Phase Transformations and Switching of Chalcogenide Phase-change Material Films Prepared by

Phase Transformations and Switching

of Chalcogenide Phase-change

Material Films Prepared by Pulsed

Laser Deposition

Der Fakultät für Physik und Geowissenschaften

der Universität Leipzig

eingereichte

D I S S E R T A T I O N

zur Erlangung des akademischen Grades

Doctor Rerum Naturalium

Dr. rer. nat.

vorgelegt

von M. Sc. Xinxing Sun

geboren am 23.12.1986 in Zhejiang/China

Gutachter: Prof. Dr. Dr. h.c. Bernd Rauschenbach (Leipzig)

Prof. Dr. Petr Němec (Pardubice/Czech Republic)

Tag der Verleihung: 13. März 2017

Bibliographische Beschreibung

Sun, Xinxing

Phase Transformations and Switching of Chalcogenide Phase-change Material Films

Prepared by Pulsed Laser Deposition

Universität Leipzig, Dissertation

171 S., 207 Lit., 76 Abb., 6 Tab.

Review:

The thesis deals with the preparation, characterization and, in particular, with the switching

properties of phase-change material (PCM) thin films. The films were deposited using the

Pulsed Laser Deposition (PLD) technique. Phase transformations in these films were

triggered by means of thermal annealing, laser pulses, and electrical pulses. The five major

physical aspects structure transformation, crystallization kinetics, topography, optical

properties, and electrical properties have been investigated using XRD, TEM, SEM, AFM,

DSC, UV-Vis spectroscopy, a custom-made nanosecond UV laser pump-probe system, in

situ resistance measurements, and conductive-AFM.

The systematic investigation of the ex situ thermally induced crystallization process of pure

stoichiometric GeTe films and O-incorporating GeTe films provides detailed information

on structure transformation, topography, crystallization kinetics, optical reflectivity and

electrical resistivity. The results reveal a significant improvement of the thermal stability in

PCM application for data storage. With the aim of reducing the switching energy

consumption and to enhance the optical reflectivity contrast by improving the quality of the

produced films, the growth of the GeTe films with simultaneous in situ thermal treatment

was investigated with respect to optimizing the film growth conditions, e.g. growth

temperature, substrate type.

For the investigation of the fast phase transformation process, GeTe films were irradiated

by ns UV laser pulses, tailoring various parameters such as pulse number, laser fluence,

pulse repetition rate, and film thickness. Additionally, the investigation focused on the

comparison of crystallization of GST thin films induced by either nano- or femtosecond

single laser pulse irradiation, used to attain a high data transfer rate and to improve the

understanding of the mechanisms of fast phase transformation.

Non-volatile optical multilevel switching in GeTe phase-change films was identified to be

feasible and accurately controllable at a timescale of nanoseconds, which is promising for

high speed and high storage density of optical memory devices. Moreover, correlating the

dynamics of the optical switching process and the structural information demonstrated not

only exactly how fast phase change processes take place, but also, importantly, allowed the

determination of the rapid kinetics of phase transformation on the microscopic scale.

In the next step, a new general concept for the combination of PCRAM and ReRAM was

developed. Bipolar electrical switching of PCM memory cells at the nanoscale can be

achieved and improvements of the performance in terms of RESET/SET operation voltage,

On/Off resistance ratio and cycling endurance are demonstrated. The original underlying

mechanism was verified by the Poole-Frenkel conduction model. The polarity-dependent

resistance switching processes can be visualized simultaneously by topography and current

images. The local microstructure on the nanoscale of such memory cells and the

corresponding local chemical composition were correlated.

The gained results contribute to meeting the key challenges of the current understanding

and of the development of PCMs for data storage applications, covering thin film

preparation, thermal stability, signal-to-noise ratio, switching energy, data transfer rate,

storage density, and scalability.

I

Table of Contents

CHAPTER 1 INTRODUCTION ........................................................................ 1

CHAPTER 2 FUNDAMENTALS: PHASE-CHANGE MATERIALS AND

THEIR APPLICATIONS ........................................................................................ 5

2.1 Principle of phase-change materials ....................................................................................... 5

2.2 Ge-Sb-Te phases ....................................................................................................................... 7 Pseudo-binary GeTe-Sb2Te3 tie-line ................................................................................. 7 2.2.1

Crystal structure of GST and GeTe ................................................................................... 9 2.2.2

Concise compendium of the crystallization process ....................................................... 12 2.2.3

2.3 Applications of phase-change materials................................................................................ 17 Optical data storage ......................................................................................................... 17 2.3.1

Electronic memory .......................................................................................................... 19 2.3.2

Frontier applications of phase-change materials ............................................................. 22 2.3.3

CHAPTER 3 EXPERIMENTAL CONDITIONS AND METHODS ........... 23

3.1 Pulsed laser deposition of thin films ...................................................................................... 23

3.2 Thermal treatment ................................................................................................................. 25 In situ thermal treatment in deposition chamber ............................................................. 25 3.2.1

Ex situ thermal treatment in an infra-red irradiation oven............................................... 25 3.2.2

3.3 Laser pulse irradiation for optical switching ....................................................................... 26 Nano- and femtosecond laser irradiation ......................................................................... 26 3.3.1

Nanosecond laser pump-probe system ............................................................................ 27 3.3.2

3.4 Arrangement for electrical switching ................................................................................... 28

3.5 Characterization methods ...................................................................................................... 30 X-ray diffraction-based analysis methods ....................................................................... 30 3.5.1

Electron microscopy........................................................................................................ 34 3.5.2

Atomic force microscopy ................................................................................................ 36 3.5.3

Spectroscopy and calorimetry methods ........................................................................... 38 3.5.4

CHAPTER 4 RESULTS AND DISCUSSION ................................................ 41

4.1 Thermal treatment-induced crystallization of GeTe films .................................................. 41 Ex situ thermally induced crystallization of GeTe films ................................................. 41 4.1.1

Effect of O incorporation on structural transitions of GeTe films ................................... 50 4.1.2

In situ thermally induced crystallization of GeTe films .................................................. 57 4.1.3

Epitaxial growth of GeTe films on BaF2(111) ................................................................ 63 4.1.4

Summary of the results on thermal treatment-induced crystallization of GeTe films ..... 67 4.1.5

II

4.2 Laser-induced phase transitions of GeTe and GST films .................................................... 69 Nanosecond laser-induced phase transition in GeTe films .............................................. 69 4.2.1

Crystallization of Ge2Sb2Te5 films by nano- and femtosecond single laser pulse 4.2.2

irradiation ........................................................................................................................ 80 Simulation of temperature distribution after pulse laser irradiation ................................ 89 4.2.3

Summary of the results on laser-induced phase transitions of GeTe and GST films ....... 93 4.2.4

4.3 Dynamic optical switching of non-volatile multi-level memory .......................................... 95 Optical switching ............................................................................................................. 96 4.3.1

Structural transformation ............................................................................................... 100 4.3.2

Morphology ................................................................................................................... 102 4.3.3

Dynamic phase transformation ...................................................................................... 106 4.3.4

Retention ........................................................................................................................ 109 4.3.5

Optical switching model ................................................................................................ 111 4.3.6

Film thickness dependence of numbers of multi-levels ................................................. 113 4.3.7

Broad wavelength response of reflectivity ..................................................................... 114 4.3.8

Summary of the results on dynamic optical switching of non-volatile multi-level 4.3.9

memory ......................................................................................................................... 115

4.4 Nanoscale bipolar electrical switching of phase-change materials ................................... 117 Background .................................................................................................................... 117 4.4.1

Polarity-dependent resistance switching of phase-change materials ............................. 123 4.4.2

Threshold of resistance switching .................................................................................. 129 4.4.3

Mechanism of bipolar resistance switching in phase-change materials ......................... 131 4.4.4

Summary of the results on nanoscale bipolar electrical switching of GST films .......... 143 4.4.5

CHAPTER 5 SUMMARY ............................................................................... 145

BIBLIOGRAPHY ................................................................................................ 149

LIST OF ABBREVIATIONS .............................................................................. 163

ACKNOWLEDGEMENTS ................................................................................. 165

CURRICULUM VITAE ...................................................................................... 167

LIST OF PUBLICATIONS ................................................................................. 169

SELBSTSTÄNDIGKEITSERKLÄRUNG ........................................................ 171

1

Chapter 1

Introduction

With the rapid growth of global networking, large data centers, and the portable

devices application-driven data storage market, the memory industry has ushered in

a comprehensive opportunity for diverse developments. The ideal new type of

memory is desirably characterized by superior all-round capabilities, including non-

volatility, long cycling life, small component size, low power consumption, fast

read/write processes, multilevel storage, 3D integration, and affordability of the

final product. One of the most promising areas in this respect is data storage

technology based on phase-change materials (PCMs) [1-3], which rely on the

reversible transformation between the disordered amorphous and ordered crystalline

states of PCMs induced by local heating/cooling either with laser pulses or

electrical pulses. The phase transformation is accompanied by a significant change

in optical reflectivity and electrical resistivity between these two solid states. The

pioneering work on the research field of PCMs was published in 1986 by

Ovshinsky [4], and considerable explorations have been made on the topic of

chalcogenides for data storage. Currently, chalcogenide-based PCMs are already

being applied in commercial optical storage media such as CDs, DVDs, and Blu-ray

disks. At the same time, the reversible electrical resistivity change of PCMs can be

applied in non-volatile electronic memory, which is considered to succeed FLASH

memory technology and is becoming the most promising candidate for future data

storage technology. However, the development of PCM-based data storage

technology gives rise to key challenges related to exploring the potential limits of

this technology, such as lowest switching energy, fastest switching speeds, largest

storage density and smallest cell sizes [5, 6].

A widely investigated chalcogenide material system in PCM based memory

technology is the Ge-Sb-Te system, particularly Ge2Sb2Te5 (GST). Regarding Ge-

Sb-Te thin films as the active layer of a memory device, magnetron sputtering is

typically employed for the thin film deposition. As an alternative to this physical

Chapter 1 Introduction

2

deposition method, pulsed laser deposition (PLD) provides several advantages: high

quality films, options for multi-compositions, quick and accurate control of the

deposition process, simple deposition system, possibility of scaling up the process

[7, 8]. To date, there exist only few studies dealing with the preparation of phase-

change films by PLD [9-12], and systematic investigations of the phase

transformation and the accompanying physical properties of PLD-prepared PCM

films are lacking. The promising PCMs possess unique properties such as high

crystallization speed, prominent optical reflectivity and electrical resistivity

contrast, excellent reversibility, large cycling numbers of reversible transitions and

high archival lifetimes of more than ten years [2]. With respect to guidance about

the optimal materials for data storage application, however, several crucial

questions exist: i) How can the thermal stability of PCMs be improved (so that they

are suitable, e.g. for use in automotive applications)? ii) How can the switching

energy be reduced? iii) What is the mechanism of phase transformation, particularly

in ultrafast switching? iv) How can the storage density be increased? v) What are

the electrical switching behavior and mechanisms when scaling down the memory

cells to the nanoscale? Undoubtedly, investigations to address the above issues are

of great importance for the ongoing development of phase-change data storage

technology.

Objectives and organization of this thesis

The thesis aims to explore the performance and mechanisms of PCMs with data

storage applications in view, including such key features as thermal stability, signal-

to-noise ratio, switching energy, data transfer rate, storage density, and scalability.

Related to this, in particular four major physical aspects, namely the crystallization

kinetics, structure transformation, optical and electrical switching have been

investigated. The main objectives are formulated as follows:

To investigate the influences of in situ or ex situ thermal treatment on

the structural transformation and the properties on the crystallization

process of PLD-deposited GeTe films, in order to optimize the thermal

stability and signal-to-noise ratio.

Chapter 1 Introduction

3

To examine laser pulse-induced phase transitions of GeTe and GST

films for attaining high date transfer rates and to improve the

understanding of rapid phase transformation kinetics.

To investigate the multi-level optical switching of single layer GeTe

films by using a custom-made nanosecond UV pump-probe system, in

order to increase the storage density and to reveal the mechanism of

rapid phase transformation.

To explore the nanoscale electrical switching and underlying mechanism

of GST films based device by using an arrangement of conductive

atomic force microscopy tips as a probe that enables write/read/erase

operations at the nanoscale.

To clarify correlations of structural phase transformation with respect to

optical or electrical property characterization by appropriate analysis.

Scope

In the thesis, Chapter 2 provides a brief introduction on the fundamentals of phase-

change materials and their application. A description of the experimental setup and

the characterization methods are presented in Chapter 3. The investigations on

various aspects of physical properties in PCMs for data storage applications and the

obtained results are presented and discussed in Chapter 4. In Chapter 4.1, the results

of thermal treatment-induced crystallization of GeTe films are presented. Chapter

4.2 deals with laser pulse-induced phase transitions of GeTe and GST films. Results

on dynamic optical switching of non-volatile multi-level memory cells are provided

in Chapter 4.3. Findings about nanoscale bipolar electrical switching of PCMs are

presented in Chapter 4.4. Finally, the contributions of the present work are

summarized in Chapter 5.

5

Chapter 2

Fundamentals: phase-change materials

and their applications

Principle of phase-change materials 2.1

Phase-change materials (PCMs) exist in two different stable states, crystalline and

amorphous, which exhibit distinctly different physical properties. The most

significant feature of this material class is the ability to repeatedly and rapidly

switch between these states through heating or cooling. The concept of using the

amorphous state of PCMs as logical ‘1’ and the crystalline state as logical ‘0’, as

well as the capability to switch between both states, has been extensively explored

for data storage applications. By utilizing the non-volatility of the optical

reflectivity switching property, PCMs have already been applied in commercial re-

writable optical storage media such as CDs, DVDs, and Blu-ray discs [2]. Recently,

based on the non-volatility of the large contrast in the electrical resistance switching

property, PCMs have been demonstrated to be promising for use in non-volatile

electronic memory devices [3, 13]. The basic principle of data storage based on

PCMs is exemplified in Figure 2.1.

The state in which the PCM is in the crystalline phase is characterized by a high

optical reflectivity and a low electrical resistance. For amorphization of the PCM,

an intense and short laser or electrical pulse is applied (Record/RESET operation).

This intense pulse provides the power to rapidly increase the temperature in the

material above the melting point Tm. Because of the very short pulse duration, high

cooling rates exceeding 109 K/s are achieved and thereby the material is quenched

into the disordered amorphous state. A moderately intense and long laser or

electrical pulse is used to crystallize the PCM (Erase/SET operation). The material

is locally heated above the crystallization temperature Tc, which lies between the

glass transition temperature Tg and the melting point Tm. The pulse-induced

Chapter 2 Fundamentals: phase-change materials and their applications

6

temperature increase causes increased mobility of atoms in the material and results

in an energetically favorable atomic arrangement in the form of the crystalline

phase. Because the differences in the optical and electrical properties between the

amorphous and crystalline states are significant, the respective state can thus be

easily read out by applying a very low intensity laser or electrical pulse.

The atomic arrangement in the material is schematically illustrated in Figure

2.1. In the crystalline phase, the atoms show a long-range order, whereas the liquid

state is characterized by complete disorder (exception: GeTe, see Ref. [14]). For the

amorphous phase, the arrangement of atoms is disordered, but has a locally limited,

short-range ordered structure extending over only a few atomic distances. The

understanding of the atomic configurations and bonding mechanisms provides an

explanation for the large contrast in optical and electrical properties between the

two phases [15]. Further details on the structural properties of Ge-Sb-Te phases in

the amorphous and crystalline states can be found in Chapter 2.2.2. It should be

noted that a high crystalline quality of the PCM in the crystalline state could be

beneficial for reduction of the required switching energy [16, 17].

In addition, it is important that the timescale of phase transitions in PCMs is

considered. For the amorphization process, a high quenching rate of more than

109

K/s is needed, because upon the quenching from a liquid into a “frozen” state,

the resulting structural configuration must be below the glass transition temperature

Tg, in order to avoid the occurrence of re-crystallization. The typical crystallization

Figure 2.1: Principle of reversible switching in phase-change materials for non-volatile memory

applications. Tm melting temperature, Tc crystallization temperature and t time.

Crystalline phase

Ordered structureHigh reflectivityLow resistance

Liquid phase

Disordered structure

Low reflectivityHigh resistance

Amorphous phase

Tm

t

Laser or current pulse

Tc Tm

t

Laser or current pulse

Tc

Erase/SET

Record/RESET

2.2 Ge-Sb-Te phases

7

process in PCMs takes place in the nanosecond regime. On the other hand,

crystallization of the amorphous region at room temperature should not happen

within the first 10 years, or even longer, in terms of data retention. This means that,

in terms of the timescale, the difference between crystallization time and retention

time spans roughly 17 orders of magnitude, whereas, regarding the temperature

scale, the corresponding temperature range spans over just a few hundred Kelvin.

Moreover, ultrafast switching behavior relying on non-thermal phase transitions in

GST or GeTe-Sb2Te3 superlattices, triggered by an ultrashort laser pulse, has been

proposed [17-19]. Nevertheless, the understanding of certain material properties and

physical mechanisms of PCMs, as well as developments based upon them, broadens

the perspective for promising future applications.

Ge-Sb-Te phases 2.2

Among PCMs, the Ge-Sb-Te system is the most interesting, in general. Here, a

description of the pseudo-binary GeTe-Sb2Te3 tie-line (Chapter 2.2.1), of the crystal

structure of GST and GeTe (Chapter 2.2.2), and concise compendium of the

crystallization process (Chapter 2.2.3) is given.

Pseudo-binary GeTe-Sb2Te3 tie-line 2.2.1

Figure 2.2 illustrates the ternary phase diagram for this material system with

intermetallic alloys that lie on the tie-line between GeTe and Sb2Te3. Along this

line, the compounds Ge2Sb2Te5, Ge1Sb2Te4 and Ge1Sb4Te7 can be found, in which

the properties change from high crystallization temperature (high stability) to low

crystallization temperature (low stability) [20]. An additional opportunity to tailor

the properties of the compounds on the GeTe-Sb2Te3 tie-line is to dope them with

elements like O, N, Si, etc. [21-23]. As a preliminary example, it was found that the

crystallization temperature Tc of 60 nm thick PLD- deposited GeTe films on Si

substrates could be increased from 240 to 280°C by incorporation of ~16 at.% O

(for details, see Chapter 4.1) [24]. PCMs with such an enhanced thermal stability

are excellent candidates for applications in various fields (i.e., automotive,

aerospace), where the use of conventional GST is limited by its relatively low Tc

(about 140°C) [25].


8

On the other hand, the materials with fast crystallization are in general of

interest, because a higher data transfer rate for data storage application can be

attained. Figure 2.2 also demonstrates the alteration of the crystallization time (in

terms of switching speed). It is known that the crystallization time can be reduced

from 50 ns for Ge2Sb2Te5 and 40 ns for Ge1Sb4Te7 to 30 ns for Ge1Sb2Te4 [20, 27].

However, stoichiometric GeTe was demonstrated to exhibit a short crystallization

time, but it is highly sensitive to deviations in the composition [28]. Even small

deviations lead to a dramatic increase of the crystallization time. The first reported

short crystallization time for GeTe is 30 ns with a high optical contrast [29];

recently, a value as low as 16 ns was found with a high electrical resistance contrast

[30].

In addition, Coombs et al. [31, 32] proposed that the compounds along the

GeTe-Sb2Te3 tie-line can be defined as materials with so-called nucleation-

dominated or so-called growth-dominated crystallization, depending on the material

properties (kinetics) during the crystallization process. According to these authors,

in a nucleation-dominated crystallization process, the majority of the total

crystallization time is taken by the nucleation process of crystallization while the

growth process of the crystallites is rather short. In contrast, in the growth-

dominated crystallization, nucleation takes place only in a very short time but the

Figure 2.2: Ge-Sb-Te ternary phase diagram illustrates the various phase change alloys.

Stoichiometric compositions lying on the pseudo-binary tie-line of GeTe and Sb2Te3 are presented

(adapted from Refs. [20, 26]).

2.2 Ge-Sb-Te phases

9

growth process takes the majority of the total crystallization time. In this respect,

the crystallization behavior in amorphous GST is, in general, characterized by a

nucleation-dominated growth process [33-35], whereas, in amorphous GeTe,

crystallization is, in general, determined by a growth-dominated growth process

[28, 30, 36]. It is also noteworthy that the interface of a complex PCM-based

memory cell structure can also play an important role with respect to the dominant

crystallization mechanism [30].

Regarding the appropriate phase-change material for the relevant application,

development of optical data storage devices is required in order to improve the

optical contrast for smaller laser wavelengths. Along the pseudo-binary line from

GeTe to Sb2Te3, the optical contrast has been shown to increase with increasing Ge

content at a wavelength of 405 nm [2]. This implies that, along the tie line, a higher

GeTe content is promising for high density optical recording applications. For

electronic memory applications, however, a good balance of the materials

properties, such as fast switching speed and high stability, is required. The most

promising and widely studied material proves to be Ge2Sb2Te5, which is located in

the middle of the tie-line in Figure 2.2. The structural properties of Ge2Sb2Te5 and

GeTe are described in the following.

Crystal structure of GST and GeTe 2.2.2

1.) Structure of GeTe

As shown in Figure 2.3a, the crystal structure of GeTe is a rhombohedrally distorted

rocksalt-type configuration (space group R3m) with Ge and Te atoms, which form

face-centered cubic (fcc) sublattices at room temperature. In this structure, the

lattice constant is a = 0.5985 nm, and a distortion angle of 88.37°, that deviates

from the usual 90° angle along the [111] direction, is present [37-39]. At

temperatures above 420°C, GeTe in the distorted rocksalt structure is transferred

into a higher symmetry rocksalt configuration. The mechanism of the distortion in

rhombohedral GeTe is connected with the Peierls effect [40]. This lattice distortion,

in conjunction with resonance bonding, determines most of the material properties

[41].

In the crystalline state, the arrangement of atoms exhibits an octahedral (six-

fold) coordination with three shorter (0.2843 nm) and three longer (0.3158 nm)


10

bond distances. Each atom possesses three valence electrons that strive to form six

covalent bonds with the nearest neighbors. However, the number of valence

electrons is insufficient to fulfill the covalent bonding state. The electrons are

delocalized in this resonantly bonded network, leading to a significant increase of

the electronic polarizability. In contrast, upon the structural transition from

crystalline to amorphous state, Ge atom positions exhibit switching from octahedral

to tetrahedral coordination through rupturing of the longer Ge-Te resonance bonds

and strengthening of shorter covalent bonds (as shown in Figure 2.4). The

amorphous GeTe structure can be represented by a random covalent network model

[2, 43]. The average Ge-Te bond length of ~0.260 nm is shorter than that in the

crystalline state [44]. Furthermore, the coordination number of the Ge and Te atoms

shows a four-fold and two-fold coordination in the amorphous state, which is

typically less than in the crystalline state.

2.) Structure of GST

GST possesses a metastable crystalline structure and a stable hexagonal structure.

The crystal structure of metastable crystalline GST can be described approximately

by a rocksalt structure (space group Fm-3m) [2, 45, 46]. Te atoms occupy one fcc

sub-lattice, whereas the other sub-lattice is randomly occupied by Ge and Sb atoms,

as well as by randomly arranged vacancies (~20%), as exemplarily shown in Figure

2.3b. In this structure, the lattice parameter is 0.602 nm. The transformation from

the amorphous to the metastable crystalline phase occurs above a threshold

temperature of 140°C and the phase transition temperature from the metastable

Figure 2.3: Crystal structure of rhombohedral GeTe (a) and cubic GST (b). The red, blue, green and

white spheres symbolize Ge, Sb, Te atoms and vacancies, respectively (adapted from Ref. [42]).

(b)(a)

2.2 Ge-Sb-Te phases

11

crystalline phase to the hexagonal phase is about 310°C [25]. It is noted that the

metastable crystalline phase shows the same type of Peierls-like distortion as in

rhombohedral GeTe. However, it is not surprising that the distorted rocksalt

structure has been found in the GeTe-Sb2Te3 pseudo-binary phase diagram with a

Ge/Sb site vacancy concentration that is proportional to (1-x)/(3-2x) for

(GeTe)x(Sb2Te3)1-x [47].

In the crystalline phase, Ge atoms exhibit an octahedral coordination, as

depicted in Figure 2.4. The crystalline to amorphous transformation process, with

breaking of the longer Ge-Te bonds and subsequent transition of the atoms from

octahedral to tetrahedral symmetry positions, is described in the framework of the

umbrella-flip model [1]. According to Liu et al. [48], 35% of the Ge atoms in the

cubic phase show a tetrahedral coordination, indicating the coexistence of

octahedral and tetrahedral sites for Ge in the GST lattice. The atomic structure of

the amorphous state is still under discussion. Nevertheless, the umbrella-flip model

is an intuitive explanation of ultrafast transition and provides a possible scenario for

the change of the bonding mechanism. Upon amorphization, the average Ge-Te and

Sb-Te bond lengths are 0.261 nm and 0.285 nm, respectively [49]. Moreover, in the

crystalline phase, GST exhibits resonance bonding, whereas in the amorphous phase

(covalent bonding), resonance bonding disappears, due to the loss of the long-range

order. In this amorphous structure (Figure 2.4, right), Ge atoms possess four nearest

Te neighbors in a tetrahedral arrangement and the positions of Sb and Te atoms

resemble those in the rocksalt structure [50].

Figure 2.4: Typical structural switching mechanism in the “umbrella flip-model” of Ge atoms in

crystalline and amorphous GST/GeTe (adapted from Ref. [1]).

Te

Ge


12

Concise compendium of the crystallization process 2.2.3

In this subchapter, the phase transition from the amorphous to the crystalline state is

shortly considered. The crystallization process in PCMs can be described with the

classical theory of crystallization, which involves the nucleation of crystal seeds

within an amorphous matrix and subsequent growth of the nuclei at the amorphous-

crystalline interface. Nucleation is the process whereby nuclei (seeds) act as

templates for crystal growth. Homogeneous nucleation occurs in the interior of the

parent phase without the involvement of a foreign substance. At temperatures below

a material’s melting point Tm, the driving force for solidification is the difference in

the Gibbs free energy ΔG between the liquid and the solid. If it is assumed that the

heat capacities of the liquid and solid are equal, then the molar enthalpy and molar

entropy of solidification will each remain constant as a function of temperature, and

ΔGhom can be calculated for a spherical seed of radius r as follows [51]

∆𝐺ℎ𝑜𝑚 = −4

3𝜋𝑟3∆𝐺𝑣 + 4𝜋𝑟2𝜎 , (2.1)

where ΔGv is the free energy of the seed crystal and σ is the surface energy. The

second term involves the increase in energy required to form a new surface. The

first term is negative and represents the decrease in Gibbs free energy upon

solidification. Because the second term is a function of the second power of the

radius, and the first term a function of the third power of the radius, the sum of the

two first increases, passes through a maximum, and then decreases as a function of

the seed radius.

The radius at which the Gibbs free energy curve is at maximum is called the

critical radius, r*, for a solid nucleus in a liquid environment

𝑟∗ = −2𝜎

∆𝐺𝑣 . (2.2)

The driving force of the Gibbs free energy will tends to cause a solid particle with a

smaller radius than r* to decrease in size. Thus, this is a particle of subcritical size

for nucleation. A viable nucleus is one with a radius larger than or equal to r*. The

critical Gibbs free energy corresponding to the radius r* is ΔG*. This term can be

shown to be

∆𝐺∗ =16𝜋

3

𝜎3

∆𝐺𝑣2 . (2.3)

2.2 Ge-Sb-Te phases

13

In practice, homogeneous nucleation is rarely encountered in solidification.

Instead, heterogeneous nucleation occurs at crevices, at interfaces, or at impurity

particles suspended in the liquid. The analysis of the energies involved in

heterogeneous nucleation follows the same method as that used for homogeneous

nucleation. In the case of a heterogeneous nucleus, the critical radius and the critical

Gibbs free energy are

𝑟ℎ𝑒𝑡∗ = −

2𝜎𝑠𝑙

∆𝐺𝑣 (2.4)

and

∆𝐺ℎ𝑒𝑡∗ =

16𝜋

3

𝜎𝑠𝑙3

∆𝐺𝑣2 ∙ 𝑓(𝜃) (2.5)

respectively, where σsl is the solid-liquid interface energy. It is important to note

that the critical radius of curvature, r*, does not change when the nucleation

becomes heterogeneous. The critical Gibbs free energy, ∆𝐺ℎ𝑒𝑡∗ , however, is strongly

influenced by the wetting that occurs at the surface of the material that catalyzes the

nucleation. A lower value of ∆𝐺ℎ𝑒𝑡∗ means a lower activation energy to be overcome

in nucleation; in other words, nucleation takes place more easily. The magnitude of

the effect can be appreciated by considering the values of f(θ). f(θ), the so-called

geometrical factor, is the ratio of the volume of the heterogeneous nucleus to the

volume of the sphere with the same radius of curvature.

In the solid-solid transformation (amorphous-crystalline transformation), a small

change in specific volume that occurs must be accommodated elastically, leading to

strain energy effects. Consequently, this strain energy must be involved, together

with the volume-free energy and the interfacial energy changes

∆𝐺(𝑟) = −4

3𝜋𝑟3∆𝐺𝑣 + 4𝜋𝑟2𝜎 + 𝑠𝑡𝑟𝑎𝑖𝑛 𝑒𝑛𝑒𝑟𝑔𝑦 . (2.6)

This strain energy can be calculated (see Ref. [52]) for a spherical seed as

4

3𝜋𝑟3∆𝐺𝐸. Then, in the case of homogenous nucleation, this contribution leads to

∆𝐺(𝑟) = −4

3𝜋𝑟3(∆𝐺𝑣 + ∆𝐺𝐸) + 4𝜋𝑟2𝜎 . (2.7)

Implications:

(a) The interfacial contribution to the nucleation barrier dominates for small seeds

and the volumetric contribution 4

3𝜋𝑟3(∆𝐺𝑣 + ∆𝐺𝐸) for large nucleus sizes. The


14

competition between these two contributions can produce a complicated

sequence of states for developing phases.

(b) Typically, for solid-state nucleation, the elastic energy contribution dominates

for small particles sizes. As the particle grows, elastic energy can be reduced by

the introduction of misfit dislocations into the interface. Such interfacial

dislocations transform a coherent interface into a semi-coherent interface at the

expense of increased interfacial energy.

(c) With exception of ideal spherical particles, this elastic strain energy is not easy

to express with simple algebraic expressions.

The rate of nucleation per unit volume and unit time, J, is proportional to

(1) the probability, 𝑤 = exp(−∆𝐺∗/𝑘𝐵𝑇), that a thermodynamic fluctuation of

the critical free energy ΔG* is given,

(2) the number of the growth species per unit volume, n, that can be used as

nucleation seed points, given by Co (in the simplest case of the homogenous

nucleation, Co is given by the initial concentration),

(3) the jump frequency of the growth species Γ from one site to another is given

by 𝛤 = 𝑘𝐵𝑇 3𝜋𝜆3𝜂⁄ , where λ is the diameter of the growth species and η is

the viscosity of the solution (for simplest case that solid precipitates are

formed in a solution).

Thus, the rate of nucleation can be described by

𝐽 = 𝑛 ∙ 𝑤 ∙ 𝛤 = 𝐶𝑜𝑘𝐵𝑇

3𝜋𝜆3𝜂∙ 𝑒𝑥𝑝 (−

∆𝐺∗

𝑘𝐵𝑇) . (2.8)

This equation indicates that a high initial concentration (supersaturation), i.e., a

large number of nucleation sites, a low viscosity, and a low critical energy barrier

favor the formation of a large number of nuclei. For a given concentration of solute,

a larger number of nuclei means smaller-sized nuclei.

The size distribution of the formed crystallites is dependent on the subsequent

growth process of the nuclei. The growth process of the nuclei involves several

processes. This multi-step process can be separated into (i) the generation of seed

points, (ii) the diffusion of the growth species, (iii) the adsorption of the growth

species onto the growth surface, and (iv) growth of the seeds through irreversible

incorporation of growth species. These steps can be further grouped into two

processes as follows. Supplying the growth species to the growth surface is termed

2.2 Ge-Sb-Te phases

15

diffusion, which includes the generation, diffusion, and adsorption of growth

species, whereas incorporation of growth species adsorbed on the growth surface

into solid structure is denoted as growth.

Growth controlled by diffusion

When the concentration of growth species declines below the minimum

concentration for nucleation, the nucleation process stops, whereas growth

continues. If the growth process is controlled by the diffusion of growth species

from the bulk solution to the particle surface, the growth rate is given by [53]

𝑑𝑟

𝑑𝑡=

𝐷(𝐶−𝐶𝑠)

𝑟𝑉𝑚 , (2.9)

where a spherical nucleus with the radius r is assumed, D is the diffusion constant

of the growth species, Cs is the concentration on the surface of the formed

crystalline precipitate, and Vm is the molar volume of nuclei. The solution of this

differential equation under the assumption that the initial size of the nucleus, ro, and

the change of the bulk concentration are negligible leads to

𝑟 = √2𝐷(𝐶 − 𝐶𝑠)𝑉𝑚𝑡 + 𝑟𝑜2 . (2.10)

For two crystalline precipitates with an initial radius difference of Δro the radius

difference Δr decreases as time increases (or as precipitates grow larger), according

to

∆𝑟 =𝑟𝑜

𝑟∆𝑟𝑜 . (2.11)

Combining Equations 2.10 and 2.11, the evolution of the radius as result of a

diffusion-controlled process is

∆𝑟 =𝑟𝑜

√2𝐷(𝐶−𝐶𝑠)𝑉𝑚𝑡+𝑟𝑜2

∆𝑟𝑜 . (2.12)

This indicates that the radius difference decreases with the increase of radius of the

precipitation and prolonged growth time. Consequently, diffusion-controlled growth

promotes the formation of uniform crystallites in an amorphous ambience.

Growth controlled by interface reaction

When the diffusion of growth species from (as here) the amorphous bulk material to

the growth surface of the crystalline precipitate is sufficiently fast, i.e., the

concentration on the surface is the same as that in the bulk, the growth rate is

controlled by the interface reaction process. The growth rate is thus proportional to

the surface area


16

𝑑𝑟

𝑑𝑡= 𝐾(𝐶) ∙ 𝑟2 , (2.13)

where K(C) is a proportionality constant that is dependent on the concentration of

the growth species. Solving this equation, the evolution of the crystalline precipitate

radius is given by

𝑟 =1

1𝑟𝑜

⁄ −𝐾(𝐶)𝑡 . (2.14)

The radius difference increases with an increasing radius of the nuclei according to

∆𝑟 =𝑟2

𝑟𝑜2 ∆𝑟𝑜 . (2.15)

Combining Equations 2.14 and 2.15 yields

∆𝑟 =𝛥𝑟𝑜

1−𝐾(𝐶)𝑡 . (2.16)

Equation 2.16 demonstrates that the radius difference increases with a prolonged

growth time, if an interface reaction-controlled process for the growth is assumed.

Consequently, this growth mechanism does not favor the synthesis of mono-sized

crystallites.

For the kinetics of phase transformation, an expression of the temperature

dependence of the growth velocity can be derived as [51, 54]

𝑢 ∝𝑇

𝜂∙ (1 − 𝑒𝑥𝑝 (−

∆𝐺

𝑘𝐵𝑇)) . (2.17)

This illustrates that a high growth velocity of phase transformation occurs for low

viscosity η and/or a high free energy ΔG to form a crystallite from the liquid state.

Self-evidently, the viscosity and the free energy are temperature dependent, as well.

Recently, J. Orava et al. investigated the kinetics of crystallization in GST by ultra-

fast DSC measurements [55, 56]. They demonstrated that the growth velocity gets

highest at an intermediate temperature between the glass transition temperature Tg

and the melting temperature Tm. At this intermediate temperature both viscosity and

free energy contribute significantly to the growth velocity. For temperatures slightly

above Tg the growth velocity is several orders of magnitude smaller than at the

above mentioned intermediate temperature, because the velocity is limited by the

still high viscosity [55]. For temperatures very close to Tm the growth velocity gets

again small. Consequently, in order to achieve the fastest crystallization, the laser

pulse induced or current pulse induced temperature rise must be optimized. In the

2.3 Applications of phase-change materials

17

case of data storage applications, this corresponds to attaining a high data transfer

rate.

Applications of phase-change materials 2.3

As mentioned above, PCMs can be rapidly and reversibly switched between an

amorphous and a crystalline state. Because both states are characterized by strong

differences in the optical and electrical properties, these materials can be employed

for data storage, including rewritable optical data storage (Chapter 2.3.1) and

electronic memory (Chapter 2.3.2). Based on this mechanism, a series of frontier

applications of phase-change materials (Chapter 2.3.3) is generated.

Optical data storage 2.3.1

Yamada et al. [20, 27] introduced the re-writable (RW) optical data storage based

on Ge-Sb-Te alloys. Within the last decades, commercial products based on PCMs

have been used in optical storage. They have progressed from compact discs (CDs)

with the operational wavelength λ = 780 nm and a storage capacity of 650 MB in

1997 over digital versatile discs (DVDs) with λ = 650 nm and 4.7 GB in 2000 to

recordable/erasable Blu-ray discs (BD-REs) with λ = 405 nm and 23.3 GB in 2003.

Based on the fixed operation wavelength of 405 nm, together with the so-called

multilayer technology, a dual-layer BD-RE with a capacity of 50 GB followed in

2004 and, later on, a triple-layer BD-RE with 100 GB storage capacity in the year

2011 [57]. These stages of development are depicted in Figure 2.5, which shows the

progress of data storage on optical discs, based on the phase-change optical

recording technology.

Basically, a higher storage density can be expected by either reducing the laser

wavelength or by increasing the numerical aperture of the objective lens (NA),

because of the storage density D ∝ (NA/λ)2, where λ is the laser wavelength. In this

thesis, the UV-wavelength of 248 nm is used for triggering the phase transition of

phase-change materials described in Chapters 4.2 and 4.3. Another approach uses

the near-field optical technique to realize an increase of the storage density,

whereby a nano-aperture or a solid immersion lens focuses the laser beam into a

very small spot on the disc surface. However, this technique is limited to a working

distance of < 40 nm, which is a challenge for practical fabrication and operation of


18

an optical disc. Nonetheless, this led to the development of optical discs based on a

super-resolution near-field structure (so-called Super-RENS), which are able to

record and read out the marked regions beyond the diffraction limit [58]. In spite of

this, the detailed working mechanism of such Super-RENS is still not completely

understood [15]. Alternatively, multi-layer recording technology has been used for

improving the commercially available recording density from 23.3 to 100 GB, as

shown in Figure 2.5. More recently, multi-level data storage has been attracting

increasing attention because of the possibility to increase the storage density even

further. This method is based on a uniquely precise control of the amorphous to

crystalline transition process, in order to realize different (non-binary) reflectivity

levels in each single layer memory cell, instead of using multi-layers. Thus, multi-

level storage techniques are not only beneficial for the increase of storage density,

but they also can potentially increase the data transfer rate and decrease the cost per

bit (for more details, see Chapter 4.3).

Another important trend is the desire for a high maximum data transfer rate.

Over the last 20 years [57], an increase of the data transfer rate from about 10 to

Figure 2.5: Progress of phase-change based re-writable optical storage media (top part adapted from

Ref. [2]): A redesign of the optics, i.e., usage of shorter wavelengths and higher numerical apertures

(NA) allowed for a reduction of the cross-section of the focused laser beam and thereby for a

decrease of the minimal mark size and a corresponding increase in recording capacity.

1995 2000 2005 2010 2015

0

50

100

Cap

acity (

GB

)

Release date (Year)

λ=780 nm λ=650 nm λ=405 nm λ=405 nm λ=405 nm

NA=0.85 NA=0.85 NA=0.85NA=0.6NA=0.5

1.2 mm

CD-RW

0.6 mm 0.1 mm 0.075 mm 0.057 mm

DVD-RAMBD-RE (Single)

BD-RE (Dual)

BD-RE (Triple)

0.018 mm0.025 mm

112 nm149 nm149 nm410 nm833 nm

CD-RWDVD-RAM BD-RE

(Single)

BD-RE (Dual)

BD-RE (Triple)


19

133 Mbit/s has been developed. The data transfer rate of a PCM is generally

associated with the amorphization and crystallization rates of its recording material

(see also Chapter 2.1). The amorphization process consists of an inherently quick

melt-quenching process, whereas crystallization is characterized by a minimum

time required for the generation of crystalline phase nuclei and their growth.

Therefore, the maximum attainable data transfer rate is mainly determined by the

crystallization kinetics of the PCM (see Chapter 2.2). According to classical

crystallization theory, nucleation and growth rates are dependent on the temperature

of the respective material. By using laser pulses as the energy source, an increase of

the temperature in the material and the induction of phase transformation in PCM

must be considered. In the case of optical switching, this energy transfer is usually

realized by supplying additional heat directly, via photon energy transfer to excited

electronic states, and subsequent relaxation [1, 59]. Coherent light is used in optical

switching to induce the phase transformation. Thereby, by reducing the irradiated

laser pulse duration, which leads to a short energy deposition time for phase

transitions in PCMs, performance at a high data transfer rate can be expected.

Various experimental studies on the property changes of PCMs upon nano-, pico-

and femtosecond laser pulse irradiation have been published in recent years [31, 60-

67]. However, comparative examinations are still lacking. Thus, the investigation of

an experimental design for a comparison of the process of GST crystallization

induced by single laser pulses in the time scale frame of ns and fs, as well as a

corresponding model of the average contribution of ns and fs laser photons to the

phase transition process, are presented in Chapter 4.2.2, with the aim of

understanding the switching process and increasing the data transfer rate.

Electronic memory 2.3.2

Similar to the optical data storage (as described in Chapter 2.3.1), PCMs are also

used as the active layer, in the form of phase-change random access memory

(PCRAM). In the year 1960, Ovshinsky observed I-V behavior in amorphous

chalcogenide glasses with a characteristic reversible switching mechanism between

high resistance state and low resistance state [4]; this principle is used for current

electronic data storage applications. Presently, the known remarkable properties of

PCRAM are non-volatility, good scalability (in terms of minimized device feature


20

size F), long endurance, short write/read (W/E) time, and multi-level cell (MLC)

capability, leading to PCRAM becoming competitive with solid-state electronic

memory. A comparison of the device characteristics of various memory

technologies is summarized in Table 2.1.

PCRAM devices use the different values of electrical resistivity upon rapid and

reversible phase transitions between the amorphous and the crystalline states of the

PCMs. The physical principle of switching PCMs relies on the atomic configuration

changes illustrated in Chapter 2.1. For practical memory applications, with logical

“0” in the cell, in which the PCM is in the crystalline state that exhibits the low

resistivity state (LRS). To write a logical “1” in the cell, a short high power pulse is

applied to the PCM cell, leading to amorphization of a partial volume of the PCM

and generating the high resistivity state (HRS), named RESET operation. In reverse,

a pulse of moderate power but with sufficiently long duration is used for

crystallization of this amorphous volume, resulting in erasing the “1” and returning

to “0” memory state, named SET operation. Due to the significant difference

between the HRS and LRS (up to five orders of magnitude), the process logical

states are easily read out through a lower power pulse. In addition, the threshold

switching effect is essential for the operation of the PCM cell, because it is

Table 2.1: Comparison of various memory technologies [3]: Dynamic random-access memory

(DRAM), NAND fresh memory, phase-change random access memory (PCRAM), ferroelectric

random access memory (FeRAM), magnetoresistive random-access memory (MRAM), resistive

random-access memory (ReRAM).

Attributes DRAM NAND PCRAM FeRAM MRAM ReRAM

Non-Volatile No Yes Yes Yes Yes Yes

Cell size 6F2 4F

2 4F

2 22F

2 20F

2 4F

2

Read time <10 ns 0.1 ms 12 ns 40 ns 35 ns <50 ns

W/E time <10 ns 1/0.1 ms 100 ns 65 ns 35 ns <1 ns

Endurance >1015

>105 >10

12 >10

14 >10

12 >10

9

Retention 64 ms >10 yrs >10 yrs >10 yrs >10 yrs >10 yrs

Write Energy ~10

-15

J/bit

~10-16

J/bit

~10-12

J/bit

~10-14

J/bit

~10-12

J/bit

~10-13

J/bit

MLC No No Yes No No Yes


21

accompanied by an abrupt transition from amorphous (HRS) to crystalline (LRS)

states. Otherwise, extremely high voltages would be required to introduce enough

power into the cell to heat it above the crystallization temperature, unless the

occurrence of threshold voltage switching [68, 69]. A corresponding typical current-

voltage (I-V) curve for a PCRAM and a schematic picture of a PCRAM in cross-

sectional view are shown in Figure 2.6.

With respect to the attributes of PCRAM (see Table 2.1), scalability is one of

the dominating issues that must be addressed in order to enhance PCRAM to next

generation non-volatile memory applications. Scaling down the feature size is not

only advantageous for increasing the storage density, but, even more important, for

the reduction of the operation energy consumption; in particular, cutting off the

RESET current. The structure design of the PCRAM cell plays an essential role,

here. When attempting to confine the current flow in the PCRAM device, Joule

heating of the PCM can be restricted and optimized. As an example, a typical

mushroom-cell structure of a PCRAM is shown schematically in Figure 2.6b.

Following an applied electric pulse (irrespective of its polarity) with a voltage

above the threshold voltage acts as Joule heating source, narrowing the contact area

between PCM and the heat contact can be used to gain an increased heating

efficiency in the PCM layer and to induce a local switching process from the high-

resistive amorphous phase to the low-resistive crystalline phase. However, scaling

down memory size below a certain dimension of < 50 nm is challenging because of

Figure 2.6: (a) Typical current-voltage (I-V) curve of SET and RESET operation of a PCM-based

device. The threshold switching behavior occurs at a threshold switching voltage (Vth). (b) A cross-

sectional schematic of a PCRAM mushroom-cell.

S

Bottom electrode

SiO2

PCM

Heater

Switchable region

Top electrode

Crysta

lline st

ate (L

RS)

Curr

ent (A

)

Voltage (V)

Amorphous state (HRS)

RESET

SET

Am

orp

hiz

ation

Cry

sta

llizatio

n

Threshold

voltage (Vth)

(a) (b)


22

the spatial limitations of lithographic device fabrication. As an alternative

technology, probe-based storage (using an AFM tip as probe for writing, reading,

and erasing operations in local PCM layers) has been attracting increasing attention.

It exhibits remarkably advantageous features [70-74]: i) higher data storage density

> Tbit/in.2, ii) lower power consumption of < 100 pJ per bit written, iii) higher data

transfer rate with large arrays of tips operating in parallel, iv) less expensive, in

comparison with the process for conventional memory device fabrication. In the

present thesis, conductive-AFM is used as an active tool for the investigation of

GST phase-change films at the nanoscale. This is not only a breakthrough regarding

the bottleneck of device scalability but, more importantly, it reveals deep insights

into the switching mechanism by simultaneously measuring electrical and

topographical properties of phase-change films at nanoscale (for more details see

Chapter 4.4).

Frontier applications of phase-change materials 2.3.3

Over the last decades, the functionality of chalcogenide-based PCMs has been

further developed with respect to the targets of optical data storage and electronic

memory applications. Simultaneously, a remarkable improvement in the

understanding of structural, optical and electrical properties of the materials with

the capabilities of nanoscale, non-volatile, prominent contrast of optical and

electrical properties between amorphous and crystalline states, fast switching speed,

multilevel storage, and 3D integration has occurred. Such a significant functionality

extension of PCMs opens the door for several promising frontier applications. For

instance, Michel et al. demonstrated that, by introducing PCM thin layers, a non-

volatile tunability of the resonance wavelength in an antenna can be achieved [75].

Similarly, the combination of surface plasmon polaritons with PCMs, with the

purpose of controlling surface plasmon waveguides, has been explored [76].

Another insight was used to uniquely combine electro-optical properties of PCMs,

in order to reveal possible applications in displays and data visualization [77].

Furthermore, a memristor capable of Boolean logic function conjunction with

memory in a single PCM cell has been proposed [78], as well as the use of PCM

devices for emulating biological synapses and neurons [79-83], which paves the

way for neuromorphic computing.

23

Chapter 3

Experimental conditions and methods

Pulsed laser deposition of thin films 3.1

Pulsed Laser Deposition (PLD) is used to deposit thin films of the phase-change

material system Ge-Sb-Te. In contrast to other physical deposition technologies, the

main advantages of PLD are (1) high quality films can be grown reliably in a very

short time and (2) the deposition of films with unusual composition and with

stoichiometric transfer of a target material to the film is possible. In addition, (3) the

vacuum system and the laser beam source are decoupled, (4) a simple vacuum

system is required, and (5) the electrical properties of the target are not crucial (e.g.,

compared to dc sputtering), and (6) PLD enables precise control of the growth rate

(sub-monolayer deposition per pulse).

The basic experimental setup is shown in Figure 3.1. A pulsed KrF excimer

laser with a wavelength of 248 nm and 20 ns pulse duration is applied. The laser

beam is focused, via a plano-convex lens (focal length = 800 mm) through a quartz

window, onto the surface of the rotating target at an incident angle of 60° with

respect to the target normal. The pulse energy is varied between 160 and 200 mJ,

Figure 3.1: Schematic of a pulsed laser deposition system.

Rotating target

Rotating

substrate

Plume

Focus lens

Pulsed laser beam

Target

Substrate

Heater

Chapter 3 Experimental conditions and methods

24

which leads to a fluence within a range between 1.0 and 1.5 J/cm2. The pulse

repetition rate is typically set to 10 Hz. The target materials are arc melted GeTe,

Ge2Sb2Te5, and LaAlOX, targets, which are placed on the multi-target holder.

Ge2Sb2Te5 and GeTe have a high absorption coefficient in the UV range (~106

cm-1

)

[84, 85] and are therefore suitable for the PLD process. The working pressure in the

PLD chamber is typically 5×10-8

mbar during deposition. The substrates (1×1 cm2)

are ultrasonically cleaned in ethanol and deionized water prior to deposition and are

positioned inside the vacuum chamber parallel to the target surface at a target-to-

substrate distance of 7.4 cm. Both target and substrates are rotated in order to avoid

deep ablation of the target and to improve the thickness homogeneity of the films,

respectively.

A resistive heater source is mounted in a specific position on the substrate

holder, near the substrates, in order to ensure an effective thermal dissipation when

increasing the substrate temperature to values required for the experiments (max.

temperature up to 1100°C).

The basic processes of PLD are the absorption of the laser radiation by the

target material, generation of a plasma plume, expansion of the plasma, and

thermalization and condensation of the target material on the substrate. In detail, the

laser beam is focused on the target material surface, causing the near-surface region

of the target to melt and then vaporize during the early stage of the ablation

processes. The continuing laser energy transfer within one laser pulse and

absorption of the laser light cause an increase in the target material vapor to high

temperatures and high pressure and, eventually, the ionized gas expands from the

target in the direction of the substrate; this is the so-called plasma plume. The

plasma plume contains a number of different species, including electrons, ions,

atoms, molecules, etc. If background gas is present in the deposition chamber, the

thermalization of the plume takes place through reduction of the kinetic energy of

the plume species. Theoretical models of interaction between plasma expansion and

background gas have been described, e.g., in Ref. [86]. Finally, the plume particles

are deposited on the substrate surface. Especially at higher substrate temperatures,

these particles are able to diffuse onto the surface and find the energetically most

3.2 Thermal treatment

25

favorable sites. For more detailed information about PLD, the reader is referred, for

example, to Refs. [7, 8, 87, 88].

Thermal treatment 3.2

The thermal treatment of the thin films is carried out by in situ annealing in the

deposition chamber during thin film deposition (Chapter 3.2.1), and by ex situ

annealing in an infra-red-irradiation oven following the thin film deposition process

(Chapter 3.2.2).

In situ thermal treatment in deposition chamber 3.2.1

To investigate the influence of substrate temperature on the structural properties of

the deposited GeTe and GST films, the substrate is heated by a resistive heater

source, which is located behind the substrate holder in the deposition chamber, as

depicted in Figure 3.1. The heater temperature is normally set to temperatures in the

range between 100°C and 500°C. The temperature on the surface of the substrate

materials is calibrated by connecting a thin (0.25 mm) type C thermocouple wire

with the substrate surface. The measurement point is pinned to the substrate surface

using ceramic cement (Omega bond CC High Temp). The results of the calibrations

on Si (100) and BaF2 (111) substrates are shown in Figure 3.2. The uncertainty of

the measurements is estimated to be ±30°C [26].

Ex situ thermal treatment in an infra-red irradiation oven 3.2.2

The post-thermal annealing of deposited GeTe or GST films is realized by an infra-

Figure 3.2: Substrate temperature calibrations.

0 100 200 300 400 500

100

200

300

400

Si (100)

BaF2(111)

Tsub

= -19.7°C+0.98Theater

Tsub

= 1.65°C+0.65Theater

Substr

ate

Tem

pera

ture

(°C

)

Temperature Heater (°C)


26

red irradiation oven system. The heating furnace features a high-accuracy

temperature control and is capable of rapid heating (> 50 K/s). As-deposited films

were annealed under high vacuum condition (5×10-6

mbar) in the temperature range

between 100 and 500°C for 20 min at a heating rate of 10 K/min. Furthermore, in

order to study the temperature dependence of the electrical resistance in vacuo, the

furnace was connected with a Keithley 2400 source measure unit in a two-point

probe setup, using gold wires for contacts, as shown in Figure 3.3.

Laser pulse irradiation for optical switching 3.3

Nano- and femtosecond laser irradiation 3.3.1

A schematic of the static ns- and fs pulse laser irradiation experiment setup is

Figure 3.3: Schematic of oven system with the possibility of temperature-dependent resistance

measurement.

Figure 3.4: Schematic of the static laser irradiation experiment setup.

Gas

Heating source: halogen lamp

Thermocouples

Keithley 2400 Source

Measurement unit

Excimer laser 248 nm

Pulse duration 20 ns or 500 fs

Convex lens

Filter

Gaussian lens

GST or GeTe

SiO2/Si

3.3 Laser pulse irradiation for optical switching

27

shown in Figure 3.4. Two laser source systems at the same wavelength of 248 nm,

but with different pulse lengths, are applied to investigate the influence of the pulse

duration on the crystallization process in the phase-change films. For the ns laser, a

KrF excimer laser with pulse length of 20 ns is utilized. The fs laser is a hybrid

excimer-dye laser with a pulse length of 500 fs. The two laser beams are focused by

a convex lens onto the surface of the deposited thin films. The average fluence is

estimated as the laser energy divided by the spot size. The laser fluence can be

varied between 0 and 200 mJ/cm2 and between 0 and 19 mJ/cm

2 for the ns pulse

laser and the fs pulse laser irradiation, respectively.

Nanosecond laser pump-probe system 3.3.2

In order to investigate the dynamics of phase transition in phase-change films, an

optical pump-probe system, based on the ns pulse laser irradiation setup, was

constructed. The schematic arrangement of the ns pump-probe system for real-time

optical reflectivity experiments is shown in Figure 3.5. The pumping source is a

KrF excimer laser with a wavelength of 248 nm and a pulse duration of ~20 ns. The

pumping laser beam is focused by a convex lens on the surface of the sample, in

order to induce local phase changes. The spot size of the rectangular laser beam is

about 24 × 6 mm2 and the laser fluence is varied between 0 and 182 mJ/cm

2 by

adjusting the laser pulse energy from 0 to 300 mJ. A continuous-wave (CW)

Figure 3.5: Schematic of the ns laser irradiation optical pump-probe system setup.

Excimer laser (248 nm, 20 ns)

Convex lens

Filter

Gaussian lens

Oscilloscope

Diode laser

405 nmGeTe

SiO2/Si

LaAlOX

Photodiode


28

probing laser beam from a laser diode with a wavelength of 405 nm is focused on

the center of the irradiated spot (d ~2 mm) at an incident angle of about 30° with

respect to the sample surface. The reflected probing beam is collected by a

high-speed Si photo detector (rise/fall time: 1 ns) equipped with a front UV filter, to

reduce the influence of the UV laser, and is connected via a coax cable to a fast

digital 500 MHz oscilloscope. The total time resolution of the pump-probe system

is about 2 ns. A second photodetector (without filter) is employed to detect the

shape of the pumping excimer laser beam as well as to calibrate the time delay

between the measured signal of the excimer laser and the diode laser. Moreover, the

influence of CW probing laser irradiation with a maximum power density

~0.7 W/cm2 on the sample is examined by a long duration (t > 60 s) analysis of the

reflection signals; this indicated that the CW probing laser irradiation does not

influence the optical signal from the sample. The layer structure of samples, as

shown in Figure 3.5, is an example with LaAlOX serving as a capping layer to

protect degradation of the active GeTe layer on SiO2/Si substrate. It should be noted

that the absorption of the LaAlOx layer with thickness ~5 nm as used in this study

upon laser irradiation wavelength of 248 nm can be neglected.

Arrangement for electrical switching 3.4

In this work, the electrical resistance switching in the memory stack (50 nm

GST/100 nm Cr/Si) is systematically investigated at the nanoscale by means of a

conductive atomic force microscope (C-AFM) based setup for write, read and erase

Figure 3.6: Schematic of a conductive atomic force microscope configuration-based setup for

electrical pulse- induced experiments.

LaserPhotodetectorpA-Amplifier

to A/D

Cr 100 nm

GST 50 nm

Si

HeaterCeramic

N2 gas

Thermocouple

type C

DC bias voltage

Conductive tip

Au wire

AFM stage

3.4 Arrangement for electrical switching

29

operations, and is shown in Figure 3.6. The measurements are performed using an

AFM equipped with a tunneling atomic force microscopy extension that is used for

C-AFM-type investigations. A commercial Pt-Ir-coated Si AFM tip, with a tip

radius of 20-25 nm as the probe, serves as a nanoscale contact (top electrode) to the

investigated phase-change film. The spring constant of the utilized cantilever is 2.8

N/m. During the experiments, a variable bias voltage is applied across the phase-

change thin films. The backside electrode (Cr layer layer) and the top electrode (PtIr

coated tip) are in series with a current amplifier in the range of 1 pA to 1 µA. The

tip can be used to perform a static measurement. Alternatively, it can be scanned in

contact mode. When the tip is used in static mode, the local nanoscale I-V curve of

the phase-change film is measured. By programming a controlled motion scheme of

the tip, an array of marks can be written into the amorphous phase-change films by

applying a voltage that exceeds a certain threshold switching voltage. During the

scan process of the tip, and maintaining a constant force between the tip and

sample, it is possible to record topographic and current images of the films

simultaneously. This enables the direct correlation of local nanoscale topography

with local electrical properties. In this study, the ramp parameters of the voltage

sweep are between -3 V to +3 V with a ramp frequency between 0.15 Hz and 10

Hz. Scanning is usually performed on areas of 500 × 500 nm² or 1 × 1 μm² with a

scanning frequency between 0.15 Hz and 0.2 Hz at a DC sample bias of 0.5 V; this

is applied to record the topography-current-images. A calibration of measured I-V

curves, using a standard resistor (100 MΩ) with a measurement sensitivity of 1

nA/V (limited current range between -12 nA and +10 nA), is shown in Figure 3.7a.

Figure 3.7: (a) Calibration of a standard resistance with value of 100 MΩ in I-V characterization by

C-AFM, (b) Calibration of substrate temperature as a function of supply power.

0 2 420

40

60

80

Measurement points

Fit by y=27+11.3xSubstr

ate

tem

pera

ture

(°C

)

Power (W)

-3 -2 -1 0 1 2 3-15

-10

-5

0

5

10

Cu

rre

nt

(nA

)

Voltage (V)

Forward

Reverse

Calib. Res.100 M

Sen. 1 nA/V

(a) (b)


30

The sample temperature is controlled by heating with a Peltier element (calibration

substrate temperature vs. heating power is shown in Figure 3.7b). In order to

minimize the oxidation of the films, the film surface is surrounded by protective

nitrogen gas. With this experimental arrangement, the measurement of I-V curves

as a function of the sample temperature can be realized.

Characterization methods 3.5

To investigate the structure, chemical composition, morphology, crystallization

kinetics and optical properties of the as-deposited films and of the films that

underwent a phase transformation induced by thermal annealing via laser irradiation

or via an electric current flow, a variety of characterization methods is used. In this

subchapter, a brief description of these techniques, including X-ray diffraction-

based analysis (Chapter 3.5.1), electron microscopy (Chapter 3.5.2), atomic force

microscopy (Chapter 3.5.3), spectroscopy and calorimetry methods (Chapter 3.5.4)

is given.

X-ray diffraction-based analysis methods 3.5.1

X-ray diffraction (XRD) is applicable to the analysis of the crystalline structure of

the GeTe and GST thin films. The principle of XRD measurements is based on the

interference between incident X-rays and periodically arranged atoms in a lattice;

i.e., the beam is diffracted only in the case of specific specular incidence angles θ

with respect to the family of lattice planes. The condition of maximum

diffracted/reflected intensity is determined by the Bragg law

2𝑑ℎ𝑘𝑙 sin 𝜃 = 𝑛𝜆 , (3.1)

where dhkl is the interplanar spacing, n is the diffraction order, λ is the X-ray

wavelength.

In this thesis, the samples are measured with a Rikagu Ultima IV Type 3

diffractometer in parallel beam geometry, using a θ-θ goniometer and parallelized

Cu K radiation (λ = 0.15418 nm). A schematic of the goniometer axes is shown in

Figure 3.8. A special feature of this arrangement is that the detector arm can be

rotated around two axes: the usual 2θ axis, as well as the additional 2θχ axis. The

latter is called in-plane axis.

3.5 Characterization methods

31

In order to obtain information about various properties of the deposited thin

films, several corresponding measurement techniques are performed, on the basis of

the utilized XRD setup, as briefly described in the following:

2θ-ω scan measurements

2θ-ω scans allow the determination of the interplanar spacing of families of lattice

planes that are oriented parallel to the sample surface, the underlying crystal

structure(s) that are present and the preferred orientation of crystallites. This can be

achieved by altering the incident X-ray beam angle ω and simultaneously altering

the detector at the angle of (2θ), while the sample remains stationary. Importantly,

ω is equal here to the angle θ between the incident X-rays and the sample surface.

The 2θ angular step width of the scans is 0.05° and an angular resolution of 0.11° is

determined by the use of a parallel slit analyzer. According to Bragg’s law

(Equation 3.1), from a distinct 2θ position of a diffraction peak, the corresponding

interplanar spacing dhkl can be deduced. In addition, the full width at half maximum

(FWHM) of the detected diffraction peak, also called Δ(2θ) can be used to estimate

the approximate average size of crystallites. This relationship is expressed by the

Scherrer formula [89]:

∆(2θ) =0.9𝜆

τ∙cos 𝜃𝐵 , (3.2)

where τ is the average size of the crystallites (average grain size), and θB is half of

the Bragg angle.

Figure 3.8: Schematic of the goniometer geometry: ω angle between incident x-ray beam and sample

surface, 2θ angle between incident x-ray beam and detector, 2θχ in-plane angle, φ sample rotation

angle around the sample normal, χ sample tilt angle around the axis perpendicular to 2θ and χ axes.

φ

χ

X-ray detectorX-ray source

Sample

ω 2θ

2θχ


32

In order to obtain a higher diffracted intensity from the polycrystalline GeTe

films, despite the small film thickness, grazing incidence diffraction (GID) at a

fixed incidence angle ω of 1° with respect to the sample surface is employed as

well. In this measurement mode, only the X-ray detector scans the angle range of

interest during the measurement.

Rocking curve measurement

Rocking curve measurements of distinct Bragg reflections are used to obtain the

average tilt distribution of the related crystallites in the films. The precondition for

this type of measurement is that the 2θ and ω axes are oriented to the sample in such

a way that Bragg’s law is fulfilled for the chosen reflection. Then, in the case of the

utilized θ-θ goniometer, the rocking curve measurement is realized by a coupled

movement of the ω axis and the 2θ axis, whereby the angle between the two axes

remains constant. This scan is performed around the position given by the Bragg

law. For typical thin films, the maximum intensity of the detected rocking curve

peak can usually be found near (2θ)/2; the corresponding FWHM of the rocking

curve is an estimate of the average tilt of crystallites around the normal of the

chosen family of planes.

In-plane pole figure measurement

Pole figure measurements are typically used to determine the orientation

distribution of crystallites in textured or epitaxial thin films. First, a distinct Bragg

reflection is chosen. Then, the pole figure measurement is realized by rotating the

sample around its normal from 0 to 360° (complete φ-Scan) for polar angles χ

ranging from 0 to 90°. A combined movement of the ω, 2θ, and 2θχ axes is

performed to realize the various polar angles. Thus, the pole figure measurement is

called in-plane pole figure measurement. The collected diffracted intensity of the

{hkl} family of planes at each position (χ, φ), is dependent on the actual crystal

structure of the sample materials and on the orientation of the crystallites in the

sample. The recorded intensities of all (χ, φ) positions are then projected on the

sample plane in stereographic projection. The angles χ and φ are also called α and β,

respectively. In the resulting pole figure, the intensity at each angle is dependent on

the number of crystallites in the films that are oriented so that they fulfil the Bragg

condition for the pre-selected reflection. For a fiber textured film, continuous rings


33

of higher intensity centered at or near the origin of the pole figure are typically

encountered. For an epitaxial film, a number of separate small spots of higher

intensity, the so-called pole density maxima, can be found in a well-defined

arrangement, and are characteristic of a highly preferred orientation of the film. In

order to evaluate actually measured pole figures, predicted pole figure models were

created by CaRIne crystallography software [90] and compared to the measured

ones.

φ-scan measurements

Instead of recording full pole figures, which include intensity data from the whole

northern pole sphere, it is, in many cases, sufficient and much more efficient to

measure partial pole figures or even single φ-scans. Such φ-scans are used here for

determining the orientation relationships between epitaxial films and the underlying

single crystal substrate. For this purpose, a distinct reflection from the film, as well

as from the substrate, is selected.

X-ray reflectivity measurement

Apart from the information about crystal structure and texture, the thickness,

density and surface/interface roughness of deposited thin films on a substrate can be

measured with an X-ray diffractometer by using X-ray reflectivity (XRR) [91]. In

terms of measurement geometry, an XRR measurement is, in principle, similar to a

2θ-ω scan, but in XRR mode, the X-ray beam is incident at a comparably small

angle relative to the sample surface. As long as the incidence angle lies below the

critical angle of total reflection (i.e., θc< 0.5°), total reflection of the X-ray beam

takes place, i.e., the X-ray reflectivity is maximal. When the incident angle

increases to values above θc, the reflectivity decreases drastically. In the case of a

homogeneous thin film on a flat substrate, the X-ray reflectivity curve exhibits an

intensity oscillation, due to constructive/destructive interference of X-rays reflected

at the surface/interfaces. The period of the oscillation is related to the film thickness

and the amplitude of the oscillation to the film density, substrate density and

surface/interface roughness. The film density can be determined from the critical

angle of total reflection. By evaluating the measured reflectivity profile, information

about film thickness, density and surface/interface roughness can be extracted. This


34

evaluation of the obtained measurements is performed by using the commercial

RayfleX software [92].

Electron microscopy 3.5.2

Scanning electron microscopy

Scanning electron microscopy (SEM) is used for the characterization of the surface

topography of the deposited thin films. In SEM, a focused beam of electrons with

an energy varying between 0.1 and 30 keV is scanned over the sample. Due to the

interaction of primary electrons with the atoms in the near surface region of the

sample, secondary electrons, bremsstrahlung and characteristic X-rays are

generated. Some of the primary electrons are rejected, as so-called backscattered

electrons. The most common use of SEM is for the detection of secondary electrons

that correspond, in their emission, with local height variations of the surface,

producing a type of surface topography images.

In the present work, a Carl Zeiss Ultra 55 field emission SEM is used for

obtaining cross-section and top-view images of the films. The instrument is

equipped with a GEMINI electron column as well as with two available secondary

electron detectors, an Everhart-Thornley detector and an in-lens detector. The SEM

is operated with an electron acceleration voltage of 15 kV and a working distance

between sample surface and electron column end of approximately 5 mm.

Energy-dispersive X-ray spectroscopy

Energy-dispersive X-ray spectroscopy (EDX) was used to determine the chemical

composition of the deposited films. The working principle of EDX is based on an

incident electron beam that knocks out an inner shell electron of an atom and,

hence, creates an electron hole where the electron was. Subsequently, the hole is

filled by an outer shell electron. The difference in energy between the inner and

outer shell produces a characteristic X-ray quantum. The emitted X-ray can be

detected by an energy dispersive spectrometer. Since each element has a unique

electron configuration of the atom, the energies are characteristic for the elements

contained in the sample and, therefore, they can be identified. The atomic

concentration can be calculated by using the signal intensities.


35

In the present work, EDX (Quantax 400-EDX, Bruker Nano GmbH, attached to

the SEM system) is used for analyzing the chemical composition of the films. An

electron beam energy of 25 keV is used. In order to increase the detected intensity

from the films, the samples are tilted to an angle of 45° and a large aperture with

60 μm diameter is used for acquisition. The acquired data are analyzed in the

Bruker ESPRIT software.

Transmission electron microscopy

Transmission electron microscopy (TEM) is used to determine the morphology,

microstructure and crystal structure of the deposited films. In TEM, an electron

beam with a high acceleration voltage of a few 100 keV is transmitted through a 20

to 200 nm thin specimen. The interaction between the incident electrons and the

atoms within the specimen results in the generation of differing signals. The basic

TEM operation modes are imaging and diffraction, which can be chosen by

adjusting the apertures and electron lenses. In imaging mode, a bright field image

can be formed using the detected direct electrons, whilst a dark field image is

formed using the electrons scattered by a particular set of crystal planes. The images

can provide information about the film morphology, appearance of grains, grain

boundaries and the defect structure of the measured films. The resolution of the

images can be improved (down to the sub-Ångstrom scale) by using spherical

aberration corrector (Cs corrector) TEM, which allows the direct characterization of

the atomic structure. The unique properties of such a TEM facility can be combined

with the capability of scanning the electron beam over the specimen, resulting in the

scanning transmission electron microscopy (STEM) technique. In diffraction mode,

the obtained electron diffraction patterns result from coherent elastic scattering and

provide information about periodicity in the materials. For instance, an amorphous

film shows broad rings in the diffraction pattern that correspond to the inter-atomic

spacing, whereas crystalline materials show sharp diffraction spots that correspond

to Bragg diffraction at the periodic lattice. Using an electron beam with a small

diameter (~5 nm) makes it possible to take diffraction patterns from precisely

selectable locations of the specimen and, therefore, facilitates determining the local

structure on the nano scale, the so-called nano beam diffraction (NBD) technique.


36

In addition, STEM can also be combined with EDX, which makes a direct

correlation of imaged region and chemical analysis of a specimen feasible.

In the present work, a FEI Titan3 G

2 60-300 analytical TEM equipped with a Cs

probe corrector and annular dark-field (ADF) detectors is used for local structure

and elemental analysis of the GeTe and GST thin films. The cross-sectional TEM

specimens of the films are prepared by using focused, high-energy Ga and

subsequent low-energy Ar ion beam milling. Typically, a small specimen thickness

of approximately 25 nm is used for STEM studies. More details about the specimen

preparation using a focused ion beam (FIB) can be found in Ref. [93]. An ultra-thin

coating of electrically-conducting Pt is used to protect the surface of a FIB lamella

during the preparation. The TEM was operated at 300 kV accelerating voltage. A

probe-forming aperture of 25 mrad is used in the STEM studies. Middle-angle ADF

(MAADF) images are taken, using annular ranges of 40-200 mrad of the HAADF

detector, whereas low-angle ADF (LAADF) images are recorded using annular

ranges of 9.3-53.6 mrad of the ADF detector. Diffraction patterns of the specimens

are obtained in the NBD mode. The analysis of local chemical compositions is

identified by a Super-X EDX system.

Atomic force microscopy 3.5.3

Atomic Force Microscopy (AFM) is used to obtain the topography of the GeTe and

GST films at the nanoscale, as well as their local phase transformation behavior.

The basic principle of AFM refers to the interactions between the probe tip and

sample surface on the atomic level. Using a very simplified picture of the tip-

surface interaction, it can be described by a superposition of the attractive van-der-

Waals forces and the Pauli repulsion, the underlying potential is the Lennard-Jones

potential [94]

𝑈(𝑑) = 4ε [(𝜎

𝑑)

12

− (𝜎

𝑑)

6

] , (3.3)

where σ and ε are constants. This potential U is schematically depicted as a function

of the tip-sample distance in Figure 3.9a. AFM measurement modes can be

distinguished as a function of the sign of resulting force (repulsive or attractive):

Contact and Non-contact Mode. In this thesis, AFM measurements in contact as

well as in intermittent contact mode (the so-called tapping mode) are used. The


37

contact mode is utilized in the conductive AFM setup, resulting in the opportunity

to register topographic and electric current images of the films simultaneously (for

details see Chapter 3.4). The tapping mode is used to gain information about the

microscopic topography of the film. In contact mode, the tip is dragged across the

surface of the sample, whereas, in tapping mode, a cantilever performs a certain

driven resonance oscillation with constant amplitude near the surface of a sample. A

schematic of an AFM setup is given in Figure 3.9b. A laser beam is reflected from

the back of the cantilever onto a position-sensitive photodetector. In general, the

position of the reflected laser beam on the detector shifts when the tip, which is

attached to the cantilever, is scanned over a sample surface with non-homogenous

topography. These shifts, caused by cantilever deflections, are recorded with help of

the photodetector and yield information about the movements of the cantilever.

These are fed into a feedback controller, which controls the cantilever’s up or down

motion, in order to compensate for the cantilever deflection (constant force mode).

A surface topography image is assembled, based on this signal. For the tapping

mode, the tip-sample interaction forces provoke changes of the cantilever

oscillation amplitude at a certain fixed resonance frequency. By maintaining the

preset constant amplitude of the oscillation during sample scanning, a surface

topography image is obtained. In this study, a cantilever with a spring constant

between 0.5 and 4.4 N/m and a resonance frequency between 315 and 366 kHz is

used for all tapping mode measurements.

Figure 3.9: (a) Tip-sample interaction and basic operation modes of an atomic force microscope, (b)

Schematic setup of an AFM.

AFM stage

Photodetector

Cantilever

(a)

Attractive

Po

ten

tia

l (V

)

Tip-sample distance (d)

Contact Mode

Non-contact Mode

Repulsive

(b)

Tip


38

Spectroscopy and calorimetry methods 3.5.4

Time-of-flight secondary ion mass spectrometry

Time-of-flight secondary ion mass spectrometry (TOF-SIMS) is used for the

analysis of the in-depth distribution of the chemical elements in the deposited thin

films. In TOF-SIMS, a focused beam of primary ions with low beam current

(commonly used energy: 10-30 keV; typical ions used in the source: Ga+, Au

+, Bi

+)

generates collision cascades near the sample surface. Due to ion sputtering,

energetic atoms and clusters are emitted from the surface. Most of them are neutral;

only the charged particles are detectable in SIMS. These emitted secondary ions are

analyzed with a time-of-flight mass spectrometer and provide qualitative

information about the presence of chemical elements. In order to obtain a depth

profile, a second ion beam source is used and operated at low ion energy (≤ 2 keV)

and a higher ion beam current, to remove surface atoms by sputtering. Typically

used ions are O2+, Cs

+, Ar

+ or Arn

+. Both ion beams are switched on alternately

during profiling. Thus, TOF-SIMS is able to obtain the in-depth composition of the

film surface as a function of the sputter time. In this study, short Ga ion pulses are

employed with an energy of 15 keV and an average beam current of about 1.5 pA

for scans in regions of 50 × 50 μm². The sputtering species are Cs ions with an

energy of 2 keV and the current range of nA is actually used. The polarity of the

system was set to detect negatively charged ions.

UV-Vis spectrophotometry

UV-Vis spectrophotometry is utilized to measure the optical reflectivity of the

films. The principle of the measurement is based on the well-known Beer-Lambert

law and on the measured intensities of the light reflected from a sample and a

standard reference material (100% absorption), respectively. In this work, the films

are measured with a Varian Cary 5000 UV-Vis spectrophotometer equipped with an

integrating sphere. The typically measured wavelength range is from 400 to 700 nm

and recorded at a step width of 1 nm.

Differential scanning calorimetry

Differential scanning calorimetry (DSC) is used for the thermal analysis of the

phase transformation occurring in the deposited films. In DSC measurements, both

sample and reference are always at the same temperature, with constant thermal


39

heating. When the heat flow changes, due to a phase transformation occurring in the

sample, the amount of heat gain or loss can be observed with differential scanning

calorimeters. Thus, DSC allows a characterization of the crystallization temperature

of the films as well as the crystallization kinetics. In the present work, a Perkin

Elmer DSC 8500 setup is employed. For the preparation of the samples, roughly

800 nm thick GeTe films are directly deposited on an Al cup as a measure sample,

while an empty Al cup is used as a reference. A temperature range between room

temperature (RT) and 300°C and at various heating rates of 5 to 100 K/min with a

temperature accuracy of 0.2°C is used for the measurement.

41

Chapter 4

Results and discussion

Thermal treatment-induced crystallization of GeTe 4.1

films

Among the phase-change materials, the binary compound GeTe is characterized by

a significantly high crystallization velocity, data retention at comparably high

temperatures, stability, and an excellent contrast in terms of optical and electrical

properties between amorphous and crystalline states [28, 30, 95-98]. However,

regarding detailed knowledge on the thermally induced crystallization mechanism

of GeTe films grown by PLD, only few reports exist [9, 24]. To improve the

understanding of the phase-change mechanism in this specific material system, the

crystallization of PLD-deposited GeTe films is investigated in the present thesis that

classifies the thermal treatment as ex situ or in situ thermal treatments. The

investigations on ex situ thermal treatment will with structural transitions of GeTe

films and the effect of oxygen incorporation into GeTe on the phase transition. The

in situ thermal treatment investigation is separated into temperature-dependent

polycrystalline and epitaxial GeTe film growth, based on different types of

substrates.

Ex situ thermally induced crystallization of GeTe films 4.1.1

Nanosecond-PLD was used to deposit GeTe films on Si substrates (for experimental

details, see Chapter 3.1). The structural properties, the topography, the

crystallization kinetics, the optical and electrical properties of these GeTe films in

the as-deposited state as well as in the crystalline state induced by ex situ thermal

treatment are presented [24].

1.) Chemical composition

The GeTe films were deposited on Si substrates with a laser pulse number of 1200

at room temperature, resulting in a typical film thickness of 60 nm at a deposition

rate of ~30 nm/min. In Figure 4.1a, the chemical composition of the films, as

Chapter 4 Results and discussion

42

determined quantitatively by EDX, is shown. The characteristic x-ray lines from

Ge, Te and Si are annotated and the composition of the films is calculated to be

Ge51.6Te48.4. Considering the experimental error of the EDX technique (± 1.6 at.%),

the composition of the films is that of Ge50Te50. In addition, the atomic ratio of

Ge/Te remained nearly constant for an increase of the sample temperature up to

350°C, as shown in Figure 4.1b.

2.) Structure analysis

Figure 4.2a presents XRD spectra of an as-deposited GeTe film and of GeTe films

annealed for 20 min in high vacuum at various temperatures. The XRD spectra of

the as-deposited GeTe film and also of the films annealed at 200 and 220°C show a

very broad peak in the 2θ range between 25° and 30°, indicating an amorphous state

Figure 4.1: (a) EDX spectrum of an as-deposited GeTe film on a Si substrate. (b) Chemical

composition ratio (Ge/Te) of GeTe films, determined by EDX as a function of annealing

temperature.

0 200 250 300 350

0.4

0.8

1.2

1.6

2.0

Ato

mic

ra

tio

Temperature (°C)

Ge

Te

(b)

(a)

0 2 4 6 8 10 12

GeK

TeM Te

L

TeL

GeK

GeL

TeL

Inte

nsity (

arb

. unit)

Energy (keV)

SiK

4.1 Thermal treatment-induced crystallization of GeTe films

43

of the film. However, when the annealing temperature reaches 240°C, diffraction

peaks at angles of 26.1°, 29.9°, 42.4° and 43.3° emerge, which are correlated to the

lattice planes (021), (202), (024) and (220) of rhombohedral GeTe, respectively.

These results illustrate the onset of the structural transformation from the

amorphous to the crystalline state, which corresponds with a remarkable optical

reflectivity increase and resistance drop, as determined by temperature-dependent

reflectivity and resistance measurements (see sections 5 and 6 of this chapter).

After a slight increase of the annealing temperature to 280°C, the maximum

reflection intensity was found, as shown in Figure 4.2a, implying optimum

crystallization conditions. With a further increase of the annealing temperature up to

350°C, the reflection intensity gradually decreases. This intensity decrease can be

Figure 4.2: (a) XRD spectra of an as-deposited GeTe film and of GeTe films with a film thickness of

60 nm annealed at different temperatures. The respective annealing temperatures are denoted. The

Miller indices correspond to the rhombohedral crystal structure of GeTe with the space group R3m

(JCPDS no. 47-1079). (b) FWHM value and calculated grain size from the crystalline phase (202)

reflection as a function of annealing temperature. The dashed lines serve as optical guides for the

eye.

20 30 40 50

Inte

nsity (

arb

.un

its)

2 (°)

380°C

350°C

300°C

280°C

260°C

240°C

220°C

200°C

(021)

Si

(220)

(202)

as-dep.

(003)

(024)

(a)

(b)

240 260 280 300 320 340

0.28

0.30

0.32

0.34

0.36

0.38

0.40

Temperature (°C)

FW

HM

(°)

(202)

30

35

40

45

50

Gra

in s

ize

(n

m)


44

attributed to partial evaporation of the films in the vacuum oven environment.

Finally, for films annealed at 380°C, no diffraction peaks could be detected, due to

complete evaporation of the film. Moreover, no separation of Ge or Te crystalline

phases could be observed within the annealing temperature range of 240-380°C.

Because the full width at half maximum (FWHM) value of the diffraction peaks

correlates with the average grain size of the crystallized films, the average grain size

of the GeTe films, expressed as a function of the annealing temperature, was

estimated from the dominant (202) reflection by using the Scherrer formula [89]. As

shown in Figure 4.2b, the average grain size is inversely proportional to the FWHM

value, and it increases from about 30 to 50 nm as the annealing temperature exceeds

the critical temperature of crystallization at 240 to 350°C. Based on the XRD results

from films annealed at different temperatures, it becomes clear that the phase

transformation from the amorphous to the rhombohedral crystalline state starts

between 220 and 240°C. This temperature for GeTe films is significantly higher

than for GST (Ge2Sb2Te5) films, indicating high data retention at higher

temperatures [99]. Furthermore, it is also higher than typical crystallization

temperatures for sputter-deposited GeTe films [95, 99, 100].

3.) Topography

Figure 4.3 shows the topography and roughness evolution of the annealed GeTe

films at various annealing temperatures under high vacuum conditions. Between RT

and 220°C, large differences in topography and roughness of the GeTe films cannot

be observed (not shown here) and the films are characterized by a smooth,

featureless topography with low roughness. However, crystallites clearly emerge

when the annealing temperature is increased to 240°C. A further increase of the

annealing temperature leads to a uniformly enhanced coarsening of crystallites, as

shown in Figure 4.3. This agrees well with the results of XRD measurements, which

only reveal diffraction peaks for films annealed at 240°C and above. The root-

mean-square roughness (rms) of the films is constant, between 1.7 to 1.9 nm in the

crystallization temperature interval between 240 and 280°C; it is related to the

homogeneous crystal growth process. A significantly higher rms value of 7.0 nm is

obtained for the film annealed at 350°C. This could be associated with increased

crystal coarsening as well as a more intense re-evaporation of the film constituents.


45

4.) Crystallization kinetics

Crystallization kinetics of the as-deposited GeTe films was investigated using DSC

at different heating rates. The results of the DSC measurements are summarized in

Figure 4.4a. With increasing heating rate, the exothermic peak corresponding to the

crystallization temperature Tc shifts gradually to higher values, i.e., from 181.4°C at

5 K/min to 198.4°C at 100 K/min. According to the homogenous nucleation theory

[101], the crystallization process of GeTe crystallites into the amorphous GeTe

matrix is diffusion-controlled. Crystalline seeds are formed as a result of this

diffusion process. For low heating rates, the crystallite dimensions are large and the

number of crystallites per unit volume small. These large crystallites are stable (the

disaggregation of crystallites by out-diffusion can be neglected). With increasing

heating rate, the probability of the formation of a larger number of smaller

crystallites per unit volume is increased. But the process of disaggregation also

becomes increasingly obvious and is contrary to the diffusion process that forms the

crystallites. This means that the crystallization is delayed, which leads to higher

measured crystallization temperatures.

Furthermore, such a time-dependent crystallization process [102] can be

expected, if the crystallization temperature Tc is influenced by the heating rate.

Figure 4.3: AFM images (500×500 nm2) of GeTe films with a thickness of 60 nm annealed at

different temperatures: (a) 240°C, (b) 260°C, (c) 280°C, (d) 350°C.

rms = 1.82 nm

Z = 18 nm

0 nm

a) rms = 1.90 nm

Z = 17 nm

b)

rms = 1.70 nmc) rms = 6.98 nmd)

Z = 16 nm Z = 48 nm


46

According to Kissinger’s formula [103], the activation energy Ea of crystallization

can be determined by

𝑙𝑛 (𝑎

𝑇𝑐2) = −

𝐸𝑎

𝑘𝐵𝑇𝑐+ 𝐴 , (4.1)

where a is the heating rate, Tc is the crystallization temperature, kB is the Boltzmann

constant, and A is a constant. The activation energy is calculated to be 3.14 eV, as

obtained from fitting the slope in the Kissinger plot, and shown in Figure 4.4b. In

the literature, measured activation energies between 1.7 and 3.9 eV for the

transition from the amorphous to the rhombohedral phase of GeTe have been

reported [104-108]. The measured activation energy obtained in this study fits

within this range. A comprehensive explanation for this relatively high activation

energy cannot be given yet. A high activation energy of crystallization could be

related to a long incubation time for crystallization.

Figure 4.4: (a) DSC diagrams of GeTe films recorded with different heating rates. (b) Kissinger plot

for determination of the activation energy Ea of crystallization.

140 160 180 200 220 240

He

at

Flo

w (

arb

.un

its)

Temperature (°C)

Heating rate

Tc=198.4°C

Tc=196.3°C

Tc=194.1°C

Tc=188.3°C

Tc=185.2°C

70 K/min

50 K/min

20 K/min

10 K/min

5 K/min

100 K/min

Tc=181.4°C

2.12 2.14 2.16 2.18 2.20 2.22-11.0

-10.5

-10.0

-9.5

-9.0

-8.5

-8.0

-7.5

ln(

2 c)

1000/Tc(K)

Ea = 3.14 eV

(a)

(b)


47

5.) Optical reflectivity

The optical reflectivity contrast between the crystalline and amorphous state is one

of the crucial optical parameters considered in phase-change optical storage. The

reflectivity measurements were performed on both as-deposited and annealed films.

Based on previous studies [10], an optimal annealing time of 20 min was applied.

For all these measurements, a GeTe film thickness of 60 nm was chosen. The

temperature-dependent reflectivity measurements are presented in Figure 4.5a. The

reflectivity of an as-deposited film varies from 72 to 66 % in the wavelength range

between the near ultraviolet and the near infrared (360-800 nm) and it increases

slightly for films annealed at 220°C. In contrast, the reflectivity of the film after

annealing at 240°C is between 83 and 88% and much higher than the reflectivity of

Figure 4.5: (a) Optical reflectivity spectra of an as-deposited GeTe film and films annealed at

different temperatures, (b) reflectivity contrast between annealed and as-deposited films at a

wavelength of 650 nm for different annealing temperatures. The dashed line in (b) serves only as a

guide for the eye.

(a)

(b)

400 500 600 700 800

60

65

70

75

80

85

90

Re

fle

ctivity (

%)

Wavelength (nm)

Si substrate

300°C240°C

350°C220°C

380°C

as-dep.

0 80 160 240 320

0

5

10

15

20

Reflectivity c

ontr

ast (%

)

Temperature (°C)


48

the as-deposited film. An additional increase of the annealing temperature from 240

to 300°C does not significantly affect the reflectivity of the GeTe films, and the

values still remain between 83 and 88%. After annealing a film at 350°C, the

reflectivity starts to decrease. The probable reason for this decrease of the optical

reflectivity of films annealed at temperatures higher than 300°C could be the much

higher surface roughness caused by crystallite coarsening and evaporation (see

Figure 4.3d). A higher surface roughness scatters light more efficiently in all

directions and the chances of absorption on nearby hillocks are therefore increased.

With a further increase of the annealing temperature up to 380°C, the reflectivity of

the film shows a dramatic decrease that results from the evaporation of film

material. In fact, almost all material has evaporated, because the reflectivity spectra

almost overlap with that of a fresh, non-coated Si substrate.

In order to obtain a high signal-to-noise ratio in an optical application, a high

reflectivity contrast is required. To evaluate the reflectivity changes in more detail,

the reflectivity contrast between the as-deposited and annealed films at a

wavelength of 650 nm were calculated and it amounted to approx. 21%, as shown in

Figure 4.5b. In this context, it should be taken into account that a phase-change film

with an optical contrast above 15% is a promising candidate for commercial

applications [109].

6.) Electrical resistivity

The resistivity measurements were carried out on 60 nm thick as-deposited GeTe

films on SiO2/Si substrates at temperatures between 0 and 300°C. In Figure 4.6, the

resistivity of the GST thin films as function of the heating rate between 5-20 K/min

is shown. Independent of the heating rate, the resistance of ~106-10

7 Ω of the as-

deposited films decreases gradually as the annealing temperature increases up to

200-220°C, due to a thermally activated conduction mechanism. The coarsening of

grains induced by the annealing process for temperatures higher than 220°C (see

Figure 4.2b) leads to a reduced grain boundary density. Based on the assumption

that the grain boundaries are the dominant scattering center for the electrical current

transport, a decrease of the grain boundary density should be connected with a

reduction of the resistivity.


49

Upon further temperature increase, a significant resistance drop, encompassing

3-4 orders of magnitude, can be observed at a transition temperature ranging from

220 to 255°C, where the decrease of the resistance is weakly dependent on the

heating rate (Figure 4.6). Here, the resistivity drop is strongly connected with the

amorphous to crystalline (rhombohedral) phase transition, which is mainly due to

the increase of mobility, rather than of carrier concentration [110, 111]. It is also

shown in Figure 4.6 that, with an increase of the heating rate, the drop of resistivity

and, therefore, the phase transition itself is shifted to slightly higher temperatures,

as follows from the discussion of the crystallization kinetics (see section 4 of this

chapter). Interestingly, the resistances in the crystalline state do not show any

dependence on the annealing temperature in the investigated temperature range.

The results are reasonably consistent with the optical reflectivity results of the

films (see Figure 4.5). This is because crystalline GeTe is a p-type semiconductor

and the Fermi level lies inside the valence band, resulting in metal-like behavior

[112, 113]. Furthermore, the transition temperature around 240°C for the film

heated at a heating rate of 10 K/min is in good agreement with the results of XRD

measurements, as discussed above, although a SiO2 buffer layer had to be added for

electrical insulation (see inset of Figure 4.6).

Figure 4.6: Resistance of 60 nm thick as-deposited GeTe films measured with heating rates of

5 K/min, 10 K/min and 20 K/min. The solid lines are only intended as guides for the eye. The inset

shows a schematic view of the samples used for resistance measurements.

0 50 100 150 200 250 30010

2

103

104

105

106

107

Re

sis

tan

ce

(

)

Temperature (°C)

5 K/min

10 K/min

20 K/min

Si (100)

SiO2 500 nm

GeTe 60 nmAu 20 nm


50

Effect of O incorporation on structural transitions of GeTe films 4.1.2

The phase-transformation of PLD-deposited GeTe films by an ex situ, thermally-

induced crystallization mechanism was described in detail in the previous section.

Based on experiences with GST alloys, O-doping may contribute to improve the

overwrite cyclability, stability and crystallization rate [114-117]. Strategies for GeTe

films doping by N, C, Cu, O were also explored [21-23, 118]. However, there are

limited reports about environmental oxidation of GeTe films [23, 119, 120]. Thus, it

is relevant to understand the influence of O incorporation on the structural

transitions of binary alloy GeTe films upon ex situ thermal treatment. This influence

is investigated in the following taking into account the properties of the pure GeTe

films described in the previous Chapter 4.1.1.

1.) Chemical composition

The Ge-Te-O (GeTeO) films were deposited by nanosecond-PLD on Si substrates

at room temperature. The number of laser pulses used was 1200, leading to a

deposition rate of 0.5 nm/s and a film thickness of 60 nm. The chemical

composition of the films was determined quantitatively by EDX. Exemplarily, an

EDX spectrum of a GeTe film with O incorporation is shown in Figure 4.7. The

characteristic X-ray lines from Ge, Te and O were detected, and the composition of

the films was calculated to be Ge34Te50O16. Considering the experimental error of

the EDX technique, the composition of the films is therefore 16 at.% O

Figure 4.7: EDX spectrum of an as-deposited 60 nm thick GeTeO film on a Si substrate. Inset: SIMS

depth profile of the film (negative secondary ions).

0 50 100 150 200

102

103

104

105

Inte

nsity (

counts

)

Sputtering time (s)

O

Si

Ge

Te

0 2 4 6 8 10 12

GeL

SiK

GeK

GeK

OK

Inte

nsity [arb

. unit]

Energy [keV]

TeL

TeL

TeL


51

incorporation in GeTe, as compared to the nominal composition of the used target

(GeTe). The O content can be partly explained by surface oxidation, but a

substantial amount is transferred from the target (as the composition of the target is

Ge46.32Te43.72O9.96). SIMS measurements, as shown in the inset of Figure 4.7, were

carried out to confirm that an O incorporation in GeTe exists throughout the

complete film. Apart from the oxidation of the film surface and the oxide at the

interface between film and substrate, the O content is almost constant over the

entire film thickness. Because the atomic radius of O (r = 66 pm) is relatively small,

compared to those of Ge (r = 122 pm) and Te (r = 137 pm) [121], O atoms in the

GeTeO films are expected to be located only at the tetrahedral interstitial sites rather

than as substitutes for Ge or Te atoms.

The series of SEM images in Figure 4.8 shows the surfaces of an as-deposited

GeTeO film and of GeTeO films after thermal treatment at different temperatures.

Starting with the as-deposited film and then annealing at up to 280°C the surfaces

of films appear to be comparably smooth. However, after annealing at 350°C, gap

formation set in and the gaps significantly enlarge with further increases in

temperature up to 380°C. Following annealing at 450°C, the films obviously

disintegrated completely by evaporation, except for a small number of pre-existing

particulates (originating from the laser ablation process itself), with a diameter

between 0.3 and 10 µm; these can still be discerned on the surface of the sample.

Figure 4.8: SEM images of GeTeO films as a function of annealing temperature.

260oCAs-dep.

350oC 380oC 450oC

280oC


52


XRD measurements were performed to obtain structural information about the

films. Figure 4.9a presents XRD spectra of an as-deposited GeTeO film and of

GeTeO films annealed at different temperatures. In each case, the films were

annealed for 20 min in high vacuum. The corresponding annealing temperatures are

shown in Figure 4.9a. The XRD spectrum of the as-deposited GeTeO film shows a

broad peak in the 2θ range between 25 to 30°, suggesting an amorphous state of the

film. XRD spectra of the films annealed at 200 and 260°C resemble that of the as-

deposited film. However, when the annealing temperature reaches 280°C,

reflections at angles of 26.1°, 29.9° and 43.3° emerge that can be attributed to the

lattice planes (021), (202) and (220) of the rhombohedral GeTe phase (JCPDS no.

47-1079), respectively. The appearance of these reflections marks the structural

transformation from the amorphous to the crystalline state. After a slight increase of

Figure 4.9: (a) XRD spectra of as-deposited GeTeO film and GeTeO films with a film thickness of

60 nm annealed at different temperatures. The corresponding annealing temperatures are denoted.

The Miller indices correspond to the rhombohedral crystal structure of GeTe with the space group

R3m (JCPDS no. 47-1079). (b) Corresponding (202) lattice plane spacing d.

280 300 320 340 360 38029.85

29.90

29.95

30.00

Temperature (°C)

2(

°)

(202)

0.2975

0.2980

0.2985

d

(nm

)

(a)

(b)

20 30 40 50

as-dep.

200°C

260°C

280°C

300°C

350°C

380°C

Inte

nsity (

arb

. u

nits)

2 (°)

450°C(22

0)

(20

2)

(02

1)


53

the annealing temperature to 300°C, peak intensities reached maximum values. The

intensities gradually decreased with a further increase of annealing temperature to

380°C. This intensity decrease can be attributed to partial evaporation of the films

in the vacuum oven environment. Finally, for films annealed at 450°C, diffraction

peaks no longer occur due to complete evaporation of the films (demonstrated by

SEM measurements in the previous chapter). It should be noted that, despite the

significant O content of the films, no additional reflections related to crystalline

oxide phase formation appeared. Thus, as confirmed by the XRD measurements

presented in Figure 4.9a, phase segregation could be detected. However, for the

crystalline state of the films, it is interesting to note that all the peak positions

shifted to lower diffraction angles with increasing annealing temperature, as shown

in Figure 4.9b. Such peak position shifts were not observed with the GeTe films

described in Chapter 4.1.1 (Figure 4.2). Regarding, for instance, the (202)

reflection, its position in 2θ is 30.02° at an annealing temperature of 280°C and

decreases to 29.87° at an annealing temperature of 380°C. This shift of peak

position with increasing annealing temperature is caused by a change of the lattice

parameter of the thin film. Using Bragg’s law, the lattice spacing d of the crystal

plane (202) was calculated and plotted as a function of the annealing temperature in

Figure 4.9b. The d value increases from 0.2973 to 0.2988 nm as the annealing

temperature increases from 280 to 380°C. Regarding the crystallization of the

investigated GeTeO films, here, phase transformation from the amorphous to the

rhombohedral crystalline state only starts at 280°C. This is a much higher transition

temperature compared to those of films with slightly Te-rich stoichiometry prepared

by magnetron sputtering and based on different stoichiometry, such as GeTe and

GeTe4. It is also higher in comparison to the transition temperature of films

discussed in Chapter 4.1.1 (a comparison of data is given in Table 4.1 in section 4

of this Chapter 4.1.2).

3.) Topography

The topography and roughness evolution of the annealed GeTeO films at different

annealing temperatures (Figure 4.10) is provided by AFM measurements. With

increasing temperature, from RT to 260°C, no large differences in the topography

and roughness of the GeTeO films appear (not shown here). However, some grains


54

clearly emerge when the annealing temperature is increased to 280°C, as in Figure

4.10a. A further increase of the annealing temperature leads to enhanced grain

growth with small protrusions appearing at 350°C and 380°C, as shown in Figure

4.10. This agrees well with the result of XRD studies for annealing temperatures

below 280°C; the GeTeO films are in the amorphous state and the crystalline state

appears after annealing at a minimum temperature of 280°C. Correspondingly, the

rms surface roughness of the films shows a gradual increase from 1.24 to 4.18 nm

with the increase of the annealing temperature from 280 to 380°C, which is related

to the heterogeneous crystal growth process. In comparison with the pure GeTe

films (in Figure 4.3), the topographies differ, although the tendencies for crystalline

grain growth with increasing temperature are similar. Rationally, the influence of

incorporated O in a GeTe film may serve to suppress the crystal growth process,

because the incorporated oxygen atoms probably inhibit the nucleation of crystals

GeTe elements clusters and it continues growing. That can be due to a smaller grain

size (in Figure 4.10) as well as to a higher critical crystallization temperature of a

GeTeO film, in contrast to a pure GeTe film. Furthermore, the surface roughness

values of GeTeO films are higher than those of pure GeTe films. This assumes that

Figure 4.10: AFM images (500×500 nm2) of GeTeO films with a thickness of 60 nm annealed at

different temperatures: (a) 280°C, (b) 300°C, (c) 350°C and (d) 380°C.

rms = 2.08 nm

rms = 2.33 nm

Z = 13 nm

Z = 17 nm Z = 25 nm

rms = 4.18 nm

a) b)

c) d)

Z = 9 nm

0 nm

rms = 1.24 nm


55

O can partially enter into the grain boundaries and exert bi-axial strain on the

crystal; this influences the growth direction of the materials, according to Ref. [23].

4.) Crystallization kinetics

In Figure 4.11a, the as-deposited GeTeO films were investigated using DSC at

various heating rates. With increasing heating rate, the crystallization temperature

Tc shifts gradually to higher values, i.e., from 258.8°C at 5°C/min to 281.1°C at

100°C/min. This indicates that the crystallization of GeTeO films proceeds in

diffusive nucleation and growth within the range of the heating rate used [23, 122].

The influence of the heating rate on the crystallization temperature Tc also shows

that the crystallization process is not independent of time [123]. According to

Kissinger’s formula (as described in Equation 4.1), an activation energy Ea of

crystallization of 3.18 eV was extracted by fitting the slope in the Kissinger plot in

Figure 4.11: (a) DSC diagrams of GeTeO films for different heating rates. (b) Kissinger plot of the

crystallization from amorphous to rhombohedral phase reveals the activation energy Ea for

crystallization of GeTeO films.

240 250 260 270 280 290 300

50 K/min

Tc= 272.5

oC

70 K/min

Tc= 275.5

oC

He

at

Flo

w (

arb

.un

its)

Temperature (oC)

5 K/min

10 K/min

20 K/min

30 K/min

Tc= 264

oC

Tc= 262.6

oC

Tc=258.8

oC

Tc= 266

oC

Tc= 281.1

oC

Heating rate

100 K/min

1.80 1.82 1.84 1.86 1.88

-11.0

-10.5

-10.0

-9.5

-9.0

-8.5

-8.0

ln(a

/T2 c)

1000/Tc(K)

Ea = 3.18 eV

(a)

(b)


56

Figure 4.11b. Comparison of the Ea and Tc values of GeTeO and GeTe films in this

study, as well as Ea and Tc values for GeTe and GST films reported in literature, are

summarized in Table 4.1. It has to be considered that a high Tc, together with a high

Ea, means high data retention. Based on the data in Table 4.1, GeTeO films exhibit

the highest Ea and Tc values, indicating that they possess the highest thermal

stability, i.e., higher data retention.

5). Optical reflectivity

Figure 4.12a shows the results of reflectivity measurements from the near

ultraviolet to the near infrared (wavelength range from 360 to 800 nm) of both as-

deposited and annealed GeTeO films. The reflectivity of as-deposited GeTeO films

varies from 55 to 58%, whereas the reflectivity of the films annealed at a

temperature of 280°C shows a significant increase and ranges from 68 to 78%.

However, a gradually decreasing reflectivity, ranging from 69 to 72%, appeared

when the annealing temperature increased from 280 to 380°C. One probable reason

for the decrease of the optical reflectivity of films annealed at increasing

temperature could be the greater surface roughness, according to AFM

measurements. On the other hand, the average film thickness slightly decreased

with increased annealing temperature due to evaporation. This may also contribute

to the observed decrease in reflectivity in Figure 4.12a. With a further increase of

the annealing temperature up to 450°C, the reflectivity of the films shows a

significant decrease, which is similar to the tendencies of a pure GeTe film at 380°C

Table 4.1: Comparison of Tc and Ea of GeTeO, GeTe and GST films.

Material Tc (°C) Ea (eV)

GeTeO 280 3.18

GeTe 240 3.14

GeTe4 a 232 2.7

GeTeb 184.8 -

GSTc 154.8 2.2

[a] H. Jiang et al., J. Appl. Phys. 109, 066104 (2011).

[b] N. Yamada et al., J. Appl. Phys. 69, 2849 (1991).

[c] L. Perniola et al., IEEE Electron Device Lett. 31, 488 (2010).


57

or a bare Si substrate (presented in Figure 4.5). The result is explained by the

evaporation of a large film fraction, as demonstrated by SEM images (section 1 of

this Chapter 4.1.2). As shown in Figure 4.12b, the corresponding reflectivity

contrasts of the GeTeO films, determined at a wavelength of 650 nm, were 18.8%

above the critical crystallization temperature within the range between 280°C and

350°C.

In situ thermally induced crystallization of GeTe films 4.1.3

The quality of phase-change films potentially benefits from the reduction of energy

consumption when switching between amorphous and crystalline states with respect

to RESET and SET operations, and from increasing the optical and electrical

contrast in terms of the high signal-to-noise ratio, leading to higher distinction

Figure 4.12: (a) Optical reflectivity spectra of as-deposited GeTeO film and of films annealed at

different temperatures, (b) reflectivity contrast at a wavelength of 650 nm of GeTeO films annealed

at different temperatures, compared to that of GeTe films. The dashed line is only intended as a

guide for the eye.

0 100 200 300 400

0

4

8

12

16

20

24

GeTe

Reflectivity c

ontr

ast (%

)

Temperature (°C)

GeTeO

400 500 600 700 800

56

60

64

68

72

76

80

84

As deposited

450°C260°C

380°C

350°C

Reflectivity (

%)

Wavelength (nm)

280°C

(a)

(b)


58

levels capability in multilevel memory applications [120]. Unlike the previously

described ex situ thermally induced crystallization, in situ thermally induced

crystallization of GeTe films was also explored in order to improve the film quality.

During PLD film growth, the growth temperature and the choice of the substrate

have a crucial influence on the growth mechanism [26].

In this subchapter, GeTe thin films PLD-deposited on Si(100) substrates (with a

very thin native surface oxide layer) at different substrate temperatures are

investigated. The effect of growth temperature on GeTe thin films’ structure,

morphology and reflectivity is characterized by using XRD, AFM and UV-Vis

spectrophotometry.


Figure 4.13 shows the XRD spectra of GeTe films deposited on Si(100) at different

substrate temperatures that ranged between 100 and 290°C. The diffraction

spectrum of the GeTe film deposited at 100°C is typical for the amorphous state.

However, the onset of polycrystalline phase formation appears at the growth

temperature of 130°C, as can be seen from the significant presence of (021) and

(220) reflections of GeTe in Figure 4.13. With a further increase of the substrate

temperature to 160°C, the dominant XRD reflections depict a transition to

preferably (003) oriented films. A continuous increase of the intensity of the (003)

reflection is clearly visible at a substrate temperature of 180°C. GeTe film growth at

substrate temperatures in the range from 200 to 240°C exhibits a significant

Figure 4.13: XRD spectra of GeTe films deposited on Si(100) substrates at different temperatures.

20 30 40 50 60

Si

(006)

200°C

Inte

nsity (

cps)

2(°)

130°C

150°C

160°C

180°C

290°C

220°C

240°C

100°C

(003)

(220)

(202)

(021)

(024)

(04

2)


59

increase of the preferred (003) reflection intensity, which is so intense that also the

corresponding second order reflection is present and of similarly high intensity.

Thus, a pronounced (003) preferred orientation has formed. For films deposited

above this temperature range (e.g., at 290°C) reflections are no longer detectable,

due to the small amount of remaining material, a consequence of the enhanced

desorption of atoms at such high substrate temperatures.

A more detailed analysis of the films at growth temperatures in the range from

200 to 240°C is shown with a zoom of the (003) reflection in Figure 4.14. It appears

that the reflection possesses an extremely high intensity of more than 104 cps,

indicating a higher crystallinity of these films. Based on the Scherrer formula, the

calculated average grain size of these highly oriented polycrystalline films is 4.7 nm

Figure 4.14: GeTe(003) reflection of the GeTe films on Si(100) deposited in the substrate

temperature range from 200 to 240°C.

Figure 4.15: GeTe(003) rocking curves of films deposited at different substrate temperatures.

23 24 25 26 27 28

0

1x104

2x104

3x104

4x104

(111

)

Inte

nsity (

cp

s)

2 (°)

200°C

220°C

240°C

(003

)

-2 -1 0 1 2

0

1x103

2x103

3x103

Inte

nsity (

cp

s)

(°)

200°C, FWHM: 1.9°

220°C, FWHM: 1.3°

240°C, FWHM: 1.6°


60

at 200oC, 7.1 nm at 220°C and 5.9 nm at 240

oC. From the corresponding rocking

curves of the GeTe (003) reflection, as shown in Figure 4.15, the FWHM was

extracted as 1.9° at 200oC, 1.3° at 220

oC and 1.6° at 240

oC. For polycrystalline

films exhibiting a preferred orientation, these values are relatively low. In order to

clarify the present crystalline texture of such highly oriented films, an in-plane pole

figure measurement was performed of a representative film deposited at a growth

temperature of 220°C. The pole density distribution in this pole figure (Figure

4.16a) exhibits a concentric ring at an α angle of 76°. Thus, the film is characterized

by a pronounced fiber texture with the [003] direction being the fiber axis. The pole

figure also contains eight distinct pole density maxima at α = 46° as well as

additional four pole density maxima at α = 76°. These are identified by the modeled

pole figure in Figure 4.16b to originate from the Si substrate and not from the film.

In conclusion, GeTe film growth in the narrow temperature window of 200-240°C

leads to a pronounced fiber texture of the films. Furthermore, at a growth

temperature of 220°C, the films possessed the highest reflection intensity and the

smallest rocking curve FWHM.

2.) Topography

In order to increase the understanding of polycrystalline GeTe film growth by PLD

in terms of the topography, AFM measurements on a series of selected films were

carried out (summarized in Figure 4.17). It is clear that different topographies result

from different growth temperatures in the range between 130 and 240°C. At a

Figure 4.16: (a) In-plane pole figure of the GeTe film on Si substrate using the {422} reflection at

2θ= 71.27°. Note that the intensity (cps) is depicted on a logarithmic scale. (b) Modeled bare Si

substrate pole figure from the Si(331) reflection.

1.5E+02

5.1E+02

8.6E+02

1.2E+03

GeTe {422}

(a) (b)

β

α

α = 76°

α = 46°

GeTe film

Si sub. Si sub.


61

growth temperature of 130°C, the film surface is relatively smooth in spite of the

several crack-like features evident in Figure 4.17a (rms roughness is 1.4 nm). With

a higher substrate temperature of 180°C, as shown exemplarily in Figure 4.17b, the

film surface is no longer smooth and it exhibits an increased rms roughness of 2.3

nm, similar to that of films deposited at temperatures of 220°C (rms = 2.6 nm) and

240°C (rms= 2.0 nm) in Figure 4.17c and Figure 4.17d, respectively. This trend of

the topography change with rising substrate temperature becomes clearer by

comparing spatial height profiles obtained from representative surface regions of

the films, as depicted in Figure 4.17e. The film surfaces obviously changed, due to

Figure 4.17: AFM images of the topography of GeTe films grown on Si(100) at different substrate

temperatures: (a) 130°C, (b) 180°C, (c) 220°C, (d) 240°C and (e) height profiles taken along the red

lines in the images of (a) to (d).

0 40 80

0

2

4

6

He

igh

t (n

m)

Length (nm)

130°C 180°C

220°C 240°C

a)

(e)

b)

c) d)rms =2.6 nm

Z=19Z=21

Z=18

Z=11

rms =2.0 nm

rms =2.3 nmrms =1.4 nm


62

the fiber-textured growth at deposition temperatures of 220°C and 240°C, in

contrast to the films obtained at 130°C. This illustrates that the higher the level of

the crystallite orientation in the films is, the higher the peak-to-valley values in the

height profiles and, consequently, the greater the surface roughness.


Figure 4.18a shows the optical reflectivity of GeTe films deposited on SiO2/Si(100)

substrates and grown at different temperatures. The reflectivity of as-deposited

films varies from 48 to 52% in the wavelength range from the near ultraviolet to the

near infrared (400-700 nm). In contrast, the reflectivity of films grown at a substrate

temperature of 130°C lies between 64 and 80%, and is thus distinctly higher than

the reflectivity of as-deposited films. A further increase of the growth temperature

Figure 4.18: (a) Optical reflectivity of GeTe films deposited on Si(100) substrates covered with

500 nm SiO2 at different growth temperatures. (b) Reflectivity contrast at a wavelength of 650 nm

between GeTe films annealed in situ during growth at different growth temperatures, compared to

amorphously deposited GeTe films post-growth annealed at different annealing temperatures. The

dashed lines are only intended as guides for the eye.

0 50 100 150 200 250 300 350 400

0

8

16

24

32

Reflectivity c

ontr

ast (%

)

Temperature (°C)

post-anneal GeTe

during growth GeTe

400 500 600 700

50

60

70

80

Optical re

flectivity (

%)

Wavelength (nm)

130°C 180°C

220°C 240°C

as-dep.

(a)

(b)


63

to 180°C obviously does not affect the reflectivity of the GeTe films. Only in the

temperature range between 220 and 240°C does the film reflectivity exhibit a

further increase, with values between 66 and 84% within the measured wavelength

range. For even higher growth temperatures, e.g. 290°C, a dramatic decline of the

optical reflectivity of films appears, caused by decomposition and partial desorption

of film components. The results agree well with those obtained from the structure

analysis presented in section 1 of this chapter. Thus, the higher optical reflectivity

of films deposited at substrate temperatures in the range from 220 to 240°C is

caused by the higher crystalline quality of the GeTe films.

In order to obtain a high signal-to-noise ratio for optical application of GeTe

films, a high reflectivity contrast is required. The reflectivity contrast, determined at

a wavelength of 650 nm, for as-deposited amorphous films and films grown at

elevated substrate temperatures, compared to post-growth annealed GeTe films, is

presented in Figure 4.18b. An optical reflectivity contrast of 28% between

amorphous and polycrystalline GeTe films was obtained at lower growth

temperatures, and of approx. 31.5% between amorphous and fiber-textured GeTe

films at higher growth temperatures. These results demonstrate that the reflectivity

contrasts of the films obtained from fiber-textured growth (in situ thermal

treatment) are approx. 10% higher than for the polycrystalline films after post-

growth annealing (ex situ thermal treatment). Moreover, the critical crystallization

temperature of films with in situ thermal treatment lies at 130°C, which is close to

100°C lower than that of amorphous films crystallized by ex situ thermal treatment

at 220-240°C (see Chapter 4.1.1). Consequently, a considerably higher optical

reflectivity contrast (31.5%) of the fiber-textured GeTe films is achieved with

relatively lower crystallization temperature. Such a high optical reflectivity is not

only beneficial for the increase of the signal-to-noise ratio, but also promising for

multilevel PCM based data storage applications.

Epitaxial growth of GeTe films on BaF2(111) 4.1.4

A further improvement of the crystalline quality can be realized by PLD deposition

of epitaxial GeTe films. A small lattice mismatch between GeTe lattice and the

lattice of the substrate is the main condition. Consequently, freshly cleaved

BaF2(111) is used as a substrate, to provoke epitaxial GeTe film growth, because


64

the lattice mismatch is only about 3.5 %. Based on the results in the previous

chapter 4.1.3 regarding the structural and optical properties of GeTe films deposited

at different substrate temperatures, the substrate temperature was selected to be

205°C, constantly, for the epitaxial growth experiments. The deposition rate in this

case was 0.2 nm/s. The typical film thickness was again 60 nm.

1.) Composition and morphology

The chemical composition of a representative PLD-deposited GeTe film on a

BaF2(111) substrate was determined from EDX, as shown in Figure 4.19. The

characteristic x-ray peaks from the elements Ge, Te, Ba and F are present and the

resulting stoichiometry of the film is determined to be Ge50.1Te49.9. As the measured

composition is almost identical to the specified target composition, this indicates

that a stoichiometric transfer from the ablation target to the film was easily

achievable. The topography of GeTe films on BaF2(111) was measured by AFM at

a scanning range of 1×1 um², as presented in the inset image of Figure 4.19. The

film surface is relatively smooth, with a rms roughness of ~1.0 nm.


Figure 4.20 presents a representative XRD spectrum of a crystalline GeTe film

deposited on a BaF2(111) substrate at a substrate temperature of 205°C. The main

reflections of the GeTe film, well distinguishable from those of the BaF2(111)

substrate, are situated at 2θ angles of 25.1° and 51.4°, which are attributed to the

Figure 4.19: EDX spectrum of a GeTe film deposited on a BaF2(111) substrate. Inset: AFM

topography image.

0 2 4 6 8 10 12

BaL

BaL

BaL

FK

GeL

TeL

GeK

Inte

nsity (

arb

.unit)

Energy (keV)


65

(111) and (222) lattice planes of rhombohedral β-GeTe, respectively. This result

suggests that the GeTe film grows in the β-phase. This is in agreement with

previous studies on GeTe films growing preferably in the cubic [111] direction

when providing (111) oriented, single crystalline substrates [124, 125].

A more detailed analysis is possible on the basis of XRD pole figure

measurements. The in-plane pole figure using the {422} reflection of a

representative GeTe film from the (111) orientation-dominated growth regime is

presented in Figure 4.21a. The intensity is plotted on a logarithmic scale. Six

intense {422} pole density maxima at a polar angle α of 76° are clearly visible, each

exhibiting a three-fold symmetry. Each couple of pole density maxima is

azimuthally separated by 40°, and the centers in between the couples are separated

by an azimuthal angle β = 120°. This is in good agreement with the modelled

GeTe{422} pole figure shown in Figure 4.21b. It is noted that the remaining, much

weaker intensity pole density maxima in the pole figure in Figure 4.21a are intensity

contributions originating from the BaF2 substrate. This is confirmed by the pole

figure of a bare BaF2(111) substrate (Figure 4.21c), measured at exactly the same 2θ

angle of 71.27° as the pole figure in Figure 4.21a. Unequivocally, the six high-

intensity pole density maxima present in Figure 4.21a are not observed in Figure

4.21c. This result is also clearly explained by selected φ scans of the GeTe film and

of the BaF2(111) substrate shown in Figure 4.21d, which were measured at the

identical 2θ angle of 71.27° (corresponding to the GeTe{422} reflection) and the

Figure 4.20: XRD spectrum of a GeTe thin film deposited on BaF2(111) at a substrate temperature of

205°C.

20 30 40 50 60

102

103

104

105

(222)

BaF2

GeTe (222)GeTe (111)

Inte

nsity (

cps)

2(°)

BaF2

(111)


66

identical polar angle α of 76°. The red-colored φ scan presents the previously

described 40° interval between couples of most intensive peaks with a three-fold in-

plane orientation symmetry of the GeTe film, as well as six much weaker peaks

consistent with the peaks from the BaF2(111) substrate appearing in the black-

colored φ scan. Therefore, a distinguishable epitaxial relationship between the α-

GeTe film and the BaF2(111) substrate is confirmed. This epitaxial relationship is

given by α-GeTe(111) || BaF2(111) and α-GeTe[1-10] || BaF2[1-10].

To estimate the average tilt distribution of crystallites in the films, rocking

curves of the β-GeTe(111) reflection were measured. One representative

measurement is shown in Figure 4.22. The very narrow peak in the middle of the

rocking curve measurement is a contribution from the β-GeTe(111) reflection. The

FWHM is 0.044°, indicating a rather small degree of mosaicity of the epitaxial

Figure 4.21: (a) In-plane pole figure of the GeTe film at the {422} reflection at 2θ = 71.27°. Intensity

(cps) is depicted on a logarithmic scale. (b) Modeled GeTe{422} pole figure provided by the CaRIne

Crystallography software. (c) Corresponding bare BaF2 substrate pole figure measured at 2θ =

71.27o. (d) φ scans of the GeTe film and of the BaF2 substrate, both measured at 2θ = 71.27° and α =

76°.

Inte

nsity (

cps)

(a) (b)

(c) (d)

BaF2 sub.

GeTe film

100

1.0E+03

1.1E+04

1.0E+05

GeTe {422}

3.5E+02

2.4E+03

1.6E+04

1.0E+05

GeTe {422}

BaF2 sub.

β

α

α = 76°

0 80 160 240 320

102

103

(°)

BaF2

102

103

104

105

80° 40°

GeTe


67

GeTe film. Compared to the FWHM in the case of fiber-textured growth of GeTe

film on Si(100) (~1.3°) and in the case of an epitaxial GST film on a BaF2(111)

substrate (~0.2°) [12], a reduction was registered.

Summary of the results on thermal treatment-induced crystallization of 4.1.5

GeTe films

The process of PLD-deposited PCM films on Si substrates at room temperature has

been established. The ex situ thermally induced crystallization of GeTe and GeTeO

phase-change films was systematically investigated, including structure analysis,

topography, crystallization kinetics, as well as optical and electrical properties. The

obtained values for the crystallization temperature Tc, the activation energy for

crystallization Ea, the optical reflectivity contrast ORa-c and the electrical resistivity

contrast ERa-c are summarized in Table 4.2. The experimental results reveal that

PLD-deposited GeTeO films exhibit a higher crystallization temperature (280°C)

and a slightly higher activation energy for crystallization (3.18 eV), compared to

pure GeTe films. The achievable higher thermal stability makes them excellent

Figure 4.22: β-GeTe(111) rocking curve of an epitaxial GeTe film on BaF2(111).

Table 4.2: Comparison of crystallization temperature Tc, activation energy Ea, optical reflectivity

contrast ORa-c and electrical resistivity contrast ERa-c between PLD-deposited GeTe and GeTeO

films.

PLD-deposited

films

Tc

[°C]

Ea

[eV]

ORa-c

[%]

ERa-c

GeTe 240 3.14 21 103 – 10

4

GeTeO 280 3.18 18.8 _

12.4 12.6 12.8

5.0x103

1.0x104

1.5x104

2.0x104

2.5x104

Inte

nsity (

cp

s)

(°)

FWHM: 0.044°

β-GeTe (111)


68

candidates for phase change applications.

The quality of PLD-deposited GeTe films on Si(100) upon in situ thermal

treatment was investigated. A structure evolution from as-deposited amorphous over

polycrystalline (130°C) to fiber textured (220-240°C) in GeTe films during growth

at elevated temperatures was demonstrated. It should be noted that, following the

in situ thermal treatment, the GeTe films showed a critical temperature of

crystallization Tc that was approx. 100°C lower than in ex situ thermally treated

films. Moreover, the higher the achieved crystalline quality of the GeTe films was,

the higher the obtained optical reflectivity. The optical reflectivity contrast of

polycrystalline and fiber textured GeTe films, compared to amorphous films, was

approx. 28% and 31.5%, respectively. The latter is 10% higher than that of ex situ

thermally induced crystalline films (21%). Finally, an improvement of the film

quality, in order to influence the crystallization behavior, was achieved by the

epitaxial growth of GeTe films on BaF2(111) at a substrate temperature of as low as

205°C. The epitaxial growth of GeTe films was confirmed by X-ray texture analysis

measuring GeTe{422} pole figures. These epitaxial films exhibited a relatively

small polar mosaicity component of as low as 0.044o, obtained from the FWHM of

rocking curve measurements.

4.2 Laser-induced phase transitions of GeTe and GST films

69

Laser-induced phase transitions of GeTe and GST 4.2

films

The application of phase-change materials for data storage relies on the fast

reversible transformation between the disordered amorphous and ordered crystalline

states, which can be achieved by local heating/cooling at high rates, either with a

laser beam or electrical pulses. As described in Chapter 2.2, distinct chalcogenides

such as GeTe and GST, which are highly attractive for such applications, exhibit

markedly different crystallization behaviors. Compared to the thermally annealing

treatment-induced crystallization discussed in Chapter 4.1, laser pulse induced

crystallization is characterized by much higher heating/cooling rates (> 109 K/s [66,

126, 127]) and resembles more closely practical crystallization kinetics in PCMs.

Furthermore, regarding the physical process of laser-induced phase transition, ns

laser pulse irradiation can be considered as a thermal process, whilst upon

excitation with ultrashort pulses, i.e., fs laser pulses, non-thermal effects are

probably involved in the phase transition process [128-130]. Therefore, it is of

extreme relevance to investigate in detail the phase transformation of phase-change

films for the particular case of laser-induced crystallization, as well as to compare

the crystallization process in the time scale frame of ns and fs single laser pulses by

variation of the laser fluence at the same wavelength.

In this chapter, results and conclusions about the laser-induced phase transition

of phase-change films are presented [127, 131]. The chapter is divided into three

sections: ns laser-induced phase transition in GeTe films (Chapter 4.2.1),

crystallization of GST films by ns and fs single laser pulse irradiation (Chapter

4.2.2), and simulation of temperature distribution after pulsed laser irradiation

(Chapter 4.2.3).

Nanosecond laser-induced phase transition in GeTe films 4.2.1

1.) Crystallization process

Crystallization as a function of pulse number

Figure 4.23a shows the optical reflectivity of an as-deposited amorphous GeTe

film and of films after 20 ns laser pulse irradiation with pulse numbers between 0


70

and 30 at a constant fluence of 130 mJ/cm2 and a repetition rate of 10 Hz. The

optical reflectivity of the as-deposited film lies in the range between 72 and 74% for

wavelengths between 400 and 700 nm. The reflectivity of the film is slightly

increased, by 3%, when the number of applied laser pulses increases to 3. A more

pronounced effect is discernable after irradiation with 5 pulses. For this film, the

reflectivity is in the range between 85 and 88%, which is significantly higher than

that of the as-deposited film. Upon further increase of the number of applied laser

pulses the reflectivity increases further, but the effect is no longer as obvious. The

corresponding results from XRD analysis are depicted in Figure 4.23b. The XRD

patterns of an as-deposited film and of a film irradiated with 3 laser pulses only

show a several degrees broad and weak peak centered around 28.3°, typical for the

Figure 4.23: (a) Optical reflectivity of as-deposited GeTe film and films after 20 ns laser pulse

irradiation with a fluence of 130 mJ/cm2 as a function of the number of pulses. The inset shows the

corresponding optical microscope image of the irradiated area after 5 pulses. (b) Corresponding

XRD spectra of GeTe films as a function of the pulse number.

20 30 40 50

20 pulses

10 pulses

5 pulses

3 pulses

Inte

nsity (

arb

. u

nits)

2(°)

(220

)

(200

)

as-dep.

(111

)

Si

400 500 600 70070

75

80

85

90 30 pulses 20 pulses

10 pulses

5 pulses

as-depositedRe

fle

ctivity (

%)

Wavelength (nm)

3 pulses

(a)

(b)

0 N

5 N


71

amorphous state of GeTe films. In contrast, when the film is irradiated by more than

5 pulses, clearly visible Bragg peaks located at 26.4°, 30.2° and 43.7° appear, which

correspond to the (111), (200) and (220) reflections of cubic phase GeTe,

respectively. Therefore, at least 5 laser pulses (not excluding the case of 4 pulses)

are required to induce the crystallization process. The XRD results are in good

agreement with the optical reflectivity measurements. Consequently, it can be

concluded that an increase in the number of laser pulses continuously increases the

crystalline fraction of the films. It can be assumed that the crystallization process is

incomplete after 20 ns single pulse laser irradiation [31], because the diffusional

processes involved in crystallization require sufficient time above the threshold

temperature typical of the GeTe material. According to this, the superposition of

laser pulses leads to an increasing crystallized volume and, consequently, to the

increase of the optical contrast.

Crystallization as a function of laser fluence

In order to examine in detail the influence of the laser fluence on the phase

transition, a series of fluence-dependent optical reflectivity and XRD measurements

of GeTe films was performed. A laser pulse number of 20 at a pulse duration of

20 ns was chosen for all samples, in order to obtain comparable data. The optical

reflectivity was evaluated at a constant wavelength of 650 nm, as presented in

Figure 4.24a. With fluences < 11 mJ/cm2, only a slight increase in reflectivity is

observed. However, the reflectivity rises abruptly to 86 % after laser irradiation

with a fluence of 14 mJ/cm2. Subsequently, the reflectivity increases only slowly

upon a gradual fluence increase from 14 and 130 mJ/cm2, indicating a relatively

large stable region with high optical contrast (up to 130 mJ/cm²). From these

results, it can be deduced that the fluence threshold for crystallization of films

irradiated with 20 pulses lies between 11 and 14 mJ/cm2. Above this threshold, the

series of laser pulses heats the films above the crystallization temperature. The

atoms in the GeTe films thus become increasingly mobile and reach the

energetically favorable crystalline state, leading to a partial crystallization of the

film. With an increase of laser fluence, the degree of crystallization increases,

corresponding to the increase of film reflectivity. The highest optical contrast

between an as-deposited GeTe film and a laser-irradiated GeTe film is calculated to


72

be ~15%, which is much higher than the optical contrast values for laser-induced

phase transition of GST [66]. From the view of application, the higher reflectivity

contrast, the higher signal-to-noise ratio and the high absolute reflectivity of the

films are independent of the irradiation fluence in the range between 36 and

130 mJ/cm2, which indicates a high stability of optical contrast and, as a

consequence, a long data retention. Also, as demonstrated in Figure 4.24a, the

reflectivity of the laser-irradiated films with fluences ≥ 130 mJ/cm2 drops to 82%.

This is attributed to ablation of the film surface, as verified by the SEM image in

the inset of Figure 4.24a. Holes in the film with sizes ranging from 200 nm to 1 μm

are clearly visible. Additionally, an influence of laser-induced crystallization on the

film surface roughness can be excluded, based on evidence provided in Figure 4.25.

Figure 4.24: (a) Optical reflectivity of GeTe films at the wavelength of 650 nm after irradiation with

20 laser pulses (pulse duration: 20 ns) as a function of the laser fluence between 4 and 162 mJ/cm2.

The films were amorphous in the as-deposited state. The inset shows an SEM image of the film

topography after laser irradiation with 20 pulses and a fluence of 162 mJ/cm2. (b) XRD spectra of

GeTe films after laser irradiation with 20 pulses as function of the laser fluence between 4 to

130 mJ/cm2. The Miller indices correspond to the rhombohedral crystal structure of GeTe with the

space group R3m (JCPDS no. 47-1079). (c) Details of the XRD patterns around the 2θ position of the

rhombohedral (202) or/and cubic (200) reflections, respectively. (d) Corresponding interplanar

spacings d of the rhombohedral (202) planes or/and cubic (200) planes as a function of the laser

fluence. The dashed line serves only as a guide for the eye.

0 40 80 120

0.294

0.296

0.298

d (

nm

)

Laser fluence (mJ/cm2)

0 40 80 120 16070

75

80

85

90

Re

fle

ctivity (

%)


20 30 40 50

Inte

nsity (

arb

.un

its)

2 (°)

(00

3)

98 mJ/cm2

4 mJ/cm2

(22

0)

(20

2)

130 mJ/cm2

78 mJ/cm2

54 mJ/cm2

36 mJ/cm2

as-dep.

(02

1)

26 mJ/cm2

22 mJ/cm2

17 mJ/cm2

14 mJ/cm2

11 mJ/cm2

(02

4)

(a) (b)

29.5 30.0 30.5 31.0

Inte

nsity (

arb

.un

its)

2 (°)

98 mJ/cm2

4 mJ/cm2

130 mJ/cm2

78 mJ/cm2

54 mJ/cm2

36 mJ/cm2

as-dep.

26 mJ/cm2

22 mJ/cm2

17 mJ/cm2

14 mJ/cm2

11 mJ/cm2

(c) (d)


73

It shows that the surface roughness of the ns laser-irradiated GeTe films is quite low

(~0.5 nm) and it remains constant, even when the laser fluence increases, in contrast

to the thermally annealed GeTe films.

Figure 4.24b presents the evolution of XRD patterns of films irradiated with

laser fluences between 0 and 130 mJ/cm2 after irradiation with 20 pulses. The

diffraction spectrum of the as-deposited film is typical for the amorphous state, as

shown before. No apparent differences in the diffraction spectra can be found up to

an applied fluence of 11 mJ/cm2. Above this value, crystalline diffraction peaks at

26.1°, 30.0°, 42.3°, 43.6°, which correspond to the (021), (202), (024) and (220)

lattice planes of the rhombohedral phase, respectively, emerge. The rhombohedral

structure can be interpreted as a slightly distorted cubic phase. This is in good

agreement with the results of the optical reflectivity measurements regarding the

threshold of the optical contrast increase. The intensity of diffraction peaks

increases slightly when the laser irradiation fluence is continuously increased from

14 to 36 mJ/cm². Simultaneously, the intensity of the (024) reflection gradually

disappears (as depicted in Figure 4.24b) because of the transformation from the

rhomhohedral GeTe phase into the rocksalt GeTe phase. Above a fluence of

36 mJ/cm2, only the latter phase is found. The measured fluence dependence

correlates with the assumption that a lower fluence is necessary for construction of

a low-symmetry rhombohedral unit cell than for a high-symmetry cubic unit cell,

because there are large atomic displacements required [132]. Nevertheless, it is

interesting that a peak shift to higher diffraction angles can be observed for fluences

Figure 4.25: Surface roughness of the GeTe films after laser irradiation at various fluences, compared

to the thermal annealing results at different annealing temperatures.

1 2 3 40

50

100

150

200

La

se

r flu

en

ce

(m

J/c

m2)

An

ne

alin

g t

em

pe

ratu

re (

°C)

RMS surface roughness (nm)

250

300

350

400


74

between 14 to 36 mJ/cm². Exemplarily, the shift of the 2θ position of the

rhombohedral phase (202) reflection at 30° towards the cubic phase (200) reflection

at 30.2° is plotted in Figure 4.24c and d. The corresponding interplanar spacing d

changes from about 0.297 nm to about 0.294 nm after this fluence increase (Figure

4.24d). The peak shift to higher diffraction angles or smaller interplanar spacings is

the result of the rhombodedral to cubic phase transformation induced by fluences

between 14 to 36 mJ/cm². It is assumed that this reduction of the interplanar spacing

is probably due to the release of film stress as a result of the transformation process

in this particular fluence range [133]. With further increase of the irradiating

fluence, from 36 to 130 mJ/cm², the interplanar spacing increases slightly, due to

the higher thermal expansion, resulting in the peak shifting back to a relatively low

angle, as shown in Figures 4.24c and d. To summarize these findings, the

crystallization phenomena require several laser pulses instead of only one, even at

the highest laser fluence used in this study.

Crystallization as a function of pulse repetition rate

The influence of the repetition rate in the range from 1 to 400 Hz on the reflectivity

of films triggered by 20 ns laser irradiation with an identical laser fluence of

78 mJ/cm2 (i.e., above the threshold for crystallization) and a fixed pulse number of

20 is shown in Figure 4.26a. Compared to the reflectivity of the as-deposited film, a

prominent increase in the reflectivity of the irradiated films of approx. 15% over the

measured wavelength range between 400 and 700 nm is obtained. As confirmed by

XRD analysis of these films (shown in Figure 4.26b), this increase in reflectivity is

due to the crystallization of the amorphous GeTe film. More interestingly, the

reflectivity increase of the films is consistently high for all chosen repetition rates

between 1 and 400 Hz, corresponding to delay times between pulses of 1 and

0.0025 s, respectively, and the resulting reflectivity curves match almost completely

over the entire measured wavelength range. Coincident with these findings, the

intensities of the GeTe phase diffraction peaks in the XRD measurements in Figure

4.26b are almost identical for the chosen laser pulse repetition rates. This indicates

that the transformation process is independent of the delay time between delivered

laser pulses in the investigated range. This is consistent with the findings of


75

Coombs and co-workers [31], who determined the probability of nucleation to be

almost unity when the delay between pulses exceeds 35 ns. In principle, each laser

pulse induces an irreversible modification in the region of the amorphous film and

contributes to the overall crystallization process; this is the so-called energy

accumulation process [81].

The initial pulses might already cause some

modifications in the form of initial crystal nucleation as well as in the early growth

of nuclei that already existed within the overall amorphous matrix of the as-

deposited film [134]. During the investigations on the crystallization as a function of

the pulse number, it was found that the threshold number of pulses for

crystallization is ≥ 5 pulses, for which a higher optical reflectivity, caused by a

higher crystalline fraction, could only be observed for a higher number of pulses

Figure 4.26: (a) Optical reflectivity of an as-deposited GeTe film and of films after 20 ns laser pulse

irradiation with a fluence of 78 mJ/cm2 and 20 pulses as a function of the pulse repetition rate

varying between 1 Hz and 400 Hz. (b) Corresponding XRD patterns of GeTe films as a function of

the pulse repetition rate varying between 1 Hz and 400 Hz.

20 30 40 50

400 Hz

100 Hz

10 Hz

Inte

nsity (

arb

. units)

2(°)

1 Hz

(22

0)

(20

0)

(11

1)

400 500 600 70070

75

80

85

90

as-depositedRe

fle

ctivity (

%)

Wavelength (nm)

Repetition rate 1 Hz

10 Hz

100 Hz

400 Hz

(a)

(b)


76

(≥ 5). Such a pulse number dependence of the optical reflectivity in GeTe films

could be beneficial for the potential application of multilevel storage. In recently

published literature, a similar characteristic has also been found with AIST and

GST materials [83, 135]. A more detailed discussion follows in Chapter 4.3.

Crystallization as a function of film thickness

Apart from the influence of the pulse laser beam parameters on the crystallization

process presented above, another important influential variable for the application

of phase-change films is the use of very thin films. A reduction of the film thickness

can affect several material properties including crystallization temperature,

crystallization rate, optical constants, etc. [136]. In order to investigate the influence

of the film thickness on laser-induced crystallization, a series of GeTe films with

different thicknesses were prepared. Film thicknesses of 20, 51, 68, and 120 nm

were obtained by precisely controlling the amount of deposited material by using

pulse numbers of 400, 1000, 1200, and 2400 pulses, respectively. As a first example

regarding film thickness dependencies, an as-deposited GeTe film with a thickness

of 51 nm and the same film after 20 ns pulse laser irradiation with 20 pulses at a

fluence of 67 mJ/cm2 were characterized by XRR, as shown in Figure 4.27a. The

film thickness-related oscillations of the XRR curve from the as-deposited film are

present in the whole displayed angular range, whereas the oscillations in the XRR

curve from the laser-irradiated film are characterized by a smaller oscillation period

caused by a shrunken film thickness upon laser irradiation. This decrease in the

thickness of the films is due to the disordered amorphous-ordered crystalline phase

transformation of GeTe films, as evidenced by the XRD patterns of the irradiated

films shown in Figure 4.27c. Moreover, a shift of the critical angle of total

reflection towards higher angles is clearly visible in the XRR curve after laser

irradiation, compared to that of the as-deposited GeTe film. This indicates that the

laser irradiated film possesses a higher density than the as-deposited one, since the

critical angle θc for total reflection is proportional to the square root of the mass

density ρ as defined by the following equation [137]: 𝜃𝑐 ∝ λ√𝜌, where λ is the

wavelength of the used X-ray beam. By corresponding modeling fitting to the

measured data in Figure 4.27a, the films thickness and mass density can be

extracted as 51 nm and 5.5 g/cm3 for the as-deposited film and 45 nm and 6.0 g/cm

3


77

for the post laser-irradiated film, respectively. This means that there is a decrease in

film thickness of about 10% upon laser-induced crystallization.

Figures 4.27b, c and d present the investigation on film thickness-dependent

optical reflectivity and XRD measurements for the 20 ns pulse laser-irradiated GeTe

films at 20 pulses and 67 mJ/cm2. The optical reflectivity contrast of the films at a

constant wavelength of 650 nm is depicted in Figure 4.27b. The reflectivity of the

as-deposited film remains as a reference (i.e., at 0% in terms of reflectivity

contrast). With an increase in film thickness from 20 to 51 nm, the reflectivity

contrast rises significantly. The highest achievable optical reflectivity contrast of

17% (with errors ± 2%) was observed for a film thickness of around 68 nm. With

further increase in thickness up to 120 nm, the reflectivity contrast shows no

obvious change and stays constant within the measurement error at this high

reflectivity contrast level. The corresponding XRD spectra of the laser-irradiated

Figure 4.27: (a) XRR measurements and the corresponding modeled curves for an as-deposited

GeTe film and the film after laser irradiation with 20 pulses and 67 mJ/cm2. The inset shows the

values of density and film thickness extracted from the measurement. (b) Optical reflectivity contrast

of GeTe films at the wavelength of 650 nm after irradiation with 20 pulses (pulse duration: 20 ns)

and 67 mJ/cm2 as a function of the film thickness that varies between 20 and 120 nm. (c) XRD

patterns of GeTe films after irradiation with 20 pulses as a function of the film thickness between 20

to 120 nm. (d) FWHM of the (200) reflection and corresponding calculated average grain sizes as a

function of the film thickness between 20 and 120 nm.

20 30 40 50

Inte

nsity (

arb

.units)

2 (°)

(220)

(200)

(111)

Si

120 nm

68 nm

51 nm

20 nm

as-deposited:experiment data

modeling data

laser-irradiated:experiment data

modeling data

0.2 0.4 0.6 0.8 1.0 1.2

101

102

103

104

105

Inte

nsity (

cps)

(°)

Si (100)

GeTe = 6.0 g/cm

3

t = 45 nm

Si (100)

GeTe = 5.5 g/cm

3

t = 51 nm

laser-irradiation

(a)

(c)

0 20 40 60 80 100 1200.20

0.25

0.30

0.35

0.40

0.45

0.50

0.55

Film thickness (nm)

FW

HM

(°)

20

30

40

50

60

70

80

Gra

in s

ize

(n

m)

(b)

(200)

(d)

0 20 40 60 80 100 120

0

5

10

15

20

Reflectivity c

ontr

ast (%

)

Film thickness (nm)


78

GeTe films with film thickness varied between 20 and 120 nm are shown in Figure

4.27c. First, it is obvious that there are reflections present for all film thicknesses

and the detected crystalline diffraction peaks of GeTe films are annotated. As

expected, the peak intensity of the crystalline phase decreases with decreasing film

thickness. It can be seen that the peak intensity from the 20 nm thick film is very

low due to the lower crystallized volume, when compared to the thicker one. It

should be noted that an increased effect of the interface on GeTe crystallization with

reduced film thickness below a critical thickness is also possibly involved. It was

reported in the literature that the crystallization temperature increases as the film

thickness is reduced below a critical thickness [138, 139]. In Figure 4.27d, the

average grain size of laser-irradiated GeTe films is plotted as a function of the film

thickness, estimated from the FWHM value of the (200) reflection by using the

Scherrer formula. The average grain size decreases with reduction of the film

thickness, with the grain size varying between 21 and 62 nm in the examined film

thickness range. Therefore, it can be concluded that the crystallization process of

GeTe films can be tailored with appropriate control by the irradiated ns pulse laser

parameters (i.e., pulse number, laser fluence) and by the film thickness.

2.) Amorphization process

It was reported in the literature that amorphization of a crystalline GeTe film can

be achieved with single laser pulse irradiation with a pulse length of 8 ns, a

wavelength of 800 nm and a fluence of 71 mJ/cm² [140]. For the present study, the

GeTe thin films, which were crystallized by ns laser irradiation as described in the

previous sections, were re-amorphized by a single 20 ns laser pulse irradiation at a

wavelength of 248 nm and at various laser fluences up to 195 mJ/cm². In Figure

4.28a, the optical reflectivity curves in the spectral region from 400 up to 700 nm of

single pulse irradiated GeTe films are shown. For comparison, the reflectivity

curves of an as-deposited GeTe film and of a GeTe film after 20 min furnace

annealing at a temperature of 300 °C are also shown. The reflectivity of this latter

crystalline film varies between 90 and 93%. Figure 4.28 also reveals that an obvious

reflectivity change after irradiation by a single laser pulse at a fluence of 98 mJ/cm²

cannot be observed. However, with an increase of the laser fluence up to about

160 mJ/cm², a continuous and gradual decrease in reflectivity is detected. Gawelda


79

et al. have also observed such a decreasing trend of the reflectivity of GeTe films

after 800 nm single pulse irradiation for fluences > 50 mJ/cm², which corresponds

to the formation of amorphous regions but also to increased ablation [140].

With further increase of the fluence from 160 to 182 mJ/cm2, a dramatic

decrease in reflectivity can be detected in Figure 4.28a. Apparently, the atoms in the

GeTe films are increasingly displaced from their lattice sites above a laser fluence

of 98 mJ/cm². The displaced atoms are quenched during the very short laser pulse

interaction with the film, which leads to a partial amorphization of the film. With

increasing laser fluence, the degree of amorphization, i.e. the amorphized volume

fraction, increases, corresponding to the decrease of film reflectivity. However,

Figure 4.28: (a) Optical reflectivity of GeTe films in the wavelength range between 400 and 700 nm

after irradiation with a single 20 ns laser pulse as function of the laser fluence. For comparison, the

optical reflectivity curves of an as-deposited GeTe film and of a crystalline GeTe film after thermal

annealing at 300°C for 20 min are shown. The inset displays an SEM image of the surface after

single laser pulse irradiation with a fluence of 195 mJ/cm2. (b) XRD patterns of ns laser pulse

irradiated GeTe films as a function of the laser pulse fluence.

20 30 40 50

Inte

nsity (

arb

. u

nits)

2(°)

182 mJ/cm2

(220)

(202)

98 mJ/cm2

130 mJ/cm2

160 mJ/cm2

(024)

(003)

(021)

400 500 600 70070

75

80

85

90

Reflectivity (

%)

Wavelength (nm)

98 mJ/cm2

160 mJ/cm2

130 mJ/cm2

182 mJ/cm2

oven annealed

as-dep.

195 mJ/cm2

(a)

(b)


80

when the fluence increases to much higher values (≥195 mJ/cm2), the reflectivity of

the laser-irradiated films shows only a slight further decrease, due to the increased

fraction of amorphization and a significant ablation, as visible in the inserted SEM

images in Figure 4.28a. In order to confirm the structural transition from the

crystalline to the amorphous state, those films irradiated with laser fluences between

98 and 182 mJ/cm2 were studied by XRD, as shown in Figure 4.28b. It can be

observed that, with an increase of the fluence from 98 to 182 mJ/cm2, the intensity

of diffraction peaks decreases gradually. Thus, the amorphous fraction in the

crystalline films increases with increasing laser fluence. These consistent results

demonstrate that amorphization of crystalline GeTe films can also be achieved with

single 20 ns UV laser pulse irradiation at sufficiently high fluences.

Crystallization of Ge2Sb2Te5 films by nano- and femtosecond single 4.2.2

laser pulse irradiation

As known from the discussion in Chapter 2.2, GST possesses a potentially higher

switching speed, compared to GeTe. Moreover, the crystallization-dominated

mechanisms differ between GST (nucleation-dominated) and GeTe (growth-

dominated). In contrast to the previous sub-chapter (4.2.1), in this sub-chapter

investigations of the crystallization characteristics of PLD-deposited GST films

with either ns or fs single pulse laser irradiation at the same wavelength of 248 nm

are compared. The details of the crystallization process are characterized by

investigating the film structure, morphology and optical reflectivity as a function of

the laser fluence [131].


XRD measurements were employed to investigate the influence of the laser fluence

on the crystallinity of as-deposited 90 nm thick GST films upon single pulse ns and

fs laser irradiation. Figure 4.29a presents the evolution of XRD patterns of films

irradiated with increasing laser fluences. The fluence was varied between 0 and

130 mJ/cm2 and between 13 and 19 mJ/cm

2 for pulse durations of 20 ns and 500 fs,

respectively. The diffraction spectrum of the as-deposited film shows that the film is

in the amorphous state. The diffraction peak at ~33° is a contribution from the Si

substrate. For 20 ns laser irradiation with fluences ≥ 36 mJ/cm2, crystalline phase

reflections at 29.6° and 42.7° are visible, which correspond to the (200) and (220)


81

lattice planes of the metastable fcc phase of GST, respectively. For laser fluences

above 98 mJ/cm2, the diffracted intensity is significantly increased and the (111)

reflection also becomes visible. It should be noted that, although partial

crystallization of the films is confirmed from optical reflection measurements (as

described in section 4 of this sub-chapter) for fluences between 15 and 36 mJ/cm2,

the XRD reflections of the crystalline phase cannot be detected, due to the low

volume fraction of the crystalline phase in this particular fluence range. With a

continuous increase of the laser irradiation fluence from 36 to 98 mJ/cm2, the

intensity of the reflections becomes gradually proportional to the volume fraction of

the crystalline phase. For the films irradiated with a single fs pulse, the diffraction

spectra are qualitatively the same with the exception that the first diffraction peaks

are already observed at a fluence of 13 mJ/cm2. A further increase of the laser

Figure 4.29: (a) XRD patterns of an as-deposited GST film and of films after 500 fs and 20 ns laser

single pulse irradiation at different fluences. (b) FWHM of the rocksalt phase GST(200) reflection

and the corresponding calculated average grain sizes as a function of the laser fluence (fs and ns

denoted by the symbols with solid and open interiors, respectively).


(a)

20 30 40 50

fcc(1

11)

fcc(2

00)

fcc(2

20)

Si(2

00)

as-dep.

15 mJ/cm2

36 mJ/cm2

67 mJ/cm2

98 mJ/cm2

13 mJ/cm2

16 mJ/cm2

19 mJ/cm2

ns single pulse

fs single pulse

130 mJ/cm2

Inte

nsity (

arb

. u

nits)

2 (°)

40 60 80 100 120

ns single pulse

30

35

40

45

50

55

60

Gra

in s

ize

(n

m)

12 15 180.24

0.28

0.32

0.36

0.40

fs single pulse

FW

HM

at

(20

0)

(°)

(b)


82

fluence promotes crystallite growth, as evidenced from the increase in diffracted

intensity. The average grain size of GST films after ns and fs irradiation as a

function of laser fluence was estimated from the FWHM of the (200) reflection by

using the Scherrer formula and the results are presented in Figure 4.29b. It follows

that the average grain size decreases from approx. 50 to 35 nm after single pulse ns

laser irradiation in the fluence range between 36 and 130 mJ/cm². In the left part of

Figure 4.29b, this tendency can also be detected for the fs irradiated films. For GST

films after fs pulse laser irradiation, the average grain size gradually increases from

about 50 to 60 nm with an increase of the fluence from 13 to 19 mJ/cm2. These

experiments demonstrate that a crystallization of the GST films is possible in both

cases: after ns single pulse and also after fs single pulse laser irradiation for laser

fluences larger than 36 mJ/cm² and larger than 13 mJ/cm², respectively.

2.) Morphology

To understand the crystallite growth process in the GST thin films after ns and fs

single pulse laser irradiation, the microstructure evolution was examined by cross-

sectional TEM. Figure 4.30 shows a cross-sectional bright-field (BF) TEM image

and a NBD pattern of an as-deposited GST thin film (Figures 4.30a and c), a film

after ns single pulse laser irradiation with fluences of 36 mJ/cm2 (Figures 4.30b, d

and e) and a film with 67 mJ/cm2 (Figure 4.30f) as well as after fs single pulse laser

irradiation with a fluence of 19 mJ/cm2

(Figure 4.30g). In Figure 4.30a, a uniform

and smooth GST film above the SiO2 / Si layer is shown; the corresponding NBD

pattern in Figure 4.30c is typical for an amorphous structure (the top layer of Pt was

only used to stabilize the samples during the TEM preparation). In contrast,

crystalline grains with an oval shape developed within the amorphous GST matrix

after laser irradiation by a 20 ns single laser pulse with a fluence of 36 mJ/cm2

(Figure 4.30b). With a further increase of the laser fluence to 67 mJ/cm2, the

crystalline precipitates in the amorphous regions grow to grains with an oval shape,

covering the film over its entire thickness. Between these grains, small angle grain

boundaries are visible (see Figure 4.30f). Similarly, the GST film occurs to be

completely crystalline and the grain diameters in the film vary between 15 and 60

nm when the laser fluence increases up to 130 mJ/cm2. In contrast, the morphology

after 500 fs single pulse irradiation of amorphous GST thin films differs from the


83

morphology of GST thin films irradiated by a ns single pulse. Figure 4.30g shows a

cross-sectional BF image of the GST thin film after single pulse fs laser irradiation

with a fluence of 19 mJ/cm2. A homogeneous polycrystalline structure with

columnar grains can be observed over the complete film thickness of ~90 nm. The

diameter of these columnar grains lies in the range between 20 and 60 nm. It should

be noted that the near surface region of the films was altered by natural oxidation in

air. However, this oxidation can be avoided by depositing a 5 nm thin capping layer

Figure 4.30: Cross-sectional BF-TEM images and NBD patterns of GST thin films: (a) and (c) as

deposited; (b), (d) and (e) after ns laser pulse irradiation with a fluence of 36 mJ/cm2; (f) after ns

laser pulse irradiation with a fluence of 67 mJ/cm2; the inset shows a magnified region; (g) after fs

laser pulse irradiation with a fluence of 19 mJ/cm2.


84

of LaAlOx onto the top of the GST films. As an example, GST films with and

without a LaAlOx layer are shown in Figure 4.31.

3.) Local atomic arrangement

High-resolution STEM (HRSTEM) investigations allow a direct insight into the

local atomic arrangement. In Figure 4.32a, a middle-angle annular dark-field

(MAADF) HRSTEM image along the [001] viewing direction of an as-deposited

GST film after single fs laser pulse irradiation is depicted. The real-space lattice

fringes in the image indicate that the GST film is characterized by a high crystalline

quality. The crystal plane spacing of the film is approx. 0.302 nm, which

corresponds to the {002} crystal plane spacing (0.302 nm) of the cubic GST

crystalline phase. It is in good agreement with the results from XRD measurements;

the interplanar distance of the {002} crystal plane is calculated to be 0.301 nm,

according to the Bragg equation. It should be noted that no additional sub-lattices

[141] are formed by Ge atoms located at tetrahedral positions or Ge atoms located at

off-octahedral positions [142] as evident from atomic-resolution low-angle annular

dark field (LAADF) HRSTEM studies of fs laser pulse irradiated GST films (see

Figure 4.32b). This is in contrast to the results of Liu et al., which claim a

coexistence of octahedral and tetrahedral sites for Ge in the GST lattice [48]. In

Figure 4.31: Contrast between the layer structures of laser-irradiated LAO/GST/SiOx and GST/SiOx

samples in cross-section TEM images.

Pt

GST ~90 nm

SiOX

GST:O

98 mJ/cm2

ns single pulse

Pt

GST ~90 nm

SiOX

LaAlO3

schematic

90 mJ/cm2


85

addition, detailed analyses of local distortions in the GST lattice showed that a

distorted octahedral atomic arrangement of GeSb is the main structural motif for the

metastable and amorphous GST phases [141]. Thus, due to laser irradiation, these

distortions of the GST lattice can even be enhanced. This will induce a breakage of

long-range order and lead to the transition from the metastable to the stable GST

phase, as was proposed by Kolobov et al. [18]. Another possibility for phase

transition is melting of the GST lattice, due to lattice heating caused by the transfer

of electron energy to the covalent backbone by applying an ultrashort laser pulse

Figure 4.32: Aberration-corrected (a) MAADF- and (b) LAADF-HRSTEM images of a

representative fs single laser pulse irradiated GST film viewed along the [001] zone axis. The insets

show a magnification of the marked regions.


86

[59]. Moreover, a non-thermal amorphization process caused by displacement of Ge

atoms from octahedral to tetrahedral sites, induced by ultrashort laser pulses, was

also recently reported [19].


Figure 4.33 shows the reflectivity of Ge2Sb2Te5 films deposited at room temperature

after 20 ns single pulse laser irradiation with fluences of 15 and 67 mJ/cm2, and

after 500 fs single pulse laser irradiation with a fluence of 15 mJ/cm2. The optical

reflectivity of the as-deposited, amorphous GST film is approximately 43% for

wavelengths between 400 and 700 nm. In contrast, the reflectivity of the film after

ns pulse laser irradiation with a fluence of 15 mJ/cm2 varies between 52 and 58%.

With increasing laser fluence, up to 67 mJ/cm2, an even higher reflectivity in the

range between 65 and 67% is observed. For the film that was fs-irradiated with a

fluence of 15 J/cm², the optical reflectivity is in the range from 64 to 67%, which is

comparable to that of the ns-irradiated film with a fluence of 67 mJ/cm2. These

results demonstrate that a phase transition from the amorphous to the crystalline

state can be achieved by single ns laser pulse irradiation with a fluence ≥ 15

mJ/cm². With an increase of the laser fluence (67 mJ/cm², Figure 4.33) the optical

reflectivity is further increased. This is an indication that the initially small

crystalline volume embedded in the amorphous matrix increased. A comparison of

the reflectivity after ns and fs single laser pulse irradiation in Figure 4.33 leads to

Figure 4.33: Optical reflectivity of GST films after deposition (as-dep.), after 20 ns single laser pulse

irradiation at fluences of 15 and 67 mJ/cm2, and after 500 fs single laser pulse irradiation at a fluence

of 15 mJ/cm2.

400 500 600 700

40

50

60

70 ns-induced, 67 mJ/cm2

ns-induced, 15 mJ/cm2

fs-induced, 15 mJ/cm2

Reflectivity (

%)

Wavelength (nm)

as-dep.


87

the conclusion that obviously single pulse fs laser irradiation is more effective

regarding the crystallization process. Similar results were also found by XRD

characterization (see section 1 of this sub-chapter).

In Figure 4.34, a series of fluence-dependent optical reflectivity contrast values

at a wavelength of 650 nm of GST films is shown. As expected, the reflectivity

contrast of the films increases with increasing laser fluence, where the reflectivity

rise after fs irradiation is very steep, up to 15 mJ/cm², in contrast to the ns laser

irradiation. A simple rescaling procedure of the laser fluence yields the number of

photons per unit area Np

𝑁𝑝 =∅

𝐸𝑝= ∅

𝜆

ℎ∙𝑐 , (4.2)

where Φ is the laser fluence, Ep the energy of the photons, h the Planck constant, c

the speed of light in vacuum and λ the laser light wavelength (here: 248 nm).

Further, the change of the optical reflectivity contrast is proportional to the amount

of the crystalline fraction implied by the laser photons [17, 83]. Consequently, it can

be assumed that the crystallization of the thin GST films is completed when no

further change of the optical contrast can be determined as a function of the laser

fluence. The results of the rescaling procedure are summarized in Figure 4.34.

In the idealized case, it might be assumed that every photon that hits the surface

produces a cylinder-like crystallized region into the surrounding amorphous matrix

with a sufficiently large crystalline area ac at the surface. Then, the formation rate of

Figure 4.34: Optical reflectivity contrast of GST films at a wavelength of 650 nm after ns and fs

laser pulse irradiation with a single pulse as a function of the laser fluence and the number of laser

photons per area.

0 20 40 60 80 100 120

0

5

10

15

20

25

ns single pulse

fs single pulse

Re

fle

ctivity c

on

tra

st

(%)


0 5 10 15

0.00

0.25

0.50

0.75

1.00

Cry

sta

llin

e f

ractio

n

No. of photon/area [

cm-2]


88

the crystalline fraction with the number of photons per area (laser fluence) can be

given by

𝑑𝐴𝑐

𝑑𝑁𝑝= 𝑎𝑐 (1 −

𝐴𝑐

𝐴𝑜) , (4.3)

where Ac is the area covered by crystalline clusters and Ao is the total irradiated

area. The integration of this simple differential equation (4.3) leads to

𝐴𝑐

𝐴𝑜= 1 − 𝑒𝑥𝑝(−𝑎𝑐𝑁𝑝) , (4.4)

where Ac/Ao represents the crystalline fraction on the total area. Equation (4.4) was

used to fit the experimental data in Figure 4.34, where the black (dash) and red

(dash dot) lines indicate the crystalline fraction according to Equation (4.4) after

single ns and fs pulse laser irradiation, respectively. Good fits are obtained, if the

circular areas of ac = 0.0025 nm² and ac = 0.0115 nm² are used for ns and fs laser

irradiation, respectively. These results correspond to the average radii of the

cylinder-like crystalline clusters of rc = 0.02 nm and rc = 0.043 nm, respectively.

The maximum length of the cylinder-like region is given by the total absorption of

the incident laser light (wavelength is 248 nm and photon energy is about 5 eV).

According to the Beer-Lambert absorption law, the absorption length can be

determined as 50 nm, when an absorption coefficient of 8·105 cm

-1 at the given

wavelength is assumed for crystalline and amorphous GST phases [143].

Consequently, the average volume that is transferred by a single laser photon with a

wavelength of 248 nm is 0.125 nm³ and 0.575 nm³ after ns and fs laser irradiation,

respectively. Assuming that the atomic density of the amorphous phase is 3.34×1022

atoms/cm³ [144], about 5 and 20 GST atoms participate in the amorphous-

crystalline phase transition per photon by ns and fs laser irradiation, respectively.

This result illustrates that the single fs pulse laser irradiation appears to be more

effective in inducing crystallization of GST thin films than the single ns pulse laser

irradiation. Explanations could be an extreme suppression of the thermal diffusion

[145, 146], because the shorter laser pulse duration is beneficial for reducing the

dwell time of the laser spot. Another possibility is the effect of multi-photon

absorption occurring during fs laser irradiation [147].


89

Simulation of temperature distribution after pulse laser irradiation 4.2.3

To correlate the thermal analysis and laser irradiation experiments in Chapters 4.2.1

and 4.2.2, the temperature distribution of the laser beam interaction with the films is

simulated and discussed in this sub-chapter. To verify the crystalline growth

mechanism of GeTe and GST films in terms of the heat flow condition after laser

irradiation, the finite element method (program package COMSOL) was used to

estimate the temperature distribution in the films.

1.) Temperature distribution after ns laser irradiation in GeTe film

The phase transitions in GeTe films are strongly determined by the temperature rise

caused by laser irradiation. The spatial distribution of the temperature T was

calculated by a finite element method [148] as function of the pulse laser fluence Φ

by solving the heat conductivity equation [149-151]

GeTeat

Siat

z)α(αΔt

)R()rΦ(t,

)rQ(t,))rT(t,(κ)r(t,Tcρ

p

opts

ssGeTe,Sisp;GeTe,SiGeTe,Si

0

exp1

, (4.5)

where Q is the laser beam induced heat source, rs is the spatial coordinate (rs =

rs(x,y,z), z is the depth, x is the distance from the laser spot center. The assumed

material parameters of GeTe are summarized in Table 4.3.

Figure 4.35 shows the calculated temperature distribution in a 60 nm thick GeTe

film on a Si substrate after 20 ns single pulse laser irradiation with a fluence of

14 mJ/cm², which represents the threshold fluence for crystallization of amorphous

GeTe. It is obvious that the temperature field in the GeTe film is comparably

uniform to a depth of about 60 nm (corresponding to the film thickness) and, in the

lateral dimension, 8 mm in the beam profile at the top heat. Consequently, it is

assumed that also for higher laser fluences, the complete films can be described by

a defined temperature. At these low laser fluences, the temperature of the GeTe film

Table 4.3: Summary of the assumed GeTe material parameters [127, 152].

density ρ 6.18 g·cm-3

heat capacity at constant pressure cp 250 J·kg-1

·K-1

thermal conductivity κ 80 W·m-1

K-1

reflectivity R 74 %

absorption coefficient α 0.053 nm-1


90

increases by only approx. 30 K. Considerably higher temperature rises can be

expected for higher laser fluences.

2.) Temperature distribution after ns/fs laser irradiation in GST film

For GST films irradiated by ns and fs single pulse laser, one- and two-temperature

models [150] are used, respectively. The calculated temperatures are for layered

samples (GST/SiO2/Si) as well as for the results of the time-dependent temperature

distribution in the surface and interface of the films. For ultrashort laser pulses

smaller than 10 ps, a disparity of the electron temperature and the lattice

temperature is supposed, and the laser-solid interaction can be physically described

by a two-temperature model [150].

Two-temperature model:

Electron temperature (Te) in GST:

)()exp()1(),(

),()),((),( ,,;,

ep

p

opts

sseeGSTseeGSTpeGST

TTzt

Rrt

rtQrtTrtTc

(4.6)

Phonon / lattice temperature (Tp) in GST:

)()),((),( ,;, pesppGSTsppppGST TTrtTrtTc (4.7)

At short laser pulses, e.g., ns laser pulses, an equality of the Te and Tp can be

assumed and a one-temperature model allows a good description of the physical

process.

Figure 4.35: Modeled two-dimensional temperature distribution of a 60 nm thick GeTe film on a Si

substrate after single laser pulse irradiation with a fluence of 14 mJ/cm2 for 20 ns. z is the depth, x is

the distance from the laser spot center.

0 1 2 3 4 5 6-200

-150

-100

-50

0

273

Z (

nm

)

X (mm)

T (K)296


91

One-temperature model (Tp = Te=T) in GST:

)exp()1(),(

),()),((),(;

zt

Rrt

rtQrtTrtTc

p

opts

ssGSTsGSTpGST

, (4.8)

where Q is the laser beam induced heat source, rs is the spatial coordinate (rs =

rs(x,z), z is the depth, x is the distance from the laser spot center), t is the time and

Δtp the laser pulse duration.

The assumed material parameters of GST are summarized in Table 4.4. At this

first estimation, the material parameters are assumed to be constant. The material

parameters of SiO2 and Si were taken from the COMSOL material database.

Further, the thermal diffusivity into the Si and SiO2 layer was calculated by a one-

temperature model in both cases. For the laser beam profile, a homogeneous beam

profile was assumed. The aim of the model is an estimation of the maximum

achievable temperature induced by laser radiation. Further energy loss mechanisms,

such as laser-plasma absorption and melting and evaporation enthalpy, were not

considered. The simulated result of ns and fs single pulse laser irradiation is

depicted in Figure 4.36.

Figure 4.36a shows the two-dimensional lattice temperature distribution in a

90 nm GST/500 nm SiO2/Si substrate system 50 ps after the single pulse

femtosecond laser irradiation with a fluence of 19 mJ/cm2, as compared to ns one

(see in Figure 4.36b). The temperature distribution shows a surface-localized heated

region, where there is a relatively rapid temperature rise and a higher temperature in

the top surface of the GST film, with a larger temperature gradient through the layer

thickness, as illustrated Figure 4.36a and in left of Figure 4.36c. Although the lattice

Table 4.4: Summary of the assumed GST material parameters [66, 143, 153, 154].

density ρ 5.5 g·cm-3

heat capacity at constant pressure cp 1.55 J·kg-1

·K-1

heat capacity at constant pressure cp;p/e 1.395/0.155 J·kg-1

·K-1

thermal conductivity κ 0.27 W·m-1

K-1

reflectivity R 65 %

absorption coefficient α 0.08× nm-1

electron-photon coupling strength γ 1017

W/(m3∙K)


92

temperature increases by about 864 K (591°C) in the top surface, which is much

higher than the conventional crystallization temperature (~140°C) for GST films,

the crystallization temperature with increasing heating rate shifts to higher

temperatures (the measured Tc of GST is 356.5° at a heating rate of 4 × 104 K/s)

[25, 55].

Figure 4.36c shows the simulation results of electron and lattice temperature

distribution at the surface of the film as a function of time. The electron temperature

immediately rises up to extremely high values (above ~6500 K) during fs pulse

laser irradiation on the GST film, then it quickly cools down (~860 K) after

40-50 ps, whereas the lattice temperature gradually increases with increasing

irradiation time, but it shows the same temperature range ~860 K at the time scale

of 50 ps. It is noted that the lattice temperature in the interface keeps almost

unchanged (only up to approximately 295 K). The heating and cooling rate of

carrier thermalization is calculated to be approx. 1.3×1016

K/s and 1.7×1013

K/s,

Figure 4.36: Modeled two-dimensional temperature distribution of a 90 nm thick GST film on a

SiO2/Si substrate: (a) after single fs pulse laser irradiation with a fluence of 19 mJ/cm2 for 50 ps, (b)

after single ns pulse laser irradiation with a fluence of 36 mJ/cm2 for 20 ns. z is the depth, x is the

distance from the laser spot center. (c) A comparison of temperature time dependence at the top

surface center of the GST film: carrier and lattice temperature as function of time after one fs single

pulse at a laser fluence of 19 mJ/cm2, (right) after one ns single pulse at a laser fluence of 36 mJ/cm

2.

(d) Corresponding comparison of the cooling rates directly after the laser pulse as well as after a

delay time of 100 ns for ns and fs single laser pulse irradiation.

0 200 400 600 800 1000

-600

-400

-200

0

293

864

T (K)

z (

nm

)

x (nm)

Si

SiO2

Ge2Sb

2Te

5

0 200 400 600 800 1000

-600

-400

-200

0

293

1288

T (K)

z (

nm

)

x (nm)

Si

SiO2

Ge2Sb

2Te

5

(a) (b)

(c) (d)

0 20 40 60 80 100

400

600

800

1000

1200

1400

8 x108 K/s

4 x108 K/s

1.5 x1011

K/s

Te

mp

era

ture

(K

)

Time (ns)

5 x1010

K/s

Cooling rates

ns single pulse

fs single pulse

10-14

10-13

10-12

10-11

10-9

10-8

10-7

1000

2000

3000

4000

5000

6000

7000

Tp

t= 500 fs

t= 20 ns

fs pulse

Tem

pera

ture

(K

)

Time (s)

ns pulseT

e


93

respectively. For comparison, the simulated temperature distribution as a function of

time in a GST film after single pulse ns laser irradiation is shown in the right part of

Figure 4.36c. Moreover, it is found that after single pulse fs laser irradiation within

approximately 100 ns room temperature is reached as shown in Figure 4.36d. In the

case of the fs single pulse irradiation, the cooling rate decreases from approx.

1.5 ×1011

K/s within the initial 50 ps to 8 ×108 K/s after a delay time of 100 ns,

whereas in the case of the ns single pulse irradiation, a decrease of the cooling rate

from initially 5 ×1010

K/s to 4 ×108 K/s after a delay time of 100 ns results (see

Figure 4.36d). The above represented results of the two-temperature model analysis

for fs laser irradiation describe the non-equilibrium process, which is more accurate

than used normal thermal diffusion equations. During laser-material interaction, the

electrons in the materials are first excited by the deposited laser energy, and these

“hot” electrons are then cooled by transferring their energy to the lattice through

electron-phonon coupling [155, 156]. Nevertheless, the temperature in the GST film

upon ns/fs laser irradiation, as calculations show, exceeds the crystallization

temperature of GST. Consequently, a crystallization process in the GST films can be

expected.

Summary of the results on laser-induced phase transitions of GeTe and 4.2.4

GST films

In this chapter, the investigation of the phase transformation of GeTe and GST films

upon pulsed laser irradiation was presented. Additionally, simulated temperature

distributions of laser irradiated samples were used to confirm the physics behind the

phase transformation after laser irradiation.

Nanosecond laser-induced phase transition in GeTe films

Phase transformation processes of PLD-grown GeTe films on silicon substrates

were initiated by irradiation with 20 ns laser pulses at a wavelength of 248 nm. The

investigations of these films revealed that a reversible phase transition was realized

by using pulse numbers ≥ 5 at fluences above the threshold fluence between 11 and

14 mJ/cm2 for the crystallization process and by using single pulses at fluences

between 162 and 182 mJ/cm² for the amorphization process. Moreover, a fluence-

dependent structural evolution from the rhombohedral to cubic phase transition was

observed. The reflectivity contrast between the laser irradiated and as-deposited


94

films was studied; it reached a maximum of 14.7% for laser fluences between 36

and 130 mJ/cm2. It should be noted that the crystallization of GeTe films could be

shown to be independent of the irradiated laser pulse repetition rate between 1 and

400 Hz. The influence of film thickness variation between 20 and 120 nm on

crystallization was investigated. A reduction of approximately 10% in the film

thickness of laser-irradiated crystallized GeTe films, compared to as-deposited

amorphous films, was found. Moreover, with decreasing film thickness, smaller

crystallized volumes were detected, because the probability of an effect of the

interface increased with smaller film thickness, it envisions an enhancement of their

crystallization temperature.

Crystallization of GST films by nano- and femtosecond single laser pulse

irradiation

Crystallization of GST thin films deposited by PLD was triggered by ns and fs

single UV laser pulses. The phase transition from the amorphous to the crystalline

state was studied as a function of the laser fluence. The appearance of a crystalline

state was found after ns-laser and fs-laser irradiation with fluences ≥ 36 mJ/cm² and

≥ 13 mJ/cm², respectively. Various crystallite shapes and growth mechanisms

following irradiation with fs and ns pulse durations were observed. GST films

irradiated with fs-laser pulses show a columnar growth, whereas in ns-laser

irradiated films, oval shaped crystallites were produced. In addition, neither

tetrahedral nor off-octahedral positions of Ge-atoms were detected in the defect-free

crystallites after fs-laser irradiation by HRSTEM examination. The optical

reflectivity contrast increased with the laser fluence for both ns and fs pulse laser

irradiation. A simple model of the irradiation-induced crystalline phase formation

was used to assess the average contribution of ns and fs laser photons to the phase

transition process. A change in the optical reflectivity contrast of 25% could be

obtained by single pulse fs-laser irradiation with a fluence ≥ 15 mJ/cm² or ns-laser

irradiation with a fluence ≥ 67 mJ/cm². This reflectivity contrast change of 25%

corresponds to a complete crystallization of the GST film.

4.3 Dynamic optical switching of non-volatile multi-level memory

95

Dynamic optical switching of non-volatile multi-level 4.3

memory

Information storage relying on chalcogenide-based alloys and performed by

switching between fully amorphous and crystalline phases was originally reported

by S. R. Ovshinsky [4]. Increasing the memory storage density achieved by

accumulated switching of phase-change materials, so called multi-level storage, has

recently been attracting extensive attention for applications in non-volatile memory

devices [67, 81, 83, 157-162]. In multi-level storage, the crystallization rate, driven

by appropriate amplitude and pulse duration of laser or electric current, reflects

distinct multi-levels of either electrical resistivity or optical reflectivity, which are

utilized for the identification of storage levels. The possibility to identify different

storage levels is determined by the maximum number of levels in a memory cell.

Therefore, an accurate control of the crystallization process and a high degree of

reproducibility are crucial for the obtained distinct intervals of intermediate levels

upon rapid phase transformation of phase-change materials. In this chapter, it is

mainly demonstrated that the optical reflectivity of GeTe films changes

accumulatively with multi ns laser pulse-induced crystallization between logic state

‘1’ (fully amorphous) and logic state ‘0’ (fully crystalline). An analogous behavior

has been observed for electrical multi-level storage [77, 81, 161].

Besides the benefit of increasing storage density, multi-level storage within a

single layer phase-change film also brings the inherent advantage of low power

consumption. Routinely, the recording (amorphization) process is obtained by

applying an intense and short pulse, either a laser or a current pulse, which induces

melting and fast quenching, whilst a moderately intense and long laser/current pulse

is used to thermally anneal the material above the crystallization temperature within

a sufficient amount of time in the erasing (crystallization) process. With the

application of multi-level storage, each appropriate excitation with a pulse

contributes to achieving a well-controlled partial crystallization/amorphization and

this allows multiple bits for recording/erasing of information in each memory cell,

based on the respective mode of energy accumulation. Consequently, the energy

consumption per bit can be reduced.


96

The work in this chapter is based on results obtained with the previously

described UV ns laser pump-probe setup system presented in Chapter 3.3.2. The

aim is to describe the key aspects of the proposed simple memory cell: 1) Non-

volatile memory. 2) Feasible and reliable control of the crystallization of GeTe thin

films for multi-level storage in a single memory cell. 3) Dynamic switching within a

few nanoseconds. 4) Detailed insights into the switching mechanism of structural

transformation and morphology.

Optical switching 4.3.1

Figure 4.37 displays the On and Off switching states operation of the optical

reflectivity as a function of laser pulse number and fluence, respectively. The PLD-

prepared GeTe films on Si substrates with a LaAlOx capping layer for this type of

Figure 4.37: Optical switching (On and Off states) operation of GeTe films as a function of the laser

pulse number and the fluence: (a) Switching On process: low to high reflectivity change of the film

as a function of the pulse number at a constant fluence of 26 mJ/cm2. (b) Switching Off process: high

to low reflectivity change of the films as a function of the laser fluence with only one single pulse,

(c) and (d) the reflected pump laser intensity with a constant fluence of 26 mJ/cm2 and various

fluences of 0-135 mJ/cm2 as a function of time, corresponding to (a) and (b), respectively.

0 2 4 6 8 10 12 80100

0

25

50

75

100

Rela

tive r

eflectivity (

%)

Number of pulses

0 1 2 3

UV

la

se

r in

tensity

Relative time (s)

26 mJ/cm2

0 20 40 60 80 100120140

Fluence (mJ/cm2)

0.0 0.2 0.4 0.6 0.8 1.0 1.2Relative time (s)

135 mJ/cm2

112 mJ/cm2

90 mJ/cm2

26 mJ/cm2

67 mJ/cm2

Off

On

(a) (b)

(c) (d)


97

experiments. The composition close to that of Ge50Te50 films was confirmed. The

samples were loaded into the custom-made ns laser pump-probe setup, as a detail

described in Chapters 3.3.2. It turns out that, as shown in Figure 4.37a, a steep

increase of the GeTe film reflectivity as a function of the number of pulses occurs,

when ns laser irradiation at a fluence of 26 mJ/cm2 is used. To evaluate the

reflectivity changes more directly, the relative reflectivity of the films (Rf) is

defined by the following equation:

𝑅𝑓 =𝑅𝑥−𝑅𝑎

𝑅𝑐× 100% , (4.9)

where Rx, Ra and Rc are the probed reflectivity of laser irradiated, amorphous and

completely crystallized GeTe films, respectively. The structurally amorphous and

completely crystallized GeTe films are referred to the as-deposited film and the

films after 20 laser pulses, confirmed by analysis with XRD and TEM (for details

see Chapters 4.3.2 and 4.3.3). During the initial four pulses, the relative reflectivity

increases gradually from 0% to almost 80% and remains constant at this high

reflectivity level for higher pulse numbers, up to 100 pulses. In the corresponding

optical microscope images of the films as show in Figure 4.38, optical reflectivity in

the large area is obviously different after irradiation with different pulse numbers.

The image of the as-deposited film shows a relatively dark region, whereas in the

images of films after laser irradiation with an increasing number of pulses, the

regions appear increasingly bright until after four pulses. The image brightness does

not change any more with larger laser pulse numbers. This result is in good

Figure 4.38: Optical microscope images of GeTe films upon ns laser pulse irradiation illustrate the

corresponding regional changes of film reflectivity upon irradiation with different pulse numbers.

as-dep. 2nd pulse

3rd pulse 20th pulse

200 μm


98

agreement with the probing reflectivity change. The evolution of reflectivity

increase is further demonstrated for phase transformation from the amorphous to the

crystalline state, as confirmed by results of both XRD and TEM (details will be

given in Chapters 4.3.2 and 4.3.3, respectively). In other words, such a change in

reflectivity, from low to high, is considered to be the logical switching On process

(crystallization). The reverse trend of optical reflectivity is the switching Off

process (amorphization), as described in detail in the following.

The amorphization of GeTe films was achieved by irradiation with a sufficiently

high laser fluence and with only one single pulse, as shown in Figure 4.37b. As an

initial stage, a high reflectivity of GeTe films by full crystallization was provided,

as in the On process described above, which means as-deposited GeTe films were

laser irradiated with 20 pulses at a constant fluence of 26 mJ/cm2. Subsequently, the

reflectivity change was investigated by applying single laser pulses at different

fluences. It seems that the reflectivity remains at a high level, despite a small

deviation after single pulse irradiation at a fluence of 26 mJ/cm2. With an increase

of the laser fluence up to 67 mJ/cm2, a decrease of reflectivity by about 50% is

detected, as shown in Figure 4.37b. For a further increase of the fluence, a gradual

decrease of the reflectivity is observed. In particular, at a fluence of 112 mJ/cm2, the

relative reflectivity of the films decreased to almost zero, meaning that the

reflectivity of the amorphized film is now equal to that of an as-deposited film. The

decrease in reflectivity upon single pulse irradiation at increasing fluences is

correlated with a reduction of the degree of crystallinity, i.e., the fraction of the

amorphized material, in GeTe films. A detailed discussion of this dynamic

crystalline-to-amorphous transition can be found in Chapter 4.3.4. With a further

increase of the fluence to 135 mJ/cm2, the reflectivity exhibits a slight increase

caused by the appearance of melting and ablation of the film (not shown here).

The reflected laser intensity of the pump laser source, measured as a function of

time, is given in Figures 4.37c and d. The pumping source consisted of an UV ns

pulse laser with a pulse duration of 20 ns and a wavelength of 248 nm. A constant

intensity of delivered laser pulses for an exemplary identical fluence of 26 mJ/cm2

is demonstrated in Figure 4.37c; this is correlated to the use of the laser as a

pumping source for the irradiation of the film in Figure 4.37a (switching On


99

process). In contrast, a gradual increase of the measured intensity with an increase

of the laser fluence from 0 to 135 mJ/cm2 was registered, as depicted in Figure

4.37d, which corresponds to the pumping source for irradiating the films in Figure

4.37b (switching Off process).

In order to clarify the accuracy of achieving distinctive levels of reflectivity

change upon the crystallization process as a function of the laser fluence and the

number of pulses, a series of pulse number-fluence-dependent optical reflectivity

measurements of GeTe films was investigated. Figure 4.39a shows the relative

reflectivity of GeTe films after laser irradiation with fluences of 0-36 mJ/cm2 as a

function of the laser pulse number. No apparent difference in the relative optical

Figure 4.39: (a) Optical reflectivity and crystalline fraction of GeTe films after ns laser pulse

irradiation at various fluences of 0-36 mJ/cm2 as a function of the pulse number, (b) Relation

between laser pulse fluence, number of pulses and the resulting relative reflectivity change for the

initial five pulses in (a).

0 4 8 12 16 20

0

20

40

60

80

100

Re

lative

re

fle

ctivity (

%)

Number of pulses

6 mJ/cm2 11 mJ/cm

2

14 mJ/cm2

17 mJ/cm2

26 mJ/cm2

36 mJ/cm2

0.0

0.2

0.4

0.6

0.8

1.0

Cry

sta

llin

e f

ractio

n

(a)

(b)


100

reflectivity after laser irradiation with 0 to 20 pulses at a fluence of ≤ 6 mJ/cm2 can

be recognized. With a slightly increased fluence of 11 mJ/cm2 and with an increase

of the number of pulses, the reflectivity gradually increases to a saturation value of

about 40%, reached after eight applied laser pulses. In contrast, the reflectivity

curve exhibits a significant rise to a level close to 85% after only four pulses, when

the laser fluence was 14 mJ/cm2. In a fluence window between 14 and 36 mJ/cm

2,

the observed trend of reflectivity change did not display very marked differences,

with the saturation level varying only by around 10%. Thus, a relatively high

reflectivity can be realized by using a higher fluence for the first 4 pulses; for higher

pulse numbers, the achieved reflectivity level remains constant. A more detailed

evolution of fluence-pulse number-dependent reflectivity of GeTe films after ns

laser irradiation within the first five pulses is shown in Figure 4.39b. It can be

observed that a continuous reflectivity increase of GeTe films can be attributed to

either the increase of the pulse number or to the increase of the laser fluence,

excluding the fluences smaller than 14 mJ/cm2, which is the threshold fluence for

crystallization (see Chapter 4.3.2). For fluences above this threshold, five distinct

levels (including the as-deposited amorphous state) of relative reflectivity of the

film can be achieved by varying the pulse number between zero and four pulses. In

addition, the possible number of different levels of reflectivity is not severely

dependent on the laser fluence in the region between 14 and 36 mJ/cm2. This

illustrates that a multilevel operation with 5 distinct levels is highly stable and also

reproducible (see Chapter 4.3.5). Consequently, these investigations demonstrate

that not only binary states of either high reflectivity or low reflectivity of GeTe

films are obtainable by using an irradiation fluence at the threshold (14 mJ/cm2), but

also non-binary states (multi-level states) that are based on accumulative

crystallization processes controlled by the laser pulse number.

Structural transformation 4.3.2

After having gained insights into the optical switching of GeTe films, the structural

transformation, as a function of pulse number and fluence, was investigated by

XRD measurements. The obtained results are given in Figure 4.40. Figure 4.40a

presents the evolution of XRD patterns of films irradiated with an increasing laser

fluence between 0 and 54 mJ/cm2 after irradiation with 20 pulses (pulse duration of


101

20 ns and wavelength of 248 nm). The chosen fluence values equal those in the

previous chapter (4.3.1). The diffraction pattern of the as-deposited film shows the

usual features that are characteristic for the amorphous state. No distinct differences

in the diffraction spectra appear for fluences up to 11 mJ/cm2. For GeTe films

irradiated with a laser fluence larger than 14 mJ/cm², crystalline phase reflections

emerge at the angular positions 26.3°, 30.2° and 43.8°, which correspond to the

(111), (200) and (220) lattice planes of the cubic GeTe phase, respectively. This is

in good agreement with the results of the optical reflectivity measurements related

to the threshold of the significant optical contrast change. It should be noted that

although the increased optical reflectivity of films irradiated with a laser fluence of

11 mJ/cm² is presumably due to the partial crystallization of the films, XRD

Figure 4.40: XRD patterns of ns laser pulse irradiated GeTe films: (a) with 20 pulses as a function of

the laser pulse fluence between 0 and 54 mJ/cm2, (b) with a fluence of 26 mJ/cm

2 as a function of the

pulse number.

20 30 40 50

20th pulse

4th pulse

3th pulse

Inte

nsity (

arb

. units)

2 (deg)

as-dep.

2nd pulse

fcc (

220

)

Si (2

00

)

fcc (

200)

fcc (

111

)

20 30 40 50

54 mJ/cm2

36 mJ/cm2

26 mJ/cm2

17 mJ/cm2

14 mJ/cm2

11 mJ/cm2

Inte

nsity (

arb

. u

nits)

2 (deg)

as-dep.

6 mJ/cm2

fcc (

22

0)

Si (2

00

)

fcc (

20

0)

fcc (

11

1)

(a)

(b)


102

reflections of the crystalline phase cannot be detected yet, due to the low fraction of

the crystallized volume induced within this fluence range. Figure 4.40b presents the

evolution of XRD patterns of films irradiated with an increasing number of pulses

between 0 and 20 at a constant fluence of 26 mJ/cm2. In contrast to the pattern of

the as-deposited amorphous film, when the film was laser-irradiated with two

pulses, weak Bragg peaks located at 26.4°, 30.3° and 43.7° appeared, which

correspond to the (111), (200) and (220) lattice planes of cubic phase GeTe,

respectively. With a further increase of the number of pulses to four and higher (up

to 20), the corresponding Bragg peaks became clearly visible and reached an

intensity saturation level, depicted in Figure 4.40b. This indicates that a continuous

increase of the crystalline fraction in the films results from the increase of the

number of laser pulses, due to a longer duration of energy absorption of the film.

The XRD results, evidencing a phase transformation in the laser-irradiated GeTe

films, are consistent with the optical reflectivity measurements presented in the

previous Chapter (4.3.2). Unequivocally, the observed higher optical reflectivity of

the films originates from a higher degree of crystallization.

Morphology 4.3.3

In order to reveal the On and Off processes of optical switching related to

crystalline phase growth and melt-quench to the amorphous state of GeTe phase-

change films on the microscopic scale, the morphology of the previously described

simple memory element was analyzed by cross-sectional TEM. Figure 4.41

summarizes cross-sectional bright-field (BF) TEM images and the corresponding

calculated 2D fast Fourier transformation (FFT) images of such GeTe thin films. A

uniform and smooth as-deposited GeTe film lying between the Si substrate and the

LaAlOx capping layer is shown in Figure 4.41a (the Pt layer is only used for TEM

specimen preparation). The entire accessible as-deposited film was confirmed to be

in the amorphous state by the FFT diagram (Figure 4.41g); this is characterized by

diffuse circular intensity distribution rings. However, it is obvious that crystalline

grains formed in the amorphous GeTe matrix after ns laser irradiation with 2 pulses

at a constant fluence of 26 mJ/cm2

(Figure 4.41b). The average grain dimensions of

the oval-shaped grains are approx. 62 nm in height and 90 nm in width, based on

the TEM results. This is because, when the crystalline grain size within the GeTe


103

film reaches the film thickness, grain growth along the direction perpendicular to

the substrate surface will be prevented by the top and bottom interfaces of the GeTe

film. Consequently, due to these geometric restrictions, any further grain growth

can only occur in the lateral direction. Furthermore, since such restricted growth of

the crystalline regions might induce mechanical stress and involve shrinkage in film

thickness upon amorphous-crystalline transition, both forces, as a result of the

altered elastic energy, possibly lead to the inhomogeneous complex height

distribution on the film’s surface. With an increase of the pulse number from two to

Figure 4.41: Cross-sectional BF-TEM images and FFT patterns of GeTe thin films on Si: (a) as-

deposited; (b), (c), (d), (e) after ns laser pulse irradiation at an identical fluence of 26 mJ/cm2

with 2,

3, 4, 20 pulses, respectively; (f) re-amorphization of 20 pulse-crystallized film by ns laser pulse

irradiation with 112 mJ/cm2

and a single pulse; (g), (h), (i) and (j) are corresponding FFT patterns to

(a), the marked regions P1 and P2 of (b) after magnification, and to (f), respectively. All GeTe films

were coated with a thin protective LaAlOx layer.


104

four pulses, the crystalline precipitates in the amorphous regions gradually grow

further in the lateral direction and the width of grains extends from 90 to 200 nm,

while, consequently, the amorphous regions between the crystalline grains are

reduced in volume, as observed in Figures 4.41b, c and d. The corresponding

structure of two selected representative regions, P1 and P2, in Figure 4.41b was

confirmed by FFT patterns to be amorphous and polycrystalline, as shown in

Figures 4.41h and i, respectively. This can explain the contribution to the

multi-level reflectivity of the film that is dependent on the pulse number chosen for

the irradiation. Since a higher volume of crystallized material is obtained upon

applying higher pulse numbers, the prolonged accumulated photon absorption time

allows the crystalline precipitates to grow and merge until a continuously

polycrystalline film is formed. Almost complete crystallization of the GeTe film can

exemplarily be observed after 20 pulses of irradiation, as shown in Figure 4.41e.

Moreover, complete amorphization of such a fully crystallized film is

achievable by ns laser irradiation with one single pulse at a fluence of 112 mJ/cm2.

This is confirmed by the corresponding results of cross-sectional TEM in Figures

4.41f and j. A homogeneous and smooth GeTe film without crystalline regions

within the film is visible that closely resembles the original as-deposited film. The

corresponding FFT pattern (Figure 4.41j) is typical for amorphous phase material,

indicating that the crystallized GeTe film was melt-quenched to a disordered

amorphous state upon the laser irradiation. These findings agree well with the

results from the optical reflectivity measurements (presented in Chapter 4.3.5).

However, it should be noted that there are differences between the as-deposited

amorphous and the melt-quenched amorphous films regarding the optical

reflectivity change upon repetitive crystallization and amorphization (see Figure

4.45). The resulting FFT images in Figures 4.41g and j are similar in that they

illustrate the amorphous nature of the thin films. It is, however, to be expected that

as-deposited and melt-quenched amorphous structures possess slightly different

short-range order that cannot be detected by FFT due to differences in TEM

specimen thickness.

To gain more information about the local structure of as-deposited, re-

amorphized and crystallized GeTe thin films, atomic resolution cross-sectional


105

TEM images of the GeTe films were recorded. The representative crystallized GeTe

film is the as-deposited film after ns laser irradiation with 20 pulses at 26 mJ/cm2

and the re-amorphized film is the crystallized film after subsequent ns laser

irradiation with single pulses at 112 mJ/cm2. The result is shown in Figure 4.42.

Compared to the as-deposited GeTe film exhibiting the presence of an

amorphous structure with a lack of long-range order, as shown in Figure 4.42a, the

HRTEM image in Figure 4.42b of the disordered structure reveals no

nanocrystalline components after melt-quenching. Note that the structural features

of melt-quenched and as-deposited amorphous are very similar and distinguishing

between them is difficult [163, 164]. Some reports [134, 165-168] have claimed the

existence of a nanocrystalline phase, and demonstrated a higher medium-range

order in the melt-quenched amorphous state, compared to the as-deposited

amorphous state in other chalcogenide materials, such as AgInSbTe or GST. Other

authors also hypothesized [169] that the ordering of Ge and Te in melt-quenched

GeTe is much higher than in as-deposited. Nevertheless, compared to a

representative crystallized GeTe film, a long-range order with a lattice fringe

Figure 4.42: Cross-sectional HRTEM images of: (a) as-deposited GeTe film, (b) re-amorphization of

the crystallized GeTe film after ns laser irradiation with one pulse at a fluence of 112 mJ/cm2, (c) a

representative GeTe film after ns laser irradiation with 20 pulses at a fluence of 26 mJ/cm2. The

corresponding FFT pattern is shown in the inset, (d) The electron diffraction patterns of GeTe films

after ns laser irradiation with 20 pulses at a fluence of 26 mJ/cm2.


106

spacing of ~0.34 nm along the [110] viewing direction can be observed in Figure

4.42c. This indicates that the crystalline phase of the GeTe film induced by ns pulse

laser irradiation possesses a high crystalline quality. A corresponding Fourier

transformed pattern obtained from this HRTEM intensity region of the GeTe thin

film is shown in the inset of Figure 4.42c and a measured electron diffraction

pattern is additionally shown in Figure 4.42d, which confirms this finding. In the

electron diffraction pattern, a local lattice spacing is identified, which is

characteristic for a typical crystalline phase of GeTe.

Dynamic phase transformation 4.3.4

In order to determine exactly how fast the phase transformations of the PLD-

deposited GeTe films take place, the dynamics of crystallization and amorphization

of as-deposited GeTe thin films driven by ns laser pulses was investigated. Figure

4.43 shows the reflectivity evaluation of a GeTe film upon laser irradiation with

each ns pulse (first to fifth pulses are shown) being at a constant fluence of 26

mJ/cm2 on a fixed irradiation region of the size of about 24 × 6 mm

2. Without the

pumping laser irradiation (zero pulses), probing the reflectivity change from an

as-deposited, amorphous GeTe film is referred to as Baseline in Figure 4.43a and it

appears as a relatively smooth line within the observed timescale range

between -200 and 200 ns. In contrast, a distinguishable level of reflectivity increase

from pre-existing to post-existing states upon each successive single pulse

irradiation for five pulses can be measured. A final total relative reflectivity of

~85% can be achieved after four pulses and an unstable increase of the relative

reflectivity with a variation of 5% after the 5th pulse to a level of more than 90% is

observed. This indicates that the crystallization process was almost completed

within the first four (or five) pulses. With a further increase of the number of pulses

to more than five, the relative reflectivity change becomes negligible, due to the

crystallization being accomplished throughout the whole film (see Figure 4.39a).

For all distinct pulse-dependent reflectivity changes in Figure 4.43, the transition

time from the lower to the higher reflectivity level upon laser irradiation lies within

an interval of ~60 ns, which is around three times longer than the pulse duration

(~20 ns) of the applied laser pulse. Possibly, this transition time is the time required

for atoms to reorganize into an adequate ordered arrangement, assuming that the


107

increased reflectivity is representative of the crystallization process in the form of a

crystalline fraction increase of the films. The time for complete crystallization of a

GeTe film after ns laser irradiation with five pulses is ~300 ns after the complete

transition process. Moreover, it is interesting to note that the gradients of the

transitions from lower to higher reflectivity levels decreased with an increase in the

number of pulses (first to fifth pulse), which implies that the velocity of

crystallization decreased with each increase of the crystalline fraction. Assuming

that the crystalline fraction is proportional to the measured reflectivity of the film, a

simple representation of the correlation between growth velocity and the number of

irradiation pulses can be extracted from the results of Figure 4.43a, as shown in

Figure 4.44. These results might be helpful for the further understanding of the fast

Figure 4.43: (a) Optical reflectivity of GeTe films after laser pulse irradiation with different pulse

numbers from 0 to 5 pulses and at a constant fluence of 26 mJ/cm2 as a function of time, (b) Optical

reflectivity of GeTe films after single laser pulse irradiation with a fluence of 112 mJ/cm2

as a

function of time. The inset displays a magnification of the curve in the time-scale range from -7 to

12 ns.

-200 -100 0 100 200

0

20

40

60

80

100 Amorphization

Re

lative

re

fle

ctivity (

%)

Time (ns)

-5 0 5 10

Time (ns)

3 ns

(a)

(b)-200 -100 0 100 200

0

20

40

60

80

100

Re

lative

re

fle

ctivity (

%)

Time (ns)

5th pulse

4th pulse

3d pulse

2nd pulse

1st pulse Baseline

Crystallization


108

phase transition mechanism and for the design of multi-level phase-change memory

elements.

The reverse process, i.e., the amorphization process (switching Off state) of a

GeTe film, characterized by a change of the optical reflectivity from the highest to

the lowest level, can be achieved by irradiation with only one single laser pulse.

Figure 4.43b shows the reflectivity change of a GeTe film with such a single laser

pulse irradiation at a fluence of 112 mJ/cm2

as a function of the time. The pre-

existing high reflectivity level of ~100% was provided by the crystallization of an

as-deposited GeTe film using ns laser irradiation with 20 pulses at a fluence of

26 mJ/cm2 (switching On process described above). When the film is irradiated

with a single laser pulse at a sufficient fluence, the film reflectivity drops sharply to

a much lower level of ~40% after about 3 ns (as seen in the inset with magnified

timescale) and then gradually decreases further to ~20% within a delay time of

~200 ns. It should be noted that the experimental setup of the pump-probe system

(as described in Chapter 3.3.2) makes it possible to provide a time resolution as fast

as about 2 ns. This would mean that the processes are maybe even faster than 3 ns

but cannot be resolved. The result of a dramatically decreased reflectivity is

assumed to substantially reduce atomic mobility caused by melt-quenching of the

crystallized GeTe film back to the stable amorphous phase, after application of an

intensive single laser pulse. Furthermore, a structural transformation was

demonstrated by the analysis of cross-sectional TEM images of such amorphized

GeTe films and this is also confirmed by the results related to the retention of

Figure 4.44: Crystal growth velocity of GeTe film after ns pulse laser irradiation with a fluence of

26 mJ/cm2 as a function of the pulse number.

0 1 2 3 4 50

50

100

150

2000 60 120 180 240 300

Crystallization time (ns)

Number of pulses

Gro

wth

velo

city (

m/s

)

0.0

0.2

0.4

0.6

0.8

1.0

Rela

tive r

eflectivity (

%)


109

reversible switching (Chapters 4.3.3 and 4.3.5, respectively).

Retention 4.3.5

The retention of the actual present state upon reversible phase transformation

between amorphous and crystalline states in phase-change films is an eminently

important parameter for data storage applications based on phase-change materials.

Repetitive optical On and Off switching processes of GeTe films with a various

number of cycles driven by ns laser pulses was investigated and is shown in Figure

4.45a. During the first cycling, crystallization (On state, high reflectivity) of an as-

deposited amorphous GeTe film was obtained at a moderate fluence of 26 mJ/cm2

with multi-pulses, i.e., a short series of five successive ns pulses; subsequently,

amorphization (Off state, low reflectivity) of the film was triggered by a single

pulse at a higher fluence (112 mJ/cm2), as discussed above. It is interesting that the

Figure 4.45: (a) Optical reflectivity of a GeTe film as a function of the number of laser irradiation

cycles (forward: with a certain number of pulses at a constant laser fluence of 26 mJ/cm2 and

reverse: a single laser pulse at a laser fluence of 112 mJ/cm2) as function of the pulse number. (b)

Representation of the optical reflectivity switching On and Off processes with various laser pulse

numbers as a function of the number of cycles.

0 5 10 15 20

0

20

40

60

80

100

Off

1,2,3

,4,5

Re

lative

re

fle

ctivity (

%)

Number of pulses

1st cycle

2nd cycle

3d cycle

4th cycle

5th cycle

On

1

On 2

,3,4

,5

0.0

0.2

0.4

0.6

0.8

1.0

Cry

sta

llin

e f

ractio

n

1 2 3 4 5

0

20

40

60

80

100

3d pulse

2nd pulse

1st pulse

0th pulse

Rela

tive r

eflectivity (

%)

Number of cycles

No. of pulses

20th pulse

5th pulse

4th pulse

0.0

0.2

0.4

0.6

0.8

1.0

Cry

sta

lline fra

ction

(a)

(b)


110

recrystallization (switching On) of a melt-quenched amorphous GeTe film requires

a smaller number of pulses (three pulses), compared to as-deposited ones (five

pulses) at the same fluence of 26 mJ/cm2. This indicates that the total amount of

energy (or crystallization time) required for the re-crystallization of a melt-

quenched film is smaller than that required to crystallize an as-deposited,

amorphous film for the first time. Thus, a difference between both amorphous

structures and the formation mechanism of crystalline nuclei can be expected.

According to previous reports [60, 170, 171], nanocrystalline precipitates were

detected in melt-quenched amorphous Ge-Sb-Te films, which should accelerate the

crystallization process. For instance, the crystallization from a melt-quenched

amorphous film was found to be an order of magnitude faster than that from an

as-deposited amorphous film [60]. Similarly, it was found that the crystallization

times for Ge-Te melt-quenched amorphous material is shorter than of as-deposited

amorphous material [28, 140]. The dependence of the crystallization kinetics of

amorphous GeTe films on their thermal history was also confirmed [169]. It has

been hypothesized that the existence of crystalline seeds is responsible for the

increase in crystallization time. An extended discussion on the local structure of as-

deposited amorphous and melt-quenched amorphous is provided in Chapter 4.3.3.

The switching behavior of On and Off processes was found to be reproducible

and stable during the investigated cycles. The intermediate levels of reflectivity

were repeatedly achieved by laser irradiation with a varied number of pulses, as

shown in Figure 4.45b. The 1st pulse of the 1st cycle did not provoke an obvious

reflectivity change, after the 2nd pulse an almost middle level of reflectivity was

reached, an then, after the 3rd trigger pulse, the reflectivity of the film increased

close to the highest (saturated) level. Compared to the 1st cycle, with the 2nd to 5th

cycle the respective 1st pulse already altered the reflectivity of the melt-quenched

amorphous GeTe film to the medium level and the following 2nd pulses were

sufficient to achieve the final high reflectivity level. Apparently, repeatedly

specified intermediate reflectivity levels are achievable in phase-change films

driven by the numbers of laser pulses, and the identification of these levels is

meaningful for multilevel storage applications.


111

Optical switching model 4.3.6

The principle of a non-volatile memory cell with multilevel optical switching of

GeTe films using ns laser pulse irradiation is illustrated in Figure 4.46. The unique

property of reversible switching between the amorphous and crystalline states upon

specific heating/cooling by applying appropriate laser pulses is the basis of the

functionality of such a memory cell. Precisely, crystallization can be considered as

accumulative switching (On states). Each laser pulse excitation can induce a partial

crystallization, and complete crystallization can be achieved by a succession of such

Figure 4.46: (a) A schematic of the multi-level optical switching of a GeTe film driven by ns laser

pulses for non-volatile memory application. (b) Corresponding multi-states evaluation of reversible

amorphous-to-crystalline phase transitions of a phase-change film using ns laser pulses. The a-GeTe

and c-GeTe represents amorphous and crystalline GeTe, respectively.

Phase-change cell

Input pulses

Cell switch On

States

Cell switch Off

Tm

Tc

t

Tm

Tc

t

I

II

III

IV

t

States

t

Input pulse

V

V

I

(a)

a-GeTeLaAlOx c-GeTe Si

States I II III IV V

Top-view

Cross-section

PCM

(b)


112

excitations. As depicted in Figure 4.46a, a moderately intense laser beam (fluence

of 26 mJ/cm2) and distinct sequences of pulse numbers (pulse duration of 20 ns for

each pulse) is used for crystallization of the phase-change memory cell. The

supplied laser pulse energy heats the material above the crystallization temperature

(but below the melting temperature) for a sufficiently long time, resulting in

transformation from the amorphous to the crystalline state. The resulting

crystallization rate of the growth-dominated GeTe film is dependent on the

irradiation time, which corresponds to the number of applied laser pulses. Thus, a

higher degree of crystallization, i.e., a higher volume fraction of crystallized

material, can be achieved by increasing the pulse number at a constant fluence,

providing an important prerequisite for realizing multi-level phase-change memory

cells.

The discrimination of these levels can be realized by reading out the optical

reflectivity of the memory cell, as shown in the example with five different

levels/states. The corresponding schematic diagram for the crystallization process of

such a phase-change GeTe film is shown in Figure 4.46b. Initially, stage I shows a

completely amorphous GeTe film. With each succeeding laser pulse excitation,

small nuclei are formed and then the crystalline regions gradually grow until the

film is completely polycrystalline (from stages II to V), induced by using ns laser

pulse irradiation with an increase of the number of pulses. The high optical

reflectivity contrast between amorphous and crystalline states is used for tuning the

distinctly intermediate levels in reflectivity as a result of the changed ratio of the

crystalline volume fraction to the amorphous fraction of the film. It should be

pointed out that the time response (~60 ns) of the phase-transition in GeTe upon

laser pulse excitation was found to be longer than the duration of the applied pulse

(~20 ns) (see Chapter 4.3.4). Further, the optical switching can be regarded as

wavelength independent from the near ultraviolet to the near infrared (400 and

700 nm) range, as confirmed in the present study (see Chapter 4.3.8).

In the reverse process, i.e., the reverse transformation from the crystalline to the

amorphous state, an intense (fluence of 112 mJ/cm2) single laser pulse is used for

amorphization of the phase-change cell, as shown in Figure 4.46a. The pulse energy

induces melting of the material, followed by rapid quenching to freeze the atoms in


113

the disordered structure. The significantly different reflectivity between the

crystalline state (high reflectivity) and amorphous state (low reflectivity) can be

easily detected. Such a sharp reflectivity drop from high to low upon amorphization

of GeTe film was determined to take 3 ns (see Chapter 4.3.4). The corresponding

evaluation of schematic state changes in the case of amorphization is also shown in

Figure 4.46b. The reversal of the state V (crystallization) back to state I (amorphous

state) is depicted. In summary, the possibility to reversibly switch a GeTe film for

practical non-volatile memory application with multi-level functionality is

provided.

Film thickness dependence of numbers of multi-levels 4.3.7

The aforementioned five distinct optical reflectivity levels of the memory cell were

obtained with a GeTe film thickness of 68 nm. By sacrificing the high signal-to-

noise ratio, an increase of the number of intermediate levels is possible through

slightly reducing the laser fluence to a subthreshold value (~11 mJ/cm2), as shown

in Figure 4.39. An alternative approach to increase the number of possible levels in

a memory cell could be to decrease the film thickness. The example in Figure 4.47

shows the optical reflectivity of an only 50 nm thick GeTe film upon ns laser pulse

irradiation at the identical fluence of 26 mJ/cm2 as a function of the pulse number.

As a fully crystallized film can be obtained only after 12 applied pulses, each pulse-

induced crystallization of GeTe film leads to up to ten different reflectivity levels

Figure 4.47: Optical reflectivity and possible indicated multi states of a 50 nm thick GeTe film as a

function of the laser pulse number.

0 5 10 15 20 80 100

0

20

40

60

80

100

Re

lative

re

fle

ctivity (

%)

Number of pulses

X

IX

VIII

VII

VI

V

IV

III

I

II


114

(as indicated in Figure 4.47) within a single memory cell, which is much higher in

comparison to those of the thicker film. For a smaller thickness of phase-change

films, more possible intermediate levels of reflectivity can be expected. On the

other hand, it should be noted that the levels must be accurately identifiable so that

they can be used for reading out stored data. A possible explanation for the effect of

film thickness on the varying number of intermediate levels is the fact that

nucleation and growth of crystalline regions in GeTe films is strongly influenced by

the film thickness, as already discussed in Chapter 4.2.1.

Broad wavelength response of reflectivity 4.3.8

As described in the sub-chapters above, the reflectivity readout of the laser

irradiated GeTe film was realized with a probing diode laser at a wavelength of

405 nm in the setup of the pump-probe system. To examine the multilevel

reflectivity of phase-change films over a wide wavelength range, a series of pulse

number-dependent optical reflectivity measurements of GeTe films in the

wavelength range from the near ultraviolet to the near infrared (400 and 700 nm)

was investigated. Each of the conditions of the ns laser-excited GeTe films was

prepared and the broad wavelength response of reflectivity was performed by UV-

Vis spectroscopy measurements. Figure 4.48 shows the optical reflectivity of an as-

deposited GeTe film and of films after 20 ns laser pulse irradiation with pulse

numbers up to 20 at a constant fluence of 26 mJ/cm2. The optical reflectivity of the

as-deposited film varies from 66 to 70% in the wavelength range between 400 and

Figure 4.48: Optical reflectivity of GeTe films as a function of the pulse number within the

wavelength range from 400 to 700 nm

400 500 600 700

65

70

75

80

85

90

20 pulses

4 pulses

3 pulses

2 pulses

Re

fle

ctivity (

%)

Wavelength (nm)

as-dep.


115

700 nm. After two laser pulses are applied, the reflectivity of the film within the

measured wavelength range is slightly increased, by around 3%. A further

enhancement by 6% appears after three pulses. A more distinct reflectivity of the

film after irradiation with four pulses is in the range from 80 to 82% (reflectivity

contrast ~14% at 650 nm). Upon further increase of the number of applied laser

pulses to, for example, 20 pulses, the reflectivity increases only slightly further by

~3%, but the effect is no longer so pronounced. Moreover, the changed optical

reflectivity of GeTe films after ns laser excited pulses is almost independent of the

wavelength, except for the maximum variation of around 4% increase from the

wavelength of 400 nm to 650 nm in the case of the films after 20 pulses irradiation.

The highest optical reflectivity contrast of 20% at a wavelength of 600 nm is

observed between as-deposited and films after 20 pulses. Therefore, in general, the

broad wavelength response (400-700 nm) of multi-states reflectivity increase in the

GeTe film is possible for the controllable crystallization process in the GeTe film

with induction by the appropriate number of applied laser pulses.

Summary of the results on dynamic optical switching of non-volatile 4.3.9

multi-level memory

Rapid and reversible amorphous-crystalline phase transformations in thin single-

layered PLD-deposited phase-change GeTe films were studied. Based on the

construction of a UV ns laser pump-probe system, non-volatile switching in GeTe

film based cells was determined by distinct changes of the optical reflectivity on the

time scale of nanoseconds. Reversible optical switching was demonstrated; the

results included a switching threshold fluence of ~14 mJ/cm2 and an energy

accumulation as function of the applied number of pulses, leading to switching On

parameters of 26 mJ/cm2

with 20 pulses and switching Off parameters of

112 mJ/cm2 with a single pulse. Thus, by appropriately applying laser fluence and

pulse number, the crystallization and amorphization processes in GST films could

be feasibly controlled, which is important for realizing multi-level phase-change

memory cells. Next, the structural transformations between amorphous and

crystalline states, as well as the local cross-sectional microstructure changes

correlated with the optical switching process, were investigated, which provided an

equally important understanding of the fast kinetics of the phase transition


116

mechanisms that were present. Upon the detailed examination of the dynamics of

phase transformation, the speed of switching On process (crystallization) was

determined to be ~300 ns and that of the switching Off process (amorphization) to

be ~3 ns. The dependency of the velocity of crystallization was found to be

inversely proportional to the number of irradiation pulses (≤ 5 pulses). Moreover,

the crystallization of a melt-quenched GeTe film required less energy than that of

an as-deposited amorphous film. On the basis of the confirmed experimental results,

a general optical switching model was established, which was helpful for

understanding the mechanism of multilevel operation. Increasing numbers of multi-

levels (up to 10 distinct levels) could be produced by choosing a GeTe film

thickness of 50 nm. The multi-level optical reflectivity was realizable at various

wavelengths between 400 and 700 nm and the highest contrast of ~20% was

obtained at a wavelength of 600 nm. Unequivocally, these experimental results

illustrate that dynamic optical switching of non-volatile multi-level memory, based

on PCMs, is not only promising for the next generation of optical data storage

memory applications, but also for improving the understanding of the physical

aspects of these unique material properties.

4.4 Nanoscale bipolar electrical switching of phase-change materials

117

Nanoscale bipolar electrical switching of phase-change 4.4

materials

The laser pulse-induced structure transition from the amorphous to the crystalline

state in phase-change materials for the optical switching process was illustrated in

the previous chapters (4.2 and 4.3). In this chapter, the phase transformation

between the amorphous and the crystalline state of the chalcogenide materials

induced by electrical pulses is demonstrated. This transformation process is the

basis of phase-change random access memory (PCRAM) devices. The

investigations on electrical switching of phase-change materials at the nanoscale

comprise, in particular, polarity-dependent resistance switching, the threshold of the

voltage for resistance switching, as well as the underlying mechanisms.

Background 4.4.1

Currently, floating-gate FLASH memory devices dominate the non-volatile

memory market, due to their low cost. However, significant challenges emerge for

the scaling of these devices, when the characteristic dimensions are to be reduced

towards the 32 nm size and below. Alternative candidates for the conventional, non-

volatile memory technology utilize the magneto-resistance (MRAM) [148],

ferroelectricity (FRAM) [172], phase-change random access memory (PCRAM) [4]

and resistive random-access memory (ReRAM) [173] technologies. Among these

memory technologies, PCRAM and ReRAM are becoming increasingly attractive

for the next generation of non-volatile memory.

PCRAM is based on dramatically different electrical resistivity states (up to five

orders of magnitude) between the amorphous and the crystalline phases of

chalcogenides materials. This large difference makes phase-change alloys ideal

candidates for solid-state memory devices. The renewed interest in PCRAM

technology was triggered by the discovery of very fast crystallizing materials such

as GST or nitrogen and/or oxygen doped GST. These materials can crystallize in

less than 100 ns. In PCRAM, the amorphous phase-change material is crystallized

by heating it above its crystallization temperature (the so-called SET operation) and

it is melt-quenched to re-obtain the amorphous phase (the so-called RESET

operation). These operations are controlled by application of an electrical current.


118

High power pulses for the RESET operation place the memory into the high

resistivity state, whilst pulses of moderate power but longer duration return the cell

to the low resistivity SET state. The remarkable, advantageous properties of

PCRAM processing include: non-volatility, long cycling life, relatively low power

consumption, fast read/write speed (ps range), and prominent resistivity contrast.

Without the threshold voltage for resistance switching, a PCRAM cannot be

used because, in the high resistivity state, extremely high voltages would be

required to deliver sufficient Joule heating power to the cell to heat it above the

crystallization temperature [68, 69]. This represents a major difference between

PCRAM and ReRAM. However, both non-volatile memory technologies – PCRAM

and ReRAM – share some similarities in function and performance (including non-

volatility, scalability and switching speed) that potentially overcome the limitations

of FLASH memory, as previously discussed in the comparison of emerging

memory technologies in Chapter 2.3.2.

ReRAM is based on an electrical resistance switching process caused by the

formation and rupture of a local conduction path in a solid electrolyte, the so-called

conductive filament, which can be formed into a conductor generated by the

application of a sufficiently high voltage. When the conductive filament is

generated, it is possible to switch towards a high resistivity state by breaking the

filament through application of an opposite voltage (RESET) and back to the lower

resistivity state through reconstruction of the filament by the original voltage (SET;

for more details see Ref. [174]). Because the respective high or low resistivity state

remains unchanged until a reverse polarity voltage is applied, this type of memory

has a non-volatile character. The write time, at approximately 10 ns, is

fundamentally shorter than those needed to change the charge state of a floating

gate in FLASH memory. Such behavior has been found in various material systems:

binary transition metal-oxides, perovskites, solid-state electrolytes and, crucially,

also in phase-change chalcogenides such as GST and AgInSbTe [72, 174-178]. The

promising properties of ReRAM are the simplicity of the structure, low switching

voltage, fast switching speeds, excellent scalability and high compatibility with the

CMOS technology.


119

With the development of PCRAM and ReRAM technologies, several key

challenges arise. For PCRAM, high density integration is of utmost importance and

it requires a small memory cell size. However, in practice scaling down is mainly

limited by the process of lithography technology during the fabrication of PCRAM.

In addition, a high RESET current is required in order to melt/quench the

crystallized PCM and bring it back to the amorphous state. Another critical problem

is the possible mechanical failure of the interface between the PCM and the heating

contact after a certain numbers of cycles that can be easily induced by the repeated

melting/quenching and recrystallization processes. For ReRAM, once they are

formed, the nature of such filaments fluctuates; this was described in the literature

on the basis of the filamentary theory [174], and leads to non-uniformities of the

high/low resistance switching states. Furthermore, the on/off resistance contrast is

relatively low, which is unfavorable for the signal-to-noise ratio in an electrical

resistance switching memory application. A concise comparison of PCRAM and

ReRAM regarding their advantages and disadvantages is presented in Table 4.5.

In order to eliminate these issues and optimize the cell performance, it is

proposed here that the concepts of the PCRAM and ReRAM could be combined

(we call such a combination “PC/Re-RAM”). As a prospective view, a unified

approach for studying correlations between phase-change and electrolytic behavior

Table 4.5: Advantages and disadvantages of PCRAM and ReRAM.

Performance PCRAM ReRAM

Advantages multi-bit cell

high on/off ratio

long term retention

simple structure

low operation voltage

Disadvantages high RESET current

cycling endurance

stable retention

low on/off ratio

Figure 4.49: Schematic of the sample structure used in C-AFM experiments.

Conductive tip

GST

Cr bottom electrodeSiO2/Si

V


120

may prove crucial for improving both device performance and reliability [179]. In a

combined PC/Re-RAM device, the chalcogenide-based amorphous GST layer is

introduced as a solid electrolyte between metal contacts, which is similar to the

typical ReRAM device of the metal/insulator/metal type. An exemplary

experimental setup using a Pt-Ir coated AFM tip as top electrode and a Cr bottom

electrode is schematically shown in Figure 4.49. Based on the C-AFM experimental

setup, a sharp tip contact on the phase-change film is not only beneficial for the

switching operation with a demanded low power consumption, due to the small

contact area and corresponding to an increase in heating efficiency, but it also

models the device switching in a nanoscale investigation, which may help to

develop a cost-effective high-density memory. During the sample preparation

process, any additional thermal treatment of the samples was avoided. The detailed

configuration of the C-AFM setup, its calibration and operation can be found in the

experimental section (see Chapter 3.4).

For clarification of the following text, low resistance state (LRS) means

crystallization of GST with atomic migration of ions forming conductive bridges,

and high resistance state (HRS) is either the initially amorphous state or the state

with nanocrystals existing in the GST film but without any conductive bridge. The

switching mechanism is depicted schematically in Figure 4.50, and the typical I-V

behavior for the combination of phase-change material as the active layer in a

memory device is shown in Figure 4.51. Figures 4.50 and 4.51 (A) represent the

pristine state. The memory cell features a HRS, which is typical for an amorphous

GST layer (as confirmed by TEM shown in Figure 4.61a) between the top electrode

(PtIr) and the bottom electrode (Cr). During the initial electroforming process (in

Figure 4.50 (B) and A B C in the I-V path in Figure 4.51), an applied electric

pulse (irrespective of its polarity) with a voltage around the subthreshold voltage

between electrodes acts as Joule heating source. This electrical pulse induces local

crystallization in the amorphous GST film. In the ideal case, the result of this

crystallization process is a crystalline low-resistivity GST filament produced in the

initial high-resistivity amorphous phase matrix. The switching process is based on

the difference between the high- and low-resistivity states. Simultaneously, an

electric pulse induced by an electrical field can lead to electrochemical reactions


121

inside the chalcogenide material, which result in a local ionic conduction, involving

Te or Ge ions. In GST, Te atoms act as anions, whereas Ge atoms act as cations

[180]. Thus, under the influence of the electrical field, migration of these ions is

along opposite directions. The mobility of the ions in the filaments of locally

crystallized material inside the GST film depends on the electrical field distribution

present between the electrodes. Only if the electric field strength is sufficiently

large, both the phase-change and the ionic migration mechanism can contribute to

form such electrically conductive filamentary pathways between the electrodes. It is

known that the crystallization in GST is nucleation-dominated in the case of formed

nanocrystallites. Such nanocrystalline GST exhibits a conductivity that is several

Figure 4.50: Schematic representation of the switching mechanism in PC/Re-RAM: (A) Pristine,

(B) Filament forming / SET state process. The red feature indicates local Joule heating. (C) ON

state (LRS), (D) RESET process, (E) OFF state (HRS).

Figure 4.51: Schematic representation of the current-voltage (I-V) curve for the combination of

phase-change material as an active layer of a bipolar resistive switching memory cell. The states

from (A) to (E) correspond to the schematic switching mechanism depicted in Figure 4.50.

A CB D E

Metal (e.g. Cr, Pt, Au, etc.)c-GSTTe- a-GST

OFF stateON state RESET processSET processPristine state

Curr

ent [a

.u.]

0 Voltage [a.u.]

Low state

high state

RESET

SET

A

C

B

D

E


122

orders of magnitude higher than the surrounding GST amorphous matrix [181, 182].

The movement of ions in the electrical field therefore follows the most favorable

diffusion path, such as the GST nanograin boundaries or/and the nanograins and

electrodes [183, 184], which might act as an electron conductive bridge, causing a

LRS of the cell, depicted as stage C (ON state) in Figure 4.50. As confirmed by

TEM and EDX analysis of PLD-deposited GST films (see section 3 in Chapter

4.4.4), the presence of generated GST nanocrystallites is evident, as well as a Te-

rich and Ge-depleted environment after electrical switching, when compared to

stoichiometric GST (225). However, if the voltage applied to the memory cell is not

high enough to form nanocrystalline GST grains, and also not sufficient to induce

migration of ions, the cell remains in the HRS.

When the polarity of the external electric field is reversed during the RESET

operation (Figures 4.50 and 4.51 D), the pre-existing electrical conductive filament

is disrupted, because of the ion repulsion from the GST nanocrystal grains

boundaries that is caused by electrostatic force. As a result, the electrical resistance

turns to the HRS (Figures 4.50 and 4.51 E). By application of a positive voltage, the

Te or Ge ions can be re-positioned between GST nanocrystallites and, thus, can

form a conductive filament again (see Figures 4.50 and 4.51 C). In conclusion, it

should be noted that a low bias voltage for the switching process is connected with

(1) the existence of small crystalline regions (grains) after the electroforming and

(2) the local movement of Te or Ge ions as a condition for the

construction/disruption of the conductive filaments between crystalline GST grains

and/or electrodes. This is the special quality of PC/Re-RAM memory cells, because

they require neither re-crystallization/re-amorphization (as is the principle of

PCRAM) nor purely an ion-built/dissolved filament (as is the principle of ReRAM)

for accomplishing the switching process. Apart from the physical switching

mechanism, possibly attributed to the combination of the phase-change of the GST

layer and the ionic migration of elements (e.g., Te, Ge) within the GST layer, as

described above, an effect of the redox reaction in the interface layer (e.g.,

formation of a thin GSTO layer associated with a filamentary conducting channel)

[177, 185], together with a phase-change in the GST layer between the electrodes,


123

could also explain the bipolar memory device switching. Detailed experiments to

elucidate the latter case are in progress.

Nevertheless, a comparison between the separated PCRAM or ReRAM and the

combined PC/Re-RAM device shows that the PC/Re-RAM provides the following

advantages:

(i) The RESET and/or operation voltage is smaller, because the resistivity

switching is dependent on the polarity and the heat loss is minimized by

the locally formed nano-contact between PCM and tip electrode, rather

than on the heat required for re-amorphization by local melt-quenching.

(ii) The On/Off resistance ratio is higher because of the confined current

flow with a higher conductivity of the generated GST nanocrystals,

reducing the non-uniformity of filament formation (owing to the

dependence of LRS on the filament) in ReRAM.

(iii) Excellent cycling endurance, since a smaller RESET/SET voltage is

required for operation, which can reduce the commonly encountered

problems in PCRAM, i.e. the largely thermal effects causing mechanical

device failure. In addition, the principle of ions migrating between

locally crystallized regions can avoid the crucial problem of randomly

forming filaments in ReRAM.

Polarity-dependent resistance switching of phase-change materials 4.4.2

The development of a concept for the combination of PCRAM and ReRAM is the

most important consequence of the presentation in the previous Chapter, in which

thin films of phase-change material were introduced as the active layer for bipolar

electrical switching memories. If this combination is realized in a PCM-based Re-

RAM (PC/Re-RAM), it can lead to applications in the fields of next generation non-

volatile memory, as well as cognitive computing, reconfigurable logic circuits, and

more.

Features of a PC/Re-RAM memory, such as excellent scalability (< 15 nm), low

SET/RESET currents (several nA), long retention (> 180 days), excellent cycling

endurance (> 1000 cycles) and high On/Off resistance ratios (about four orders of

magnitude), as well as a high estimated data density (0.7 Tbit/inch2), should be

ultimately attainable. Based on conductive AFM as an active switching tool, it is


124

possible to examine and to quantify the switching process on the nanoscale in a

laboratory environment. A schematic of the prototype cell (GST/Cr/SiO2/Si) and the

configuration based on C-AFM was already shown in Figure 4.49 and discussed in

detail in the experimental section (see Chapter 3.4). As an example, a measured I-V

characteristic of GST films is shown in Figure 4.52. The bias voltage is applied

between the bottom Cr electrode and the top electrode consisting of a PtIr coated

AFM tip, where the top electrode is connected to the ground and an amperemeter is

attached between the latter two, in order to measure the electric current flowing

through the GST film. By sweeping the voltage from -3 V→ 0 → +3 V → 0 → -3

V at a sweeping rate of 1.8 V/s (as demonstrated in the top left inset of Figure 4.52),

the I-V curve of the PtIr/GST/Cr stack shows typical bipolar switching with a

pinched hysteresis loop. There are two segments that correlate with the two

distinguishable resistance levels of a low ohmic resistance state (LRS) and a high

ohmic resistance state (HRS), characterized by the different slopes of the curves.

Figure 4.52: Measured I-V characteristic of GST/Cr/SiO2/Si arrangement with PtIr coated C-AFM

tip (diameter = 20 nm) as top electrode. The inset in the top left of the figure shows the

corresponding applied sweeping voltage as a function of time (at a sweep rate of 1.8 V/s) and the

inset in the bottom right of the figure represents the corresponding I-V curve with respect to the

absolute norm of the current at a logarithmic scale. The arrows indicate the directions of the voltage

sweep.

-3 -2 -1 0 1 2 3

1E-12

1E-11

1E-10

1E-09

1E-08

|Cu

rre

nt| (

A)

Voltage (V)

0 1 2 3 4 5 6 7-3

-2

-1

0

1

2

3

Vo

lta

ge (

V)

Time (s)

swee

p ra

te 1

.8 V

/s

-3 -2 -1 0 1 2 3-15

-10

-5

0

5

10

RESETCurr

ent (n

A)

Voltage (V)

HRS

current limit

Vth

current limit

LRS

Vh

SET


125

Initially, the cell stage of the as-deposited GST film is HRS (see black curve) and

sharp switching to LRS occurs after the bias voltage exceeds the threshold voltage

Vth, suggesting that the conduction is due to the formation of a filament in the GST

film. For the negative voltage polarity (black curve in Figure 4.52), the decreased

current indicates that the cell switches to the HRS and remains in HRS until the bias

voltage becomes positive again. With a further voltage sweep in this direction, the

current of the cell increases abruptly at the value of SET or of the threshold voltage

(about +1.55 V) caused by switching from the HRS to the LRS. If the voltage

exceeds +1.55 V, the sample current rises dramatically up to >10 nA, which is the

current limit for preventing electrical breakdown. Following reversal of the voltage

sweeping (red line in Figure 4.52), the cell initially maintains the LRS at the

positive voltage polarity, then the current gradually declines with the decrease of

the bias voltage with an almost linear I-V behavior, as observed for the voltage

sweep in the range between around +0.9 V and 0.8 V. For voltages below 0.8 V,

the current is limited to a constant level of 12 nA, similar to the positive current

limit. It should be noted that with the very low current at the voltage below the so-

called holding voltage (Vh), the HRS is reached. It is possible for the state of the

cell to return back to HRS by the RESET operation in the negative voltage polarity

regime. The representation of the I-V curve on a logarithmic scale, using the

absolute value of the current, is shown in the inset of Figure 4.52 (bottom right).

The I-V curve loops in a pinch to zero current at zero bias voltage and it changes in

the bias-polarity-dependence current of the curves, which is identical to a common

capacitor-like cell structure. Moreover, the current through the GST film rises

suddenly up to a level of three orders of magnitude higher than the prior state, when

the voltage is swept above the threshold voltage for crystallization. Similar pinched

hysteresis I-V curves were also found during investigations of the I-V curve as a

function of the repeated switching, voltage sweeping rate, and film thickness

(shown in the following), indicating that the experimental results for polarity-

dependent switching coincide well with the proposed switching mechanism as

introduced in Chapter 4.4.1.

The intrinsic memory effect of the cell was found to be reproducible. The I-V

curves of the device for the representative first ten sequences of repeatable


126

switching are plotted together in Figure 4.53a. It shows that the bipolar switching

characteristics of the PtIr/GST/Cr cell are well maintained during repeated cycling.

The cell can be SET (HRS) by applying a positive bias and RESET by applying a

negative bias. Extended endurance characteristics were determined by repeated

sweeping of the voltage between -3 V and +3 V at a sweeping rate of 1.8 V/s and at

a temperature of 85°C up to 1000 cycles, as shown in Figure 4.53b. The resistance

of the HRS shows an enhanced scattering after 600 cycles, whereas the resistance of

the LRS remains mostly constant. However, a considerable difference between HRS

and LRS is clearly maintained during the 1000 cycles. The two levels of resistance

switching range between about 8 MΩ and 100 GΩ, which corresponds to

approximately 3-4 orders of magnitude between the LRS and the HRS.

The switching probability of the prototype cell is also obtained from data in

Figure 4.53b and is shown in Figure 4.53c. The probability is the calculated

cumulative resistance distribution probability of the different cycle numbers (up to

Figure 4.53: (a) Repeatable switching I-V curves of a Pt-Ir/GST/Cr cell. Only the first ten repeated

switching cycles are shown, for the sake of clarity. Insert: corresponding log-scale switching I-V

curves show LRS/HRS ratio of ~factor 5×103. (b) Cycling endurance characteristics of a Pt-

Ir/GST/Cr cell at a cell temperature of 85°C for 1000 cycles. (c) Corresponding cumulative

probability of resistance distribution of the cell. (d) Data retention characteristics of the cell at room

temperature for more than 180 days.

(a) (b)

(c) (d)

-2 -1 0 1 2-15

-10

-5

0

5

10

current limit

current limit

RE

SE

T

Cu

rre

nt

(nA

)

Voltage (V)

SE

T

Repeated switching

-2 -1 0 1 210

-12

10-11

10-10

10-9

10-8

0 200 400 600 800 100010

7

108

109

1010

1011

1012

Resis

tance (

)

Number of cycles

LRS

HRS

~ 5

1E7 1E8 1E9 1E10 1E11 1E12

0

20

40

60

80

100

Cum

ula

tive p

robabili

ty (

%)

Resistance ()

1 10 10010

8

109

1010

1011

1012

Ron

Roff

Re

sis

tan

ce

(

)

Retention time (days)


127

1000 cycles). The resistance of the HRS shows a slight fluctuation between 1011

Ω

and 1012

Ω. The resistance of the LRS is quite uniform at 8×108 Ω. The ratio of the

HRS to the LRS remains without remarkable degradation over the range of 3-4

orders of magnitude of the resistance.

Furthermore, to investigate the data retention characteristics of the PtIr/GST/Cr

cell, the electrical properties have been measured for a time interval exceeding 6

months. The variation of the resistance as a function of the retention time is

presented in Figure 4.53d. It shows that in such a cell both HRS (Roff) and LRS

(Ron) can be maintained at a high ratio of 3-4 orders of magnitude for several tens of

days. Subsequently, the values of HRS are slightly reduced, but are still clearly

distinguishable from the LRS (one order of magnitude difference) after a retention

time of more than 180 days. In general, a higher ratio of HRS to LRS can be of

advantage not only for the signal-to-noise ratio in data transformation but it is also

promising for multilevel memory technology applications (analogous to the

principle of optical multi-level storage; see Chapter 4.3). The achieved ratio of HRS

to LRS of 3-4 orders of magnitude in this work is higher than in the conventional

ReRAM [186]. Such a high HRS/LRS ratio can possibly be explained by the

proposed electrical resistance switching mechanism in GST films that is described

in detail in Chapter 4.4.4.

Figure 4.54 shows the measured I-V curves of typical features of pinched

hysteresis loops and sweeping rate-dependent hysteresis in the Pt-Ir/GST/Cr cell.

The bipolar switching behavior of the cell, in principle, seems to be almost

independent of the voltage sweeping rate in the investigated range between 1.8 and

120 V/s (see Figures 4.54a-e). Moreover, it is worth noting that, with respect to the

resistance values of the cell obtained from Figures 4.54a-e, the values of the LRS

(Ron) lie very close on a common resistance level, whereas the level of the HRS

(Roff) can be adjusted by changing the sweeping rate, as demonstrated in Figure

4.54f. At a sweeping rate ≤ 3.6 V/s, the resistances (Roff) are approximately

5.2×1011

Ω, which is ~4×103 times higher than the LRS (Ron). In contrast, at a

higher sweeping rate (e.g., ≥ 12 V/s), the Roff values are approximately 1.8×109 Ω,

which is ~15 times higher than the HRS (Roff) values. Such distinct levels (I and II)

of Roff obtained at different sweeping rates can allow for possible multi-bits per cell


128

(multi-level storage), as indicated in Figure 4.54f. The reduction of the HRS (Roff)

at higher sweeping rates, as compared to lower sweeping rates, may be explained by

the shorter duration of the applied pulse, progressively narrowing the filaments

within the GST layer that connects both electrodes and, consequently, resulting in a

lower resistance.

Figure 4.54: (a) to (e) representative I-V curves of a Pt-Ir/GST/Cr cell at different voltage sweeping

rates from 1.8 V/s to 120 V/s. (f) corresponding On/Off resistance as a function of the sweeping rate.

The arrows in (a) to (e) indicate the directions of voltage sweep. The dashed lines in (f) are only

intended as a guide for the eye.

-3 -2 -1 0 1 2 3-15

-10

-5

0

5

10

Curr

ent (n

A)

Voltage (V)

1.8 V/s

-3 -2 -1 0 1 2 3-15

-10

-5

0

5

10 3.6 V/s

Curr

ent (n

A)

Voltage (V)

-3 -2 -1 0 1 2 3-15

-10

-5

0

5

10 12 V/s

Curr

ent (n

A)

Voltage (V)

-3 -2 -1 0 1 2 3-15

-10

-5

0

5

10 60 V/s

Curr

ent (n

A)

Voltage (V)

-3 -2 -1 0 1 2 3-15

-10

-5

0

5

10 120 V/s

Curr

ent (n

A)

Voltage (V)

(a) (b)

(d)

(f)

(c)

(e)

1 10 100

108

109

1010

1011

1012

~factor 4 103

Level I

Roff

Ron

Resis

tance (

)

Sweeping rate (V/s)

Level II

~factor 15


129

Threshold of resistance switching 4.4.3

The threshold of resistance switching plays an essential role in the operation and

performance of resistance switching memory cells. As shown in Figure 4.52 (in

Chapter 4.4.2), the resistance of the HRS is suddenly switched to the LRS at a

specific electrical voltage, known as the threshold voltage (Vth). The resistance in

the LRS up to this critical voltage is called the holding voltage (Vh). As long as the

sweeping voltage is lower than Vh, the current across the GST layer tends towards a

value close to zero, due to the change of resistance from the LRS to the HRS.

Figure 4.55 summarizes the statistic values of Vth and Vh measured at 12 different

points on the sample surface, as marked by crosses in the topography image (500 ×

500 nm²) of the ~50 nm thick GST film. On average, a threshold voltage Vth of 1.55

± 0.4 V and a holding voltage Vh of 0.5 ± 0.1 V are obtained by sweeping the

voltage from 0 up to +3 V and from +3 V down to 0, with a sweep rate of 1.8 V/s.

Thus, the results for Vth and Vh are stable and reproducible for a homogeneous GST

film of a given film thickness.

However, it is worth noting that the Vth values are tunable by varying the GST

film thickness. These experimental results are shown in Figure 4.56. Values of Vth

in the range between 1.2 and 4.2 V are obtained for GST film thicknesses varied

between 10 and 300 nm. Figure 4.56a shows the dependence of Vth on the film

thickness. For film thicknesses > 50 nm, the threshold voltages are approximately

constant. However, once the film thickness is less than 50 nm, the values of the

threshold voltage increase exponentially with decreasing film thickness. It is noted

Figure 4.55: Distribution of threshold voltage and holding voltage measured at various probe points,

marked in the inset, showing the surface topography of the investigated GST film. The numbers in

the inset mark the first, fifth and ninth measuring points.

4 nm

0 2 4 6 8 10 120.0

0.4

0.8

1.2

1.6

2.0

Number of probe point

Th

resh

old

vo

lta

ge

Vth (

V)

0.0

0.4

0.8

1.2

1.6

2.0

Ho

ldin

g v

olta

ge

Vh (

V)

1

5

9


130

that the positive polarity Vth is slightly different from the negative polarity Vth. This

is possibly related with the formation of an asymmetric trap-band structure within

GST under both positive and negative electrical fields, because structural defects in

the interface of GST films and electrodes are involved [187]. The values for the

threshold voltage in Figure 4.56a were extracted by fitting the data points of the

current at the critical sweeping voltage, as described in Figures 4.56b and c. It is

obvious that, above Vth, I-V characteristics of GST films at various film thicknesses

show a linear, i.e., ohmic relationship, indicating the respective resistance states of

the films switched to LRS. Consequently, the relationship of resistance and film

thickness, as shown in Figure 4.56d, illustrates that the LRS is uniform with a

change of the film thickness, whereas the HRS is slightly increased when the

thickness of GST films is reduced to ≤ 50 nm. Nevertheless, the divergence of the

ratio of HRS and LRS is within less than one order of magnitude.

In a previous report, it was shown that the crystallization temperature of GST

films is dependent on the film thickness in an inverse exponential relationship

Figure 4.56: (a) The threshold voltage as a function of film thickness. (b), (c) Positive and negative

I-V curves as a function of film thickness. (d) The electrical resistance of HRS and LRS as a

function of film thickness.

(a)

(d)0 1 2 3 4 5

0

2

4

6

8

10 10 nm

20 nm

80 nm

150 nm

300 nm

Cu

rre

nt (n

A)

Voltage (V)

Vth

-5 -4 -3 -2 -1 0-10

-8

-6

-4

-2

0

10 nm

20 nm

80 nm

150 nm

300 nm

Curr

ent (n

A)

Voltage (V)

Vth

0 50 100 150 200 250 300

1

2

3

4

Th

resh

old

Vo

lta

ge

(V

)

Film thickness (nm)

Positive Vth

Negative Vth

Fit by y=1.56+7.31e-0.082x

Fit by y=1.65+6.48e-0.081x

0 50 100 150 200 250 300

108

109

1010

1011

1012

Re

sis

tan

ce

(

)

Film thickness (nm)

HRS

LRS

(b)

(c)


131

[136]. The crystallization temperature was measured by using either in situ

resistivity or in situ XRD during annealing at identical heating rates. On the basis of

these experiments, it was concluded that the role of interfaces will certainly increase

with reduced film thickness, which leads to a larger influence on the crystallization

process of PCMs. This is in agreement with the trend of increasing Vth with reduced

GST film thicknesses in this study.

Mechanism of bipolar resistance switching in phase-change materials 4.4.4

The conventional physical mechanism for resistive switching based on phase-

change materials is Joule heating of the materials, resulting in a phase

transformation between amorphous and crystalline states. Correspondingly, the

current driven through the device switches the thin GST film to LRS when T > Tc

(SET), and switches back to HRS when T > Tm (RESET). This is unipolar resistive

switching, which is not dependent on the polarity of the bias voltage, because

negative voltages can also be applied for SET or RSET. However, the experimental

results for the I-V characteristics of a PC/Re-RAM device (see Chapter 4.4.2)

demonstrate that the resistive switching of the GST film presented here is, indeed,

dependent on the polarity of the bias voltage. Thus, bipolar resistive switching is

found here; this has previously been attributed to electrochemical metallization

(ECM) or to solid-state electrolytic behavior. Therefore, it is suggested that the

switching mechanism is a combination of both crystalline phase change in the GST

film and solid-state electrolytic behavior. The investigation is structured in the

following sections according to: 1) the Poole-Frenkel conduction mechanism, 2)

visualization of the polarity-dependent resistance switching processes provided by

C-AFM as an active switching tool, and 3) correlation of microstructure and

chemical composition of the C-AFM tip-induced marker by analysis of cross-

sectional HRTEM, NBD and focused EDX.

1.) Poole-Frenkel conduction mechanism

The Poole-Frenkel emission is one common conduction mechanism used to explain

the HRS current [188-190]. Defect-related localized states in amorphous phase-

change materials (here a-GST) act as electronic traps that are able to capture

electrons (so-called Anderson localization). However, when thermally excited, the


132

trapped electrons can be emitted into the conduction band, as depicted in Figure

4.57. The rate of this excitation of electrons into the conduction band is dependent

on the applied electrical field. The electron emission in the framework of the Poole-

Frenkel mechanism [191] is given by

𝐽 ∝ 𝐸 exp [−𝑞(∅𝐵−√

𝑞𝐸

𝜋𝜀)

𝑘𝑇] , (4.10)

where J is the current density, E is the electric field across the a-GST film, k is the

Boltzmann constant, T is the temperature, øB is the voltage barrier, q is the electron

charge, and ε is the dielectric constant of the material. Considering E = V/d and

J = I/d, where d is the film thickness, V is the voltage, and I is the current, equation

(4.10) can be rearranged to

𝑙𝑛(𝐼/𝑉) ∝ −𝑞∅𝐵−√ 𝑞3

𝜋𝜀𝑑 √𝑉

𝑘𝑇 . (4.11)

Figure 4.58 shows that this behavior fits well with the experimental results under

the assumption of the Poole-Frenkel effect being present in this case. Figure 4.58a

shows the relation of ln(I/V) versus √V in the HRS under negative bias for different

temperatures. The sweeping voltage is below -1.0 V, which is smaller than the

threshold switching voltage (~ -1.5 V). For an a-GST film thickness of ~18 nm (as

determined from the TEM image in Figure 4.61), the electric field is derived as

~105 V/cm. This is a sufficiently high applied electric field to drive the carriers

from Poole-Frenkel emission across the entirely amorphous GST matrix [188, 189,

192]. Here, the fit of the experimental results indicate that the Poole-Frenkel

Figure 4.57: Schematic representation of the Poole-Frenkel emission. Electrons at the valence band

edge inside amorphous GST are localized in trap states. Due to an external voltage (EF (left) –

EF(right)), the emission barrier for a jump towards the next trap state (or the electrode) is lowered

and can be overcome by thermal excitation. Electrons can thus flow across the insulating layer

without the external voltage exceeding the breakdown voltage.

qøB

a-GST

EF EC

EF

EV


133

emission is the dominant current conduction mechanism below Vth.

To gain more insight into the conduction mechanism, the conduction in the

amorphous GST film, as function of the bias voltage and temperature, is shown in

Figure 4.58: Analysis of the experimental results on the conductivity in the HRS, based on the

Poole-Frenkel effect: (a) ln(I/V) vs. sqrt(V) plots for a negative bias voltage, (b) ln(I/V) vs. 1/kT to

extract the activation energy Ea for carrier emission, (c) the activation energy plotted as a function of

sqrt(V) to determine the dielectric constant of the material.

0.2 0.4 0.6 0.8 1.030

31

32

33

34

negative bias

80°C

70°C

60°C

50°C

PF fit

ln (

I/V

) (S

)

sqrt (V) (V)

33 34 35 36-30

-31

-32

-33

-34

ln (

I/V

) (S

)

1/kT (eV-1)

bias - 0.2 V

- 0.3 V

- 0.4 V

- 0.5 V

0.4 0.5 0.6 0.70.40

0.45

0.50

0.55

0.60

Ea (

eV

)

sqrt (V) (V)

slope= -0.53

Intercept = 0.81 eV

(a)

(b)

(c)


134

Figure 4.58b. Low bias voltages between -0.2 and -0.5 V and temperatures between

50 and 80°C are investigated, as shown in the ln(I/V) versus 1/kT plots. Due to the

Arrhenius type of equation (4.10), the activation energy Ea for Poole-Frenkel carrier

emission of the amorphous GST film can be given as (considering again E = V/d)

𝐸𝑎 = 𝑞∅𝐵 − √𝑞3

𝜋𝜀𝑑 √𝑉 . (4.12)

The value of Ea can be obtained as extracted from the slopes of the Arrhenius plots

in Figure 4.58b, and is in the range from 0.43 to 0.56 eV for external voltages

between -0.5 and -0.2 V. Such values of Ea are consistent with the case in which the

Fermi level is pinned around the middle of the bandgap, as predicted by previously

reported bandgap energies of amorphous GST [68, 192-194]. Moreover, the

activation energy decreases as the bias voltage increases, as represented in the plot

of Ea vs. √V in Figure 4.58c. This is in agreement with previous reports [195],

indicating a hopping transport of carriers by thermal emission over a potential

barrier, which is generally lowered by the applied electric field [194, 196]. The

slope of the curve in Figure 4.58c provides the dielectric constant 𝜺 of the material.

With a very small measuring frequency (scan rate of 10 Hz), the dielectric constant

was accordingly estimated to be ~1.15. This value is quantitatively smaller than the

one predicted in the literature [197, 198]. By extrapolating to the zero bias point, the

trap energy (ET= qΦB) in the a-GST film is extracted to be around 0.81 eV. The

Fermi level is pinned close to the valence band edge in a degenerate p-type

semiconductor (i.e., close to EV, but electrons cannot occupy EF). Since the GST

conductivity is p-type, acceptor-like trap states determine the electronic conduction.

ET is located between EF and EV [68]. In summary, the conduction mechanism of

the HRS in a-GST films, based on the I-V characteristics as dependent on

temperature and applied electrical field, is identified as being based on the Poole-

Frenkel effect. For more details about this electrical conduction mechanism, see

Refs. [192, 194].

2.) Polarity-dependent resistance switching

Conductive-AFM is used to determine the surface topography, as well as the local

electrical conductivity of the phase-change thin films, as described above. Arrays of

marker points at the surface of amorphous phase-change films can be written by a


135

computer-controlled movement of the tip in the scanning contact mode and by

application of a voltage (exceeding the threshold switching voltage). Moreover, it is

also possible to record topography and current images simultaneously (see

description in Chapter 3.4) in order to visualize the resistance switching process.

Figure 4.59a shows an 8×10 array of marker points in the amorphous GST film

with a distance of 80 nm between the marker points. Figure 4.59b shows the current

profile of a chosen marker point row (dashed line in Figure 4.59a). The tip-induced

current markers in the amorphous GST films were realized by application of a

sweeping voltage of +3 V at a sweeping rate of 1.8 V/s. The local marker points are

evidenced by the electric current image that was obtained from the scanning contact

mode over the region (1×1 μm²) with an applied bias voltage of +0.5 V. The current

Figure 4.59: (a) A current image of a C-AFM tip-induced 8×10 array of markers in an amorphous

GST film and (b) a current profile plotted along the dashed line in (a).

0.4 nA

(a)

(b)

0 200 400 600 800 10000.0

5.0E2

1.0E3

1.5E3

2.0E3

2.5E3

Cu

rre

nt

(pA

)

Lateral distance (nm)


136

image exhibits a considerably high contrast between the tip-induced local marker

points (bright spots indicate a high conductivity) and the surrounding amorphous

material. It is possible to change the local electrical resistance/conductivity of the

amorphous GST film significantly by applying a writing current to the AFM tip.

The bright spots represent the contact area of the current-induced crystalline

filaments at the GST surface (see Chapter 4.4.1). The spots on the conductive

filaments at the surface have a diameter between 25 and 50 nm and they are

frequently non-homogeneous. This inhomogeneity of the markers at the surface

may be caused by improper tip-sample electrical contact, including a rough surface

topography, wearing-off of the conductive tip coating, or lateral spread of the Joule

heating in the presence of small inhomogeneities in the film. The current profile

extracted along a row of a current image, as marked in Figure 4.59a, demonstrates

that the tip-induced points are characterized by an approximately three times higher

current response, in comparison with the amorphous surroundings. The size of the

observed bright spots (in-plane filament dimension) is approximately proportional

to the probing current across the GST film. From the current profile at the

respectively formed in-plane filament dimension in Figure 4.59b (as plotted from

Figure 4.59a), the localization of the switching spots of, on average ~15 nm, is

reproducibly observed in the full width at half maximum. This implies that the

minimum dimension of the locally formed atomic conductive pathways within the

intra-connection of the GST film may be less than 15 nm. Details of the local

geometry of a conductive filament are described in the subsequent section.

Considering an in-plane filament diameter dimension of 15 nm, the resistance

switching behavior potentially allows for very high data storage densities up to

0.7 Tbit/inch².

The next step in this study was to analyze the switching behavior of these tip-

induced markers by changing the electrical polarity. The electrical resistance of the

tip-induced points can be switched between a low and a high resistance level by

reversing the polarity of the applied electric field. The experimental results are

summarized in Figure 4.60. For an operation bias voltage of ±0.5 V, topography (a-

c) and current images (d-f) were simultaneously recorded over the region of


137

500×500 nm². The corresponding current profiles (g-i) were plotted from the

current images along the dashed lines in (d-f). One exemplary tip-induced marker

point was selected to study the polarity-dependent resistance switching. The

topography (see Figures 4.60, a-c) remains unchanged during a switching process

between +0.5 V and 0.5 V and the reverse. By applying a positive bias voltage of

+0.5 V, the LRS of the marked region is clearly distinguished from the surrounding

amorphous phase, as shown in the current image (Figure 4.60d). Upon the polarity

change of the applied bias voltage to 0.5 V, the current flow through the GST film

at the marked point is in the range of several pA only (Figure 4.60e); i.e., it is much

lower in comparison to the previously produced LRS (several nA). This indicates

that the electrical resistance of the marked point switches from the LRS to the HRS.

The electrical contrast between LRS and HRS is about three orders of magnitude

Figure 4.60: Polarity-dependent resistance switching of a GST film at the nanometer scale: (a-c)

topography images for bias voltages of +0.5 V, 0.5 V and +0.5 V, respectively; (d-f) current

images, corresponding to the images (a-c); (g-i) current profiles extracted from the marked lines in

the current images (d-f), respectively.

+0.5 V -0.5 V +0.5 V

0 100 200 300 400 500

0

1

2

3

4

Cu

rre

nt

(nA

)


0 100 200 300 400 500

0

1

2

3

4


Cu

rre

nt

(nA

)

0 100 200 300 400 500

0

1

2

3

4

Cu

rre

nt

(nA

)


Z = 5 nm

Z = 4 nm Z = 3 nm

Z = 2 nA

Z = 2 pA Z = 2 nA

(a) (b) (c)

(d) (e) (f)

(g) (h) (i)


138

(see corresponding current profiles in Figures 4.60g and h, respectively).

Furthermore, it is also clear from Figure 4.60f that the marker point has been

brought back to the LRS after another positive bias scanning with +0.5 V. The

current profile in Figure 4.60f indicates that the current across the GST film is in the

order of several nA; i.e., in the same range as before the process of switching

(compared with Figure 4.60g). The magnitude of the bias voltage in contact-mode

scanning for polarity switching is 0.5 V, which is lower than the threshold voltage

of about 1.55 V. This indicates a resistance switching of the GST film, as discussed

in the case of the I-V characteristic in Chapter 4.4.1. This implies that the observed

bipolar resistance switching is due to the change of the voltage polarity, causing

ionic migration, rather than to thermal heating of the material, resulting in a phase-

change process. In detail, only sufficient power is required to drive the local Te or

Ge ion movement as the functionality for bridging or rupturing conduction channels

between GST nanocrystal grains or grains and electrodes in the amorphous GST

film. Consequently, the bipolar electrical switching of the memory cell can be

realized with a lower bias voltage than expected. This required voltage is less than

the voltage needed to provoke a reverse phase-change in GST.

3.) Structural changes connected to the switching process

In order to understand the correlation between the local switching process and the

microstructure of phase-change GST films, thin lamellae of the phase-change

memory films were prepared with a FIB and analyzed by cross-sectional HRTEM,

NBD and EDX analysis. To locate the conductive channel regions in the GST

memory, the vicinity of tip-induced regions was tracked using pre-prepared FIB

markers. Results of the cross-sectional HRTEM studies for both as-prepared and

tip-induced regions of the GST film are shown in Figures 4.61a and b, respectively.

For the as-prepared GST films, including metal electrodes, a uniformly amorphous

and smooth GST thin film was found that had a thickness of ~20 nm between the Pt

top layer (that was deposited during the FIB process in order to protect the GST

surface from subsequent ion irradiation during FIB preparation) and the Cr bottom

electrode. A thin natural surface oxide of the GST layer, as well as at the surface of

the Cr layer, in the form of CrOx was detected (the effect of oxygen in this system is

being investigated further), as shown in Figure 4.61a. The long-range disordered


139

amorphous structure in the GST layer was confirmed from the 2D fast Fourier

transformed (FFT) image region shown in the inset of Figure 4.61a, corresponding

to the region marked by the square. In sharp contrast, after introducing an electrical

current locally into the tip-induced region, the GST presents nanocrystals with

lattice planes clearly visible in HRTEM imaging mode (see Figure 4.61b).

Figure 4.61: Cross-sectional TEM images, NBD and EDX analysis of GST based memory cell: (a)

HRTEM image of as-prepared GST film. (b) HRTEM image of tip-induced region. The insets are

the fast Fourier transformation patterns corresponding to the marked square regions in the HRTEM

images. (c) NBD images 1-6 corresponding to the numbered positions in the HAADF-STEM image

above. The diffraction pattern (4) can be identified as the [434] zone axis of the cubic Ge2Sb2Te5

phase with a tilt angle of ~1°. There are some residual diffraction spots in (5), resulting from the

electron beam running across the edge of the crystalline area. (d) Local EDX analysis corresponding

to the marked rectangle regions of (a) and (b); the dashed lines mark the composition of Ge2Sb2Te5.

(d)

(c)

(a) (b)

As-prepared Tip-induced0

10

20

30

40

50

60

70

Co

nce

ntr

atio

n(a

t.%

)

Ge

Sb

Te

1 2 3 4 5 6

1

4 5 6

2 3

Pt-Ir tip

GST

Cr

Pt-Ir tip

CrCr

GSTGST

As-prepared Tip-induced

CrOx CrOx

5 nm 5 nm


140

The spacing of these planes (~0.21 nm) corresponds well with the (022) planes

of cubic GST225. The corresponding FFT image (inset in the top left panel) is

extracted from the square region of the HRTEM image marked by the dashed box.

The presence of the spots and their specific arrangement in the FFT image indicates

a metastable cubic structure (Fm-3m). The corresponding local tip-induced

resistivity switching of the GST film with a bias voltage (> Vth) is illustrated by the

resulting I-V characteristic, as detailed in Chapter 4.4.2. It is clear that a sufficiently

high voltage applied between the conductive tip and the Cr bottom electrode leads

to the formation of a conductive nano-filament within the amorphous GST thin

film. Consequently, a high current flows through this conductive channel, which

generates Joule heating of its surroundings, and once the temperature and the

holding time that is required for GST crystal nucleation and growth is reached, the

formation of the crystalline phase occurs in the amorphous GST film. The geometry

of the conductive channel that accompanies the crystalline region is strongly

dependent on the distribution of the electric field [199-202]. The shallow valley

visible in the film surface of the tip-modified marker in Figure 4.61b may be caused

by the effect of the tip loading force and Coulomb attraction. From the cross-

sectional view, it appears that the contact radius of the tip and sample is around 10

to 15 nm and the penetration depth into the film is ~4 nm. Because of the shape of

the apex at the tip-film contact, the application of a voltage results in a radial

distribution of the electric field towards the bottom electrode, leading to crystallite

nucleation and successive growth starting from the tip-film contact (as sketched in

Figure 4.50). A quantitative investigation of the temperature profile in the presence

of an electrical nanoprobe, using a theoretical model, can be found in Ref. [200].

The estimated electric field of ~150 MV/m in this study exceeded the critical

electric field (~50 MV/m) for inducing crystallization of the GST film [202, 203].

In this scenario of a local, electric field-induced conductive channel and subsequent

thermal crystallization, the primary conductive channel may have measured 10 nm

or less in width. The NBD measurements further corroborate the information given

above.

A broadening in the lateral direction of the tip-induced crystallization is also

confirmed by the NBD measurement shown in Figure 4.61c. The displayed NBD


141

images 1 to 6 correspond to different positions in the STEM image marked by red

circles. For the NBD measurements, the STEM beam was adjusted to a very small

convergence semiangle of 1.6 mrad by means of the smallest available condenser

aperture (20 um) and appropriate lens settings. The resulting diffraction patterns

were recorded with a camera length of 145 mm, while a HAADF-STEM image

under the same instrument conditions was also acquired (upper part of Figure

4.61c). These NBD patterns thus individually carry electron diffraction information

from a region in the sample roughly 5 nm in diameter along the entire TEM sample

thickness of approximately 40 nm. The diffuse halo patterns obtained from spot

positions 1 to 3, which are located to the left side of the tip-film contact region,

demonstrate the presence of an amorphous phase. More specifically, the first diffuse

ring maximum is at 3.232 nm-1

, the second at 5.114 nm-1

, both with significant

broadening. As a reference, the octahedral sites in cubic GST225 have nearest-

neighbor distances of 3.316 nm-1

, whilst hypothetical tetrahedral sites would have

nearest neighbor distances of 3.829 nm-1

. It thus seems likely that the short range

order is of octahedral coordination (also in the light of the second diffraction ring),

as expected. Nevertheless, the structures from spot positions 1 to 3, as the NBD

patterns show, are all uniformly amorphous. Position 4 corresponds to the tip-

induced area at which the electron diffraction reveals the presence of the cubic

phase of GST. The positions and intensities of the diffraction spots suggest a [434]

zone axis orientation of the largest grain contributing to the diffraction pattern, with

a tilt misorientation of ~1°. Along the [434] zone axis, the dominant diffraction

spots from the (‒202), (‒3‒35), (‒1‒33), (‒35‒1) (‒551), (‒4‒68), (‒2‒66), (4-82)

and (6-80) planes are clearly identified, as indicated in Figure 4.61c. When the

beam position is moved to the right side of the tip-film contact region (position 5),

the amorphous ring contribution to the diffraction pattern reappears, with only a

small crystalline component running across by the NBD beam. By moving the beam

position further towards the right, no more crystalline diffraction is observed, as

seen for position 6. Thus, it can be concluded that the GST nanocrystals formed by

tip-induced switching probably form a cylindrical region across the film. This is

consistent with the observations shown in Figure 4.61b.


142

The local compositional changes between as-prepared and after tip-induced

modification of the film were analyzed by STEM-EDX. The quantification results

shown in Figure 4.61d are extracted from the regions in Figures 4.61a and b that are

enclosed by dashed lines. The measured relative atomic concentrations of elements

Ge, Sb and Te in the as-prepared and tip-induced region are Ge22.1Sb21.5Te56.3 and

Ge17.7Sb20.3Te62.0, respectively. Considering the experimental error of the EDX

technique (up to ± 4 at.%, depending on S/N ratio and peak overlap), the

composition of the as-prepared film can be assumed to be that of Ge22.2Sb22.2Te55.5,

whereas the composition of the tip-induced region is slightly over-stoichiometric,

regarding the Te content, and under-stoichiometric, regarding the Ge content. The

results suggest that, during the voltage application in the tip-induced local switching

experiment, either Te or Ge atoms might migrate, due to ionic migration that

depends on the electric field (i.e., electropositive Ge with lowest ionization potential

among GST and electronegative Te ions) or to segregation of material components

during local Joule heating. This could be attributed to the movement of Te or Ge

ions forming bridges or ruptures of conduction channels between GST crystal

grains or between grains and electrodes in the amorphous GST film upon polarity

switching by voltage application, resulting in resistance switching (for more details,

see Chapter 4.4.1). Very recently, it has been reported [72, 178, 204-207] that, in

some compositions of the ternary system Ge-Sb-Te, bipolar resistive switching was

also observed, which was attributed to either Sb or Te ion migration, based on solid-

state electrolytic transport. Moreover, if the entire current in the I-V characteristic

flows along the ~10 nm wide channel formed by the GST nanocrystallites between

the Pt-Ir tip and the Cr electrode, as found, the current density and resistance would

be 104 A/cm² and 10

8 Ω, respectively. These values are two orders of magnitude

lower than those at which electromigration typically occurs in metal interconnects

(e.g., of Cr, Pt, etc.). Therefore, it should be possible to exclude an influence of in-

diffusing electrode material in the present study. In other words, bipolar resistance

switching behavior of a PC/Re-RAM device should, in this regard, be independent

of the metal electrode materials.


143

Summary of the results on nanoscale bipolar electrical switching of 4.4.5

GST films

A concept for a PC/Re-RAM memory device was developed that was based on the

combination of both phase change mechanisms and the electrolytic properties of

PCMs. For this purpose, the active layer of a GST phase-change film based memory

was prepared by PLD. The properties and mechanism of the simple device were

verified for actual bipolar electrical switching. Based on using advanced C-AFM

techniques, bipolar electrical switching on the nanoscale (~15 nm) could be

demonstrated, and high data storage densities (0.7 Tbit/inch2) could be estimated on

the base of the present device setup. A high On/Off resistance ratio of ~5×103 was

obtained and the switching process could be influenced by voltage sweeping rates in

the regions ≤ 3.6 V/s and ≥ 12 V/s, showing distinct resistance values that reveal the

possibility of realizing several multi-levels per cell. A retention time of > 180 days

and a cycling endurance of > 103 times of such memory cells was measured.

Regarding the switching energy, above the obtained threshold voltage of ~1.5 V,

the current increased to a level in the range of only a few nA; this limited the

required power to a relatively small amount. Moreover, the threshold voltage was

influenced by film thickness. The conduction mechanism in amorphous GST films

was explored by an analysis based on the Poole-Frenkel conduction model. In

addition, the operation of a polarity-dependent switching process was revealed by a

local visualization using simultaneously measured current and topography images

of the GST film. The investigation of local conductive AFM tip-induced switching,

in correlation with a local microstructure and chemical composition analysis by

STEM, suggested the formation of cubic GST(225) nanocrystals within a

cylindrical region with a diameter of ~10 nm across the film, as well as a slight

deviation from stoichiometric GST (225) to a Te-rich and Ge-depleted composition.

From these results, the present switching mechanism was proposed to consist of a

combination of the usual phase changes in GST films and the migration of mobile

Ge/Te ions. The demonstrated creation of this novel type of PC/Re-RAM memory

does not only open up new opportunities for exploring physical mechanisms but

also proves PCMs to be promising candidates for next generation non-volatile

memory applications.

145

Chapter 5

Summary

This thesis covers the key challenges of PCMs in science and technology

development, regarding usage of PCMs in data storage applications. For this

purpose, PCM thin films of GST and GeTe were prepared with the PLD technique.

These films were the basis of a thorough investigation of phase transformations of

PCMs with respect to structural change, crystallization kinetics as well as local

microscopic mechanisms. The structural, optical, and electrical properties of these

films were examined with a particular focus on the switching processes. Finally, the

thesis explores topics such as controllable multi-level storage, scalability, and new

switching mechanism of PCMs for next generation non-volatile data storage

applications. Briefly, the contributions of this thesis include the following:

1. The ex situ thermally induced crystallization process of PLD-prepared GeTe

films was characterized regarding relevant film properties such as structure,

chemical composition, topography, optical reflectivity, and electrical resistivity.

The crystallization temperature for pure GeTe film was found to be 240°C.

Investigation of the crystallization kinetics revealed the activation energy to be

3.14 eV. Moreover, a higher crystallization temperature Tc (280°C) and a higher

activation energy for crystallization Ea (3.18 eV) could be obtained with oxygen

incorporation into GeTe films. The correspondingly higher achievable thermal

stability makes them excellent candidates for phase change application.

2. The in situ thermally induced crystallization process of GeTe films was explored

for dominant growth temperature regimes related to the amorphous, polycrystalline,

and fiber textured states. It was found that the optical reflectivity increased with the

corresponding degree of crystalline quality. An even higher degree of crystalline

quality was achieved by realization of epitaxial growth of GeTe films on BaF2(111)

substrates by PLD.

Chapter 5 Summary

146

3. The ns UV laser-induced phase transformations of PLD-prepared GeTe films

were identified by optical reflectivity and structural evaluation. It could be

demonstrated that the processes of crystallization and amorphization are

controllable as functions of the laser irradiation parameters pulse number, fluence,

and pulse repetition rate. With a threshold fluence of as low as 11-14 mJ/cm2 with

pulse numbers ≥ 5, the amorphous films crystallize in the rhombohedral GeTe

phase. In contrast, with higher threshold fluences of above 36 mJ/cm2 with pulse

numbers ≥ 5, the amorphous films crystallize in the cubic GeTe phase. The re-

amorphization of the crystallized films could be achieved by using single pulses at

fluences between 162 and 182 mJ/cm². By using these determined parameters,

reversible phase transformations of GeTe films were realized.

4. The crystallization of GST thin films induced by either ns or fs single laser pulse

irradiation was compared at the same wavelength of 248 nm. The appearance of a

crystalline state was found after ns-laser and fs-laser irradiation with fluences

≥ 36 mJ/cm² and ≥ 13 mJ/cm², respectively, by evaluation of structural and optical

reflectivity. Local microstructure analysis revealed that GST films irradiated with

fs-laser pulses exhibit a columnar growth, whereas, in ns-laser irradiated films, oval

shaped crystallites were produced. The corresponding growth mechanisms are

clearly different. Based on the fluence dependence of the optical reflectivity of GST

films, a model of irradiation-induced crystalline phase formation was used to assess

the average contribution of ns and fs laser photons to the phase transition process.

The obtained results demonstrate that single fs pulse laser irradiation appears to be

more effective in inducing crystallization of GST thin films than single ns pulse

laser irradiation, enabling the attainment of a higher data transfer rate. In particular,

the results gained from the fs single laser pulse irradiation provide a better

understanding of the rapid phase transformation mechanism in practical PCM

applications.

5. Inspired by the multi-level storage technique to increase storage density and

reduce energy consumption, non-volatile multi-level phase-change memory based

on single-layer GeTe films was developed. Distinct multi-levels, instead of the

commonly used binary states, could be achieved by generating intermittent states,

using ns UV laser irradiation with a controlled combination of laser parameters.

Chapter 5 Summary

147

Each laser pulse excitation induced a partial crystallization, and complete

crystallization was achieved by a succession of such excitations. Various

reflectivity states could be realized by producing different volume fractions of

crystallized material. A fluence of 26 mJ/cm2 and a specific sequence of pulse

numbers were used for the crystallization of the films. In the reverse process, a

single laser pulse at a higher fluence of 112 mJ/cm2 was used for amorphization of

the films. The formation of these distinct levels was shown to be reproducible.

Regarding the dynamics of the described phase transformations, the switching times

for the ON operation (complete crystallization) and OFF operation (amorphization)

were found to be ~300 ns and ~3 ns, respectively. The local microstructure phase

transformation processes were correlated to the observed reversible switching of the

optical reflectivity, which are important for understanding the underlying

mechanisms of the PCM-based multilevel phase-change memory.

6. A concept for a PC/Re-RAM memory device was developed that was based on

the combination of both phase-change mechanisms and the electrolytic properties of

PCMs. The memory cell exhibited excellent scalability (< 15 nm), low

RESET/SET current (several nA), small RESET/SET operation voltages (0.5 V),

long retention (> 180 days), excellent cycling endurance (> 1000 cycles) and high

On/Off resistance ratios (~5×103), as well as a high estimated data density (0.7

Tbit/inch2). Bipolar electrical switching of PCMs at the nanoscale was performed

by using C-AFM techniques for write, read, and erase operations. The physical

switching mechanism of PC/Re-RAM was validated by the Poole-Frenkel

conduction model. The polarity-dependent resistance switching processes could be

visualized by simultaneous topography and electric current images. The local

microstructure on the nanoscale of such memory cells and the corresponding local

chemical composition were correlated. The formation of cubic GST nanocrystals

within a cylindrical region with a diameter of ~10 nm across the film was observed,

as well as a slight deviation from stoichiometric GST (225) to a Te-rich and Ge-

depleted composition. From these results, the switching mechanism was proposed

to consist of a combination of the usual phase changes in GST films and the

migration of mobile Ge/Te ions. The demonstrated creation of this type of PC/Re-

RAM memory does not only open up new opportunities for exploring physical

Chapter 5 Summary

148

mechanisms but also proves PCMs to be promising candidates for next generation

non-volatile memory applications.

149

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163

List of Abbreviations

ADF

AFM

BF

C-AFM

CW

DRAM

DSC

EDX

FeRAM

FFT

FIB

FWHM

GeTeO

GID

GST

HAADF

HRS

HRTEM

HRSTEM

I-V

LAADF

LRS

MAADF

MRAM

NBD

PCM

PCRAM

PLD

ReRAM

SEM

SIMS

STEM

TEM

XRD

XRR

annular dark-field

atomic force microscopy

bright-field

conductive atomic force microscopy

continuous-wave

dynamic random access memory

differential scanning calorimetry

energy-dispersive X-ray spectroscopy

ferroelectric random access memory

fast Fourier transformation

focused ion beam

full width at half maximum

Ge-Te-O

grazing incidence diffraction

Ge2Sb2Te5

high-angle annular dark-field

high resistivity state

high-resolution transmission electron microscopy

high-resolution scanning transmission electron microscopy

current-voltage

low-angle annular dark-field

low resistivity state

middle-angle annular dark-field

magnetoresistive random access memory

nano beam diffraction

phase-change material

phase-change random access memory

pulsed laser deposition

resistive random access memory

scanning electron microscopy

secondary ion mass spectrometry

scanning transmission electron microscopy

transmission electron microscopy

X-ray diffraction

X-ray reflectivity

165

Acknowledgements

This dissertation would not have been possible without our team which contributed

to and gave valuable assistance during the period of this study. I would like to

acknowledge all of them. In the case I forgot someone to mention, I do apologize

for sincerely.

First and foremost, my utmost gratitude goes to my supervisor, Prof. Dr. Dr. h.c.

Bernd Rauschenbach, for having offered me the opportunity to work on this

interesting and challenging topic at the IOM and for his consistent support and

guidance throughout my Ph.D. study. Prof. Rauschenbach has been leading me to

the right direction to go when I faced various obstacles during the completion of

this research work. I really appreciate him propelling me in my research efforts,

being engaged in the deep physics regarding the experimental findings and

conclusions, as well as encouraging me to think creative and critical, which I am

benefitting from throughout my academic career, significantly. In addition, I would

also like to acknowledge my second supervisor, Prof. Dr. Stefan G. Mayr, for his

general support and valuable advices on my research.

I especially thank Dr. Jürgen W. Gerlach, for his friendly and patient offer of all

kinds of help with his broad knowledge, from teaching me how to use XRD to

assistance in vacuum technology, from results discussion to all manuscripts

corrections as well as the proof-reading of the present thesis. Moreover, his rigorous

style of scientific work has certainly influenced me.

I would like to thank Dr. Erik Thelander, as a senior member of the PLD-group, for

training me how to utilize the PLD technique, for discussing experiments and

results regularly, for having introduced me to the surroundings and for valuable

suggestions, which I benefitted from to quickly adapt to the working environment.

The other PLD-group members, Isom Hilmi and Philipp Schumacher, are

acknowledged as well.

I would like to thank Dr. Andriy Lotnyk and Ulrich Ross for performing the STEM

characterization, for the subsequent valuable discussions during the data analysis

and for proof-reading the manuscript. I also thank Agnes Mill for TEM lamellae

preparation.

Dr. Martin Ehrhardt I would like to thank for sharing his extensive knowledge on

laser applications with me and for always being willing to help me. I would also

like to thank Dr. Pierre Lorenz for performing COMSOL simulations and for

https://www.google.de/search?q=TEM-Lamella&spell=1&sa=X&ved=0ahUKEwiv-fjkja_OAhUKXRQKHVmsADQQvwUIGygA

166

valuable discussions. I thank Dr. Joachim Zajadacz for performing E-beam

lithography, Eva Salamatin for performing optical interference microscopy, Dr.

Klaus Zimmer for comments on laser experiments. Dr. Tomi Smausz (University of

Szeged, Hungary) is acknowledged for help on the fs laser system.

I would also like to acknowledge the other collaborators, though some results have

not been presented in the final thesis: Prof. Dr. Bernd Abel and Dr. Tilmann Häupl

for performing the Raman spectroscopy measurements, Prof. Dr. Marius

Grundmann and Dr. Holger von Wenckstern (University of Leipzig) for performing

the impedance spectroscopy measurements, Dr. Carsten Bundesmann for

conducting the ellipsometry measurements, Dr. Heidemarie Schmidt and Tiangui

You (Technische Universität Chemnitz) for initially trying C-AFM measurements,

Dr. Wolfgang Knolle for trying electron beam irradiation of several samples.

I am also grateful to many more colleagues at the IOM for either contributing to this

work or for providing me with many useful suggestions: Dr. Ulrich Decker and

Nadja Schönherr for their help regarding UV-Vis spectroscopy; Ingrid Reinhardt for

performing DSC measurements; Dr. Frank Frost, Dr. Jan Lorbeer, Marina

Sarmanova and Dr. Ariyan Arabi-Hashemi for tips on AFM measurements. Dietmar

Hirsch for conducting SIMS analyses and teaching me how to operate EDX/EDS;

Petra Hertel for performing the magnetron sputtering, providing the sufficient

sample boxes and helping me in the clean-room; Toni Liebeskind for supplying me

with wafer material; Marco Müller for tips on electronic circuit design.

I feel very grateful for my former and current officemates: Dr. Annemarie Finzel,

Susann Liedtke, Michael Mensing, who warmly offered all kinds of help, and in

particular Christoph Grüner for his patient and consistent help, as well as his

assistance on experimenting with the annealing system and the resistance

measurements.

I also want to acknowledge all members of the mechanical workshop and of the

administration, as well as all other PhD candidates and staff at the IOM for

providing me with a convenient work atmosphere from their lots of times being

helpful. I gratefully acknowledge the “BuildMona” graduate school for providing

various training activities, which benefitted me with the concept of an

interdisciplinary approach in conducting research.

Last, but not least, I want to deeply express my gratitude to my family members and

all my friends. They gave me continuous support and encouragement during the

time I was pursuing my Ph. D. degree.

https://woi.chemie.uni-leipzig.de/mitarbeiter/prof-dr-bernd-abel/

http://research.uni-leipzig.de/hlp/staff/

http://research.uni-leipzig.de/hlp/staff/

167

Curriculum Vitae

Person

Name: Xinxing Sun

Date of Birth: 23. December 1986

Place of Birth: Zhejiang, China

Nationality: China

Education

2013-2017 Ph.D. in Physics, Leibniz Institute of Surface Modification and

Leipzig University, Germany

Doctoral candidate and member of the graduate school

“BuildMoNa”

Title of the thesis: Phase Transformations and Switching

of Chalcogenide Phase-change Material Films Prepared

by Pulsed Laser Deposition

Supervisor: Prof. Dr. Dr. h.c. Bernd Rauschenbach

2009-2012 Master’s degree (M.Sc.) in nanophysics, East China Normal

University, China

Title of the thesis: Scattering and Antireflection Coatings

for Opto-electronic Device Application

Supervisor: Prof. Dr. Zhuo Sun

2004-2008 Undergraduate diploma in Electronic Engineering, Hangzhou

Dianzi University and Self-taught Higher Education Examinations

of Zhejiang Province, China

Work Experience

2007-2008 Engineer, State Key Lab of Silicon Materials and Department of

Materials, Zhejiang University, China

Working project: Transmission Electron Microscopy and

its Applications

Supervisor: Prof. Dr. Xiaobin Zhang and Youwen Wang

169

List of publications

1. X. Sun, U. Ross, J. W. Gerlach, A. Lotnyk and B. Rauschenbach, Nanoscale

bipolar electrical switching of Ge2Sb2Te5 phase-change material thin films,

submitted to Nature nanotechnology.

2. X. Sun, A. Lotnyk, M. Ehrhardt, J.W. Gerlach and B. Rauschenbach,

Realization of multi-level states in phase-change thin films by fast laser

pulse irradiation, Adv. Optical Mater. 1700169 (2017). DOI:

10.1002/adom.201700169.

3. X. Sun, M. Ehrhardt, A. Lotnyk, E. Thelander, J.W. Gerlach, T. Smausz, U.

Decker and B. Rauschenbach, Crystallization of Ge2Sb2Te5 thin films by

nano- and femtosecond single laser pulse irradiation, Sci. Rep. 6, 28246

(2016).

4. A. Lotnyk, S. Bernütz, X. Sun, U. Ross, M. Ehrhardt and B. Rauschenbach,

Real-space imaging of atomic arrangement and vacancy layers ordering in

laser crystallized Ge2Sb2Te5 phase change thin films, Acta Mater. 105, 1

(2016).

5. X. Sun, E. Thelander, J.W. Gerlach, U. Decker, B. Rauschenbach,

Crystallization kinetics of GeTe phase-change thin films grown by pulsed

laser deposition, J. Phys. D: Appl. Phys. 48, 295304 (2015).

6. X. Sun, E. Thelander, P. Lorenz, J.W. Gerlach, U. Decker and B.

Rauschenbach, Nanosecond laser-induced phase transitions in pulsed laser

deposition-deposited GeTe films, J. Appl. Phys. 116, 133501 (2014).

7. X. Sun, X. Chen, Z. Zhang and Z. Sun, Plasmon based antireflection

coatings containing nanostructured Ag and silica medium, Appl. Surf. Sci.

258, 3785 (2012).

8. L. Zhou, X. Chen, F. Zhu, X. Sun and Z. Sun, Improving temperature-stable

AZO-Ag-AZO multilayer transparent electrodes using thin Al layer

modification, J. Phys. D: Appl. Phys. 45, 505103 (2012).

9. B. Xue, R. Liu, Z. Xu and X. Sun, Synthesis, characterization and

photocatalytic activity of hierarchical spherical TiO2 nanostructures, Chinese

J. Inorg. Chem. 25, 1 (2009).

10. X. Sun, Z. Sun, L. Pan, M. Zhang, X. Chen, P. Guo. A measuring device for

thermal conductivity of solid materials, China Invention Patent, Publication

number: CN101949873A.

170

Conference contributions (talks and posters):

1. X. Sun, A. Lotnyk, M. Ehrhardt, P. Lorenz, J.W. Gerlach, B. Rauschenbach,

Accurately controllable phase transitions in PLD-deposited GeTe films for

multi-level memory application, talk, DPG spring meeting, Dresden (2017).

2. X. Sun, M. Ehrhardt, A. Lotnyk, E. Thelander, J.W. Gerlach, T. Smausz, U.

Decker, B. Rauschenbach, Phase transition on crystallization of PLD-

deposited Ge2Sb2Te5 films induced by nano- and femtosecond single laser

pulse irradiation, poster, DPG spring meeting, Regensburg (2016).

3. X. Sun, E. Thelander, J.W. Gerlach, B. Rauschenbach, Nanosecond laser-

induced phase transitions in PLD-deposited GeTe films, talk, DPG spring

meeting, Berlin (2015).

4. X. Sun, E. Thelander, J.W. Gerlach, U. Decker, B. Rauschenbach, Pulsed

laser deposited GeTe phase-change films on Silicon, poster, European

Phase Change and Ovonics Symposium, Amsterdam (2015).

5. X. Sun, E. Thelander, J. W. Gerlach, U. Decker, D. Hirsch, B.

Rauschenbach, Effect of O-incorporation on structural transitions of PLD-

deposited GeTe films, poster, Annual BuildMoNa Conference, Leipzig

(2015).

6. X. Sun, E. Thelander, H. Lu, J.W. Gerlach, B. Rauschenbach, Laser-induced

phase transitions of PLD-deposited GeTe films, poster, DPG spring

meeting, Dresden (2014).

7. X. Sun, E. Thelander, H. Lu, J.W. Gerlach, B. Rauschenbach, Ultraviolet

laser-induced phase transitions of PLD-deposited GeTe films, talk, Annual

Conference of the Graduate School BuildMoNa, Leipzig (2014).

8. X. Sun, E. Thelander, H. Lu, J.W. Gerlach, U. Decker, D. Hirsch, B.

Rauschenbach, Structural investigation of pulsed laser deposited phase-

change films, poster, 18th International Summer School on Vacuum,

Electron and Ion Technologies, Sozopol (2013).

171

Selbstständigkeitserklärung

Hiermit versichere ich, dass die vorliegende Arbeit ohne unzulässige Hilfe und ohne

Benutzung anderer als der angegebenen Hilfsmittel angefertigt, und dass die aus

fremden Quellen direkt oder indirekt übernommenen Gedanken in der Arbeit als

solche kenntlich gemacht wurden.

Ich versichere, dass alle Personen, von denen ich bei der Auswahl und Auswertung

des Materials sowie bei der Herstellung des Manuskripts Unterstützungsleistungen

erhalten habe, in der Danksagung der vorliegenden Arbeit aufgeführt sind.

Ich versichere, dass außer den in der Danksagung genannten, weitere Personen bei

der geistigen Herstellung der vorliegenden Arbeit nicht beteiligt waren, und

insbesondere von mir oder in meinem Auftrag weder unmittelbar noch mittelbar

geldwerte Leistungen für Arbeiten erhalten haben, die im Zusammenhang mit dem

Inhalt der vorliegenden Dissertation stehen. Außerdem versichere ich, keinen

Promotionsberater in Anspruch genommen zu haben.

Ich versichere weiterhin, dass die vorliegende Arbeit weder im Inland noch im

Ausland in gleicher oder in ähnlicher Form einer anderen Prüfungsbehörde zum

Zwecke einer Promotion oder eines anderen Prüfungsverfahrens vorgelegt und in

ihrer Gesamtheit noch nicht veröffentlicht wurde.

Ich versichere außerdem, dass keine früheren erfolglosen Promotionsversuche

stattgefunden haben.

Leipzig, 12. May 2017

Xinxing Sun

Documents

Phase Transformations and Switching of Chalcogenide Phase-change Material Films Prepared by