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Insight into the Growth and Control of Single-Crystal Layers of Ge Sb Te Phase-Change Material Ferhat Katmis,† Raffaella Calarco,*,† Karthick Perumal,† Peter Rodenbach,† Alessandro Giussani,† Michael Hanke,† Andre Proessdorf,† Achim Trampert,† Frank Grosse,† Roman Shayduk,†,z Richard Campion,‡ Wolfgang Braun,†,§ and Henning Riechert† †
Departments of Epitaxy and Microstructure, Paul-Drude-Institut f€ur Festk€orperelektronik, Hausvogteiplatz 5-7, D-10117 Berlin, Germany ‡ School of Physics & Astronomy, University of Nottingham, University Park Nottingham, NG7 2RD, England ABSTRACT: Employed for a long time in optical disks, Ge2Sb2Te5 is nowadays considered the most promising material also for phase-change nonvolatile memories. In the current paper, Ge Sb Te phase-change material thin films with a nominal composition of Ge2Sb2Te5 were grown by molecular beam epitaxy on slightly mismatched GaSb and InAs substrates with (001) and (111) orientations. In situ quadrupole mass spectrometry and reflection high-energy electron diffraction allowed tight control of the growth parameters, revealing that Ge2Sb2Te5 grows in epitaxial fashion only within a narrow window of substrate temperatures around 200 °C. Smooth surfaces were achieved solely on (111)-oriented substrates. Rough surfaces and interfaces were observed by transmission electron microscopy for films grown on (001)-oriented substrates. Whereas films deposited on (001) substrates possess two different vertical epitaxial orientations, single-crystalline layers exclusively (111) oriented were achieved on (111) substrates, as shown by synchrotron radiation X-ray diffraction. All the results point to the superior crystalline quality and morphology of Ge2Sb2Te5 layers grown on (111) surfaces, independent of the substrate material.
1. INTRODUCTION Phase-change materials discovered by S. Ovshinsky1 did not immediately find a technological application. Later, with successful employment in the field of optical data storage2 they experienced a renaissance. Nowadays, they are being investigated by several semiconductor companies for nonvolatile memory applications.3 The most promising materials are the alloys along the GeTe Sb2Te3 pseudobinary line (GST), which are characterized by fast switching speed, excellent endurance, and high scalability.4 7 A fundamental research subject is the detailed investigation of structural changes that occur during phase transitions, which are employed in the storage of data in memories. Recent results8 10 demonstrate that the dramatic differences in physical properties, such as the refractive index between the amorphous and the crystalline phases of GST, are essentially related to a change in the character of the chemical bonds. If one could design a material that retains a high degree of structural perfection during the phase change while allowing a complete change in the character of the bonds, it would be possible to address fundamental questions related to the nature of bonding. However, thin films of GST are usually prepared via sputtering, which results in formation of polycrystalline textured layers. The presence of orientational disorder complicates interpretation of the results. We therefore grow GST epitaxially on slightly mismatched substrates to obtain films with a unique crystalline orientation. The technique of our choice is molecular beam epitaxy (MBE), r 2011 American Chemical Society
which offers several in situ investigation techniques, such as ellipsometry11 and electron diffraction12,13 to monitor the epitaxially grown layers. Epitaxial growth of GST alloys is a challenging topic because the temperature window available for growth, i.e., from 170 to 210 °C, is quite narrow, making accurate and reproducible temperature control hard to achieve.14 In previous studies we showed that epitaxial layers can be reversely laser switched between the ordered and the disordered state in a way comparable to polycrystalline films15 and that a rhombohedral distortion occurs in the GST films epitaxially grown on GaSb(001), suggesting that epitaxial growth on the GaSb(111) surface would be more promising.16 In the following, we compare the MBE growth of GST on slightly mismatched GaSb(001), GaSb(111), InAs(001), and InAs(111) substrates.
2. EXPERIMENTAL PROCEDURE The layers were grown in a MBE apparatus dedicated to chalcogenides at the PHARAO beamline at BESSY II (Helmholtz Center for Materials and Energy, Berlin).17 The system is equipped with separate dual-filament hot lip effusion cells for evaporation of elemental Ge, Sb, and Te, in situ reflection high-energy electron diffraction (RHEED), and a line of sight quadrupole mass spectrometer (QMS). Moreover, the Received: July 8, 2011 Revised: August 17, 2011 Published: August 18, 2011 4606
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during the growth process. A schematic representation of growth monitoring via QMS is given in Figure 1a. A Hiden HAL IV QMS is used as it is capable of scanning masses up to 512 amu, thus enabling one to detect atoms and molecules like Ge, Sb, Te, Sb2, Sb3, Sb4, Te2, and GeTe. The QMS is located at a distance of 30 cm from the substrate at an angle of 20° to the substrate normal. During the measurements the shutter in front of the quadrupole is repeatedly (every 20 s) closed and opened to alternately measure the total signal and the background, which is subsequently subtracted from the signal. In order to reduce the amount of doubleionized molecules and splitting of large parent molecules into daughter fragments, the electron energy of the ionizer is kept at 20 eV and the emission current at 800 mA. The local structure of the films was analyzed by a JEOL 3010 transmission electron microscopy (TEM), operated at 300 kV with a point resolution of 0.17 nm. In-plane SR-XRD measurements were performed using a X-ray beam energy of 9.34 keV.
Figure 1. (a) Schematic of molecules impinging on and desorbing from the surface during growth. (b) GST growth rate (red) and desorbed GeTe molecules (black) as a function of substrate temperature.
Figure 2. RHEED images of the growth of GST with nominal Ge2Sb2Te5 composition at a substrate temperature of 200 °C on InAs(001) (a c) and InAs(111)A (d f). Images were taken before, during ((b) after 15 min deposition, (e) after 30 s), and at the end of the deposition ((c) after 120 min, (f) after 180 min). chamber can be connected in situ to a six-circle diffractometer for synchrotron radiation X-ray diffraction (SR-XRD) analysis. Undoped GaSb (001) and (111) epi-ready substrates were prepared for growth by desorbing the native oxide under an Sb flux, followed by deposition of a 100 nm thick GaSb homoepitaxial buffer layer at 350 °C. InAs(001) substrates were loaded in a dedicated MBE system, annealed up to 460 490 °C to remove the native oxide, and then overgrown with a 100 nm thick homoepitaxial buffer at 400 °C. During the cooling process, the As-stabilized (2 4) reconstructed InAs(001) surface pattern changed displaying a diffuse (1 1) reconstruction.18 The InAs(001) substrates were finally transferred to the PHARAO system using a vacuum shuttle to avoid air contamination. InAs(111) substrates were directly introduced into the GST growth chamber and the native oxide removed at 490 °C without a preceding buffer growth. Substrate preparation as well as the growth process were monitored by RHEED. In-line QMS, an instrument developed for growth of ternary III V alloys,19 allows one to monitor the species desorbed from the surface
3. RESULTS AND DISCUSSION GST alloys possess the fcc rocksalt structure with lattice parameters in the 6.01 6.04 Å range. Since GaSb and InAs have a zinc blende structure with lattice constants of 6.09 and 6.06 Å, respectively, they appear to be suitable substrates for performing GST heteroepitaxial growth. Epitaxial growth of GST presents different growth regimes depending on the chosen growth temperature. At room temperature and up to about 100 °C, asgrown films are amorphous with a smooth morphology. The growth rate as a function of temperature is shown in Figure 1b, in which the red line is a guide to the eye and the black points represent the experimental results obtained for GST grown on GaSb(001). At these temperatures the growth rate is about 0.5 nm/min. The growth rate was evaluated measuring the sample thickness using either secondary electron microscopy (SEM) cross sections or X-ray reflectivity (XRR) thickness fringes for films on (001)- and (111)-oriented substrates, respectively, and dividing the obtained value by the growth time. At a slightly higher temperature of around 150 °C, the films become polycrystalline. Only at 200 °C a continuous epitaxial film is obtained. At these substrate temperatures, however, the growth rate decreases dramatically to about 0.3 nm/min (Figure 1b), and therefore, we assume that only a fraction of the supplied flux incorporates into the film. To investigate this behavior, we made use of the QMS. In Figure 1b, the QMS intensity of the desorbed GeTe molecules (in black) is plotted as a function of the growth temperature. At low temperatures up to 200 °C almost all the GeTe molecules stick on the surface, as we do not observe any signal of desorbed species from the surface. Between 200 and 250 °C the QMS signal slowly increases. Above 250 °C, no growth is observed and the QMS intensity reaches a plateau, revealing a complete desorption of the GeTe molecules. Additionally, no GenSbm or SbnTem molecules are detected in the desorbing flux up to our detection limit of 512 amu; this result indicates a strong bond between Ge and Te in GST, in agreement with the findings of Kolobov et al., who could not detect Sb Ge bonds by means of X-ray absorption fine structure spectroscopy.20 Monitoring via QMS the kinetics of desorption in the course of growth allows one to achieve equivalent surface conditions (temperature) on different substrates even in the presence of varying thermal coupling to holders and thermocouples. Indeed, by applying the same control method we could fabricate epitaxial GST layers on InAs substrates as well. In Figure 2 the RHEED patterns acquired during growth of Ge2Sb2Te5 on InAs are shown. Figure 2a shows the diffraction 4607
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Crystal Growth & Design pattern of the (1 1)-reconstructed InAs(001) surface. Formation of the interface layer with GST takes several minutes; in fact, in Figure 2b we could still observe the streaky RHEED pattern of the substrate, indicating that the surface is not yet fully covered with a complete layer. The subsequent stage of growth is characterized by development of a bright spotty pattern typical for bulk diffraction (see Figure 2c). Thus, we conclude that the surface is rough enough to allow for electron transmission. Moreover, we observe a crystal diffraction pattern; therefore, we additionally infer that the film is epitaxially grown. Further growth proceeds without significant changes in the RHEED pattern, as discussed by R. Shayduk et al.16 for growth of GST InAs(111)A. RHEED patterns of the In-terminated InAs(111) (InAs(111)A) surface are shown in Figure 2d. The surface exhibits a (2 2)
Figure 3. Cross-sectional bright-field TEM image of epitaxial GST on GaSb(001).
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diffraction pattern, commonly observed for both In-terminated InAs(111)A and As-terminated InAs(111)B surfaces.21 The starting substrate surface appears smooth; it produces a streaky pattern with lines of intensity normal to the surface that vary only slowly in intensity, as presented in Figure 2d. During deposition of Ge2Sb2Te5 (Figure 2e), the streaks change but the pattern does not transform into a transmission pattern as in the case of growth on the (001)-oriented substrate. The diffraction spots slightly decrease in intensity, indicating an initial roughening of the surface. At the end of growth (Figure 2f) the pattern again gains in intensity, revealing a smoothing of the surface. To better understand the origin of the roughening of the layers grown on the surface with (001) orientation, as observed by RHEED,14 TEM investigations were carried out. For this purpose, representative images of a GST layer grown on a (001) surface, here GaSb(001), are shown in Figure 3. The layer was grown at 200 °C, as previously pointed out, which ensures the best crystallinity. During transfer through air and sample preparation a surface oxide has formed on the layer. The surface of the layer is fairly rough; however, the layer is continuous and a morphologically rough interface with the substrate is found. This is surprising since the RHEED of the pregrowth GaSb(001) surface indicates large atomically flat terraces separated by monolayer high steps.15 Obviously, the interface becomes unstable in contact with GST at the onset of deposition. This morphological instability is not due to intermixing, since no diffraction peaks of mixed compounds are detected in X-ray diffraction. The interface instability does not impair the epitaxial orientation of the film. The strong buckling of
Figure 4. (a) In-plane reciprocal space map of a GST film on GaSb(111). High-resolution line scan along the high-symmetry Hhex (b) and Hhex = Khex (c) directions. The scans are taken in situ in the MBE chamber after completion of growth. 4608
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Crystal Growth & Design the interface, however, creates local facets that approach the (111) orientation. It suggests that the material prefers to grow in this direction; in fact, we also observed formation of (111) facets on the film surface. In addition, the rhombohedral distortion present in the films is most easily accommodated when the film grows along (111).16 The samples grown on substrates with (111) orientation were further characterized by SR-XRD. Figure 4 shows representative examples from the X-ray diffraction results for GST on GaSb(111). The experiments were performed in grazing incidence geometry with the scattering vector lying inplane, parallel to the sample surface. The intensity is given in a normalized logarithmic scale. A large area reciprocal space map (Figure 4a) was taken around the GST (202)cub peak. The coordinate system is given in hexagonal coordinates (H,K) in units of the reciprocal lattice (rlu) of the GaSb(111) surface, while the coordinates of each individual peak are given with respect to the bulk cubic unit cell. The map clearly shows the GST (202)cub and GST (404)cub peaks located close to the GaSb (202)cub and GaSb (404)cub peaks, respectively. Figure 4b and Figure 4c show high-resolution in-plane scans along the highsymmetry directions of the map shown in Figure 4a, i.e., the Hhex || [ 211]cub and the Hhex = Khex || [ 101]cub directions, respectively. In both directions the GST layer peak is clearly separated from the substrate; no unexpected additional peaks are visible. The GST peak positions are slightly different from the values measured for sputtered samples.22 A detailed investigation of this discrepancy will be the subject of further studies. Truncation of the real crystal at its surface produces a distribution of intensity along the surface normal direction, denoted as crystal truncation rods (CTR), which intersect with the scan axis at Hhex = 1 and Hhex = 2, as shown in Figure 4b. These findings prove the single-crystalline nature of the GST film. The GST peaks are surrounded by a much broader intensity than the perfect, sharp substrate peaks, which is attributed to the lower crystalline quality of the film, compared to the substrate. This is also evident from the high-resolution line scans, in which the layer peaks are much broader than the substrate peaks. The distribution of the intensity is circular around the center of each peak, showing the morphological and structural in-plane homogeneity of the film. Furthermore, there is no distribution of intensity along circles of a constant scattering vector, strongly indicating the absence of polycrystalline grains. For the case of GaSb(001) substrates, in contrast, GST was shown to grow with two main orientations in the vertical direction, both [001] and [111], and in addition, strong diffraction from polycrystalline grains was observed.16 The results clearly demonstrate that growth of GST on (111)-oriented substrates is superior to that on (001)-oriented substrates, without depending significantly on the substrate material.
4. CONCLUSIONS Molecular beam epitaxy was employed to grow epitaxial Ge2Sb2Te5 thin films on nearly lattice-matched (001)- and (111)-oriented GaSb and InAs substrates. The growth behavior was found to be influenced by the substrate orientation rather than by the substrate material. In situ QMS allowed for determination of the window of substrate temperatures, within which Ge2Sb2Te5 grows epitaxially. Indeed, at temperatures slightly lower than 180 °C, polycrystalline layers were obtained, whereas for temperatures above 220 °C thermal decomposition of the film was observed. Epilayers deposited on (001) surfaces resulted in rough films even on the scale of the coherence length of
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RHEED (∼63 nm), as indicated by the presence of spots in the diffraction pattern. In contrast, in the case of (111) substrates, RHEED streaks developed from the beginning of growth, pointing to a rather smooth film surface. SR-XRD measurements demonstrated that Ge2Sb2Te5 layers on GaSb(111) grow in the cubic phase with a unique orientation. Epilayers grown on GaSb(001), instead, were shown to consist of two main domains, characterized by the vertical epitaxial relationships Ge2Sb2Te5[001]//GaSb[001] and Ge2Sb2Te5[111]//GaSb[001], respectively, with a high degree of azimuthal misalignment.15 The different growth behavior on (111) and (001) surfaces could be explained by the presence of a rhombohedral distortion in the Ge2Sb2Te5 cubic lattice along [111], which would be more easily accommodated for (111)-oriented substrates.16 It is noteworthy that such a distortion was reported for GeTe, one of the extremes of the GeTe Sb2Te3 pseudobinary tie line.23 In conclusion, a comparison of the growth of Ge2Sb2Te5 thin films on (001)- and (111)-oriented GaSb and InAs substrates highlighted the superior properties of the epilayers grown on (111) surfaces, in terms of both surface roughness and crystallinity. These results are promising, as achievement of highly ordered, epitaxial Ge2Sb2Te5 is expected to simplify the understanding of the complex amorphous to crystalline switching mechanism, which makes phasechange materials unique.
’ AUTHOR INFORMATION Corresponding Author
*E-mail:
[email protected]. Present Addresses z
Institut f€ur Physik und Astronomie, Universit€at Potsdam, Karl-Liebknecht-Straße 24 25, 14476 Potsdam, Germany. § CreaTec Fischer Co. GmbH, Industriestr. 9, 74391 Erligheim, Germany.
’ ACKNOWLEDGMENT We would like to acknowledge S. Behnke, C. Stemmler, and C. Herrmann for technical support and also thank O. Bierwagen, P. Fons, and A.V. Kolobov for critical reading of the manuscript. Support for this work was provided by the Deutsche Forschungsgemeinschaft (BR 1723/3-1) and the Japan Science and Technology Agency through an international research corporation grant. One of the authors (K.P.) was supported by Deutscher Akademischer Austausch Dienst. Synchrotron radiation was supplied by the Helmholtz Center Berlin. ’ REFERENCES (1) Ovshinsky, S. R. Phys. Rev. Lett. 1968, 21, 1450–1453. (2) Yamada, N.; Ohno, E.; Nishiuchi, K.; Akahira, N.; Takao, M. J. Appl. Phys. 1991, 69, 2849–2856. (3) Atwood, G. The Evolution of Phase Change Memory: Why PCM is Ready for Prime Time as a Next-Generation, Nonvolatile Memory [online]. Published online: July 26, 2010. http://www. micron.com//get-document/?documentId=5539&file=evolution_of_phase_change_memory.pdf (accessed August 17, 2011) (4) Greer, A. L.; Mathur, N. Nature 2005, 437, 1246–1247. (5) Wuttig, M. Nat. Mater. 2005, 4, 265–266. (6) Lankhorst, M. H. R.; Ketelaars, Bas W. S. M. M.; Wolters, R. A. M. Nat. Mater. 2005, 4, 347–352. (7) Hamann, H. F.; O’Boyle, M.; Martin, Y. C.; Rooks, M.; Wickramasinghe, H. K. Nat. Mater. 2006, 437, 383–387. 4609
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