Transmission electron microscopy investigationsof the CdSe based quantum structures

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Transmission electron microscopy investigations

of the CdSe based quantum structures

Elena Roventa

University of Bremen

2006

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Transmission electron microscopy investigations

of the CdSe based quantum structures

Ph.D. thesis

to obtain the academic degree

Doctor in Natural Sciences

— Dr. rer. nat. —

of the Fachbereich 1

at the University of Bremen

submitted by

Dipl. Phys. Elena Roventa

1. Examinator: Prof. Dr. rer. nat. Andreas Rosenauer 2. Examinator: Prof. Dr. rer. nat. Detlef Hommel Day of examination: 22.09.2006

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Introduction

Optical communication systems have progressed very rapidly from the research labs into commercial applications. They have already been established within transport networks as point to point links, broadcast distribution and interconnecting electrical nodes. Currently the progress of this technology is significant in the diffusion of multi -wavelength extended capacity links with -wavelength routing at the nodes and add-drop operation of the high data flowing in the domain. Future optical communications networks for terabyte transmission rates require the use of optical routing to cope with ever increasing capacity demand due to growing Internet traffic. In addition, another important role plays the data storage, where light also can provide a solution. Optical data storage systems (ex. DVD) permit to store the information with a high density and fast access.

The special requirement for both applications is the use of a proper light source as monochromatic as possible and of a defined color. Moreover, light with a high inten-sity and easy to focus even over long distances is needed. Additionally, the source has to be small, efficient, robust and cheap. Semiconductor lasers, in which the semicon-ductor serves as photon source, provide such a monochromatic coherent light of high intensity as a result of light amplification by stimulated emission of radiation (LASER). Furthermore, the rapid development of semiconductor technology allows the produc-tion of small devices with high reliability. This leads to the conclusion that, in our days, the semiconductor laser diodes are one of the best light sources used in communication technology.

However, the application range of semiconductor lasers goes beyond the commu-nication area. In fact, our everyday life is surrounded by laser applications which are included in laser pointers, (used to make a bright spot to point with), navigation system for aircraft navigation (semiconductor laser-driven gyroscope about the size of a com-puter chip), laser sights for rifles and guns, health science (use in diagnosis, surgery and imaging), biotechnology etc. One of the most interesting application of lasers is the dis-play technology. Using lasers, new, sharp and brilliant disdis-plays with a color spectrum can be fabricated in a very small size.

Since the realization of the first electrically pumped laser diode in 1962, in the last 40 years these devices transformed from pure academic curiosities to commercial pro-ducts that can be found almost everywhere. Semiconductors lasers are most commonly made from the III-V semiconductors GaAs/AlGaAs (for light with a wavelength around 850 nm) or InP/InGaAsP (for 1.5 µm). The current injected into a forward-biased p-n junction creates pair of free electrons and holes, which then emit and recombine. The emission spectrum from red to infrared is already covered by the commercially avai-lable laser diode. The blue-violet emitting laser diodes based on gallium nitride (GaN)

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are also commercialized (since 2000). The emission is not really blue (≈480 nm) and the In incorporation in the active region is still limited, so these devices could only be optimized for data storage systems (Sony presented a DVD recorder with a capacity of writing of 22.5 GB on a 12 cm disc [1]).

In comparison, for display applications a wavelength emission above 500 nm is needed. Electrically pumped blue-green-yellow emitting laser diodes can in principle be fabricated using ZnSe based semiconductor material. Although continuous wave (cw) ZnSe laser diodes working at room temperature have been achieved, the lifetime is still limited to about 400 h [2]. As a consequence, the most pressing problem re-garding the ZnSe laser diodes is to understand the degradation mechanism behind the insufficient stability. The instability of the active region, which is in general a Cd-based quantum well, compels these laser diodes to be still in a research stage.

According to these circumstances, the interest of this thesis was focused on studies of quaternary CdZnSSe quantum wells with high Cd content which show emission at 560 nm, which is one of the longest emission wavelengths obtained for II-VI laser diodes so far [3]. Furthermore the use of self-assembled CdSe/ZnSSe quantum dot stack struc-tures as active region of the ZnSe laser diode at the same emission wavelength [4] is expected to improve the lifetime of the device. For such quantum dot lasers a lower laser threshold is expected and defect induced degradation effects can be diminished. For optimal device performance a narrow QD distribution is necessary. It follows that the ordering and correlation effects play an important role, therefore a closer look on CdSe/ZnSSe stack structures was performed as well.

Using mainly Transmission Electron Microscopy, the attention was drawn first to the microstructural changes and possible composition fluctuations of the CdZnSSe quater-nary quantum well due to the degradation process and secondly to the the chemical composition and ordering phenomena in the CdSe dots embedded in ZnSSe material. The choice of the electron microscopy method is connected with the electron wave-length which is in the picometer range, and gives information about microstructural changes at nanometer scale. The presence of the chemically sensitive (002) reflection in zincblende CdSe crystals offers the possibility to gain information about chemical com-position of quantum well or quantum dot structures. Using the (002) dark-field as well as high resolution imaging technique, composition-related information can be obtained. On the other hand, the presence of strain sensitive reflections can provide details about structural defects (stacking faults, dislocations etc) present in these quantum structures. Employing Compositional Evaluation by Lattice Fringe Analysis (CELFA) and Digital Analysis of Lattice Images (DALI) in connection with other experimental methods (such as photoluminescence), information about outdiffusion and segregation effects can be obtained. Whereas Atomic Force Microscopy and Scanning Tunneling Microscopy pro-vide chemical, structural and strain information of crystal surfaces, electron microscopy yields information averaged along the thickness of the specimen in the electron beam direction, which is in the order of a few nanometers to approximately 50 nm.

The following synopsis explains the structure of the thesis. Chapter 1 contains the basic theoretical information necessary for the understanding and interpretation of the experimental results. It starts with an overview of the principal properties of II-VI com-pounds followed by a short explanation of the epitaxial process used for the growth of such complicated structures as laser diodes. In particular the two growth method: Molecular Beam Epitaxy (MBE) and Migration Enhanced Epitaxy (MEE) are described

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Introduction

and compared. At the end of the chapter, an overview of the principal defects that can be found in a zincblende crystal structure is presented.

Several experimental methods have to be employed for a complete characteriza-tion of such laser diodes. However, we focus mainly on TEM results here. Therefore, Chapter 2 presents the basic principles of electron microscopy, a description of electron diffraction and imaging formation theory. Since the sample preparation is an important issue for an accurate investigation of the crystal structure, the sample preparation tech-niques are also briefly presented. To confirm the findings by transmission electron mi-croscopy, a comparison with the results obtained using other experimental techniques is necessary. Therefore, a very short description of the complementary techniques: pho-toluminescence, electroluminescence and high resolution X-ray diffraction is given.

More then ten years of research deals with degradation process of ZnSe based lasers. From optical investigations to electrical and structural ones, the acquired informations allowed continuous improvement of the ZnSe heterostructure based devices. However, studies of microstructural changes of the quantum well area are few and not really con-clusive. Therefore, Chapter 3 is focused on the microstructural studies of the degraded CdZnSSe quantum well. Using an advanced TEM sample preparation method it was possible to investigate the degraded quantum well underneath the injection stripe. Fur-thermore, electro-optical properties and structural properties are compared in order to obtain a complete view of the degradation process.

Chapters 4 and 5 are dedicated to analysis of CdSe/ZnSSe quantum dot structures. Chapter 4 addresses the topic of elemental distribution in the structure. An estimation of the chemical composition in the CdSe/ZnSSe can be obtained using the (002) dark-field technique by comparing experimental intensity profiles with simulated profiles. Chapter 5 presents TEM studies of the ordering phenomena in the CdSe/ZnSSe struc-tures. The dependence of the ordering process on ZnSSe spacer thickness was verified on a five-fold CdSe/ZnSSe series. The results were compared with the data obtained for a ten-fold stack in order to observe the influence of the number of CdSe layers. The TEM findings are also compared with grazing incidence small angle X-ray scattering results.

The results of this work will be summarized in Chapter 6. A short outlook will also be given, concerning possible experimental work that can be performed regarding to study the degradation in quaternary quantum wells as well as the ordering phenomena in quantum dot stack structures.

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Chapter 1 Theoretical background

The ZnSe and CdSe-based alloys offer a great potential for application in opto-electronic semiconductor devices due to the wide energy spectral range covered by their ternary and quaternary compounds from blue to green. Recently (during the last decade), by employing advanced techniques such as Molecular Beam Epitaxy (MBE) or Migration Enhanced Epitaxy (MEE), low dimensional structures with high crystalline quality has been achieved.

In this chapter some basic properties of II-VI compounds are discussed. The struc-tural and optical properties are given, followed by an introduction to the principal growth modes, MBE and MEE growth techniques. In the last part, structural defects that can be found in the zincblende crystal structures are shortly described.

1.1 II-VI compounds based on ZnSe

1.1.1 Crystal structure

The II−VI compounds crystallize in such a manner that the atom of one element is in the center of a regular tetrahedron, the vertices’s of which are occupied by atoms of the other element. In general, different crystal structures can be obtained: the zincblende (cubic type) wurtzite (hexagonal type) and rocksalt structure [5]. The zincblende struc-ture is, in general, the stable phase for the II-VI semiconductor compounds (ZnSe, ZnS, CdS, etc). The only exception is CdSe, which typically has wurtzite as stable phase. On substrates with zincblende structure, as GaAs, it has been demonstrated that CdSe can also be grown in zincblende configuration up to a few micrometers thickness (≈ 4 µm) [6]. During the present work only the zincblende configuration will be discussed.

The zincblende structures can be described, as well, in terms of a close-packing of equal spheres (atoms), held together by interatomic forces. It is considered as a stacking of atoms (Fig. 1.1 a), starting with a close-packed triangular lattice (A) as the first layer. The next layer is formed by placing an atom in the depressions left in the center of every other triangle in the first layer, thereby forming a second triangular layer (B), shifted with respect to the first. The third layer (C) is formed by placing the atoms directly above those interstices in the first that are not covered by the atoms from the second layer and then the sequence ABC is repeated. This leads to an ABC stacking sequence along the [111] direction. An important aspect of the zincblende geometrical arrangement is that it has a lower degree of symmetry than diamond. In a view of the

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Figure 1.1: a)ABC stacking sequence for zincblende structures b) Zincblende polarity and c) Non-primitive unit cell

.

absence of a center of symmetry, the h111i directions are polar axes. This means that the [111] and£111¤directions are different. This applies to the (111) and¡111¢planes as well. Figure 1.1 b shows a model of the zincblende structure illustrating the difference between (111) and¡111¢planes. The (111) planes consist of atoms of group II, while the ¡

111¢planes consist of atoms of group VI.

The unit cell of a zincblende crystal structure consists of two face-centered cubic (fcc) lattices, displaced from each other by one-quarter of the body diagonal. The metal atoms (Zn,Cd) are placed on one fcc lattice and non-metal atoms (S,Se) on the other fcc lattice. The conventional unit cell of a zincblende structure is shown in Fig. 1.1 c where a1 = a [100], a2 = a [010], a3 = a [001] are the translation vectors and a is the

lattice parameter. Each unit cell contains 4 metal atoms which occupy positions at: ri

= [0,0,0], [1

2,12,0], [12,0,12,], [0,12,12] and 4 non-metal atoms with positions at rj = [14,14,14],

[3

4,34,14], [34,14,34], [14,34,34].

The direct lattice, defined in real space, enables us to calculate the distances between atoms in a crystal and angles between the bonds connecting those atoms. The concept of reciprocal lattice was then introduced for a direct computing of the spacing between successive lattice planes in a crystal lattice and therefore provides appropriate space for the description of diffraction (see chapter 2). For a given set of the direct basis vectors in real space, a set of reciprocal lattice basis vectors b1, b2, b3 can be defined such that:

b1 = a2× a3 (a1× a2) · a3 , b2 = a3× a1 (a1× a2) · a3 , b3 = a1× a2 (a1 × a2) · a3 (1.1)

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1.1 II-VI compounds based on ZnSe

Thus, a1· b1 = 1, a1 · b2 = 0, etc. The reciprocal lattice for zincblende structures has a

body centered cubic (bcc) structure. Each lattice vector in the direct lattice is described by:

r = ua1+ va2+ wa3 (1.2)

with u, v, w being integers numbers. A reciprocal reciprocal lattice vector is given by:

g_hkl= hb1+ kb2+ lb3 (1.3)

where h, k, l are also integer numbers and known as Miller indices of a set of lattice planes. The scalar product of the lattice vector r and a reciprocal lattice vector g_hklis

g_hkl· r = uh + vk + wl = N, (1.4) where N is an integer. The reciprocal lattice vector g with Miller indices (h, k, l) is perpendicular to the plane with Miller indices (hkl) and the following relationship is valid:

ghkl =

1

dhkl

, (1.5)

where dhkl is the distance of adjacent planes of set (hkl). This means that the length of

a reciprocal lattice vector is equal to the inverse of the spacing between the correspon-ding lattice planes dhkl. As will be seen, the concept of the real and reciprocal space is

important for the interpretation of electron diffraction patterns [7].

1.1.2 Physical properties and material parameters

Table 1.1 summarizes some of the physical parameters of the II-VI binary semiconductor compounds used in this thesis. The precise knowledge of the lattice parameter of diffe-rent compounds plays an important role, especially in the process of epitaxial growth. Experimental values combined with theoretical calculations, established the values of lattice constants of II-VI binary compounds [6, 8, 9].

Moreover, by mixing two different metal atoms in the metal sublattice or two non-metal atoms in the non-non-metal sublattice,

• ternary alloys (e.g. ZnSSe, CdZnSe, ZnSeTe) are obtained.

Using different atoms in both sublattices,

• quaternary alloys (e.g. CdZnSSe, MgZnSSe)

are formed. The change in composition influences the lattice constant and the band gap energy of the formed crystal. The lattice constant of a ternary alloy AxB1−xC, aABC can

be expressed by a linear combination of the lattice constants of binary compounds: AC and BC (aAC and aBC) involved. This is known as Vegard’s law [10]:

aABC = x · aAC + (1 − x) · aBC (1.6)

An example is given here for CdxZn1−xSe and ZnSySe1−y (Fig. 1.2). Adding Cd (S)

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Figure 1.2: Linear variation of lattice constant with composition for CdxZn1−xSe and ZnSeyS1−y ternary compounds varying the Cd and S content, respectively.

Material ZnSe ZnS ZnTe CdSe CdS GaAs Lattice constant a at RT [nm] 0.5669 0.5409 0.6104 0.6077 5.811 0.56533 Lattice mismatch relative to GaAs f (%) -0.27 +4.32 -7.9 -7.4 -2.5

-Elast. const. c11(N/m2) 8.50 9.81 7.12 7.41 7.7 11.9

Elast. const. c12(N/m2) 5.02 6.27 4.07 4.52 5.4 5.38

Elast. const. c44(N/m2) 4.07 4.48 3.12 1.32 2.32(w) 5.95

Poisson ratio 0.371 0.39 0.363 0.381 0.412 0.311 Stacking fault energy (mJ/m2₎ _12.9 ₆ ₁₆ ₁₂ _- ₅₅

Thermal conductivity [W/(K cm)] 0.19 0.251 0.108 0.09 - 0.44 Band gap Egat 300 K [eV] 2.69 3.68 2.394 1.66? 2.5 1.42

Ex. binding energy EB

X[meV] 20.8 40 12.4 - - 4.2

Table 1.1: Some physical parameters of the II-VI compounds used in this thesis in com-parison with GaAs. The values are taken from the latest edition of theLandolt-B¨onstein

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varies linearly if x respectively y is varied from 0 to 1, which is confirmed also by experi-mental data [12, 8, 13]. To large extent, Vegard’s law can be also applied for quaternary compounds [14, 15, 16, 17].

A heteroepitaxial structure implies the growth of compounds with different lattice constants on different substrates. The physical parameter which assigns the difference between lattice constants of different materials is called lattice mismatch f . Due to the fact that all structures investigated are grown on GaAs(001) substrates the lattice mis-match for the II-VI/GaAs heterostructures is defined according to:

f = aGaAs− aII−V I aII−V I

, (1.7)

where aGaAs and aII−V I are the lattice constants of the substrate and lattice constant

of II-VI layers bulk material, respectively. As a consequence of this lattice mismatch, the epilayers are grown under strain, if no relaxation takes place. If f has a negative value, the epilayer is under compressive strain and for f positive it is under tensile

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strain. GaAs substrates are favorable because the lattice mismatch relative to II−VI compounds is relatively small, especially in the case of ZnSe (f = −0.27).

For a strained crystal structure a precise description of the elastic properties is ne-cessary. In this respect, the elastic moduli give a measure of the ratio of stress εkl to

Figure 1.3: The components of the stress tensor for a body crystal.

associated strain σij in dependence of the crystal, which can be express as:

σij =

X

cijklεkl (1.8)

where σij and εklare second order tensors, and consequently cijklmust be a fourth order

tensor. Fig. 1.3 explains the components of the stress tensor. The definition of the strain tensor is analogous. Since the σij and εkl are symmetric tensors, the Hooke’s tensor

cijkl is also symmetric. Therefore the elastic coefficients cijkl are often represented as a

contracted 6×6 matrix in notation cnmwhere the indicies m and n are indicies

correspon-ding to a pair of ij or kl indicies, e.g.: c1111= c11, c1122 = c12 and c2323 = c44. Moreover,

for cubic crystal systems, the crystal symmetry has to be taken into account. As a con-sequence, the elastic moduli can be reduced to only three independent constants: c11

modulus for axial compression, c44 shear modulus and c12 modulus for dilation

com-pression [18, 19]. The values of those constants for binary compounds described here are listed in table 1.1. In cubic materials, the ratio between transversal and longitudinal strain, called Poisson’s ratio, gives also useful information about elastic behavior of the crystal structure.

ν = c12 c11+ c12

. (1.9)

Another important physical parameter regarding elastic properties of crystalline ma-terial is the elastic strain energy, which defines the energy stored in a strained solid material equal to the work performed in deforming the solid from its unstrained state. Therefore, when the elastic strain energy exceeds a critical value, plastic deformation occurs. This gives rise to defect formation in the heterostructure, where stacking faults (see defects section) play an important role. Especially for the II-VI material system, the low energy for stacking fault formation and high dislocation mobility of II-VI com-pounds are some of the big disadvantages for device processing [9, 6, 8]. Furthermore,

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the fabrication of semiconductor heterostructures requires the stacking of layers with different band gap energies and refractive indices.

Figure 1.4: Variation of bandgap energy of ZnSSe depending on S content

Semiconductor alloys allow to vary the band gap energy and consequently the re-fractive index of the material depending on the composition. If the lattice constant of a compound shows a linear dependence over composition x, the variation of the band-bap energy with composition x is conventionally described by a quadratic dependence [21, 22]:

Eg(x) = x · EAC+ (1 − x)EBC − b · x · (1 − x) (1.10)

where b is the bowing parameter and EAC and EBC band gap energies of binary

com-pounds from which the ternary alloy was formed. For example, in the case of ZnSxSe1−x

b is 0.68 eV [9]. A graphical representation of the ZnSSe band gap variation over sulfur

concentration is shown in Fig. 1.4.

The electronic properties of semiconductors depend on their band structure. In gene-ral, the electronic band structure of a semiconductor has numerous bands with asy-mmetric shapes and several energy maxima and minima. For optical transitions in the visible spectrum, only a small part of the band structure has to be considered, due to two factors: first, optical transitions are direct transitions in which the momenta of the initial and final electronic state are essentially equal and second, the completely filled bands do not contribute directly to the optical transition in the frequency range of in-terest. In the particular case of II-VI compounds studied here, which are direct band gap semiconductors, it is reasonable to consider the energy bands as symmetric and parabolic in shape. In a simple model, the band structure contains one conduction band and three valence bands. The lowest valenceband is called spin-orbit band and the up-per are called heavy-hole and light-hole bands corresponding to their curvature.

ZnSe, CdSe, etc are semiconductor compounds with direct band gap, which means that the valence band maximum and the conduction band minimum occur at the same

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Figure 1.5: Lattice constant versus bandgap for binary compounds investigated in this thesis [20]. The range of wavelength for visible spectrum covered by this compounds is also depicted.

value of momentum ~·k = 0 of the reciprocal space, called Γ point. Since no momentum change is necessary during direct recombination, such a recombination process occurs readily, producing a photon with the energy difference of the initial electron and hole (photon energy equal to the the band gap energy). Therefore, direct bandgap semicon-ductors are efficient sources of light. Fig. 1.5 shows the distribution of bandgap energy versus lattice constant for the semiconductor compounds discussed.

The discussion regarding the optical properties of those compounds requires the understanding of a quasiparticle like the exciton. An exciton is a bound state of an electron and a hole, or in other words, a Coulomb potential correlated electron-hole pair. It is an elementary excitation, or a quasiparticle of a solid. The exciton Bohr radius

aB in the case of a direct semiconductor can be derived within the framework of the

effective mass approximation [23]. In this case, electron and hole are treated as particles of effective masses mvand mc(the valence and conduction band effective masses, which

are taken for simplicity to be isotropic) which interact through an attractive Coulomb interaction screened by the dielectric constant ε of a crystal. The mass of the exciton is described by the reduced mass formed by the effective masses of electron and hole:

mc,v = h2· · d2_E dk2 ¸−1 (1.11) 1 m∗ = 1 mc + 1 mv (1.12)

For ZnSe-based material, the resulting effective masses for electron and hole are high, in comparison with III-V semiconductors as GaAs, which leads to small values for the exciton Bohr radius aB . The exciton Bohr radius aBis given by :

aB =

4πε0~2ε

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Furthermore, this leads to the possibility of obtaining quantum structures (quantum wells or quantum dots) in the dimension range of exciton Bohr radius aB (ex. aCdSeB =

5.4 nm ). The binding energy for an exciton in the ground state follows the relationship [24]: EB X = − m∗_{· e}4 32π2_ε2 0~2ε2 (1.14)

where ε and ε0are the dielectric constant of the material and vacuum, respectively.

In the case of zincblende compounds, with tetrahedral configuration, metal atoms have sufficiently high ionization potentials and do not completely loose their electrons, but share them with the neighboring non-metal atoms. Since the elements of group VI have a stronger electronegativity than the elements from group II, a two-electron cloud shifts from atoms of group II to atoms of group VI. Therefore, in the II-VI compounds, the binding is partly ionic, partly covalent, and they exhibit properties typical for both types of binding [5]. Detailed studies regarding II-VI compounds can be found else-where [25, 26, 20, 27].

1.2 Epitaxial growth

Semiconductor separate confinement heterostructure (SCH) used for fabrication of light emitting devices (LED) or laser diodes (LD) require the deposition of thin layers of semi-conductor material with well defined composition and thickness and high crystalline perfection. During such epitaxial process, the semiconductor material is deposited on a seed crystal, the so-called substrate. If the substrate is made from the same material as the deposited layers, i.e., ZnSe layers grown on a ZnSe crystal, a homoepitaxial growth process is performed. Since it is very difficult to grow ZnSe bulk crystals (also costly!), the ZnSe layers and device growth is performed on foreign substrate material and then the growth process is heteroepitaxial. The most frequently used substrate material for ZnSe-based device growth is GaAs due to the availability and well known properties. Furthermore, the substrate crystal orientation determines the crystallographic orienta-tion of the deposited layers so, in particular here the [001] growth direcorienta-tion was chosen. The big advantage of these substrates is the relatively small lattice mismatch f between GaAs and ZnSe at room temperature: f = −0.27% (see Table 1.1).

1.2.1 Growth modes

The individual atomic processes which govern the film growth are in principal: conden-sation of the new material from the molecular beam followed immediately by redesorb-tion or diffusion along the surface. This diffusion process might lead to adsorpredesorb-tion, particularly at special sites like edges and kinks or the diffused particle may desorb. In all these processes, characteristic activation energies have to be overcome. During growth, interdiffusion is often an important process. Substrate and film atoms can in-terchange places and the film/substrate interface is smeared out. In the following phe-nomenological description of the growth process the quasi-equilibrium conditions are assumed. A simple formal distinction between the condition for the occurrence of the various growth modes can be made in terms of surface or interface energy γ, i.e the

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char-1.2 Epitaxial growth

acteristic free energy (per unit area) to create an additional piece of surface or interface.

Figure 1.6: Simplified picture of an island of a deposited film

The growth morphology of the epitaxial layers can be then determined from the balance between different surface energies involved: γsfor the substrate, γf for the film,

and γs/f for the surface/film interface. In the heteroepitaxial growth process, where

lattice mismatch plays an important role, the elastic strain energy accumulated with increasing thickness must also be considered. Since γ can also be interpreted as a force per unit length of boundary, force equilibrium where the substrate and, for example, a 3D island deposited touch the film (Fig. 1.6 ) requires:

γs= γs/f + γfcos(φ). (1.15)

In this way, three growth mechanisms can be distinguished: -Frank-van-der-Merwe [28, 29] or layer by layer mode - Stranski-Krastanov [30] or layer plus island growth mode, -Vollmer-Weber [31] or island growth mode.

Figure 1.7: Layer by layer epitaxial growth

In the Frank-van-der-Merwe growth mode the interaction between the substrate and layer atoms is stronger than that between neighboring layer atoms. Each new layer starts to grow only when the last one has been completed. Taking into account the surface energies defined above and the eq.(1.15), the following relationship is valid:

γf + γs/f ≤ γs independent of the thickness. Furthermore, in this case the strain energy

is considered almost zero. This means, that this kind of growth is applicable only for homoepitaxial growth or material systems where the lattice mismatch is extremely low. The mixed Stranski-Krastanov growth mode (layer plus island) can be explained by assuming a lattice mismatch between the deposited film and substrate (usually the la-ttice constant of a grown layer is by 2-4 % higher compared to the substrate). The lala-ttice

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Figure 1.8: Layer by layer plus island epitaxial growth.

of the film adjusts to the substrate lattice through elastic deformation. The transition from layer to island growth occurs when the elastic strain field exceeds the range of adhesion forces with the deposited material. Also, many other factors can contribute to this growth mode, such as the symmetry and orientation of the overlayers with respect to the substrate.

Figure 1.9: Layer by island epitaxial growth.

For the case of Volmer-Weber growth mode the interaction between the neighboring film atoms exceed the overlayer/substrate interaction: γf + γs/f ≥ γs. The island

de-posits always by a multilayer conglomerate of adsorbed atoms.

Nevertheless, these thermodynamic criteria are frequently non-valid, first because the growth takes place far-from equilibrium conditions, and second because the surface free energies are macroscopic values, and are therefore not well defined for atomic-scale objects. Kinetic effects are extremely important as they allow metastable structure formation. For this reasons, studies of growth in real time are crucial in order to fully appreciate the relevance of the different processes.

1.2.2 Molecular Beam Epitaxy

During the research on ZnSe-based structures for light emission in the short-wavelength range of the visible spectrum, it was found that the Molecular Beam Epitaxy (MBE) is the only practically relevant epitaxy method to obtain n- and in particular p-type doped material. Up to now, only one electrically pumped ZnSe-laser diode – operating at 77 K – was fabricated by a Metal Organic Chemical Vapour Deposition (MOCVD) epitaxy process [33]. During the MBE process, the elements that form the semiconductor crystal are evaporated from individual cells, with a very small orifice. Such a cell is known as an effusion (or Knudsen) cell. The evaporation process is carried out in an ultra-high-vacuum (UHV) chamber where the pressure is usually in the order of 10−11_{Torr. Thus it}

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1.2 Epitaxial growth

Figure 1.10: Schematic drawing of the MBE system [32]. The source material is coming from the Knudsen cells as molecular beams on the substrate. Here the crystal grows.

of the chamber – each cell produces a beam of material [34]. Those material beams are directed onto the surface of the heated substrate crystal, where a reaction between the different materials occurs and thus the semiconductor crystal grows. Since the whole process usually is carried out far away from the thermodynamical equilibrium, it is only controlled by the reaction kinetics between the participating elements. Hence, the growth process – and consequently the growth rate and composition – of the crystal can be controlled by the temperatures of the cells on the one hand, and the substrate crystal on the other hand. Typical growth rates are in the order of 300-600 nm/h in the case of ZnSe growth. Each cell is equipped with a shutter which enables the abrupt turn on and off of the molecular beam. Given a shutter movement time of 1 s, a layer thickness control on the order of atomic monolayers is possible. A more detailed and comprehensive description of MBE can be found in the literature [26, 25, 20].

The MBE system used at University of Bremen consists of two connected EPI930 growth chambers: one for III-V and one for II-VI material systems. The II-VI cham-ber is equipped among other with Zn, Cd, S and Se with very high purity (6N) and the doping is performed with ZnCl2 and N-plasma source for n- respectively p-doped

semiconductors. To monitor the surface and crystal quality, during the growth, light and electron based techniques are used: Reflection High Energy Electron Diffraction (RHEED), Ellipsometry and X-ray Photospectroscopy (XPS).

1.2.3 Migration Enhanced Epitaxy

In addition to the conventional MBE growth mode, techniques based on pulse-mode supply of the reactant species to the growing epilayer surface play an important role. By this method, an improvement of grown surface quality can be achieved. Here the

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in-terest is focused on the Migration Enhanced Epitaxy (MEE) growth technique. With this method, monolayers of pure constituent elements are alternatively deposited onto the substrate, produced by molecular beam bunches from effusion cells. As a consequence, the film grows stepwise at relatively low growth rate [25].

Figure 1.11: Analogy of MBE (a) with MEE (b) growth mode for ZnSe. The dash lines delimit the beginning and the end of the MEE cycle.

The principal idea of the MEE growth technique and the comparison with the MBE is shown in Fig. 1.11 for ZnSe compound. If in the case of MBE growth, the Zn and Se atoms impinge simultaneously on the surface of the substrate and the ZnSe compound is formed, in the case of MEE growth, the Zn and Se atoms are supplied alternatively onto the surface. Moreover, between the Zn and Se supply a short break is provided, as it it seen in Fig. 1.11. A succession of Zn and Se deposition including the break is a MEE cycle. When in one cycle a monolayer (ML) is deposited, then so-called Atomic Layer Epitaxy (ALE) is performed. For II-VI compounds, in general, at temperatures above 220 the maximun growth rate that can be obtained is in the order of 0.5 ML/MEE cycle [35].

The film thickness is determined by the total number of cycles of pulses than rather the beam intensity or source temperature. In this way, at least one complete coverage of a reactant component is formed on a substrate. The experimental results [36, 37] indicate flat growth surfaces, which is one of the most important characteristic of MEE. For a II-VI material system, another very important effect is the rapid surface migration of the atoms during growth interruption. Because of the enhanced migration of the surface ad-atoms, this technique can be applied to lattice-mismatched systems for the realization of self-organized dots controlled by a modulated MEE deposition sequence. This was realized for InAs/GaAs system [38, 39] and also applied successfully for the growth of Cdse/Zn(S)Se self-organized quantum dots in Bremen [25, 27].

1.3 Defects in crystal structure

In general, real crystals are not perfect, they always contain imperfections. An imper-fection or ”fault” in the regular periodic arrangement of atoms in a crystal is called a

defect. Crystal lattice defects (defects in short) are usually classified according to their

dimensions: Fig. 1.12 illustrates some of the 0, 1, 2 and 3 dimensional defects that can be present in a crystal lattice.

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1.3 Defects in crystal structure

Figure 1.12:a) Interstitial impurity atom, b) Edge dislocation, c) Self interstitial atom, d) Vacancy, e) Precipitate of impurity atoms, f) Vacancy type dislocation loop, g) Interstitial type dislocation loop, h) Substitional impurity atom (taken from www.tf.uni-kiel.de)

• 0-dimensional defects: point defects (vacancies and interstitial atoms).

A vacancy is formed by the removal of an atom from an atomic site (Fig.1.12 d) and an interstitial by the the introduction of an atom into a non-atomic site (Fig.1.12 c). They can appear in crystalline materials by plastic deformation, high energy parti-cle irradiation or by rapid quenching from high temperature. Moreover, impurity atoms are also point defects which can be placed in different types of sites:

substi-tutional (one atom of the parent lattice is replaced by the impurity atom) (Fig.1.12

h) and interstitial (impurity atom located at a non-atomic site) (Fig.1.12 a) For ex-ample, in ionic crystals, point defects are charged and come in pair to maintain charge neutrality. Therefore, two types of defects can be found: Frenkel defects (charge interstitial-vacancy pair), and Schottky defects (differently charge pairs of vacancies) (not shown here).

• 1-dimensional defects: all types of dislocations: perfect dislocations (Fig.1.12 b),

partial dislocations (always in connection with a stacking fault), dislocation loops (Fig.1.12 f, g), grain boundary and phase boundary dislocations. Due to greater importance of linear defects in crystal structures, a detailed description will be given in the next paragraph.

• 2-dimensional defects as: stacking faults, grain boundaries for crystals of one

mate-rial or phase; phase boundaries and a few special defects as example boundaries between ordered domains. These defects will be also discussed in combination with linear defects.

• 3-dimensional defects: Precipitates (Fig.1.12 e), usually involving impurity atoms

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be filled with a gas.

Defects, even in very small concentrations, can have a dramatic impact on the properties of a semiconductor material. However, the interest is focused here on crystal structures grown using epitaxial methods as MBE or MEE. During the process of semiconductor crystal growth, the formation of such defects is unavoidable, due to the relative differ-ences in physical properties between crystal semiconductors involved and device im-perfections. Furthermore, the existence of one type of defects can lead to the appearance of another type, it is an interconnected process.

Furthermore, they are influencing the structural, optical and electrical properties of the final semiconductor heterostructure depending on the type and density of the defects. Several investigation techniques are suitable to study such defects. Since they have the dimensions in the range of the atomic scale, one of the most powerful tool used for crystal defect investigations is transmission electron microscopy. Using different imaging techniques, particularly one- and two-dimensional defects can be visualized and analyzed.

1.3.1 Dislocations and stacking faults

1.3.1.1 Dislocations

Dislocations are linear defects around which some of the atoms of the crystal lattice are misaligned. The most useful definition of a dislocation is given in terms of the

Burgers circuit. A Burgers circuit is an enclose clockwise circuit around the core of a

dislocation, going from lattice point to lattice point. Such a path (ABCD) is illustrated in Fig. 1.13. If the same atom to atom sequence is made in a dislocation free crystal and the circuit does not close then the first circuit, Fig. 1.13 a), must enclose one or more dislocations. The special vector needed for closing the circuit in the reference crystal is

Figure 1.13: Burger circuit in a crystal which contains an:a) edge dislocation, b) screw dislocation (taken from www.tf.uni-kiel.de)

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by definition the Burgers vector (b) (Fig. 1.13 b). Moreover, the presence of dislocations in a crystal leads to lattice distortion. The line along which the lattice is distorted is called dislocation line and can be described by a vector l. There are two basic types of dislocations: the edge dislocations and the screw dislocations. Edge dislocations are caused by the termination of a plane of atoms inside the crystal. In such a case, the adjacent planes are not straight, but instead bend around the edge of the terminating plane so that the crystal structure is perfectly ordered on either side (Fig. 1.13a). Thus, the Burgers vector b is perpendicular to the dislocation line l. Edge dislocations could (in principle) be generated by the agglomeration of point defects: self-interstitial on the extra half-plane, or vacancies on the missing half-plane.

The screw dislocations are more difficult to visualize, but basically comprise a structure

in which a helical path is traced around the linear defect (dislocation line) by the atomic planes of atoms in the crystal lattice (Fig. 1.13). The Burgers vector b is in the case of the screw type parallel to the line of the dislocation l.

Furthermore, in the most general cases, the dislocations situated in real crystalline solids are rarely of a pure edge or pure screw nature, rather they exhibit aspects of both types, therefore are termed ”mixed” dislocations, which means that the Burgers vector can have particular orientations.

Since, by definition, the Burgers circuit is an atom to atom movement, the closure failure must be between two atom sites in perfect crystal and will be a lattice vector. In semiconductor materials, b is aligned with close-packed crystallographic directions and its magnitude is equivalent to one interatomic spacing. Moreover, a dislocation defined in this way is a perfect or unit dislocation. Another important consideration regarding dislocations is that dislocation lines can end at the surface of the crystal and at grain boundaries, but never inside the crystal. Thus, dislocations must either form closed loops or branch into other dislocations. When three or more dislocations meet at a point (node), it is a necessary condition that the sum of the Burgers vectors equals zero.

A dislocation also introduces a certain amount of elastic energy into the crystal per unit length of the dislocation line [40]. The elastic energy, for screw dislocation, can be calculated taking into account the work done in displacing the faces of a crystal of a unit length by a distance b along the slip direction relative to each other and it is:

Escdis = µb2 4πln µ r1 r0 ¶ (1.16) where r1 is the external diameter of the crystal and µ the shear modulus. In the same

way, the elastic strain energy for an edge dislocation can be derived:

Eeddis = µb2 4π(1 − ν)ln µ r1 r0 ¶ (1.17)

with ν the Poisson ratio. In the same way for mixed edge-screw dislocation,taking into account that the Burgers vector b of a mixed type can be decomposed into a edge compo-nent be = bsinθ and a screw component bs = bcosθ, the strain energy can be calculated.

Moreover, from these expressions, it can be seen that the elastic energy is relatively in-sensitive to the dislocation character, therefore, taking realistic values for r1 and r0, a

general equation can be written approximately as:

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where α ≈ 0.5-1. This leads to a very simple rule (Frank rule) for determining whether or not a dislocation reaction will occur. In general, a dislocation of strength b will tend to dissociate into dislocations b1and b2when E1+ E2 < E or b2 > b21+ b22.

Figure 1.14: a) Dislocation glide b) Dislocation climb

Dislocations can move if the atoms from one of the surrounding planes break their bonds and re-bond with the atoms at the terminating edge. There are two basic types of dislocation movement: gliding and climbing (see Fig. 1.14). The dashed line shows the glide plane. The Gliding process consist of conservative movement of dislocations, on the plane defined by the Burgers vector and the dislocation line direction. The move-ment takes place without a transport of matter and it is particularly effective for de-formation propagation. The problem can be approach considering the concept of slip. When the plastic deformation occurs in a crystal then the atoms of one plane slide over another plane of atoms and this planes are called slip planes. The dashed line shows in the Fig. 1.14 the glide plane. Normally, the slip plane is the plane with the highest density of atoms and the direction of slip is the direction in the slip plane in which the atoms are most closely spaced. Therefore in f cc and zincblende crystals respectively, the slip take place on {111} planes in h110i directions and these two form the so-called

slip system.

At low temperatures where diffusion is difficult, and in the absence of a non-equili-brium concentration of point defects, the movement of dislocations is restricted almost entirely to glide. However, at higher temperatures, an edge dislocation can move out of its slip plane by climbing. The process of climbing requires movement in a different plane than the glide plane of a dislocation, and usually is accompanied by the pre-sence of stacking faults. To move a dislocation apart from its slip surface requires a movement of the atoms at a long distance (not conservative process). In this climbing process, diffusion of the vacancies or interstitial atoms can take place. However, this process is thermally activated and usually appears at high temperatures. The most common climb processes involve the diffusion of vacancies either towards or away from the dislocation.

Stacking faults

In section 1.1.1 it was emphasized that the perfect lattices can be described as a stack of identical atom layers arranged in a regular sequence. A stacking fault is a local region

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in the crystal where the regular sequence has been interrupted and consists of one or

Figure 1.15: a) Intrinsic stacking fault where the fault sequence is ABCABABCA b) Extrinsic stacking fault where the fault sequence is ABCABACABCA (taken from www.tf.uni-kiel.de)

two layer interruption in the stacking sequence. Here the interest is focused on close packed structures. For example, in f cc crystals (see Fig. 1.15) where the stacking se-quence is ABCABC along h111i direction, the interruption can be produced by vacancy agglomeration with the fault sequence ABCABABCA (intrinsic stacking faults) or by interstitial agglomeration with the fault sequence ABCABACABCA (extrinsic stacking faults). Stacking faults have a characteristic energy per unit area called stacking fault

energy. Furthermore, due to the additional lattice discontinuity associated with an

ex-trinsic fault it is expected to have a higher stacking fault energy than an intrinsec fault. The stacking faults which occurs in zincblende crystals are described in the section be-low.

1.3.2 Dislocations and stacking faults in zincblende structures

The deformation behavior of the zincblende structure is closely related to the structure of dislocations. The principal lattice vectors, and therefore the most likely Burgers vec-tors for dislocations in face-centered cubic structures are of the type a

2 [110] and a [001]

and the dislocations are called perfect dislocations. Then, due to the fact that the energy of dislocation is proportional to the square of its Burgers vector, b2_{, the dislocations with}

Burgers vector of a

2 [110] are more energetically favorable.

Furthermore, when a stacking fault ends inside the crystal, the boundary in the plane of the fault, separating the faulted region from the perfect region of the crystal, is a

partial dislocation (displacement less than a unit lattice vector). Two important partial

dislocations have been recognized in f cc crystal structures: Shockley partial (associated with the slip) and Frank partial (associated with the climb).

An explicit understanding of the different types of dislocations and their correspon-ding Burgers vectors in zincblende crystal structures is given by the construction of the Thompson tetrahedron (Fig. 1.16). As mentioned before, the glide planes for dis-locations in an f cc lattice are the four {111} planes and the corresponding h110i slip directions.

The Thompson tetrahedron is simply formed by the {111} planes with consistently indexed planes and edges. By convention the vertices’s are enumerated A,B,C and D and the centers of their triangles by α, β, γ, δ. The lines between vertices’s are h110i directions, they are used to define the Burgers vector of the perfect dislocations. The faces are {111} planes and they show the positions of the potential stacking faults.

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Figure 1.16: The crystal orientations in the opened unfold Thompson tetrahedron. The thicker and thinner arrows can be interpreted as Burgers vectors of perfect or partial dislocations, respectively.

The vectors AB, AD, DC etc represent the Burgers vectors of the type 1

2ah110i, which

are primitive translation vectors in the f cc structure. In consequence, dislocations which have these Burger vectors are called perfect dislocations. In terms of moving a perfect dislocation 1

2ah110i, it can be shown (using close-packed model) that it is energetically

favorable to dissociate into two partial dislocations according to the reaction of the type:

AB = Aδ + δB (1.19) or a 2[−110] = a 6[−12 − 1] + a 2[−211]. (1.20)

These dislocations are Shockley partial type and can be visualized by the vectors that point from the corner to the center of a face in Thompson tetrahedron representation. The corresponding Burgers vector is equal to b = a

6h112i.

In general the dissociation of a perfect dislocations is independent of the character (edge, screw, mixed) of the dislocation. Geometrically the Frank partial dislocation is formed by inserting or removing one close-packed layer of atoms. The Burgers vector is normal to the (111) plane of the fault and is equal to the change in spacing produced by one close packed layer, i.e. b = (a

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and since the Burger vector is not contained in a close-packet plane it can move only by climb.

Burger vectors Burgers vectors Dislocation type Thompson notation in crystallographic notation

AD a1

2h011i Perfect dislocation

Aβ a1

6h112i Shockley partial

A α a1

3h111i Frank partial

βα a1

6h110i Stair-Rod dislocation

Table 1.2:Different types of dislocations and corresponding Burgers vectors inf cc crys-tal structures, whereaindicates the lattice constant of the material.

Strain hardening of the semiconductors can be attributed to the progressive intro-duction during straining of barriers to the free movement of dislocations. One specific barrier observed in f cc crystals is the Lomer Cotrell dislocation. Using Thompson’s no-tation, the reactions leading to the Lomer Cotrell dislocation are as follows: The perfect dislocations DA(β) and BD (α) dissociate in the β and α planes, respectively, to form Shockley partial dislocations.

DA(β) → Dβ + βA (1.21)

BD(α) → Bα + αD (1.22)

Two Shockley partial dislocations react to form the new partial dislocation

αD + Dβ → αβ (1.23)

In general the partial dislocations that have the dislocation lines represented by the edges of the Thompson tetrahedron and their Burgers vectors by the line joining the centers of the faces of the tetrahedron (βα, αβ, γα etc.) are called stair-rod dislocations.

1.3.3 Defects in heterostructures

ZnSe-based heterostructures investigated here are grown on GaAs substrates as it is described in section 1.2.2. Therefore, a description of defects that are formed due to the lattice mismatch present at the interface between the substrate and II-VI epilayer (e.g. ZnSe/GAs) or different epilayers (e.g. CdSe/ZnSe) is necessary. This is in principal connected with the degradation behavior of ZnSe-based devices under current injection. It has been shown that even a small lattice mismatch present at interfaces, gives rise to defect formation inside the epitaxially grown crystal. In the case of ZnSe on GaAs, the latter has the smaller lattice constant. Thus, the ZnSe layer is compressively strained, as indicated in Fig. 1.17. The same situation is valid for CdSe and ZnTe, where the lattice mismatch regarding GaAs is higher (see Table 1.1).

In the opposite case, for instance ZnS, the layer is under tensile strain. Moreover, when the lateral lattice constant of the ZnSe layer equals that of the GaAs substrate and the layer is elongated along the growth direction-then the layer is fully strained

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growth direction _GaAs ZnSe pseudomorphic GaAs ZnSe relaxed critical critical d < d d > d

Figure 1.17: Schematic representation of pseudomorphic and relaxed ZnSe grown on GaAs. Relaxation occurs when the layer thicknessdexceeds the critical thicknessdcritical

.

or pseudomorphic. With increasing layer thickness, the strain energy accumulates, and as soon as the critical thickness (300 nm) [41] is exceeded, the strain is released via the generation of dislocations. Accordingly, the lattice constant of the layer relaxes to its bulk value. Finally, the fully relaxed (layer thicknesses beyond 1000 nm) [41] case is reached and the whole layer has the ZnSe lattice constant, as schematically shown on the right-hand side of Fig. 1.17. The critical thickness does not only depend on the lattice mismatch, but also on the substrate preparation and the growth start procedure.

Figure 1.18: Schematic representation of the nucleation of misfit dislocation networks (D) in the quantum well by a stacking fault (S) and its associated threading dislocations (T). The network is generated under current injection [42].

The different types of crystalline defects that appear in ZnSe-based heterostructures and their influence on the degradation of the processed devices has been studied

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inten-1.3 Defects in crystal structure

sively (e.g. [43, 42, 44, 26]). Two different types of defects are mainly important in this context:

• stacking faults • misfit dislocations.

Stacking faults (SF) usually nucleate at the GaAs/ZnSe heterointerface. They are di-rectly related to the substrate surface morphology, which is affected by the substrate preparation, the deoxidation process, and the growth start procedure. Such a SF is bound by partial dislocations, which have a Burgers vector b = 1/3h111i in the case of a Frank-type SF, and a Burgers vector b = 1/6h121i for a Shockley-type SF [43]. These SF propagate upwards during growth and form a V-shaped defect in cross section, as shown in Fig. 1.18. In some situations, the partial dislocations of a SF can react and form a perfect 60◦_{-dislocation that threads upwards. Such a dissociation process often}

occurs at a layer interface [43, 42]. The lifetime of a light emitting device depends on the density of stacking faults in the structure, since they are electrically active and act as non-radiative recombination centers [43].

Another negative feature of SFs is that they act as nucleation centers for misfit dis-locations (MD). Misfit disdis-locations are generated, when the layer and the substrate do not have the same lattice constant. For ZnSe-based devices, this is especially relevant in the case of Cd-containing quantum wells. When a SF (or its associated threading dislocation) intersects the highly strained quantum well, the dislocation can dissociate and form a MD in the quantum well. Under current injection, these MD can multiply and finally form a network of dislocations as indicated in Fig. 1.18 [42]. These networks are responsible for the formation of the so-called dark defects [43]. The defects which are correlated with the Cd-based quantum well degradation will be discussed in chapter 3.

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Chapter 2 Experimental techniques

The interaction of the accelerated electrons with the crystal atoms offers the great po-ssibility of obtaining useful information about crystalline materials [45, 46]. Using an assemble of electromagnetic lenses, a real image of the structure can be formed and magnified. With the CM20 UT electron microscope used in this work, point resolution of 0.19 nm can be obtained. This can be used for the determination of crystal properties with an atomic resolution. Therefore, Transmission Electron Microscopy (TEM) inves-tigation methods were mainly used in the frame of this work, providing information about crystal quality and chemical composition which can be then correlated with epi-taxial parameters.

This chapter starts with a short description of electron diffraction theory (2.2). Then, for a proper interpretation of TEM micrographs, the process of image formation is shortly explained (2.3) and the different imaging techniques employed for TEM inves-tigations described (2.4). Moreover, due to the fact that the electrons are absorbed in the specimen, thicknesses of the sample less then 200 nm have to be used. There-fore, appropriate sample preparation techniques, pointing out the importance of the specimen preparation for TEM (2.5), are also described. For extensive and descrip-tive presentations of electron microscopy several standard books are available (e.g.: [47, 48, 49, 50, 51]).

The complexity of the semiconductor structure investigated, requires the compari-son with other investigation methods. Therefore, TEM findings were compared with X-ray diffraction and Grazing incidence X-ray diffraction results (crystal structure in-formation as well as Photoluminescence and Electroluminescence. A brief description of these methods is also given (2.6).

2.1 Basic principle of an electron microscope

During this work, the analyzes of the II-VI semiconductor heterostructures were per-formed using a CM 20 UT (Ultra twin) Philips microscope. The essential components of this electron microscope (and in general) are: the illumination system, the objective lens and the imaging system (Fig. 2.1).

The illumination system comprises the electron gun and the system of condenser lenses. Its task is the formation of an electron beam and focusing it onto the speci-men. The entire electron gun is a triode, where the LaB6 source is used as the cathode.

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Figure 2.1: Optical column of the electron microscope

In addition to this cathode, there is a Wehnelt cylinder, and an anode at earth potential with a hole in its center. The cathode is attached to the high tension cable, which it is connected to the high-voltage power supply. When the LaB6 source is heated to high

enough temperature, the electrons have enough energy to overcome the natural barrier (“work function“ (Φ)). The cathode has a negative potential (-200 kV) with respect to the anode. The electrons are accelerated towards the anode and pass it through the anode hole, with a velocity greater than half the speed of light.

To get a controllable beam of electrons through the hole in the anode and into the microscope, a small negative bias is applied to the Wehnelt cylinder. The electrons com-ing from the cathode are converged to a point called gun crossover located between the Wehnelt cylinder and the anode. Therefore, the Wehnelt acts as a simple electrostatic lens: the first lens in the microscope. The gun crossover is considered as the object for the condenser system.

The condenser system consists of two magnetic lenses: the first condenser lens C1

and the second condenser lens C2, and an aperture. This can operate in two principal

modes: parallel beam and convergent beam. The first mode is used for TEM imaging and diffraction, while the second is used for scanning (STEM) imaging, microanalysis

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2.1 Basic principle of an electron microscope

and microdiffraction. The C1 lens is located close to the anode and forms a

demagni-fied image of the gun crossover, with a diameter of 10-100 µm, depending on the lens excitation. This is called spot size (smallest illuminated area).

In order to focus the beam onto the specimen a second condenser lens (C2) is used. The

C2uses the C1crossover as the object. This lens has a variable current control and hence

variable focal length. As a result, C2 affects the convergence of the beam and also the

diameter of the illuminated area on the specimen. When a parallel beam is required (high-resolution observations), C2 is adjusted in a way to produce an underfocused

image of the C1 crossover. The C2 condenser aperture (situated below) controls the

fraction of the beam which is allowed to hit the specimen. It therefore helps to control the intensity of illumination and the width of the illuminated area. In conclusion, the condensor system prepares the electron beam for the interaction with the specimen.

Parameter Value High-tension 200 keV de Broglie wavelength 2.51 pm Spherical abberation Cs 0.5 mm Chromatic abberation Cc 1 mm Point resolution 0.19 nm Beam convergence θc 1.3-2.5 mrad

Table 2.1: Important parameters of CM20 UT electron microscope used during this work

Objective lens The most important lens in an electron microscope is the objective lens

(OL). The CM20 is equipped with an ultra twin (UT) objective lens, which is constructed as a symmetric lens. In the symmetric lens design, the electron experiences nearly half of the objective lens magnetic field before reaching the specimen, and this upper half of the field is known as objective condensor lens. Responsible for this is the upper polepiece of the objective lens. It can be used to demagnify further the incident probe, down to the nanometer size range.

The first function of the objective lens is to bring the various diffracted electron beams, produced after the interaction of the electron beam with the specimen, to a crossover while introducing minimal aberrations and second function is to form the first intermediate image of the specimen in the image plane. The influence of the lens aberrations will be discussed in the section devoted to theory of imaging formation. Then, this image is subsequently magnified by the intermediate lens and the projector lens system.

The objective lens is characterized by two different planes: the back focal plane (BFP) and image plane. In the back focal plane of the objective lens the diffraction pattern (DP) is formed. In this plane the objective aperture (OA) can be inserted. Its function is to select electrons which are diffracted by specific lattice planes to contribute to the image, and thereby affects the appearance of the image, as the contrast of the final image. By inserting the aperture and tilting the beam, different types of images can be formed (see Sec. 2.4). To form diffraction patterns from small areas of the specimen, a “selected area“ aperture is used. This operation is called selected-area diffraction and consists of inserting

Transmission electron microscopy investigationsof the CdSe based quantum structures

Transmission electron microscopy investigations

of the CdSe based quantum structures

Elena Roventa

University of Bremen

2006

Transmission electron microscopy investigations

of the CdSe based quantum structures

Ph.D. thesis

to obtain the academic degree

Doctor in Natural Sciences

— Dr. rer. nat. —

of the Fachbereich 1

at the University of Bremen

submitted by

Dipl. Phys. Elena Roventa

Contents

Abbreviations

Introduction

Chapter 1

Theoretical background

1.1

II-VI compounds based on ZnSe

1.1.1

Crystal structure

1.1.2

Physical properties and material parameters

1.2

Epitaxial growth

1.2.1

Growth modes

1.2.2

Molecular Beam Epitaxy

1.2.3

Migration Enhanced Epitaxy

1.3

Defects in crystal structure

1.3.1

Dislocations and stacking faults

1.3.2

Dislocations and stacking faults in zincblende structures

1.3.3

Defects in heterostructures

Chapter 2

Experimental techniques

2.1

Basic principle of an electron microscope