Influence of Partial Substitution of Te by Se and Ge by Sn on the

Article pubs.acs.org/cm

Influence of Partial Substitution of Te by Se and Ge by Sn on the Properties of the Blu-ray Phase-Change Material Ge8Sb2Te11 Saskia Buller,† Christine Koch,† Wolfgang Bensch,*,† Peter Zalden,‡ Roland Sittner,‡ Stephan Kremers,‡ Matthias Wuttig,‡ Ulrich Schürmann,§ Lorenz Kienle,§ Thomas Leichtweiß,∇ Jürgen Janek,∇ and Boyke Schönborn⊥ †

Institute of Inorganic Chemistry, Christian-Albrechts-University, Max-Eyth Str. 2, 24118 Kiel, Germany I. Institute of Physics, RWTH Aachen, Sommerfeldstr. 14, 52056 Aachen, Germany § Institute for Materials Science, Synthesis and Real Structure, Christian-Albrechts-University, Kaiserstr. 2, 24143 Kiel, Germany ∇ Institute of Physical Chemistry, Justus-Liebig-University, Heinrich-Buff-Ring 58, 35392 Gießen, Germany ⊥ Institute of Physical Chemistry, Christian Albrechts University, Max-Eyth Str. 2, 24118 Kiel, Germany ‡

ABSTRACT: The opto-electronic properties of the Blu-ray phase-change material Ge8Sb2Te11 were investigated and compared to partially substituted compounds. Investigations were performed on Ge8Sb2Te6Se5 (I) and SnGe7Sb2Te7Se4 (II). To monitor the influence of substitution on both cation and anion sites onto the properties, in Ge8Sb2Te6Se5, only Te was substituted by Se, while in SnGe7Sb2Te7Se4, Ge also was partially replaced by Sn, yielding an equivalent degree of substitution. The crystallization of the amorphous compound to a rhombohedral phase was observed. The first transition temperature (Tc1) is 210 °C for I and 185 °C for II, while the second transition temperature (Tc2) cannot be clearly determined, because of the coexistence of both phases. Thin-film properties were examined and an increase of density and roughness was observed for both materials upon crystallization. In situ electrical resistance measurements show a huge electrical contrast between the amorphous samples and the crystalline samples, which amounts to 4 orders of magnitude for II and 5 orders of magnitude for I. Optical properties were characterized with variable incident-angle spectroscopic ellipsometry (VASE) and Fourier Transform infrared (FTIR) spectroscopy. Substitution leads to an enormous increase of the reflectivity contrast for blue light: The reflectivity contrast between the different phases increases from 20% to 50% for I and from 20% to 74% for II. KEYWORDS: phase-change materials, blu-ray, optical contrast, electrical conductivity, structure, thin films

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INTRODUCTION Data storage is one of the most important technologies for information-based societies. Nowadays, more and more information must be stored faster on even smaller devices. Commercially used optical storage media (CD, DVD, Blu-ray) and electrical storage media (PC-RAM) fulfill these demands in a suitable way.1−4 Requirements for commercial application are high reflectivity contrast, cyclability, long-term stability, and fast crystallization speed.5 Materials that are used in these applications are so-called phase-change materials (PCM), based, e.g., on pseudo-binary compounds of (GeTe)x(Sb2Te3)1−x or on Sb2Te:AgIn (AIST).6 These materials can be reversibly switched between an amorphous and a crystalline phase. The different states exhibit unique physical properties, characterized by pronounced optical and electrical changes used as the logical values of a memory device. Switching and reading is performed by a variable (power, duration) laser pulse for optical devices while variable-current pulses are performed for switching and reading in PC-RAM. Amorphous phases are characterized by low reflectivity and © 2012 American Chemical Society

high resistivity while crystalline phases show contrary properties.7 To understand the phase transition upon the reversible switching between the amorphous and the crystalline phase, the origin of the optical contrast8,9 and the electronic and atomic local structures of the different phases10−15 were intensively studied, applying a large variety of analytical tools.16,17 The crystalline structure of PCM devices based on Ge−Sb− Te phases is, in most cases, the NaCl-type structure, where Te atoms are located on the anion site while Ge, Sb, and vacancies occupy the cation site. The amount of vacancies is fixed by the ratio of GeTe to Sb2Te3 and decreases with increasing GeTe content. The structural details of the amorphous phase of PCMs are still not well understood, despite intense experimental and theoretical research. Extended X-ray absorption fine structure (EXAFS) analysis is a powerful analytical method to study crystalline as well as amorphous Received: June 11, 2012 Revised: August 16, 2012 Published: August 20, 2012 3582

dx.doi.org/10.1021/cm301809g | Chem. Mater. 2012, 24, 3582−3590

Chemistry of Materials

Article

Figure 1. In situ XRD patterns of (a) Ge8Sb2Te11, (b) SnGe7Sb2Te7Se4, and (c) Ge8Sb2Te6Se5. The transition from amorphous to crystalline phase at the different temperatures is clearly seen by the occurrence of diffraction peaks. (Asterisk (*) denotes the substrate.)

chemical bonds. The influence on the optical/electrical behavior and the bonding characteristics, which are crucial for the properties for storage media, are reported here.

samples, yielding information about bond lengths, coordination numbers, and neighbor atoms of the absorber atom. Several EXAFS studies were undertaken on PCMs, but the interpretations are, in part, contradictory, and, until now, a comprehensive picture of the real structural and bonding situation in the amorphous state could not be provided.10,12,15,18−20 In addition, theoretical studies were performed to identify the real structures and to understand the phase transitions.20−22 Very recently, a combination of experimental X-ray scattering and theoretical data led to the assumption that GeTe forms a “core” network with ring formation in Ge2Sb2Te5, while Sb is involved in a network with Te atoms. The GeTe network stabilizes the amorphous phase and also persist in the crystalline phase, which enables the fast crystallization and stability on the other hand.23 The pseudo-binary compound Ge2Sb2Te5 is one of the best established and most investigated PCMs.19,24−30 This material combines all requirements for commercial application in CD or DVD devices. Further improvement of the properties was achieved by doping with other metals31−33 or with oxygen and/ or nitrogen.34−37 The substitution resulted in a severe alteration of the properties such as crystallization rate, archival lifetime, and optical and electrical contrast.33,38,39 The substitution of elements in a structure is always accompanied by a change in bonding characteristics. Recently, it was demonstrated that resonant bonding is essential for the optical contrast in PCMs.8,9,16,40 The knowledge about resonant bonding and substitution effects may pave the way to design PCMs with an excellent compromise of the required properties, i.e., fast crystallization time, reliable read/write process, long-term stability. One way to increase the storage density is the use of laser light with shorter wavelengths. One candidate for commercial application with blue laser light is the material Ge8Sb2Te11 for so-called Blu-ray discs. The compound crystallizes as lowtemperature modification (rhombohedral symmetry) and as a high-temperature modification (rock salt structure type).41,42 The rhombohedral structure can be explained as a distorted rock salt structure in which all Ge atoms are displaced equally along the ⟨111⟩ direction.43 In the crystalline modifications, ∼9% of the cationic sites are empty (vacancies). In the present study, the influence of cation/anion substitution by Sn and Se, respectively, in Ge 8Sb 2 Te 11 (Ge8Sb2Te6Se5 (I) and SnGe7Sb2Te7Se4 (II)) on the properties were investigated. Selenium (Se) has a smaller atomic/ionic radius, less metallic character, and higher electronegativity than tellurium (Te), while tin (Sn) has a larger atomic/ionic radius, is more metallic, and is less electronegative than germanium (Ge). Both substitutions increase the ionic character of the

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EXPERIMENTAL PROCEDURES

Thin film samples of ∼50 nm and 1 μm thickness were deposited on silicon wafers or glass substrates by DC-magnetron sputtering, using stoichiometric targets. The base pressure in the deposition chamber was ∼2 × 10−6 mbar. Sputtering was performed using an Ar pressure of ∼5 × 10−3 mbar. Ge8Sb2Te6Se5 was deposited by cosputtering targets of Ge3Sb2Te6 (Umicore) with diameters of 10 cm at a power of 10 W and GeSe (Umicore) at a power of 8 W. SnGe7Sb2Te7Se4 was also deposited by cosputtering targets of SnSb2Se4 (Umicore) with diameters of 10 cm (power = 5 W) and GeTe (Evochem) at 20 W. Unfortunately, film samples sputtered with a SnSb2Te4 target are crystalline in the as-deposited state. Therefore, samples of Sn8Sb2Te11 were not prepared, although the analysis of the sole influence of Ge substitution by Sn would be very interesting. The compositions of the films were determined by energy-dispersive X-ray spectrometry (EDX) and wavelength-dispersive X-ray spectrometry (WDX), yielding compositions very close to the desired stoichiometry. In both cases, an acceleration voltage of 15 kV was used to measure the L-series of the elements in the 1-μm-thick samples. Penetration depths of the electron beam were calculated using CASINO44 Monte Carlo simulation and determined to be preponderate in the layer and not in the substrate. Optical properties have been determined using variable incident-angle spectroscopic ellipsometry (VASE, J. A. Woolam Co.) using energies from 0.7 eV to 5.2 eV. Furthermore, Fourier transform infrared (FTIR) spectra were measured from 0.05 eV to 1 eV. Investigations were performed on amorphous and crystallized samples. Samples for optical measurements were deposited on Ag-coated glass substrates. The structural phase transitions were determined by in situ X-ray diffraction (XRD) measurements on an X’Pert Pro PANalytical diffractometer equipped with an Anton Paar HTK 1200N high-temperature chamber, a Göbel mirror, and a PIXcel detector using Cu radiation. The patterns were collected within one hour at constant temperature on 1-μm-thick samples on Si wafers. For temperature-resolved experiments, the temperature was gradually increased, at a heating rate of 5 K/min, from room temperature up to 360 °C in steps of 10 or 5 K, respectively. X-ray reflectivity (XRR) measurements were performed on the same diffractometer, using an incident angle of 0.5° and receiving slit mode. To prevent oxidation, a protective He atmosphere was used. Temperature-dependent sheet resistance measurements were performed on 1-μm-thick films on glass substrates. A four-point probe setup was employed following the procedure proposed by Smits.45 The values of the sheet resistance (Rs) are dependent on the setup and are corrected by the factors ascertained from ref 46. Therefore, four 200-nm thin gold contacts were sputtered onto 3-nm thin chromium contacts on the film samples with a distance of 1 mm. The samples were heated continuously from room temperature to 360 °C under a protective N2 atmosphere, using a heating rate of 5 K/min. The temperature of 360 °C was held constant for 30 min. The sheet resistance (Rs) can be converted, in principle, to a resistivity (R) if the thickness of the film is known, but 3583



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opposite trend was reported for GeSb2Te4.38 The influence of Sn on Tc1 seems to be less pronounced. The bond energies of heteronuclear bonds can be calculated by using the homopolar bond energies and the electronegativities of the elements.48,49 According to this approach, SnTe has a lower bonding energy and GeSe has a higher bonding energy than GeTe. This trend of bonding energies correlates to the trend of transition temperatures. Ge8Sb2Te6Se5, where parts of GeTe were substituted by GeSe shows transition temperatures at higher values. SnGe7Sb2Te7Se4 in which GeTe is partially substituted by GeSe and SnTe, the transition temperatures are shifted to higher values (because of the stronger GeSe bonds) and, on the other hand shifted, to lower transition temperatures (because of the softer bond energies of SnTe). Because of the higher amount of GeSe bonding interactions than SnTe bonds in SnGe7Sb2Te7Se4, the resulting transition temperature is shifted to higher values. The transition temperatures of the substituted samples should be comparable or even identical if the influence of Se or Sn substitution would be the same, which is not the case, considering that the number of atoms substituted in the samples is identical. The transition temperatures from the rhombohedral to the cubic phase cannot exactly be determined because the transition seems to be a slow process where the two phases coexist. For SnGe7Sb2Te7Se4, the ⟨110⟩ (rhombohedral) and ⟨220⟩ (cubic) diffraction peaks were analyzed because the other reflections almost do not change their position. However, the transition from the low-temperature phase to the hightemperature phase in Ge8Sb2Te11 is clearly seen in Figure 2a. The experiments yield crystallization temperatures of Tc1 ≈ 185 °C for SnGe7Sb2Te7Se4 and ∼210 °C for Ge8Sb2Te6Se5. For pure Ge8Sb2Te11, a transition temperature of 165 °C was determined. However, for the Se-substituted sample a more distinct transition to the cubic phase, compared to the pristine material and the Se/Sn substituted sample, was observed. The results reported in ref 43 indicate that the transition from the rhombohedral to the rock-salt-like structure is an order− disorder transition, rather than a displacive phase transition. The sample investigated in this study was first heated at ∼200 °C in a furnace and X-ray scattering experiments were done between −185 and 209 °C. A ferroelectric-to-paraelectric transition occurs at ∼112 °C. Matsunaga et al.41 investigated the phase transition with in situ XRD, using a different setup and starting also with crystalline Ge8Sb2Te11. The material was placed in a quartz capillary and was heated with a nitrogen stream at distinct temperatures, starting at ∼92 K and the heating rate from measurement to measurement was 10 K min−1. The results indicate a sudden phase transition from the low-temperature phase to the high-temperature phase at ∼112 °C and a first-order transition was suggested. In the first step of such a first-order phase transition, a free-energy barrier must be overcome, which is the work of formation of a small embryo or nucleus of the new phase. The nucleus can only emerge from random thermal fluctuations within the original phase. In the present investigation, the films were amorphous in the asdeposited state, and they were heated stepwise, leading to different phase transition behavior. Rietveld refinements of the patterns of the crystalline samples of SnGe7Sb2Te7Se4 (360 °C) and Ge8Sb2Te6Se5 (360 °C) exhibit a slight decrease of the lattice parameters, compared to the literature values of Ge8Sb2Te1141 (Tables 1 and 2, Figure 3). This observation can be explained based on the atomic radii,

during the structural phase transition the thickness changes and was not determined simultaneously. Atomic force microscopy (AFM) images (Park Scientific Instruments), using a 5-μm scanner in contact mode, were recorded to monitor the differences in surface topography. High-resolution transmission electron microscopy (HRTEM) and selected-area electron diffraction (SAED) measurements were performed in a Tecnai F30 STwin microscope (300 kV, field emission gun (FEG) cathode, spherical aberration coefficient Cs = 1.2 mm) on 30-nm thin film samples deposited on nickel grids. The samples were continuously heated with a double-tilt heating sample holder (Gatan) in the electron microscope with 5 K/min until the transition occurred, then the temperature was held constant for 30 min. Ge8Sb2Te11 was not additionally prepared, but data or samples for comparison were either provided by the coauthors or were compiled from the literature.

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RESULTS AND DISCUSSION In Situ X-ray Diffraction (XRD). In situ X-ray diffraction (XRD) patterns of the 1-μm-thick samples on Si wafers are displayed in Figure 1. The as-deposited films are in the amorphous state and, upon heating, crystallize first in a rhombohedral (R3m) and then in a cubic phase (Fm3m ̅ ). The transition from the amorphous phase to the rhombohedral phase is clearly seen, but the transition from the rhombohedral phase to the cubic phase is less clear. For a more accurate determination of the latter transition temperatures for Ge8Sb2Te6Se5 the intensities of the ⟨012⟩ diffraction peak of the rhombohedral phase and the ⟨200⟩ reflection of the cubic phase were monitored as a function of temperature (Figure 2b). Hence, increasing the selenium content leads to an increase in the transition temperature. The same trend was observed for samples of the phase-change system Ge2Sb2Te5,47 while the

Figure 2. (a) Temperature-induced phase transitions in Ge8Sb2Te11 observed from the diffraction peaks of the low-temperature rhombohedral phase (⟨012⟩ reflection) and the high-temperature cubic phase (⟨200⟩ reflection). The dashed line corresponds to 165 °C where crystallization begins. (b) Comparison of the intensities of the rhombohedral and cubic diffraction peaks at the different temperatures for the prepared samples. [Data for Ge8Sb2Te11 sample on Si, from P. Zalden, Aachen, Germany.] 3584



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measurements were performed on glass substrates with a constant heating rate while XRD investigations were done on Si substrates heating the samples stepwise. The electrical contrast is ∼4 orders of magnitude higher for SnGe7Sb2Te7Se4 and Ge8Sb2Te11, and is even higher for Ge8Sb2Te6Se5 (∼5 orders of magnitude). The transition temperatures are shifted to higher values with increasing Se content and the resistances of Ge8Sb2Te11 and SnGe7Sb2Te7Se4 of the as-deposited films are more or less identical, while the resistance of Ge8Sb2Te6Se5 is larger, by ∼1 order of magnitude. The behavior after annealing of Ge8Sb2Te11 cannot be directly compared to the substituted samples because different annealing temperatures were used. Investigations of Se substituted GeSb2Te4 showed no influence on the electrical contrast.38 Upon cooling from 360 °C to room temperature, the samples show semimetallic behavior, which is observable by the constant increase of the sheet resistance from ∼50 Ωsq to ∼140 Ωsq for SnGe7Sb2Te7Se4 and ∼170 Ωsq for Ge8Sb2Te6Se5. Optical Properties. As reported in refs 8, 16, and 40, resonant bonding in the crystalline phase was identified as the origin of the large optical contrast in PCMs. This type of bonding is influenced by the lattice structure and its distortion. In PCMs based on (GeTe)x(Sb2Te3)1−x, three p-electrons per site are available for covalent bonds. However, the octahedral coordination in the crystalline phase requires six covalent bonds, which are not saturated by the number of p-electrons. The lack is compensated by resonant bonding, as depicted in Figure 7. The larger the distortion, the less pronounced the resonant bonding in the structure. Density functional theory calculations showed the influence of distortion on the properties of the GeTe prototype system. The displacement of Te atoms along the ⟨111⟩ direction and a distortion of the unit cell along the cubic ⟨111⟩ direction lead to the formation of a rhombohedral structure such as that in the Ge8Sb2Te11 system. It was reported that the distortion yields a gain of energy of the structure and a reduction of the resonant bonding, but the resonant bonding characteristics prevail and the optical dielectric tensor (ε∞) remains large. The dielectric constant ε∞ describes the degree of polarizability of the electronic system and is an additional fingerprint characteristic of PCMs. A small degree of hybridization of orbitals and low ionicity are found for all known PCMs, and some selected examples are shown in Figure 8. Details of the calculation procedure can be found in refs 40 and 52. Substitution of Te by Se and Ge by Sn in Ge8Sb2Te11 has a distinct influence on the bonding characteristics with an increase of the degree of ionicity and increased hybridization. Hybridization is more pronounced in Ge8Sb2Te6Se5 than in SnGe7Sb2Te7Se4, while the degree of ionicity is altered in a more or less identical way. FTIR spectroscopy has been identified as a powerful tool to identify the presence of resonant bonding and to quantify its

Table 1. Atomic Positions Obtained by the Rietveld Refinement for Ge8Sb2Te11, Ge8Sb2Te6Se5, and SnGe7Sb2Te7Se4 at Room Temperature cubic, Fm3̅m x

y

Ge, Sb Te

1/2 0

1/2 0

Sn, Ge, Sb Te, Se

1/2 0

1/2 0

Ge, Sb Te, Se

1/2 0

1/2 0

rhombohedral, R3m z

x

y

Ge8Sb2Te11 1/2 0 0 0 0 0 SnGe7Sb2Te7Se4 1/2 0 0 0 0 0 Ge8Sb2Te6Se5 1/2 0 0 0 0 0

z

R_Bragg (%)

0.544(2) 0

1.83

0.526(1) 0

3.62

0.5394(8) 0

2.53

because Se is smaller than Te. Rietveld refinement was only possible assuming the coexistence of a rhombohedral and a cubic phase, which is in accordance with the slow transitions observed in the temperature-dependent XRD measurements. High-Resolution Transmission Electron Microscopy (HRTEM). The HRTEM experiments support the results obtained by temperature-dependent XRD measurements. At low temperatures, the materials are in the amorphous state and, in the selected area diffraction patterns (SAED), no diffraction peaks are observed. During annealing, the sudden appearance of diffraction peaks in the SAED indicate the transition from the amorphous phase to the rhombohedral phase. While the temperature was held constant, the grains of the crystalline phases grow with increasing time, as depicted in the sequence of SAED patterns in Figure 4. The diameters of the diffraction rings correlate well with the expected rhombohedral structure. As known for many phase-change materials, Ge8Sb2Te11 also shows the typical vacancy layers (VLs).50,51 The VL or van der Waals gaps are developed by internal arrangements of the atoms resulting in well-defined layers that are not occupied by atoms. The VLs are observed in the high-resolution micrographs as lines with bright spots (Figure 5), representing the space between neighboring Te layers. Temperature-Dependent Sheet Resistance Measurements. Temperature-dependent sheet resistance measurements are shown in Figure 6. The sheet resistance data (Rs) for the amorphous films at room temperature are ∼0.21 GΩsq for SnGe7Sb2Te7Se4 and Ge8Sb2Te11 and ∼1.49 GΩsq for Ge8Sb2Te6Se5. These values slightly decrease with increasing temperature until a sudden drop is observed at ∼185 °C for Ge8Sb2Te11, 204 °C for SnGe7Sb2Te7Se4 and 220 °C for Ge8Sb2Te6Se5. The somewhat-larger values for Tc1, compared to those obtained by the in situ XRD measurements, can be explained by the different experimental conditions. Resistance

Table 2. Results of the Lattice Parameter and the Ratio of the Two Phases of the Rietveld Refinement for Ge8Sb2Te11, Ge8Sb2Te6Se5, SnGe7Sb2Te7Se4 at Room Temperature, and Ge8Sb2Te11 (at −181 °C) cubic Ge8Sb2Te11 SnGe7Sb2Te7Se4 Ge8Sb2Te6Se5 Ge8Sb2Te11a a

rhombohedral

a (Å)

a (Å)

c (Å)

ratio, cubic/rhombohedral

5.950(3) 5.946(1) 5.9868(8) 6.05800

4.2099(6) 4.181(6) 4.1574(5) 4.20340

10.247(4) 10.476(4) 10.370(4) 10.4578

41/59 23/77 29/71

Data taken from ref 41. 3585



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Figure 3. Results of the Rietveld refinements performed on (a) Ge8Sb2Te11, (b) SnGe7Sb2Te7Se4, and (c) Ge8Sb2Te6Se5. Points rerpesent experimental data. Line represents calculated data. Vertical bars denote Bragg positions (top: rhombohedral phase, bottom: cubic phase). Bottom trace represents the difference between calculated and experimental data.

Figure 4. Sequence of temperature- and time-dependent SAED measurements of Ge8Sb2Te6Se5 (left, 120 °C, in situ; middle, 220 °C, in situ; right, 230 °C, in situ after 25 min).

Figure 5. High-resolution micrographs of Ge8Sb2Te11. Vacancy layers (VLs) are formed by internal arrangement of the atoms.

magnitude.8 The FTIR spectra were simultaneously analyzed in the energy range from 0.05 eV to 4.5 eV using the SCOUT software.53 To model the dielectric function of the material under investigation, ε̃(ω) is decomposed into the following parts: (a) a constant that accounts for the polarizability in the higher energy range; (b) a Tauc−Lorentz oscillator,54 which describes the optical properties of the interband transitions and additionally for crystalline samples; (c) a Drude-type contribution, which describes the free carriers. ε1(ω) and ε2(ω), which represent the real and imaginary part of the dielectric function ε̃(ω) = ε2 + iε2 were determined for amorphous and crystalline samples. The dielectric functions are

related with the index of refraction n and the extinction coefficient k as follows: ε1 = n2 − k2 and ε2 = 2nk. The dielectric functions (ε1 and ε2) are displayed in Figure 9, showing a single peak that is caused by bonding−antibonding transitions. The optical constant ε∞ was determined at energies close to zero from the dielectric function ε1 at ε1(0.05 eV) = ε∞. At low energies, the Drude-like absorption dominates the dielectric functions of crystalline samples (see Figure 9). To enable the comparison of the optical constants, the Drude term was subtracted. The typical behavior of PCMs is observed, where amorphous samples exhibit values of ∼11, while the crystalline samples show higher values of ∼25, depending on the degree of polarizability (for detailed values, see Table 3). 3586



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Figure 6. Sheet resistance measurements of film samples on glass substrates. The sheet resistance decreases significantly upon the transition from the amorphous to the crystalline phase. Substitutions lead to a pronounced increase of transition temperature. [Data for Ge8Sb2Te11 supplied by R. Sittner, Aachen, Germany.]

Figure 9. Dielectric functions ε1(ω) and ε2(ω) of the amorphous and crystalline samples. Crystalline samples were annealed under an Ar atmosphere at 250 °C (Ge8Sb2Te11), 267 °C (Ge8Sb2Te6Se5), and 208 °C (SnGe7Sb2Te7Se4). Dashed lines represent the functions without the Drude term. [Data of amorphous Ge8Sb2Te11 taken from M. Woda as reported in ref 55. Measurements of crystalline Ge8Sb2Te11 were made on a sample supplied by P. Zalden.]

Figure 7. Scheme of two distorted covalent bonding situations (left/ right) and the resonance bonding (middle) as known from, e.g., benzene.

Table 3. Dielectric Function ε∞ (Degree of Polarization) and Eg (Optical Bandgap) of the Amorphous (amo) and Crystalline (cry) Samples ε∞ Ge8Sb2Te11 Ge8Sb2Te6Se5 SnGe7Sb2Te7Se4

Eg

amo

cry

amo

cry

13 11 9

26 27 22

0.82 0.89 0.97

0.58 0.66 0.62

has more metallic character than Ge, and it seems that Sn influences the band gap of the amorphous phase in a more pronounced way. From the dielectric functions ε1 and ε2, the optical constants n and k can be determined (see Figure 10). At energies below the band gap, the extinction coefficient k becomes zero and ε1 is constant (ε1 = n2). With these optical constants, the reflectivity of the materials can be calculated. Multiple scattering occurs, as described by ref 56. The amplitude r of the reflected waves is evaluated using the Airy formula, with the Fresnel coefficients r01 and r12 for the air/layer and layer/ substrate boundaries, respectively. Taking the square of the absolute values of the obtained amplitudes yields the total reflected intensity R. Calculations were carried out with a film thickness of 1 μm. Optical constants of the Si substrate were fixed to n = 3.66 (ref 57) and k = 0.019. The reflected intensities of the amorphous samples are less than these of the crystalline samples (see Figure 10).

Figure 8. Calculated proportions of the hybridization and the degree of ionicity for some known PCMs. Substitution of Se and Sn in Ge8Sb2Te11 has the effect of slightly increasing the two bonding characteristics.

Comparison of the values demonstrates that the influence of substitution on ε∞ is larger for SnGe7Sb2Te7Se4 than for Ge8Sb2Te6Se5, compared to the pristine material. The optical band gap Eg was determined at the energy where the absorption α (not shown) exceeds 10 000 cm−1, and, for crystalline samples, after subtracting the Drude contribution. Compared to pure Ge8Sb2Te11, the values for Eg are larger, which can be explained by the less-metallic character of Se, compared to Te. For Ge8Sb2Te6Se5, a larger influence on the crystalline band gap is observed while, for SnGe7Sb2Te7Se4, the alteration is more pronounced for the amorphous phase (see Table 3). Sn 3587



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Figure 10. Index of refraction (n) and extinction coefficients (k) of the amorphous and crystalline samples. At energies below the band gaps, k becomes zero and n becomes a constant.

Figure 11. (a) Calculated reflectivity, depending on the wavelength of amorphous (filled symbols) and crystalline (empty symbols) samples; vertical lines represent laser light of 405 nm (Blu-ray) and 830 nm (DVD, CD). (b) Relative reflectivity (SnGe7Sb2Te7Se4 has the highest contrast and Ge8Sb2Te11 has the lowest contrast).

One important trait for data storage is a huge contrast between the two phases, to guarantee a proper readout of the data. Therefore, the ratio R cryst − R amo R amo

=

|rcryst| −

Table 4. Experimental Densities of the Substituted Samples, Compared to the Pristine Material in .

|ramo|

|ramo|

density (g/cm3)

of the reflected intensities of the crystalline Rcryst and the amorphous Ramo phases is also plotted in Figure 11. Ge8Sb2Te11 has a contrast of ∼20% (405 nm), while substitution of Te by Se leads to an enormous increase of the contrast to ∼50% (405 nm) for Ge8Sb2Te6Se5. The highest contrast of ∼74% (405 nm) is obtained for the SnGe7Sb2Te7Se4 sample. Such a high optical contrast for PCMs was, to the best of our knowledge, never determined previously. Calculations were also made for 80-nm thin samples on SiO2 substrates (n = 1.46), which is more realistic for commercial application. The same trends and high values for the optical contrast were obtained for the thin film. X-ray Reflectivity (XRR) Measurements. X-ray reflectivity (XRR) measurements of the thin-film samples were performed for the as-deposited and the crystallized samples. This method allows to determine the thicknesses and densities of the films by measuring the intensity of the reflected beam, depending on the angle.58,59 The results of the XRR analysis are listed in Table 4. During crystallization, the densities increase by ∼6(±3)%, which is a known phenomenon for PCMs.60,61 Atomic Force Microscopy (AFM). Contact-AFM images of the amorphous and crystallized samples are displayed in Figures 12a−d. The as-deposited films show no significant surface morphology and an average roughness of ∼10 and 20 Å for SnGe7Sb2Te7Se4 and Ge8Sb2Te6Se5, respectively. Crystallization of the amorphous films leads to formation of small

phase

Ge8Sb2Te11

I

II

amorphous crystalline

5.5 6.0

5.56 5.89

5.63 5.94

a

Estimated values of Ge8Sb2Te11 from ref 62, Ge8Sb2Te6Se5 (I), and SnGe7Sb2Te7Se4 (II).

crystallites, which results in larger surface roughness. The rootmean-square (rms) roughness increases to ∼18 and 21 Å. A similar trend of increasing roughness after annealing was observed in XRR measurements.

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CONCLUSION The partial substitution of Te by Se and Ge by Sn in the Bluray phase-change material (PCM) Ge8Sb2Te11 leads to pronounced changes of the structural, electrical, and optical properties. In the as-deposited state, the films are amorphous, as evidenced by X-ray diffraction (XRD) and transmission electron microscopy (TEM). In situ XRD experiments demonstrate that the transition temperature from the amorphous to the metastable state increases with the Se content, compared to the pristine material. In high-resolution TEM images, nanosized crystallites are seen and, in the image of a sample heated at 400 °C, the typical vacancy layers are observed. The results of in situ XRD suggest that the substitution of Te by Se has a more pronounced effect on 3588



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ACKNOWLEDGMENTS

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REFERENCES

Article

The authors would like to acknowledge the financial support of the State of Schleswig-Holstein and the Deutsche Forschungsgemeinschaft (DFG, BE 1653/18-1 and SFB 917 (Nanoswitches)).

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Figure 12. Contact AFM images of (a/b) amorphous/crystalline SnGe7Sb2Te7Se4, and (c/d) Ge8Sb2Te6Se5 films. The average roughness increases from the amorphous phase to the crystalline phase.

the transition temperature than the replacement of Ge by Sn. According to the concept of hybridization and ionicity of the chemical bonds, the two substituted samples exhibit a higher degree of ionicity and hybridization, compared to Ge8Sb2Te11. These changes are in accordance with the bonding energies of GeTe, GeSe, SnTe, and SnSe, and higher bonding energies correlate with the larger values for the crystallization temperatures. It is interesting to note that Se substitution in GeSb2Te438 does not lead to higher transition temperatures, while the same trend like in the title compounds was observed for Ge2Sb2Te547 and Ge3Sb2Te663 systems. Whether there is a relationship between the vacancy concentration on the cationic sublattice or the degree of substitution on the transition temperatures cannot be answered yet, because of the small data basis. The substitution in Ge8Sb2Te11 also significantly alters the optical band gap. For the amorphous state, a significant increase is found for SnGe7Sb2Te7Se4, while the difference from Ge8Sb2Te11 in the crystalline state is more pronounced for Ge8Sb2Te6Se5. A drastic change is also observed for the reflectivity of the materials in the amorphous and metastable crystalline states. At 405 nm, the maximal optical contrast amounts to ∼20% for Ge8Sb2Te11, while the Sn-substituted sample exhibits ∼74% contrast and Ge8Sb2Te6Se5 exhibits a contrast of ∼50%. Electrical sheet resistance measurements evidence an increased electrical contrast for Ge8Sb2Te6Se5 of ∼5 orders of magnitude, compared to the pristine material. Such a large contrast is beneficial to avoid read-out errors in PC-RAM devices. The influence of Sn on the transition temperatures and on the electrical properties is less pronounced than that of Se.

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Influence of Partial Substitution of Te by Se and Ge by Sn on the

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