Microstructural characterization of sputter-deposited Pt thin film electrode

Transcription

1 Microstructural characterization of sputter-deposited Pt thin film electrode Ji-Eun Lim, Jae Kyeong Jeong, Kun Ho Ahn, Hyeong Joon Kim, and Cheol Seong Hwang a) School of Materials Science and Engineering and Inter-university Semiconductor Research Center, Seoul National University, Seoul , Korea Dong-Yeon Park and Dong-Su Lee Inostek Incorporated, Gasan-dong, Keumchun-gu, Seoul , Korea (Received 21 May 2003; accepted 7 October 2003) Pt thin films of various thicknesses (30 nm 200 nm) were deposited on Si wafers with SiO 2, Ti, TiO 2,orIrO 2 buffer layers at various temperatures (room temperature 200 C) by a direct current magnetron sputtering process. The Pt films showed a strong (111)-preferred texture irrespective of the thickness, under-layer, and growth temperature. The authors previously reported [J-E. Lim, D-Y. Park, J.K. Jeong, G. Darlinski, H.J. Kim, and C.S. Hwang, Appl. Phys. Lett. 81, 3224 (2002)] that the films were composed of three kinds of grains with slightly different (111) lattice parameters (bulklike, 1.0% and 2.1% larger). This study details the microstructural variations of the Pt films according to the variations of experimental parameters. The different deposition conditions produced slightly different crystalline structures, but the three different (111) lattice parameters were always found. Epitaxial (200) Pt films on a (200) MgO substrate and a highly (111) textured Au thin film on a SiO 2 /Si did not show the same splitting in the lattice parameter. The grains with 1.0% and 2.1% larger (111) lattice parameter almost disappeared after postannealing at 1000 C. However, surface chemical binding of the Pt film before and after annealing was unchanged. Therefore, it is believed that the lattice parameter splitting in the (111) textured Pt film originated from the interfacial grains with the distorted crystal structure due probably to growth stress. I. INTRODUCTION Noble metal thin films such as Pt, Ir, and Ru, especially Pt thin films, have been extensively used as the electrodes of the ferroelectric and high-dielectric functional devices due mainly to their chemical inertness. The crystallization behavior and, thus, the electrical properties of these functional thin films are critically dependent on the structure of the underlying Pt electrode. Therefore, understanding and control of the Pt electrode are critically important to obtain the desired performances of the devices. There have been numerous reports on the growth and characterization of Pt films on various substrates using various methods. The reported growth methods include on- and off-axis magnetron sputtering, 1 3 electron beam evaporation, 4,5 metal organic chemical vapor deposition 6,7 and electroplating. 8 Among them, magnetron sputtering might be the most suitable method for most a) Address all correspondence to this author. cheolsh@plaza.snu.ac.kr applications (excluding applications requiring conformality of the electrode such as dynamic random access memory capacitors) for its good quality control, large area deposition, and mature deposition tool technologies. The sputtered-pt thin films usually show a highly (111) textured growth behavior (the 111 direction of the film is parallel to the substrate surface normal direction, whereas the in-plane orientation is random) because the (111) crystal plane has the largest atomic packing density thereby the smallest surface energy, as is the case for most face-centered-cubic (fcc) metal films. 1 3,9 The growth conditions have a profound influence on the properties and structure of Pt films such as the stress, grain size, degree of texturing, surface roughness, and defect concentration, although they are all (111) textured. All these factors influence the performance of functional thin films in various ways. In particular, the stress of the electrodes is of interest here because the functional properties of the ferroelectric and high-dielectric films, such as the remanent polarization and dielectric constant, are quite dependent on the stress state of the films and electrodes. In addition, it is believed that the electrode film 460 J. Mater. Res., Vol. 19, No. 2, Feb Materials Research Society

2 stress also influences the crystallographic behavior and crystallization of the films themselves. The authors reported that a x-ray diffraction (XRD) study, in Bragg Brentano geometry using a highresolution diffractometer, of sputter-deposited Pt films with a strong (111) texture showed the (111) peak splitting into three separate peaks suggesting that there are three different grains with slightly different (111) lattice parameters. 10 The relative intensity of the peak corresponding to the bulk value of the (111) lattice parameter increased with the increasing Pt film thickness, and the ferroelectric performances of the overgrown Pb(Zr,Ti)O 3 (PZT) thin films were improved with the increase in the Pt thickness due to the enhanced crystallization of the PZT layer. However, normal XRD analysis of the Pt films using the input x-ray with large divergence angle showed that the (111) peak was composed of a single peak, and its position was merely shifted to a higher diffraction angle. This could lead to the incorrect conclusion that the tensile stress increased as the Pt film thickness increased. Therefore, it was necessary to investigate the crystallographic evolution of the Pt films in more detail to elucidate the reason for the (111) peak splitting; that is, for the formation of three types of grains with different (111) lattice parameters. In this study, detailed XRD analyses of the magnetron sputter-deposited Pt films were performed as a function of varying deposition parameters, including the film thickness, substrate temperature during the film growth, postannealing conditions, and types of underlying layers. Through detailed experiments, this study attempted to show the factors that affect the formation of the three kinds of Pt grains. II. EXPERIMENTAL Four major experimental parameters were changed to investigate the formation mechanism of the three different Pt grains: the film thickness, deposition temperature, types of underlying layers, and postannealing temperatures and atmospheres. In addition, a highly (111) textured fcc Au film and epitaxial Pt films on a MgO substrate, where the epitaxial relationship was (200) Pt //(200) MgO, were also prepared. Through oblique-incident direct current magnetron sputtering, (111) textured Pt thin films (thicknesses ranging from 30 nm to 200 nm) were deposited on 10-nmthick Ti/SiO 2 /Si substrates (4-inch diameter). The Pt and Ti layers were grown sequentially in the same chamber without a vacuum break to prevent Ti oxidation. The detailed deposition conditions are given in Table I. The TiO 2 layer was obtained by oxidizing the deposited Ti layer at 650 C for1hinair. The Pt films grown on the TiO 2 layer by the same deposition conditions as the Pt films on Ti were postannealed at 650 C in air for 1 h to improve the adhesion. A 100-nm-thick Pt film was grown on a 30-nm-thick IrO 2 /SiO 2 /Si wafer (6-inch diameter) using a face-to-face type sputtering system. For comparison, a 100-nm-thick Pt film was grown directly on the SiO 2 /Si wafer (8-inch diameter) using another face-to-face type sputtering system. The growth temperature of all the Pt thin films was set to 200 C. Here, the Ti, TiO 2, and IrO 2 served as the adhesion layer between the Pt film and the SiO 2 /Si substrates, and the SiO 2 layers were grown by thermal oxidation to prevent Pt silicidation. To observe the effects of the growth temperature, 50- nm-thick Pt films were directly grown on the SiO 2 /Si substrate at room temperature, 130 C, and 200 C, respectively. The heat-treatment effects on the microstructure were tested by annealing the 50-nm-thick Pt/SiO 2 /Si samples at temperatures ranging from 700 C to 1000 C under air or N 2 atmosphere using a vertical tube furnace. The XRD study was performed using a conventional x-ray diffractometer (M18XHF-SRA, MAC Science Co., Tokyo, Japan) using Cu K radiation in the Bragg Brentano geometry and high-resolution XRD (HRXRD). The HRXRD, which were conducted in the 2 and rocking modes, respectively, were performed using a double-crystal [asymmetrically cut single-crystal Si(220)] diffractometer (Bede Scientific Instruments, Durham, UK, HRXRD with 2-kW Cu radiation source) with a full width at half-maximum (FWHM) spread of 12 arcsec of the incident x-rays (Cu K 1). To align accurately the sample position with respect to the detector, a Si(400) reflex was used at each measurement. The absolute values of 2 for Pt(111) peaks were calibrated using the 2 value of the Si(400) peak. 2 value of [2 value of bulk Pt(111) peak] was taken as 0 in the 2 HRXRD. Reciprocal space mapping (RSM) around the (111) reciprocal lattice point of the Pt film was performed using the triple-crystal diffraction geometry (FWHM < 5 arcsec). Figure 1 shows a schematic diagram showing the reciprocal lattice positions of the Pt(111) and Si(400) and the diffraction geometry. It should be noted that a variation in the G vector TABLE I. Summary of Ti and Pt deposition conditions. Deposition conditions Ar flow rate (sccm) Deposition pressure (torr) Sputtering power (watt) Deposition temperature ( C) Target size (inch) Target rotation Ti yes Pt yes J. Mater. Res., Vol. 19, No. 2, Feb

3 FIG. 1. The schematic diagram showing the reciprocal lattice positions of the Pt(111) and Si(400) and diffraction geometry. length (variation in 2 value) indicates a variation in (111) interplanar spacing, and a variation in corresponds to a divergence in the 111 directions of the Pt grains. The surface morphology and grain size of the Pt films were observed using atomic force microscopy (AFM; Park Scientific Instrument Co., Sungnam, Korea) and scanning electron microscopy (SEM; Philips XL30, Eindhoven, The Netherlands). The detailed grain morphology and Pt adhesion-layer interface were investigated by high-resolution transmission electron microscopy (HRTEM) equipped with a field emission electron source (JEOL-3000F, Tokyo, Japan). III. RESULTS AND DISCUSSIONS Figure 2 shows the conventional XRD patterns of the Pt films with thicknesses of (a) 30, (b) 50, (c) 100, and (d) 200 nm, respectively, deposited on the Ti adhesion layer at a wafer temperature of 200 C. Only the Pt(111) peak can be observed irrespective of the film thickness suggesting that the Pt films are highly (111) textured. There is hardly any change in the (111) peak shape and position except for the increase in the peak intensity with increasing thickness due to the increased x-ray scattering. HRXRD patterns around the (111) peak of the Pt films grown on the Ti adhesion layer with thicknesses ranging from 30 nm to 200 nm in 2 mode were reported in Fig. 1(a) of Ref. 10. It was observed that the (111) peaks are composed of three peaks. Reference 10 also showed the same three (111) peaks from a 100-nm-thick Pt film directly on SiO 2 and from the 50-nm-thick Pt films on the TiO 2 and IrO 2 adhesion layers. The intensity variations of the three peaks with increasing Pt film thickness were discussed in Ref. 10, and it was suggested that the peak near 0 originated from the growing grains. It was also FIG. 2. Conventional XRD patterns of the Pt films with thicknesses of (a) 30, (b) 50, (c) 100, and (d) 200 nm, respectively, deposited on the Ti adhesion layer at a wafer temperature of 200 C. suggested that the first two peaks might have originated for some reason from the Pt nuclei that have a distorted lattice parameter. When the adhesion layer was changed to TiO 2, there was little difference in the Pt HRXRD pattern from a comparison between Figs. 1(a) and 1(b) of Ref. 10. However, when the sputtering geometry was changed from the oblique type to face-to-face type, the relative separation between the first and third peaks was slightly increased. 10 The sputtering geometry usually has a large influence on the stress of the deposited films through the peening effect. 11 Therefore, it is reasonable to assume that the origin for the formation of the first and second peaks have something to do with the stress state of the film. This will be discussed in detail later. Figures 3(a) and 3(b) show the two-dimensional reciprocal space mapping (RSM) results around the (111) Pt reciprocal lattice points of the 30- and 200-nm-thick Pt films, respectively. The figure clearly shows that the (111) Pt reciprocal lattice points are composed of three subsidiary components with a relatively large elongation along the rocking direction due to the relatively poor alignment of the 111 directions of each grain. One further interesting observation from Fig. 3 is that the distance between the first and second peaks became smaller with almost constant distance between the first and third peaks as the rocking angle became negative irrespective of the Pt film thickness. The crystallographic implication of this variation is unclear. 462 J. Mater. Res., Vol. 19, No. 2, Feb 2004

4 The HRTEM images of the Pt/Ti interface [Figs. 7(b) and 8] show that there was no interfacial reaction that would result in a Pt x Ti y alloy formation. Therefore, the extra two peaks are certainly not from any alloy phase. In addition, the appearance of the two extra peaks from the Pt/SiO 2 /Si [Fig. 1(b) of Ref. 10], where there can hardly be any interfacial reaction between the Pt and SiO 2, clearly shows that those originated from the Pt film itself. To further check if this peak split occurs merely from some artifacts related to the equipment, the HRXRD pattern of Si(400) was obtained (data not shown). It could be clearly understood from the single and sharp Gaussian peak shape that there are no artifacts with the equipment. Figures 4(a) and 4(b) show the wide-angle XRD spectrum using conventional XRD and HRXRD of the Au(111) peak of the Au/SiO 2 /Si sample, respectively. Figure 4(a) clearly shows that the film is highly (111) textured but Fig. 4(b) shows that there is no (111) peak splitting. Therefore, the (111) peak splitting is not a general phenomenon of sputtered fcc metals. Figure 5 shows the HRXRD of the Pt(200) peak of the epitaxial (200) Pt/MgO sample, respectively. This clearly shows that the Pt film is (200) oriented but there is no (200) peak splitting. More details about this epitaxial Pt film can be found in Ref. 12. From the comparison between the three cases above, polycrystalline (111) textured Pt, polycrystalline (111) textured Au, and epitaxial (200) Pt, it can be inferred that the two extra (111) Pt peaks originated from some of the (111) oriented Pt grains with a 1.0% and 2.1% larger (111) lattice spacing. Furthermore, from the nonvarying peak positions and intensities of the first and second (111) peaks in Fig. 1(a) of Ref. 10 with increasing film thickness, grains with the increased (111) lattice spacing exist either at the interface or surface. It was concluded from the following experimental observations that these grains exist at the interface. The columnar growth manner of the growing Pt grains with the bulklike lattice parameter [corresponding to the third (111) peak] was accompanied by an increase in the lateral dimension of each grain, as can be seen from the AFM images of the Pt films with the different thicknesses shown in Fig. 6. The root-mean-square surface roughness of the films increased with increasing thickness from approximately 0.68 nm at 30 nm to 2.6 nm at 200 nm. Apart from this change in surface structure of the Pt films with the increasing thickness, the first and second (111) peaks in Fig. 1(a) of Ref. 10 did not change at all, suggesting that the Pt grains responsible for the occurrence of the first and second (111) peaks do not exist on the film surfaces. FIG. 3. The two-dimensional RSM results around the (111) Pt reciprocal lattice points of the (a) 30- and (b) 200-nm-thick Pt films. FIG. 4. (a) Wide-angle XRD spectrum using conventional XRD and (b) HRXRD of the Au(111) peak of the Au/SiO 2 /Si sample. FIG. 5. HRXRD of the Pt(200) peak of the epitaxial (200) Pt/MgO sample. J. Mater. Res., Vol. 19, No. 2, Feb

5 FIG. 6. AFM images of the Pt films with thicknesses of (a) 30, (b) 50, (c) 100, and (d) 200 nm. More direct observations of the internal grain structure of the Pt films on the Ti adhesion layers can be found from the cross-section HRTEM. Figure 7(a) shows the low-magnification TEM micrographs of the 30-nm-thick Pt film on the 10-nm-thick Ti film and Fig. 7(b) shows the high-resolution TEM micrographs of the interface between the 200-nm-thick Pt film and the 10-nm-thick Ti film. The Pt grains have columnar shapes with an average column diameter of approximately the same as the film thickness. The Pt Ti interface shows no interfacial reaction. Because this film showed the split (111) peak in the HRXRD, some of the grains in Fig. 7(a) should have increased (111) lattice parameters by 1.0% and 2.1% in addition to the grains with the bulklike lattice parameter. The volume fraction of these grains with the three different (111) lattice parameters appears to be roughly the same from the similar intensities of the three diffraction peaks. As the film thickness increases, most of the grains show a columnar growth behavior extending from the Pt Ti interface to the surface. However, some of the grains stop growing, and a second grain nucleates on top of the stopped grains, as typically shown in Figs. 8(a) and 8(b). Figures 8(a) and 8(b) show the bright-field and the dark-field images, respectively, of the 200-nm-thick Pt FIG. 7. (a) Low-magnification TEM micrographs of the 30-nm-thick Pt film on the 10-nm-thick Ti film, and (b) high-resolution TEM micrographs of the interface between the 200-nm-thick Pt film and the 10-nm-thick Ti film. FIG. 8. (a) Bright-field and (b) dark-field images of the 200-nm-thick Pt film on the Ti adhesion layer. 464 J. Mater. Res., Vol. 19, No. 2, Feb 2004

6 film on the Ti adhesion layer. The grain with the dark contrast in Fig. 8(a), indicated by a white arrow, shows a different contrast from the grain with the light contrast, indicated by the black arrow, on top suggesting that their crystallographic orientations are different. This clearly shows that there is the possibility that the grains with the increased (111) lattice parameter stopped growing during the deposition run and another Pt grain with a normal (111) lattice parameter grew on top. This might be due to the increased volume free energy of the grains with the nonequilibrium lattice parameter with increasing grain size. An attempt was made to obtain direct evidences of this hypothesis by utilizing the convergent-beam electron diffraction of the various grains and making a direct comparison of the lattice fringes in HRTEM. Unfortunately, the blurred shapes of the higher-order Laue zone (HOLZ) lines, due probably to the high density of point defects, and the difficulty in simulating the HOLZ lines based on the assumption of the arbitrarily distorted unit cell shapes did not allow a reliable estimation of lattice parameters. A direct comparison between the lattice fringes from the different grains was also not reliable due to the small difference in the (111) lattice parameters (only 1.0% and 2.1%). Therefore, any direct evidence for the existence of grains with an increased (111) lattice parameter could not be obtained. However, Fig. 8, at least, shows that there were some small growth-stopped grains at the interface and these might be the grains with the larger (111) lattice parameters. The approximate diameters of the grains with the increased (111) lattice parameters can be estimated from the FWHM values in the 2 mode [Fig. 2(a) of Ref. 10] using the Scherrer formula, as shown in Eq. (1), 13 t = 0.9 B cos B, (1) where t,, B, and are the grain diameter, x-ray wave length, FWHM in radians, and Bragg diffraction angle, respectively. The FWHM of the first and second peaks converged to approximately 500 arcsec. ( radians) in Fig. 2(a) of Ref. 10, which corresponds to a t of approximately 65 nm. This value coincides very well with the size of the growth-stopped grains in Fig. 8. Figure 9 shows the variations in the peak intensities of the three peaks as a function of the postannealing temperatures ( C). Here, the 50-nm-thick Pt films were grown directly on the SiO 2 /Si wafer at room temperature to alleviate any possible interfacial reaction during the annealing in air atmosphere for 1 h. The postannealing greatly increased the third peak intensity where the higher the temperature, the higher the intensity. However, the intensities of the first and second peaks remained almost constant or slightly reduced up to 900 C and then decreased to very low levels at 1000 C. The increase in the third peak intensity was due to the increase in the surface grain size, as shown in the SEM pictures in Fig. 10. Another interesting observation is the relative shift of the positions of the three peaks. The peak positions generally shift to higher 2 values with increasing annealing temperature as a result of thermal stress. The thermal expansion coefficients of Pt and Si are / C and / C, respectively. 14 Therefore, tensile stress should develop in the film during cooling. The peak shift to higher 2 values originates from this tensile stress. However, interestingly, the shifts of the first and second peak positions are much smaller than that of the third peak, as shown in Fig. 11. This suggests FIG. 9. HRXRD patterns of the 50-nm-thick Pt/SiO 2 /Si samples at the (a) as-deposited state and annealed at (b) 700, (c) 800, (d) 900, and (e) 1000 C. FIG. 10. SEM micrographs of 50-nm-thick Pt films annealed at (a) 700, (b) 800, (c) 900, and (d) 1000 C. J. Mater. Res., Vol. 19, No. 2, Feb

7 that the growing grains (the grains with the bulklike lattice parameter) exert some additional compressive stress to the grains with the higher (111) lattice parameters. Here, it was implicitly assumed that the grains with the higher (111) lattice parameters were produced by compressive stress during the initial stages of film formation. The extra compressive stress might also suppress the growth of these grains with the increased (111) lattice parameters up to 900 C. These grains might be dissolved into the bulk Pt grains at 1000 C owing to the excessive thermal energy at this temperature. Gallego et al. 15 investigated the surface structure of the Pt(111) plane using a low-energy electron diffraction technique. They carefully prepared a clean Pt(111) surface by the in situ baking of the sample in an ultra high vacuum condition and at a high temperature (1300 K). They found that the interplanar spacing of the first and second surface layers of the (111) plane increased and decreased, respectively, by approximately 1% due to surface reconstruction from the careful fitting of the diffraction data. 15 According to this report, the surface states of the current Pt films before and after postannealing at 1000 C were investigated by x-ray photoelectron spectroscopy (XPS). It was assumed that the XPS signal comes from the Pt surfaces of only 2-3 nm in depth. It should be noted that the extra two (111) peaks disappeared after high-temperature annealing (Fig. 9). Figure 12 show the XPS spectra of the Pt 4f core-level binding energy shift of the as-deposited and postannealed samples. The C 1 s peak position was taken as the reference to calibrate the energy scale of the spectra. They were almost identical, suggesting that there were negligible changes in the structure during annealing. If the Pt grains with the increased (111) lattice parameters existed on the surface, there should be certain changes in the XPS peak positions after annealing because the peak position must be sensitive to the chemical binding status and thus the interplanar spacing. There is other evidence showing that the (111) peak splitting did not originate from the surface grains. Figure 13 shows the HRXRD of the 100-nm-thick Pb(Zr 0.52 Ti 0.48 )O 3 film on the 30-nm-thick Pt/Ti/SiO 2 /Si substrate. The PZT film showed a (111) textured crystallization behavior, but the PZT (111) peak was not detected in this HRXRD due to the relatively small intensity of the peak. Instead, the three (111) Pt peaks were clearly detected after annealing at 650 C for 30 min. Therefore, even the chemical modification of the Pt(111) surface by the deposition of a PZT FIG. 12. XPS spectra of the Pt 4f core-level binding energy shift of the as-deposited and postannealed at 1000 C samples. FIG. 11. The variations in the positions of the first, second, and third peaks as a function of the annealing temperatures. FIG. 13. HRXRD of the 100-nm-thick Pb(Zr 0.52 Ti 0.48 )O 3 film on the 30-nm-thick Pt/Ti/SiO 2 /Si substrate. 466 J. Mater. Res., Vol. 19, No. 2, Feb 2004

8 film did not induce any changes in the (111) peak split. This serves as further evidence of the absence of the grains with an increased (111) lattice parameter on the film surface. In fact, if the one or two monolayers of the modified surface by the reconstruction, as reported by Gallego et al., 15 were responsible for the (111) peak split, the FWHM of the extra peaks in the 2 mode should be approximately 44, according to Eq. (1), which is certainly not the case. The effects of the deposition temperatures on the HRXRD of the 50-nm-thick Pt/SiO 2 /Si samples are shown in Fig. 14 where the films were grown at (a) room temperature, (b) 130 C, and (c) 200 C. The surface grain size increased with the increasing deposition temperature as shown in Figs. 15(a), 15(b), and 15(c). The HRXRD shows several interesting results. The peak intensity of the third peak (bulklike peak) increases with the increasing temperature coinciding with the increase in the surface grain size shown in Fig. 15. The third peak position of the room temperature grown sample was shifted to a negative 2 value, suggesting that the film was under the compressive stress due to the peening effect. 16,17 When the growth temperature was increased to 130 C and 200 C, the third peak position shift changed to a positive value, indicating that the film stress became tensile due to the thermal stress effect as in the case of the thermal annealing shown in Figs. 9 and 11. The magnitude of the third peak position shift was approximately arcsec when the growth temperature was changed from room temperature to 130 C or 200 C. The first and second peaks also shifted to a positive 2 value, but the magnitude was only approximately +300 FIG. 15. SEM micrographs of the 50-nm-thick Pt/SiO 2 /Si samples grown at (a) room temperature, (b) 130 C, and (c) 200 C. FIG. 14. HRXRD of the 50-nm-thick Pt/SiO 2 /Si samples grown at (a) room temperature, (b) 130 C, and (c) 200 C. arcsec. Furthermore, the peak intensities of the first and second peaks even decrease, whereas that of the third peak certainly increases when the deposition temperature increases from 130 C to 200 C. This is qualitatively coincident with the variations in the peak positions and intensities of the three peaks as a function of the annealing temperature shown in Figs. 9 and 11. This also suggests that the (111) Pt grains that produced the first and second HRXRD peaks are produced by biaxial compressive stress whereas the grains that produced the third HRXRD peak are relatively free of compressive stress. It is believed that the many experimental observations up to now suggest the following experimental procedure occurred during the sputter-deposition of the (111) textured Pt films. Phillips et al. 16 recently showed that the J. Mater. Res., Vol. 19, No. 2, Feb

9 (111) textured Pt film grew on a SiO 2 surface through an island growth manner. At the very early stages of the film growth, the film stress changed from compressive (up to the nominal thickness of approximately 0.5 nm) to tensile (up to the nominal thickness of 3.3 nm) then to compressive again after approximately 3.3 nm. They ascribed the second tensile stress and the third compressive stress to the coalescence of Pt nuclei and peening effect. It is certain that the experimental thickness range of the current Pt films (>30 nm) is in the third compressive stress range. Now it can be assumed that the compressive stress by the peening effect 17 caused the (111) lattice parameter of the initial grains to increase in the early stage of film growth. However, with the increase in film thickness (or the grain height), the accumulation of internal strain energy due to a nonequilibrium lattice distance renders the grains unstable to further growth, and grains with bulklike (111) lattice parameters nucleate and grow further. Moreover, the initially nucleated grains with the bulklike lattice parameters kept growing on the film surface. However, this sort of renucleation is not a general feature of (111) textured fcc metal films, as shown in the case of the (111) textured Au film shown in Fig. 4. This is because the mechanical characteristics of the materials, such as Young s modulus and Poisson ratio, should also be involved in determining the strain energy. One further unsolved question is why the 1.0% and 2.1%? It might be reasonable to assume a continuous variation in the lattice parameter or a single value of lattice distortion corresponding to the given stress magnitude. It is possible that the magnitude of the compressive stress changes with an increase in film thickness, and this change induces the two different lattice distortions. Regarding this hypothesis, HRXRD experiments on samples with a much smaller Pt thickness ( 10 nm) would be very interesting. However, the scattering intensities from the thinner films (<30 nm) were so weak that the diffraction study was not possible using the current XRD equipment. A brighter x-ray source, such as synchrotron x-rays, may be useful. Even when the varying stress assumption is accepted, still the question as to why the specific values of 1.0% and 2.1% lattice distortions are induced remains. To solve this question, full quantum mechanical calculations based on the first principle are currently underway. The basic idea of this calculation is to determine if there is any specific lattice distortion that corresponds to the local energy minima in the interplanar spacing versus the system total energy curve. A preliminary result using a model system composed of 16 layers of Pt (111) crystal plane, which is laterally compressed, showed that approximately 0.9% out-of-plane tensile strain is more stable than other strain values. 18 IV. CONCLUSIONS A HRXRD study of sputter-deposited Pt films with a (111) texture showed that the films were composed of three kinds of grains with slightly different (111) lattice parameters (bulklike, 1.0% and 2.1% larger). It is believed that the lattice parameter splitting in the (111) textured Pt film originated from interfacial grains with a distorted crystal structure due probably to growth stress produced by the peening effect. However, the precise causes for the specific (111) lattice distortion of 1.0% and 2.1% are still unclear. ACKNOWLEDGMENTS This work was supported by Korea Research Foundation (Grant No. KRF D00348), the Korea Ministry of Science and Technology through the National Research Laboratories program and the National R&D Project for Nano Science and Technology (M B ), and the System IC 2010 program of the Korean government. REFERENCES 1. D-S. Lee, D-Y. Park, M.H. Kim, D-I. Chun, J. Ha, and E. Yoon, in Thin Films-Structure and Morphology, edited by S.C. Moss, D. Ila, R.C. Commarata, E.H. Chason, T.L. Einstein, and E.D. Williams. (Mater. Res. Soc. Symp. Proc. 441, Pittsburgh, PA, 1997), p B.M. Lairson, M.R. Visokay, R. Sinclair, S. Hagstrom, and B.M. Clemens, Appl. Phys. Lett. 61, 1390 (1992). 3. H. Tokura, B. Window, D. Neely, and M. Swain, Thin Solid Films 253, 344 (1994). 4. P.C. McIntyre, C.J. Maggiore, and M. Nastasi, J. Appl. Phys. 77, 6201 (1995). 5. P.C. McIntyre, C.J. Maggiore, and M. Nastasi, Acta Mater. 46, 879 (1997). 6. J.M. Lee, C.S. Hwang, H.J. Cho, and H.J. Kim, J. Electrochem. 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