towards the top of the structure.
Similar observations were already reported [101], [119]. Interestingly, besides typical
hexagonal columns, the non-hexagonal microrods growth is also observed (Fig. 9.1 c)).
Figure 9.1: 45° tilted SEM image of GaN nanowires grown on Si(111) substrate. a) GaN rods on Si(111) substrate.
b) GaN nanocolumns with 3 pairs of InGaN/GaN core/shell MQWs grown on Si(111) substrate. c) Non-hexagonal
microrod.
Figure 9.2 shows the PL spectra measured from the GaN nanostructures grown without
MQW incorporation. Left graph is a summary of all three types of structures studied. There are
standing and laying NWs as well as unfavorable not-activated GaN nucleation sites (small
pyramids). Right graph shows only well aligned vertical nanocolumns perpendicular to
the surface.
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Figure 9.2: Low-temperature (10K) photoluminescence spectra of GaN nanostructures. Left graph a) shows
a summary of measured structures: standing and lying NWs as well as parasitic not-activated nucleation. Right
graph b) shows only the spectra measured for well aligned vertical NWs.
All standing NWs show similar spectra in terms of spectral shape. Depending on
the NW, the emission is centred between 3.478 and 3.488 eV. At low temperature (10 K),
the emission is dominated by the recombination of donor bound excitons and is centred at
3.471 eV. This means that investigated NWs are compressively strained. The emission from
the parasitic islands (not activated GaN nucleation) is easy to recognize as it is centred at
3.37 eV. Such a redshift may be due to the fact that this material is full of stacking faults.
Coming to the shoulder at 3.41-3.43 eV, it could be also related to the stacking faults (emission
at 3.41-3.42 eV in strain-free GaN). The lower energy shoulders (below 3.40 eV) can be related
to LO-phonon replica or to donor-acceptor pairs, since the Mg-doping was used a few growths
runs before growing the presented samples. The NWs lying on the substrate show a spectrum
slightly broader than the standing ones. This might come from the strain due to some bending
of the NWs. In addition, these NWs look more intense, which is a consequence of the antenna
effect [130]. The antenna behaviour originates from small size and geometry of
the nanostructures. The NWs appeal to act like a classical antenna – they are essentially
responding only to radiation polarized parallel to the antenna (NW) axis [131].
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9.2 InGaN distribution in GaN/InGaN core/shell heterostructures by nano-scale
Cathodoluminescence mapping
In this paragraph, the optical properties and InGaN distribution in GaN/InGaN
heterostructures will be discussed. The set of three samples grown with different TMIn flow is
studied by nano-scale cathodoluminescence mapping. Sample A refers to 2.6, Sample B to 5.2
and Sample C to 10.4 µmol/min of TMIn molar flow.
The InGaN/GaN samples exhibit similar morphology to the pure GaN rods: there are
standing and lying wires as well as parasitic islands. Only standing core/shell heterostructures
were measured by µPL. To avoid collecting PL from the parasitic islands, the set-up was aligned
so that the laser was focused on the top of the wires. For statistical reasons 6 NWs per sample
were measured. Figure 9.3 depicts the PL and CL peak positions, obtained from the InGaN/GaN
nanocolumns. Both PL and CL measurements reveals the same trend. There is a clear red-shift
of InGaN emission due to higher TMIn-flux. The difference of the peak position for µPL and
CL characterization originates from the different excitation energies used for
the measurements.
Figure 9.3: PL (square) and CL (triangle) wavelength as function of TMIn supply of InGaN/GaN core-shell
heterostructures.
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9.2.1 Nano-scale cathodoluminescence mapping of Sample A
A representative GaN microrod is depicted on Fig 9.4 a). It is about 30 µm high and has
a diameter of about 6 µm in the top part of the microwire, giving an aspect ratio of 5. The flat
top of the microstructure suggests the N-polarity of the rod. However, one can also observe
small pyramid grown on the c-plane of the microrod. Such a feature might be considered as
a polarity inversion domain (see chapter 3.1.2). Therefore, most probably the studied
microstructure is of mix-polarity with dominant N-polar domain.
The integral intensity of the recorded cathodoluminescence signal is shown in
Fig 9.4 b). One can observe the locally high CL intensity at the edges of the microwire.
Interestingly, there is a striation-like CL contrast in the bottom part of the microrod. Such
a pattern might be related to the stacking faults or some cubic inclusions in the material.
Moreover, locally high CL intensity is also observed for the small parasitic microstructures in
between the big vertical microrods.
The wavelength image – Fig. 9.4 c) and wavelength maps – Fig. 9.4 d) are very useful
in terms of In incorporation investigations. Considering the CL-linescan along the wire we
observed a blue-shift from 436 to 375 nm of the InGaN MQW emission. The highest CL
intensity coming from the GaN near band edge (NBE), with a broad FWHM, originates from
the bottom part of the microrod. The intensity modulation of GaN NBE along
the microstructure and the onset of InGaN emission in the upper part of the microrod are clearly
visible in the wavelength maps – Fig. 9.4 d). The broad InGaN MQW emission is centered at
405 nm with the highest emission intensity in the very top of the microrod. The yellow
luminescence is situated at 550 nm.
75
Figure 9.4: a) SEM image of a single GaN nanowire with InGaN/GaN MQWs. b-c) SEM-CL mapping. d) CL-
linescan along the InGaN/GaN microrod. Two micrographs from image d) refer to the same GaN microrod.
The position on both linescan maps was measured from 0-35 µm. The MQWs were deposited with the lowest
investigated TMIn flow of 2.6 µmol/min.
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9.2.2 Nano-scale cathodoluminescence mapping of Sample B
9.2.2.1 Non-hexagonal microrod
On the sample grown using 5.2 µmol/min TMin flow, besides typical microrods
exhibiting hexagonal symmetry, some non-hexagonal structures were found. A SEM image
with an example of such a microwire is shown of Fig. 9.5 a). Studied microrod is about 70 µm
high and has a diameter of about 6.5µm in the top part giving an aspect ratio of 10.8. There is
a typical hexagonal pedestal, acting as a base for microrod growth. However, after several
micrometers the cross section of the structure is changing from hexagonal to the rectangle. Such
a transformation might originate from the high amount of defects, which are pronounced as
a striation like pattern on the SEM image. Furthermore, this rectangle symmetry might be
a consequence of growth rate suppression of the two opposite m-planes. The detailed
investigation of this structure shall be conduct by subsequent TEM characterization.
The integral intensity, wavelength image and the CL spectra coming from the whole
microstructure are depicted on the images b), c) and d) of Fig 9.5, respectively. Considering
the CL wavelength image – Fig. 9.5 d) one can distinguish two distinct wavelength regions.
The GaN NBE dominates the bottom part of the microrod, whereas the upper part is attributed
to the InGaN MQW. The yellow luminescence peak from the microrods is situated at 550 nm.
The broad InGaN emission band has the highest intensity at 431 nm. The onset of InGaN
emission from the upper part of microrod is also visible from the CL map – Fig. 9.5 e). There
is a blueshift of InGaN CL from 460 nm to 428 nm. The clear dominant InGaN emission from
the nanowire top can be also observed from the wavelength image of the upper part of the rod
and CL spectra from this region - Fig. 9.5 f) and g), respectively.
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Figure 9.5: a) SEM image of the non-hexagonal GaN nanowire with InGaN/GaN MQWs. b-d) SEM-CL mapping
e) CL-linescan along the InGaN/GaN microrod. e) CL spectrum of the studied microrod. f) SEM-CL mapping from
upper part of microrod. g) CL spectrum of upper part of the microrod. The MQWs were deposited with the middle
investigated TMIn flow of 5.2 µmol/min.
9.2.2.2 Typical hexagonal microrod
A typical SEM image of a representative GaN microrod, exhibiting typical hexagonal
symmetry, is shown on Fig. 9.6 a). The flat top of the microwire suggest N-polarity of
the structure. The CL wavelength images from the upper part of the microrod, CL spectra and
CL map from investigated region are depicted on Fig. 9.6 b-d), e) and f), respectively.
The CL characteristics of the typical hexagonal micowire are similar to these obtained
for a non-hexagonal structure. CL wavelength image clearly reveals a blue shift of the InGaN
78
MQW emission along the microrod in the upper part. Interestingly, a slight additional blue shift
of InGaN emission can be distinguished from the c-plane top facet in comparison to
the microrods side-walls.
79
Figure 9.6: a) SEM image of the upper part of an individual hexagonal GaN nanowire with InGaN/GaN MQWs.
b-d) SEM-CL mappings in different spectral regions, e) CL spectrum of the studied microrod. f) CL linescan along
microrod. The MQWs were deposited with medium TMIn flow of 5.2 µmol/min.
9.2.3 Nano-scale cathodoluminescence mapping of Sample C
The investigated sample with MQWs deposited using the highest TMIn flow of
10.4 µmol/min exhibits very similar properties to the previous structures. The SEM image of
the representative microstructure is shown on Fig. 9.7 a). Studied microrod is about 75 µm high
and has a diameter of 5 µm in the top part giving an aspect ratio of 15. The structure morphology
duplicates previously studied microrods. There is a hexagonal pedestal at the base of
the microwire, the diameter of the rod is increasing towards the top and some striation like
patterns are visible on the side walls.
The integral intensity characteristics from Fig. 9.7 b) once again show locally higher CL
intensity signal at the edges and the base part of the rod, as well as increased CL signal detected
from the GaN parasitic islands, located around the big microwires.
On the wavelength image – Fig. 9.7 c), one can observe two distinct wavelength regions.
First one is the bottom part of the microrod with the highest GaN NBE contribution. The InGaN
MQW emission, centered at 498 nm, originates from the upper part of the structure. The most
intense emission is at the very top of the microrod. The intensity modulation of GaN NBE along
the microrod as well as the onset of InGaN MQW emission can be seen in the Cl-linescans in
Fig. 9.7 d). The InGaN coverage of the sidewalls of the microrod is the biggest in comparison
to the previously studied samples with the lower TMIn flow. Here, the InGaN clusters are
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covering about 80% of the m-planes, whereas in the previous set of the samples the coverage
ratio was found to be about 40% (In incorporation only on the uppermost part of the rod).
Figure 9.7: a) SEM image of the single GaN nanowire with InGaN/GaN MQWs. b-c) SEM-CL mapping. d) CL-
linescan along the InGaN/GaN microrod. Two micrographs from image d) refer to the same GaN microrod.
The position on both linescan maps was measured from 0-35 µm. The MQWs were deposited with the highest
investigated TMIn flow of 10.4 µmol/min.
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Comparison of all three investigated samples, grown with different TMIn flow, lead to
the following observations:
the µPL and SEM-CL measurements clearly show a red shift of the InGaN
MQW emission of the three samples with an increasing TMIn flow due to
the higher indium incorporation,
the SEM-CL mappings and CL-linescans of the microrods reveal the most
intense emission of the InGaN MQW from the upper parts of the microrods,
a dominant emission of the GaN near band edge emission is observed from
the bottom parts of the microwires,
a blue shift of the InGaN MQW emission in the upper part of the micrords was
found.
The reason for the two distinctive regions: bottom part of the microrod with highest
GaN NBE signal and upper part with strong InGaN incorporation could be attributed to
the conditions of the MOVPE reactor. There are two different growth conditions in the bottom
and upper part of the microstructure during the MOVPE growth. Most likely, the origin of these
two different growth conditions is an antisurfactant role of silane used in the process (see
Chapter 5). Similar observations were reported by Tessarek et al. [119].
The SiNx stabilizing layer covers the side walls of the rods. However, there is a gradient
of the SiN layer coverage along the NW. Since the bottom part is exposed to SiN formation
much longer than the top part, the SiN is complete in the lower part of the rod. As
a consequence, no InGaN/GaN growth takes place in this area. Contrary, at the upper part of
the NW the SiN coverage is not completed and InGaN might be deposited there. The result of
InGaN/GaN MQW growth is a spotty pattern observed in the top part of the microrod, as shown
on discussed SEM images above. The InGaN coverage of the sidewalls can be enlarged by
supplying higher TMIn flow (see 9.2.3), however it also changes the InGaN material
composition and as a consequence, changes the optical properties of the wells (redshift of
the InGaN emission peak).
Additionally, in our case the different growth behavior in the upper and lower part of
the rod is strengthened due to relatively high structure dimensions – in average the rods are
above 40 µm high.
The broad GaN NBE with a high FWHM might be attributed to many defects within
the microrods as basal plane stacking faults and some cubic inclusions. In terms of future full
assembled device application, this aspect should be optimized. The defects in the material
structure, especially visible in the bottom part of the microrods (see Fig. 9.4), shall be
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eliminated in order to improve the crystal quality and thus to improve the optical and structural
properties of the NWs. Since the defects are mostly observed in the bottom part of the wires,
they are originating from the nucleation step. It means that the applied conditions are still not
optimized. There is plenty of room for further investigations, e.g. working pressure or V/III
ratio influence on the NWs properties (see chapter 3.4).
Furthermore, a blue shift of the InGaN MQW emission in the upper part of
the microrods observed for the first two samples could be attributed to a lower indium
incorporation, smaller quantum well thickness or different strain conditions.
9.3 Advanced structural characterization of GaN microrods grown under
different conditions by TEM
The first investigated group of the rods consists of pure GaN microcolumns without
MQW incorporation. Figure 9.8 depicts the SEM images of two studied samples. The rods
shown in Fig. 9.8 a) were grown following the optimized growth conditions, as described in
Chapter 8 (additional SEM image – Fig. 8.7). This sample will be called afterwards sample A.
The second image – Fig 9.8 b) shows the group of multicolumnar GaN microrods. Here,
the microstructures were grown with two main differences in comparison to optimized process.
First of all, the SiNx in-situ masking layer was deposited for 500 s instead of 250 s. The second
difference was shorter deposition time. After 240 s of vertical growth with SiH4 support,
the microstructures continued to grow for subsequent 240 s, instead of 720 s as in optimized
process. This sample will be called afterwards sample B.
Figure 9.8: SEM micrographs showing two groups of GaN microrods: a) optimized growth conditions lead to
the formation of individual vertical columns, b) multicolumnar formation due to different growth conditions.
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The TEM cross section of sample A, grown under optimized conditions, is shown on Fig 9.9.
Only hexagons or similarly faceted cross sections were observed. There is a fraction of rods
with a thickness of 1 – 1.3 µm, however also some thicker structures are found. Interestingly,
there are voids within several microwires. It suggests that investigated structures are hollow
with a nanopipe inside them.
Figure 9.9: TEM cross section of GaN microrods bases from sample A: a) overview of a group of microstructures,
b) detailed view of smaller microwire group.
Figure 9.10 depicts the bottom parts of GaN microrods of sample B, grown under non-
optimized growth conditions. The sample preparation resulted in very thick sections due to
thicker microrods found in this sample. The structures 8-10 µm wide were found. Similarly to
sample A, here also some regular hexagonal cross sections were found – Fig. 9.10 a-b).
Interestingly, also different morphology was observed – Fig. 9.10 c-e). The non-hexagonal
microrods’ base was formed due to coalescence of a few crystallites. As a consequence, not
regular shape of the pedestal facets was grown. Moreover, in studied microstructrures more
than one void was found. This multi holes are also most likely a result of a few grains
concrescence.
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Figure 9.10: TEM cross section of GaN microrods bases from sample B: a-b) solid hexagonal structures, c-e) not
regular hexagonal base of the microcolumn is formed by coalescence of a few crystallites.
The heterostructures containing InGaN MQWs were also characterized. The microrods
were grown following the conditions as discussed in chapter 8.3 (see also the optical
characterization of the rods by µPL and nano-scale CL). The three samples correspond to
2.6 µmol/min, 5.2 µmol/min and 10.4 µmol/min of TMIn flow supplied during MQW
deposition. The mentioned samples will be called afterwards, sample C, D and E, respectively.
Figure 9.11 presents the TEM cross sections of microrods bottom parts from samples C, D and
E. In each of the studied samples two different rod morphologies were found. First, the typical
hexagonal cross sections of the structures were observed, as presented in the upper row of
the TEM micrographs from Fig 9.11. Additionally, in many cases, faceted crystallites were
observed on the surface of the rods – bottom row of the TEM micrographs from Fig. 9.11.
The facets alignments suggest that the crystallites are epitaxial to the rod. The nucleation started
at the edges of the microrods. Such snow flake cross section morphology is becoming more
dominant as the TMIn flow increases (transition from C to E). The sample grown using
the highest TMIn flow, exhibits most pronounced snow flake cross section (E).
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Figure 9.11: TEM cross sections of microrod bottom parts from samples C, D and E. The left column corresponds
to regular hexagonal microrods, whereas right column represents the snow flake morphology of the microwire
cross section.
Besides cross sections of the microrods another features from the samples C, D and E
were studied. Figure 9.12 depicts the deitaled TEM view of a horizontal GaN microrod from
86
sample C, a nanopipe within the GaN microrod from sample D and dark field of the microrod
basis from sample E.
The laying microrod from sample C was ion milled from both sides. The crystallites on
the sidewalls are seen from a side view showing their distribution along the rod. Typically,
sharp edges are likely to attract and nucleate more crystallites then sidewalls. However, here it
is not possible to tell whether the prepared section contains crystallites grown on plane sidewalls
or sharp edges of the rod.
The detailed view of cross section of microrod from sample D reveals a presence of
a nanopipe within the microwire. Interestingly, in that case, the nanocavity exhibits regular
hexagonal cross section. In other studied structures (samples A-E) such holes were also
observed, but with irregular shape and bigger size.
The substrate of sample E is covered with a textured film with a crystallite size of ~ 50-
100 nm. The texture means close to single crystalline orientation grains as seen on the Fig 9.12
c) dark field image.
Figure 9.12: Detailed TEM view of the a) GaN microrod laying on the substrate, b) nano-pipe within the GaN
microrod (detailed view of cross section depicted in the Fig. 9.11 D above), c) dark field of the microrod basis
(pair of bright field depicted in the Fig. 9.11 E above).
The advanced TEM characterization of self-assembled GaN NW allowed to understand
more deeply the growth of such nanostructures on Si(111) substrates. The first very important
finding was importance of initial nucleation phase of GaN rods. This stage is crucial for
the structural properties of GaN wires. If the nucleation sites are too dense, then coalescence of
a few crystallites might occur. As a consequence, the non-hexagonal pedestal is formed which
acts as a base for microrod growth. Moreover, due to the grain coalescence the non-regular
shape configurations are present what enhances the creation of defects and voids within
the GaN structures. The second important finding was also attributed to the microrod
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morphology. Two different cross-sections were found: regular hexagonal and flake-like one.
Interestingly, the fraction of two morphologies was found to be TMIn-flux related. The bigger
the TMIn-flux was applied during the MQWs growth, the more dominant snow-flake
morphology was observed. Apparently, higher TMIn-flow supported the selective growth on
the vertical edges resulting in cluster-like InGaN formation and consequently snow-flake like
morphology.
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Chapter 10
Summary and conclusions
A comprehensive study on MOVPE of InGaN/GaN nanowire heterostructures has been
conducted. Epitaxial growth of GaN rod structures, based on three different approaches was
investigated. The grown samples were analyzed in detail by several measurement techniques to
provide information about structural and optical properties of the nanostructures. The ultimate
goal was to develop the growth process for GaN NW on silicon substrates and qualify their
properties in terms of the future nanoLED applications.
The first investigated growth approach was the vapor-liquid-solid (VLS) growth mode.
Due to the high reactivity between the Au and Si, the growth of GaN NW on Si(111) was not
successful. There is a room for improvement for a buffer development and investigation of Au-
initiated VLS mode on new interlayers on silicon. In our case, the focus was set to deeply
investigate the structural and optical properties as well as elemental characteristics of single
InGaN/GaN NW grown on sapphire substrates to qualify the VLS approach as a fabrication
method for nanoLED applications. The studied structures were characterized by means of
different complementary non-destructive techniques. In particular, the QWs emission was
studied by photoluminescence (PL), the crystal quality and strain through Raman scattering,
X-ray diffraction (XRD) and X-ray absorption near edge structure (XANES) spectroscopy with
hard X-ray nanoprobe and the elemental distribution thanks to X-ray fluorescence (XRF) and
energy dispersive spectroscopy (EDS).
XRF intensity maps revealed a homogeneous distribution of Ga elements along
the c-axis of a GaN NW. On, the other hand, In distribution was inhomogeneous. The coaxial
InGaN QWs formed at the lateral surfaces of the NW, but they did not completely form at
the base of the NW. The root cause of this event was assign to the nature of VLS growth mode.
The inhomogeneous incorporation of In is most probably driven by the nanodroplet localized
on top of the nanowire. Interestingly, Au catalyst was detected only on top of the NWs,
implying no incorporation of Au along the NW. This observation is very important in terms of
future full assembled device, i.e. nanoLED. The unfavorable incorporation of Au elements may
strongly decrease the properties and functionality of the device. However, presented results
89
proof that Au-initiated VLS growth of GaN NWs may be considered as an alternative for SAG
or self-organized growth of GaN nanostructures. EDX technique provided information about
approximate size of the heterostructure. It was found that the radius was approximately equal
to 80 nm, the thickness of the QWs and the barrier were around 2 nm and 4 nm, respectively
and the concentration of In was around 20 %. Based on XRD and XANES characterization one
could conclude that the inner GaN core was free of strain. Additionally, no mixture of phases
in the NW was revealed. All of the studied NWs exhibited wurtzite crystal structure. The low
temperature PL spectra collected at the bottom and upper part of the nanowire showed
an inhomogeneous distribution of the InGaN and the In concentration along the core-shell
structure, as already reported by XRF.
The second investigated approach to synthesize GaN nanowires was selective area
growth (SAG) on Si(111) substrates. To study the growth mechanism of selectively grown GaN
microrods on silicon substrates a special template was prepared. A mask consisted of different
shapes: hexagonal, circle and square openings as well as different diameters of the windows,
between 1-8 μm and separation distances of the opening between 2 – 16 μm. In order to simplify
the process and limit the necessary technology steps Si(111) wafers were directly structured.
A MOVPE growth process was designed, which allowed formation of GaN microrods
vertically aligned to the silicon substrate. The GaN microcolumns were rotated by 30° angle
with respect to the Si(111) flat, which suggested the formation of m-plane {101̅0} GaN facets.
Additionally, the top of the structures was of pyramidal shape – six {101̅1} facets. This
morphology is attributed to the metal polarity of GaN wires. The Al-face polarity of the AlN
seeding layer determined the Ga-polarity of the GaN layer. Consequently, GaN pyramids were
formed. The adapted growth time was necessary for different opening diameters to obtain well
developed, selectively grown vertical GaN rods. Too long growth time resulted in a loss of
selectivity due to lateral growth and parasitic nucleation on the mask. As a result, neighboring
structures tended to coalesce. Additionally, the morphology of the structures changed. Growth
continued in different directions resulting in randomly shaped structures. The optimized time
for microrod structures growth was found to be 1 h and 3 h for opening diameters of 4 μm and
8 μm, with the spacing of 8 μm and 16 μm, respectively. The structure development was
investigated for a microrod grown in the opening of a nominal diameter of 8 μm and separation
distance of 16 μm. The height of the m-plane and distance of two opposite side facets was
measured as a function of growth time. The dependence was not proportional in time. There
was a very pronounced lateral growth over the mask during nucleation increasing the diameter
of the microrod from 8 µm (nominal size of the openings) to more than 16 µm (after 30 min of
90
growth). The aspect ratios for a microwire grown for 3 h were: 0.34 and 1.38 considering height
of the m-plane and total height of the rod, respectively. In addition, the supporting role of silane
during vertical growth of the GaN columns was shown experimentally. Structure morphology
for samples grown without and with SiH4 injection was very different. In the first case, without
silane support, there was a formation of only truncated pyramidal structures. Opposite, once
the silane was introduced as an antisurfactant, the initial pyramids tend to grow vertically and
formed GaN microrods. The structural properties of the nanowires were investigated by Raman
spectroscopy. The frequency and FWHM of the E2h mode served as good indicators for
the assessment of strain fields and the crystal quality of the material. The values obtained for
the microrods grown for 3 h with the nominal diameter of 8 µm revealed nearly strain-free
structures (central frequency of 567 ± 1 cm-1) of crystal quality (FWHM = 4.2 ± 1 cm-1)
comparable to strain-free bulk crystals. Additionally, the evolution of the peak frequency and
FWHM was studied along the full microrods length. The estimated value of strain was in range
of -0.1%, which suggested that the microrods grew nearly free of strain during most of
the stages of the growth.
In order to achieve the functional full-assembled nanoLED the further research shall be
conducted. The studied structures exhibited good crystal quality, however the aspect ratio of
the rods shall be improved. Moreover, the core/shell MQWs shall be incorporated and
intensively studied. Especially, the In incorporation aspect shall be investigated in detail.
The third and last investigated approach was self-organized growth of GaN nanowires
on Si(111) substrates. The innovative, novel approach was developed during this project to
meet the goal: enable the fast and cost-efficient growth of NW on Si(111) and qualify
the structures as building blocks for the future nanoLED.
At the beginning, the AlN buffer as a basis for subsequent GaN nanowire growth on
silicon substrates was investigated. The impact of asynchronous injection of precursors before
AlN layer deposition was studied. It was found that both predose steps, including Al preflow
or nitridation led to mixed polarity of AlN layer. However, Al predose resulted in development
of more pronounced metal domain in comparison to NH3 predose in which case the AlN layer
consisted mostly of N-polar domains. The observation was based on the KOH etching
experiment and subsequent comparison of AFM micrographs of the etched AlN buffers.
The influence of the AlN buffer polarity on the nanowire morphology was shown. Metal-polar
AlN buffer resulted in growth of only not activated GaN nucleation sites and some bigger GaN
pyramidal structures with very high density. No vertical nanostructures were observed.
Contrarily, N-polar AlN buffer ensured formation of GaN nanowires. Besides vertical and tilted
91
structures, also some not activated nucleation sites were found, but with much less density in
comparison to Al-polar buffer. The nanowire growth optimization model based on
the nanostructure density as a function of the key process parameters was proposed. Impacts of
SiNx in-situ masking layer deposition time, growth temperature and silane injection time were
investigated. Based on the observations and optimization steps the following optimized
parameters were defined: 5 s of ammonia preflow before AlN buffer, 250 s of SiNx in-situ
masking layer deposition, 1020 °C GaN growth temperature and 240 s of silane injection time.
Such growth parameters led to the growth of GaN nanowires vertically aligned to the substrate.
There was a very small fraction of tilted structure and the multicolumnar structure formation
was suppressed. Moreover, the unfavourable not activated nucleation seed density was limited.
Besides technological aspects also maintenance was discussed. The unfavourable effect of AlN
susceptor coating on nanowire growth homogeneity was underlined. It was found that highly
coated deposition plate may strongly affect the uniform and homogeneous deposition of AlN
with desired polarity. The material was unintentionally decomposed from the susceptor surface
and transported onto the wafer during initial stages of growth promoting Al-polar buffer growth.
The centre part of the wafer exhibited more metal dominant character in comparison to the anti-
flat region.
The properties of self-organized grown InGaN/GaN nanowires were investigated in
detailed by a number of characterization techniques. The structural properties and In
incorporation was characterized by microphotoluminescence (µPL) and nano-scale
cathodoluminescence (CL) mapping. The comparison of three samples grown with different
TMIn flow during MQWs deposition, led to four main observations. The both measurement
techniques clearly showed a red shift of the InGaN MQW emission of the three samples with
an increasing TMIn flow due to the higher indium incorporation. Two distinctive regions within
the microrod structures were observed. The SEM-CL mappings and CL-linescans revealed
the most intense emission of the InGaN MQW from the upper part of the microwires.
A dominant emission of the GaN near band edge (NBE) emission was observed from
the bottom parts of the structures. The origin of these segregation effect was attributed to the
antisurfactant role of silane, used in the process to enhance the vertical growth of the rods.
The SiNx stabilizing layer formed on the side walls of GaN microrods with the coverage
gradient. The lower parts of the rods were exposed to SiN formation much longer in comparison
to upper parts. Thus, SiN masking layer was complete in the bottom part of the microwires. As
a consequence, no InGaN/GaN growth took place in this region of the structure. Opposite, in
the upper part of the NW, where the SiN coverage was not completed, InGaN could have been
92
deposited. The different growth behaviour in the bottom and upper part of the microrod was
also strengthened due to relatively high structure dimension – in average the investigated rods
were above 40 µm high. The last finding dealt with a presence of a blue shift of the InGaN
MQW emission in the upper part of the microrods observed for the samples grown using
the smallest and moderate indium flow. The blue shift could be a result of lower Indium
incorporation, smaller quantum well thickness or different strain conditions.
The advanced structural properties of self-organized GaN microrods grown under
different conditions were characterized by TEM. The first investigated group of structures
grown under optimized conditions exhibited hexagonal or similarly faceted cross sections.
There was a fraction of rods with thicknesses of about 1-1.3 µm, but also some bigger structures
were found. Some single voids within several microwires were observed, which suggested that
structures were hollow with a nanopipe inside them. The cross sections of another group of
GaN microrods grown under non-optimized conditions revealed different properties. First of
all, the bigger structures with a diameter ranged between 8 and 10 µm were found. Besides
typical hexagonal cross section also different morphology was observed. The irregular, non-
hexagonal shape of the microrods’ base originated from the coalescence of a few nucleation
crystallites. Consequently, pedestals consisting of irregular facets were formed. Additionally,
in studied structures more than one void was found. The multi holes structures were the result
of a few grains coalescence. It is therefore suggested to further optimize the growth process of
self-organized GaN nanowires on Si(111) by investigating the impact of other growth
parameters (i.e. pressure and V/III ratio) with the major focus on the nucleation part in order to
improve the crystal quality as well as to minimise the detrimental not activated nucleation
between NWs. Moreover, the process window for InGaN MQW deposition shall be defined in
order to solve the InGaN inhomogeneity issue. The higher TMIn flux applied during MQW
deposition resulted in higher coverage of side walls with InGaN. On the other side, higher TMIn
flow was a root cause of snow-flake cross section morphology of the structures. Additionally,
higher density of GaN NW is expected and therefore new set of parameters shall be also
determined based on the existing optimization model.
93
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List of Figures
Figure 1.1: Bandgap of binary InN, GaN, and AlN and their ternary alloys as a function of in-
plane lattice constants (no bowing assumed). ............................................................................ 2
Figure 2.1: Perspective views along [0 0 0 1] direction of wurtzite and cubic zincblende GaN,
a) and b) respectively [9]. The large circles represent gallium atoms and the small circles
nitrogen. c) The hexagonal unit cell of GaN defined by the lattice parameters: the length of
the hexagon’s side (a), (b) and the height (c) of the hexahedron. .............................................. 7
Figure 2.2: Atomic arrangements in two possible GaN polarities: Ga-faced and N-faced [10].
.................................................................................................................................................... 7
Figure 2.3: Schematic draft of axial and core/shell nanowire, a) and b) respectively. ........... 10
Figure 2.4: Calculated internal quantum efficiency versus current density for c-plane [a) and
b)] and m-plane [c) and d)] growth, under zero bias [a) and c)] or a 3.5 V applied voltage
[b) and d)]. Figure adapted after [26]. .................................................................................... 11
Figure 2.5: Nanowire-based multicolour LED: a) Schematic of the heterostructure cross-
section and energy band line-up. b) Optical microscopy images collected from around the p-
contact of nanowire LEDs in forward bias, showing different colour of emitted light: purple,
blue, greenish-blue, green and yellow. c) Normalized electron-luminance spectra recorded
from five representative forward-biased NW LEDs with 1%, 10%, 20%, 25% and 35% In (left
to right), respectively. Figure adapted after [27]. ................................................................... 12
Figure 2.6: GaN/AlN/AlGaN NW-based transistor. a) Left: cross-sectional, high-angle annular
dark-field scanning TEM image of a radial nanowire heterostructure. Scale bar is 50 nm. Right:
Band diagram illustrating the formation of am electrpm gas (red region) at the core-shell
interface. b) Intrinsic electron mobility of a transistor as a function of temperature (after
correction for contact resistance). c) Logarithmic scale Ids-Vg curve recorded at Vds = 1.5V
(channel length 1 µm, 6 nm ZrO2 dielectric). Inset shows the linear scale plot of the same data.
Figure adapted after [35] ........................................................................................................ 13
Figure 3.1: a) Meltback etching of Si by Ga [52], b) Ga-rich, Si-rich and SiNx formation after
GaN deposition directly on Si substrate [53]. .......................................................................... 15
Figure 3.2: a) GaN NW with two different polarity domains – the pattern on the c-plane top
facet induces the existence of both polarities within the structure, b) the same GaN NW after
KOH etch; flat region on the c-plan facet represents Ga-polarity, whereas rough surface is
attributed to N-polar regions [65], c) GaN NW after KOH etch, arrows indicate the remaining
Ga-polar regions [68]. ............................................................................................................. 17
Figure 3.3: Schematic model of the VLS growth of GaN NWs utilizing Au nanoparticles [91].
.................................................................................................................................................. 18
Figure 3.4: Typical SEM images of GaN NWs grown on the sapphire substrate by the VLS
technique. The metal catalyst used in the growth experiment was a very thin Au film. ........... 19
102
Figure 4.1: Schematic of an AIXTRON 3x2” MOCVD reactor. A: thermocouple, B: tungsten
heater, C: showerhead, D: reactor Lid, E: optical Probe, F: showerhead water cooling,
G: double O-ring seal, H: susceptor, I: quartz liner, J: susceptor support, K: exhaust. ......... 27
Figure 4.2: Graphical visualization of three different growth regimes for MOCVD process:
A – kinetic limited regime, B – mass transport limited regime, C – reduced growth rate due to
desorption from surface instead of incorporation. ................................................................... 28
Figure 4.3: Schematic of a synchrotron facility in Grenoble including an injection system,
a storage ring and beamlines. The injections system consists of an electron gun, a linac and
a booster. The parts of storage ring are radio frequency cavities, bending magnets and
undulators or wigglers. Figure adapted from [108]. ............................................................... 28
Figure 4.4: Draft of the experimental setup for recording XRF and XANES using a synchrotron
X-ray nanobeam at the beamline ID22. Figure adapted from [108]. ...................................... 29
Figure 4.5: XRF and Auger electron yields for K-shell as a function of atomic number, solid
and dotted curve, respectively. Figure adapted from [112]. ................................................... 30
Figure 5.1: Comparison of GaN microrods grown on sapphire substrate by self-organized
mode. Left image refers to the process without silane support, whereas right one depicts GaN
microrods grown with SiH4 injection [119]. ........................................................................... 33
Figure 5.2: The schematic model of GaN NWs growth [119]. ................................................ 34
Figure 6.1: Process sequence steps involved in the Au-initiated VLS growth of GaN NWs on
sapphire substrates by MOCVD. .............................................................................................. 37
Figure 6.2: V/III ratio influence on the GaN NW morphology during the Au-initiated growth of
GaN NWs. The growth temperature was 870 °C and the total pressure applied was 100 mbar.
.................................................................................................................................................. 38
Figure 6.3: Average NW diameter (nm) as a function of total working pressure (mbar). ...... 39
Figure 6.4: XRF intensity maps showing the distribution of Ga, In and Au along the GaN NW
– a), b) and c), respectively. ..................................................................................................... 40
Figure 6.5: EDS In profile perpendicular to the diameter of the NW. The inset shows a HR-
TEM of a single dispersed NW with the magnification of the bottom part of the same NW. ... 41
Figure 6.6: a) Representative (210) and (211) XRD peaks along with their best Gaussian fit.
The peak position and the FWHM are indicated for both fits. b) Evolution of the lattice
parameters along the c-axis. .................................................................................................... 42
Figure 6.7: Representative unpolarized Raman spectra of a single NW taken with 514 nm laser
line. ........................................................................................................................................... 43
Figure 6.8: Representative LT-PL spectra of a single NW taken at 5 K. The signal from top and
bottom of the NW – black and red lines, respectively. ............................................................. 44
103
Figure 7.1: The utilized mask designed for SAG GaN microrods growth on Si(111) substrates.
Red circled regions showed the investigated mask units. The labels under each mask unit stand
for: H – hexagonal, C – circle, S – square opening; first number refers to window diameter,
second number refers to spacing between two neighboring windows. For example, Hx8x16
stands for hexagonal opening of 8 µm diameter and distances between openings of 16 µm. . 46
Figure 7.2: The process steps for SiNx/Si(111) mask preparation for GaN microrods selective
area growth. ............................................................................................................................. 47
Figure 7.3: The scheme of SAG of GaN microrods on Si(111). ............................................... 48
Figure 7.4: SEM image of SAG GaN rods grown on a SiNx/Si(111) template. The nominal
diameter of the openings was 8 µm and the distance between the openings was 16 µm. Samples
were grown for 30 min, 1 h, 3 h – images: a), b), c), respectively. (d) Sketch representing
the orientation of m-plane GaN sidewalls of the microrod with respect to Si(111) flat of
the substrate. The nominal diameter of the openings was 4 µm and the distance between
the openings was 8 µm. Sample was grown for 1 h. ................................................................ 49
Figure 7.5: Graph showing height of m-plane side facets (black) and distance between two
opposite side facets (blue) of GaN microrods as functions of growth time. The nominal diameter
of the opening was 8 µm and the opening separation was 16 µm. .......................................... 51
Figure 7.6: 45 degrees tilt view SEM image of SAG GaN rods grown on a SiNx/Si(111)
template. The nominal diameter of the openings was 2 µm and the distance between the
openings was 4 µm. Samples were grown for 1 h without (image a) and with silane support
(image b). ................................................................................................................................. 52
Figure 7.7: Room temperature PL spectra of single microrods grown with and without silane
support. The spectra correspond to single microrods of 8 µm nominal diameter and 16 µm
distance between openings. ...................................................................................................... 53
Figure 7.8: Comparison of Raman spectra measured for time series samples. Microrods were
grown for 30 min, 1 h and 3 h. Each spectrum corresponds to a single microrod of 8 µm nominal
diameter and 16 µm distance between openings. ..................................................................... 54
Figure 7.9: (a) Raman peak corresponding to the E2h phonon of an 8 µm diameter microrod
measured with the excitation light focused at different depths (z) along the microrod axis.
The apex of the pyramidal tip of the microrod defines z=0 and for increasing depth z<0. (b)
Evolution of the E2h frequency and FWHM as a function of the depth. The frequency of the bulk
E2h is plotted as a dashed line [128]. ....................................................................................... 55
Figure 8.1: AFM images of the AlN samples after 60 min of etching in 10% KOH at 30 °C. Left
image refers to 5 s of TMAl predose, right image to 5 s of NH3 predose. ............................... 58
Figure 8.2: 45°-tilted SEM view of self-assembled GaN nanowires grown on Si(111) substrate
using different polarity AlN buffers. The preflow applied before AlN deposition was 5 s of TMAl
or NH3 for sample a) and b), respectively. ............................................................................... 60
Figure 8.3: 45°-tilted SEM view of self-assembled GaN nanowires grown on Si(111) substrate
showing the description of the measured structures used to determine the density: not activated
104
GaN nucleation seeds (blue), random structures (orange), vertical wires (red), tilted wires
(green), and multicolumnar wires (yellow). ............................................................................. 61
Figure 8.4: Density of the structures as a function of SiNx deposition time. Graph a) shows
density of total structures (black), sum of the rods (red) and not activated nucleation sites
(blue). Total structure density stays for the sum of the rods density and density of the randomly
shaped structures. Graph b) depicts the density of vertical (black), titled (red) and
multicolumnar (blue) wires. ..................................................................................................... 63
Figure 8.5: Density of the structures as a function of GaN growth temperature. Graph a) shows
density of total structures (black), sum of the rods (red) and not activated nucleation sites
(blue). Total structure density stays for the sum of the rods density and density of the randomly
shaped structures. Graph b) depicts the density of vertical (black), titled (red) and
multicolumnar (blue) wires. ..................................................................................................... 65
Figure 8.6: Density of the structures as a function of silane injection time ratio (time with SiH4
supply divided by total GaN NW growth time). Graph a) shows density of total structures
(black), sum of the rods (red) and not activated nucleation sites (blue). Total structure density
stays for the sum of the rods density and density of the randomly shaped structures. Graph b)
depicts the density of vertical (black), titled (red) and multicolumnar (blue) wires. ............... 67
Figure 8.7: 45°-tilted SEM view of self-assembled GaN nanowires grown on Si(111) substrate
using optimized growth conditions. .......................................................................................... 68
Figure 9.1: 45° tilted SEM image of GaN nanowires grown on Si(111) substrate. a) GaN rods
on Si(111) substrate. b) GaN nanocolumns with 3 pairs of InGaN/GaN core/shell MQWs grown
on Si(111) substrate. c) Non-hexagonal microrod. .................................................................. 71
Figure 9.2: Low-temperature (10K) photoluminescence spectra of GaN nanostructures. Left
graph a) shows a summary of measured structures: standing and lying NWs as well as parasitic
not-activated nucleation. Right graph b) shows only the spectra measured for well aligned
vertical NWs. ............................................................................................................................ 72
Figure 9.3: PL (square) and CL (triangle) wavelength as function of TMIn supply of
InGaN/GaN core-shell heterostructures. ................................................................................. 73
Figure 9.4: a) SEM image of a single GaN nanowire with InGaN/GaN MQWs. b-c) SEM-CL
mapping. d) CL-linescan along the InGaN/GaN microrod. Two micrographs from image d)
refer to the same GaN microrod. The position on both linescan maps was measured from 0-35
µm. The MQWs were deposited with the lowest investigated TMIn flow of 2.6 µmol/min. ..... 75
Figure 9.5: a) SEM image of the non-hexagonal GaN nanowire with InGaN/GaN MQWs. b-d)
SEM-CL mapping e) CL-linescan along the InGaN/GaN microrod. e) CL spectrum of the
studied microrod. f) SEM-CL mapping from upper part of microrod. g) CL spectrum of upper
part of the microrod. The MQWs were deposited with the middle investigated TMIn flow of 5.2
µmol/min................................................................................................................................... 77
Figure 9.6: a) SEM image of the upper part of an individual hexagonal GaN nanowire with
InGaN/GaN MQWs. b-d) SEM-CL mappings in different spectral regions, e) CL spectrum of
the studied microrod. f) CL linescan along microrod. The MQWs were deposited with medium
TMIn flow of 5.2 µmol/min. ...................................................................................................... 79
105
Figure 9.7: a) SEM image of the single GaN nanowire with InGaN/GaN MQWs. b-c) SEM-CL
mapping. d) CL-linescan along the InGaN/GaN microrod. Two micrographs from image d)
refer to the same GaN microrod. The position on both linescan maps was measured from
0-35 µm. The MQWs were deposited with the highest investigated TMIn flow of 10.4 µmol/min.
.................................................................................................................................................. 80
Figure 9.8: SEM micrographs showing two groups of GaN microrods: a) optimized growth
conditions lead to the formation of individual vertical columns, b) multicolumnar formation
due to different growth conditions. .......................................................................................... 82
Figure 9.9: TEM cross section of GaN microrods bases from sample A: a) overview of a group
of microstructures, b) detailed view of smaller microwire group. ........................................... 83
Figure 9.10: TEM cross section of GaN microrods bases from sample B: a-b) solid hexagonal
structures, c-e) not regular hexagonal base of the microcolumn is formed by coalescence of
a few crystallites. ...................................................................................................................... 84
Figure 9.11: TEM cross sections of microrod bottom parts from samples C, D and E. The left
column corresponds to regular hexagonal microrods, whereas right column represents the
snow flake morphology of the microwire cross section. .......................................................... 85
Figure 9.12: Detailed TEM view of the a) GaN microrod laying on the substrate, b) nano-pipe
within the GaN microrod (detailed view of cross section depicted in the Fig. 9.11 D above),
c) dark field of the microrod basis (pair of bright field depicted in the Fig. 9.11 E above). ... 86
106
List of Tables
Table 2.1: Main physical properties of III-N nitrides [12]. ....................................................... 8
Table 2.2: Main optoelectronic properties of III-N nitrides [17]. ............................................. 8
107
List of Abbreviations
AFM .................................................................................................... Atomic force microscopy
AlN ................................................................................................................. Aluminium nitride
Au ......................................................................................................................................... Gold
CCS ................................................................................................. Close-Coupled Showerhead
CL .............................................................................................................. Cathodoluminescence
EDX ................................................................................. Energy-dispersive X-ray spectroscopy
ESRF .......................................................................... European Synchrotron Radiation Facility
EXAFS ........................................................................ Extended X-ray absorption fine structure
FWHM ............................................................................................ Full width at half maximum
Ga .................................................................................................................................... Gallium
GaN ...................................................................................................................... Gallium nitride
H2 .................................................................................................................................. Hydrogen
HRTEM ........................................................ High-resolution transmission electron microscopy
In ........................................................................................................................................ Indium
InGaN ....................................................................................................... Indium gallium nitride
InN ......................................................................................................................... Indium nitride
IQE ................................................................................................... Internal quantum efficiency
KOH ........................................................................................................... Potassium hydroxide
LED ............................................................................................................. Light emitting diode
LPCVD ........................................................................ Low-pressure chemical vapor deposition
LT-PL ................................................................................ Low temperature photoluminescence
MOCVD ...................................................................... Metal organic chemical vapor deposition
MOVPE ................................................................................. Metal organic vapor phase epitaxy
MQW ............................................................................................................ Multi quantum well
N2 .................................................................................................................................... Nitrogen
NH3 ............................................................................................................................... Ammonia
NW ............................................................................................................................... Nanowire
PL .................................................................................................................. Photoluminescence
QCSE .......................................................................................... Quantum-confined Stark effect
RIE .............................................................................................................. Reactive ion etching
SAG ............................................................................................................Selective area growth
SEM .............................................................................................. Scanning electron microscopy
Si ........................................................................................................................................ Silicon
SiC ........................................................................................................................ Silicon carbide
SiH4 .................................................................................................................................... Silane
SiN ......................................................................................................................... Silicon nitride
SRH ..............................................................................................................Shockley-Read-Hall
TDs ........................................................................................................... Threading dislocations
TEGa ................................................................................................................... Triethylgallium
TEM ...................................................................................... Transmission electron microscopy
TMAl ........................................................................................................... Trimethylaluminium
TMGa ............................................................................................................... Trimethylgallium
VLS ................................................................................................................ Vapor-liquid-solid
XANES ...............................................................................X-ray absorption near edge structure
XRD .................................................................................................................. X-ray diffraction
XRF ................................................................................................................ X-ray fluorescence
108
Scientific appendix
Patents and Publications:
1) B. Foltynski, M. Vallo, C. Giesen, M. Heuken, GaN wires on Si, patent Ai 2014/22
2) B. Foltynski, N. Garro, M. Vallo, M. Finken, C. Giesen, H. Kalisch, A. Vescan,
A. Cantarero, M. Heuken, The controlled Growth of GaN Microrods on Si(111)
Substrates by MOCVD, Journal of Crystal Growth 03/2015; 414:200-204
3) B. Foltynski, C. Giesen, M. Heuken, Self-organized growth of catalyst-free GaN nano-
and micro-rods on Si(111) substrates by MOCVD, Physica Status Solidi B 05/2015;
252(5)
4) A. Hospodková, M. Nikl, O. Pacherová, J. Oswald, P. Brůža, D. Pánek, B. Foltynski,
E. Hulicius, A. Beitlerová, M. Heuken, InGaN/GaN multiple quantum well for fast
scintillation application: radioluminescence and photoluminescence study,
Nanotechnology 25 455501
5) A. Kovács, M. Duchamp, R. E. Dunin‐Borkowski, R. Yakimova, P. L. Neumann,
H. Behmenburg, B. Foltynski, C. Giesen, M. Heuken, B. Pécz, Graphoepitaxy of high-
quality GaN layer on graphene/6H-SiC, Advanced Materials Interfaces 12/2014; 2(2)
6) B. Pécz, L. Tóth, G. Tsiakatouras, A. Adikimenakis, A. Kovács, M. Duchamp, R. E.
Dunin-Borkowski, R. Yakimova, P. L. Neumann, H. Behmenburg, B. Foltynski,
C. Giesen, M. Heuken, A. Georgakilas, GaN heterostructures with diamond and
graphene, Semiconductor Science and Technology 11/2015; 30(11):114001
7) To be submitted:
a. B. Foltynski, M. Müller, G. Z. Radnoczi, C. Giesen, A. Dempewolf, F. Bertram,
B. Pécz, J. Christen, M. Heuken, Cathodoluminescence and
µPhotoluminescence of InGaN/GaN Nanowire-based core/shell
heterostructures
b. E. Secco, B. Foltynski et. al., Elemental distribution of coaxial InGaN/GaN
quantum wells in nanowires
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Conference contribution:
1) B. Foltynski, N. Garro, M. Vallo, M. Finken, C. Giesen, H. Kalisch, A. Vescan,
A. Cantarero, M. Heuken, The controlled Growth of GaN Microrods on Si(111)
Substrates by MOCVD, 17th International Conference on Metalorganic Vapor Phase
Epitaxy (ICMOVPE), Lausanne, Switzerland, 13th – 18 July 2014, poster session
2) B. Foltynski, C. Giesen, M. Heuken, Self-organized growth of catalyst-free GaN nano-
and micro-rods on Si(111) substrates by MOCVD, International Workshop on Nitride
Semiconductors, Wroclaw, Poland, 24-29 August 2014, poster session
3) B.Foltynski, M. Müller, P. Corfdir, C. Giesen, A. Dempewolf, J.Christen, M. Heuken,
Cathodoluminescence and µ-Photoluminescence of InGaN/GaN Nanowire-based
core/shell heterostructures, Deutsche Gesellschaft für Kristallzüchtung und
Kristallwachstum Conference (DGKK), Magdeburg, Germany, 11-12 December 2014,
talk
4) M. Heuken, Preparation and Properties of Nanostructures, 6th International
Conference on Nanomaterials - Research & Application (NANOCON), Brno, Czech
Republic, 5-7 November 2014, talk
5) B. Reuters, B. Foltynski, D. Fahle, H. Hahn, B. Hollander, M. Heuken, H. Kalisch,
A. Vescan, Epitaxial Growth of AlInGaN layers on AlN templates for backbarrier
application in DHFET, Compound Semiconductor Week 2014 (CSWEEK),
Montpellier, France, 11-15 May 2014, talk
6) W. Witte, B. Foltynski, M. Finken, M. Heuken, H. Kalisch, A. Vescan, Vertical Field
Effect Transistors with p-GaN Current-Blocking Layer, Workshop on Compound
Semiconductor Devices and Integrated Circuits (WOCSDICE), Delphi, Greece, 15-18
June 2014, talk
7) E. Secco , N. Garro, A. Cantarero, M-H. Chu, J. Segura-Rui; G. Martinez-Criado,
Bartosz Foltynski, H. Behmenburg , C. Giesen, M. Heuken, Elemental distribution of
coaxial InGaN/GaN nanowires grown by metalorganic chemical vapor deposition, 10th
International Conference on Nitride Semiconductors (ICNS-10), Washington, USA, 25-
30 August 2013, poster session
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8) B. Foltynski, M. Vallo, C. Giesen, M. Heuken, Catalyst-free growth of GaN
nanostructures on Si(111), Advanced School on Semiconductor Nanowires, Alghero,
Italy, 6-12 October 2013, poster session
9) M. Vallo, B. Foltynski, C. Giesen, M. Heuken, Gold-assisted growth of GaN
nanostructures on Si(111), Advanced School on Semiconductor Nanowires, Alghero,
Italy, 6-12 October 2013, poster session
10) B. Foltynski, H. Behmenburg, C. Giesen, M. Heuken, MOCVD growth of InGaN
nanowires for optoelectronics and energy harvesting device application, Marie Curie
Actions Conference at EuroScience Open Forum (ESOF), Dublin, Ireland, 11-15 July
2012, poster session
11) E. Hulicius, A. Hospodková, M. Nikl, O. Pacherová, J. Oswald, P. Brůža, D. Pánek,
B. Foltynski, A. Beitlerová, M. Heuken, Radioluminescence and photoluminescence of
InGaN/GaN multiple quantum well nanoheterostructure, Olomouc, Czech Republic,
18th Conference of Czech and Slovak Physicists, 16-19 September 2014, poster session
111
Acknowledgements
First and foremost I would like to gratefully and sincerely thank my advisor Prof.
Michael Heuken for his guidance, understanding, patience and most importantly the faith he
put in me. It has been an honour to be his Ph.D student. I appreciate all his contributions of
time, ideas, and funding to make my Ph.D experience productive and stimulating. He
encouraged me to not only grow as an experimentalist but also as an independent thinker.
I suspect that not that many graduate students are given the opportunity of self-development by
being allowed to work with such independence.
I would also like to give a heartfelt, special thanks to Dr. Christoph Giesen for managing
my Ph.D project within AIXTRON SE. He has been motivating, encouraging and most of all
he was always willing to help.
I thank Prof. Wilfried Mokwa for kindly accepting to referee this work.
I thank Prof. Angella Rizzi and Dr. Joerg Malindretosfor for managing the Marie Curie
Nanowiring project. I thank all of the fellows and researchers for making our network efficient
and for all of the fruitful discussion during our meeting and workshops.
In regards to the external collaboration, I thank colleagues from University of Valencia,
Otto-von-Guericke University Magdeburg, Hungarian Academy of Science, European
Synchrotron Radiation Facility (ESRF), Paul-Drude-Insitute and Academy of Science of
the Czech Republic. I thank Prof. Nuria Garro and Eleonora Secco for their efforts to perform
Raman spectroscopy measurements, Prof. Jurgen Christen, Dr. Frank Bertram and Marcus
Müller for nano-scale cathodoluminescence study, Prof. Bella Pecz for TEM characterization,
Dr. Gema Martinez and Dr. Manh-Hung Chu for their support during measurements in ID-22,
Dr. Pierre Corfdir for photoluminescence analysis and Dr. Alice Hospodková for fruitful
discussions.
I thank colleagues from AIXTRON SE for their input, valuable discussions
and accessibility. In particular, I would like to thank Dr. Hannes Behmenburg for giving me
an introduction to GaN growth by MOCVD, technician Waldemar Fischer for keeping Trine
(3x2” CSS reactor) always in good conditions, Mustafa Öztürk, Dr. Olivier Feron, Dr. Holger
Grube and Martin Vallo for good discussions, Beate Sahl for her help in the chemical labor.
I thank my colleagues from RWTH Aachen University, especially – Matthias Marx (Finken)
for his contribution to selective area growth experiments, Wiebke Witte for her support in
the chemical labour and Benjamin Reuters for fruitful discussions about MOCVD growth.
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Of course no acknowledge would be complete without giving thanks to my parents.
Both have instilled many admirable qualities in me and given me a good foundation with which
to meet life. Both have always expressed how proud they are of me and how much they love
me. I too am proud of them and love them very much.
