Phase Formation and Morphology of Nickel Silicide Thin Films

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Phase Formation and Morphology of Nickel Silicide Thin Films Synthesized by Catalyzed Chemical Vapor Reaction of Nickel with Silane Antony Premkumar Peter,* Johan Meersschaut, Olivier Richard, Alain Moussa, Johnny Steenbergen, Marc Schaekers, Zsolt Tőkei, Sven Van Elshocht, and Christoph Adelmann IMEC, Kapeldreef 75, B-3001 Leuven, Belgium ABSTRACT: The synthesis of nickel silicide thin films via a vapor− solid reaction has been studied by exposing thin (10 nm) Ni films to silane (SiH4). The crystalline phases, the Ni/Si stoichiometric ratios, as well as the surface and interface properties of the resulting silicide films were investigated as a function of the growth parameters such as the SiH4 partial pressure, the reaction temperature, and the exposure time. At low temperature (300 °C), SiH4 exposure led to the self-limiting deposition of Si on Ni by catalytic decomposition of SiH4 but not to silicate formation. Between 350 and 400 °C, phase pure orthorhombic NiSi films were obtained that were formed directly without any apparent intermediate Ni-rich silicide phases. A transformation to NiSi2 occurred at 450 °C and above, and at 500 °C phase pure NiSi2 was obtained. Here, the transient formation of NiSi was observed that transformed into NiSi2 for prolonged SiH4 exposure. The results indicate that the Si solubility governs the phase formation sequence whereas kinetics are determined by Ni diffusion and the reaction rate. Resistivity values of 21 and 36 μΩ cm were found for the NiSi and NiSi2 thin films, respectively, corresponding to the values reported for films obtained by solid-state reactions.

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focused on the development of nickel silicide nanowires.6,16−23 Despite the attention silicide nanowires have elicited, tailoring the desired morphology with controlled dimensions, phase, and composition with good reproducibility presents still a formidable challenge. In previous studies on the vapor−solid reaction of SiH4 with Ni, it was shown that process parameters, including the SiH4 partial pressure (the supersaturation) and the reaction temperature, as well as the surface preparation, were highly critical in determining the morphological characteristics of nickel silicide nanostructures.21−24 No clear and unambiguous picture has emerged that explains the process windows for nanowire or thin film growth and the phase sequence. In the present study, we report on the observation of several nickel silicide phases as thin films after vapor−solid silicidation of Ni in SiH4. We describe in detail the process parameter dependence, the phase sequence, as well as the resulting thin film morphology and properties. For all studied conditions (that overlap strongly with previously reported process conditions21−24), we observe thin film formation without any sign of nanowire growth unless the Ni films are transformed into nanoclusters by high-temperature annealing. This demonstrates that the vapor−solid synthesis route of nickel silicides is

INTRODUCTION Nickel silicide (NixSiy) thin films are of central importance for numerous technological applications. As an example, nickel monosilicide (NiSi) and nickel disilicide (NiSi2) have been widely used in modern microelectronic and optoelectronic devices. NiSi, due to its low resistance, lower silicon consumption during formation, and midgap work function, has become the material of choice for contact applications in complementary metal−oxide−semiconductor (CMOS) devices where an ultrashallow penetration of the contact into the substrate is mandatory.1−4 Nickel-rich silicides, such as Ni2Si and Ni31Si12, show high work functions and have been considered as electrode materials for p-type metal−oxide− semiconductor devices.5−7 Furthermore, the Si-rich silicide NiSi2 has received much attention for low-power memory devices and in optical communication applications.5 NiSi2 has also been investigated in the context of Schottky barrier height studies since the band alignment depends strongly on the orientation of the Si substrate.8 In the past, nickel silicide thin films have been mainly synthesized via diffusion-mediated solid-state reactions of deposited thin Ni films with the Si substrate, a process that has often been termed self-aligned silicidation (salicidation).3,9−15 In this article, we report on a different strategy to synthesize NixSiy thin films that is based on the reaction of Ni thin films with SiH4. While this process has been previously reported for the Ni−Si system, these studies have mainly © XXXX American Chemical Society

Received: October 16, 2014 Revised: December 11, 2014

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highly promising for thin film applications in nanoelectronics and that the nanowire growth requires further consideration to understand its mechanism.

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EXPERIMENTAL SECTION

All nickel silicide thin films were synthesized by reacting Ni films with silane (SiH4). In a first step, 50 nm of SiO2 was grown on 300 mm Si(100) wafers by rapid thermal oxidation to avoid salicidation reactions with the Si substrate. Subsequently, Ni films were deposited by physical vapor deposition (PVD) at room temperature in an Applied Materials Endura sputtering chamber; all studies were carried out using a fixed Ni thickness of 10 nm, unless noted otherwise. Nickel silicide thin films were then synthesized by exposing the Ni films to SiH4 in an Applied Materials Producer BLOk plasma-enhanced chemical vapor deposition (PECVD) chamber. SiH4 was undiluted, and thus the chamber pressure was equivalent to the SiH4 partial pressure. All silicide deposition−cum−chemical vapor reactions (also referred to in the following as silicidation reactions) were carried out at a chamber pressure of 1.0 Torr, unless noted otherwise, and at temperatures between 300 and 500 °C. We remark that all wafers and chambers were contamination controlled. The maximum metallic contamination on the wafer surface was below 1011 atoms/cm2, and the maximum particulate contamination was below 0.05/cm2. The phase evolution of the nickel silicide films was examined by grazing incidence X-ray diffraction (GIXRD) with an incidence angle of ω = 0.7°. The film thickness was determined by X-ray reflectivity (XRR). Both GIXRD and XRR were performed in a MetrixL diffractometer from Jordan Valley Semiconductors using Cu Kα radiation. The sheet resistance (Rs) was measured at 49 points across the wafer using a 4-point probe KLA-Tencor Rs100 system. The microroughness of the films was monitored by haze analysis in a KLA Tencor SP2 SurfScan instrument with a laser wavelength of 355 nm. Absolute haze values typically depend on the material and surface properties, in particular on the film reflectivity; however, relative trends with growth parameters can be reliably interpreted.25 Haze measurements were complemented by scanning electron microscopy (SEM) and atomic force microscopy (AFM) analysis using a Hitachi SU8000 microscope and a Nanoscope dimension 3100 instrument, respectively. Transmission electron microscopy (TEM) images were acquired in a FEI Tecnai F30 electron microscope operating at 300 kV after focused ion beam sample preparation. High-angle annular darkfield scanning transmission electron microscopy (HAADF-STEM), energy dispersive spectroscopy (EDS), and electron energy loss spectroscopy (EELS) data were acquired in parallel. The stoichiometry of the deposited nickel silicide thin films was determined by Rutherford backscattering spectrometry (RBS) using a 1.5 MeV He+ beam with sample tilt and backscattering angles of 36° and 170°, respectively.

Figure 1. (a) GIXRD pattern and (b) SEM image of a 10 nm thick asdeposited Ni film before SiH4 exposure.

known from previous studies that the detailed solid-state reaction path is characterized by a complicated phase formation sequence, which shows a strong dependence on the thickness of the Ni film, the Si substrate type (amorphous or crystalline), the surface preparation, the temperature, and the annealing conditions.11,27 Thus, it is imperative to understand the phase and growth evolution of the Ni−SiH4 reaction to control the film properties. This not only concerns the stoichiometry but also the morphology and crystallinity of the reacted silicides, including, within the framework of emerging vapor−solid reaction studies, the dimensionality of the formed structures.1,6,7,16,17,21−24 Figure 2a shows the dependence of both Rs and haze values of nickel silicide thin films on the silicidation reaction chamber pressure (1.0−6.0 Torr). The reaction time and temperature were kept constant (60 s; 500 °C). The Rs and haze of the initial Ni film were 19.5 Ω/□ and 500 °C were required to thermally decompose SiH4.21,29,32 Furthermore, the absence of any reaction between thick NiO (grown by oxidation of Ni films in O2 at 400 °C) and SiH4 (results not shown) even at high SiH4 partial pressures (9 Torr) and very long exposure times (45 min) confirm that the silicide formation was selective to the Ni catalyst. Hence, this process allows for a selective formation of silicides, very similar to the salicide processes using solid-state reactions. Cross-sectional TEM images of films grown at 350 and 500 °C are shown in Figure 5. For both samples, a discontinuous interlayer exhibiting a bright contrast was observed within the bottom part of the NiSix films (Figures 5a and d). Note that the crystallographic orientation of the NiSi2 grown at 500 °C (Figure 5e) was continuous across the interlayer, corroborating

appearing at 2Θ = 28.5° that can be assigned to the (111) reflection of NiSi2. This phase transformation from NiSi to NiSi2 can be seen clearly at 500 °C. Several reflections corresponding to NiSi were found either to be absent [(102), (202), (103), (020)] or to undergo a strong reduction in intensity [(002), (112)]. In parallel, a number of reflections appeared that were due to the formation of NiSi2 [(111), (220), (311), (400)]. Hence, almost phase pure NiSi2 was obtained after silicidation at 500 °C. Although the phase transformation from NiSi to NiSi2 resulted in a negligible increase in Rs, it did strongly increase the film roughness, as can be deduced from the haze increase from ∼70 ppm at 450 °C to ∼150 ppm at 500 °C. However, it should be noted that even though the Rs value remained approximately constant during the NiSi−NiSi2 phase transformation, the different volume expansion during the formation of NiSi and NiSi2 will lead to a thickness variation and thus to differences in the resistivity of NiSi and NiSi2 thin films. This will be discussed further below. Figures 4a−c show RBS spectra of the nickel silicide thin films synthesized at three different temperatures (300, 350, and 500 °C). The background Si signal from the substrate and the thermal SiO2 is illustrated in the figure together with the signal due to reacted film. The RBS Ni/Si stoichiometric ratios of the reacted films at 350 and 500 °C were 1.05 and 0.65, respectively. This confirms the GIXRD observation that

Figure 4. RBS spectra of Ni films after silicidation in SiH4 at different temperatures: (a) 300, (b) 350, and (c) 500 °C. The reaction time and SiH4 partial pressure were kept constant (60 s, 1.0 Torr). The red lines indicate best simulations based on layered structures. D

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Figure 5. TEM cross-sectional images of nickel silicide thin films grown by the chemical vapor reaction of Ni with SiH4 at different temperatures: (a), (b) 350 °C and (d), (e) 500 °C. The insets in (b) and (e) show fast Fourier transforms of high-resolution lattice images inside grains in films synthesized at 350 and 500 °C, respectively. (c) and (f) show HAADF-STEM images of the films grown at 350 and 500 °C, respectively. The SiH4 chamber pressure (1.0 Torr) and exposure time (60 s) were kept constant during the formation of all films.

Figure 6. (a) GIXRD patterns (offset for clarity); (b) sheet resistance (Rs) and haze values; (c)−(f) SEM top-view morphologies of nickel silicide thin films processed at various SiH4 reaction times, as indicated. The reaction temperature and chamber pressure were kept constant (500 °C, 1.0 Torr).

the discontinuity. This was less obvious for the NiSi film grown at 350 °C (Figure 5b), possibly because of smaller grain sizes. The Fourier transforms of high-resolution lattice images inside well-defined grains in the films grown at 350 and 500 °C (insets of Figures 5b and 5e) were in agreement with the crystalline structures of NiSi and NiSi2, respectively. The HAADF-STEM images (Figures 5c and 5f) showed dark contrast of the interlayer with respect to the NiSix layers, indicating the average atomic number was lower. In addition, there was little contrast in the NiSix layers both above and below the interlayer, indicating that the same silicide phases were obtained. The EDS quantification performed at five

different locations of both films resulted in compositions close to NiSi and NiSi2 for the films grown at 350 and 500 °C, respectively. Few NiSi crystals exhibiting a brighter contrast in the HAADF-STEM images (not shown here) were also observed in the NiSi2 film grown at 500 °C. This was confirmed by EDS, in keeping with the RBS and GIXRD results that show the presence of excess Si and a NiSi minority phase in the NiSi2 layer. The analysis of the interlayers inside both the NiSi and NiSi2 films by electron energy loss spectroscopy (EELS) indicated the presence of O. Note that O was absent elsewhere in the silicide layers within the sensitivity of EELS. Although the exact chemical nature of the interlayer is difficult E

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These results were further confirmed by the AFM morphology and rms roughness analysis in Figure 7. The

to establish due to projection effects, it clearly establishes that the interlayer is an oxide. As reactions with the underlying SiO2 layer are not expected at temperatures of ≤500 °C,33 the thin oxide interlayer most probably stems from a thin surface oxide of the initial Ni film. Upon exposure to SiH4, it can be transformed into NiSiOx-like material.34 If we assume that the interlayer itself does not diffuse (much), it can be considered as an insoluble marker. Then, its presence close to (but not right at) the interface of the NiSix layers with SiO2 is a strong indication that both Si and Ni can diffuse at the reaction temperatures, but Ni outdiffusion through the silicide is much faster than Si indiffusion and kinetically determines the reaction rate.21,35,36 Hence, the silicidation reactions occur via catalytic decomposition of SiH4 at the surface in combination with Ni outdiffusion. TEM also indicated that the surface of the films was rather rough for both samples. The average thickness of the NiSi2 layer was 36 nm, whereas the average thickness of the NiSi layer was about 21 nm. This was in agreement with the XRR analysis (data not shown). The observed thickness change with respect to the initial Ni thickness (10 nm) is in line with the NiSi and NiSi2 volume expansion (∼2.2× and ∼3.6× , respectively) found by previous solid-state reaction studies.9,10,13 Based on the measured thickness and Rs values (see Figure 3a), thin film resistivity values of 21 and 36 μΩ cm could be determined for the NiSi and NiSi2 phases, respectively, which are close both to the values reported in the literature for self-aligned mono- and disilicide thin films and to bulk values.5,30,37−40 This is remarkable given the fact that the formation of morphologically stable NiSi2 films has been reported only for much higher temperatures (≥750 °C) for solid-state reaction routes.3,10,13−15,30,41 We now turn to the dependence of the silicidation process on the SiH4 exposure time at different temperatures. Figure 6 shows results of GIXRD, Rs, and haze, as well as SEM morphologies for different SiH4 exposure times in the range between 15 and 120 s. Here, the temperature was 500 °C, and the SiH4 partial pressure was fixed at 1.0 Torr. GIXRD (Figure 6a) indicated that the predominant phase formed after the shortest SiH4 exposure (15 s) was NiSi. After a longer SiH4 exposure of 30 s, an additional (111) reflection due to NiSi2 started appearing. Subsequently, the NiSi reflections were strongly reduced in intensity close to the detection limit in parallel with an increase in peak intensities for NiSi2 reflections. Phase pure (within the sensitivity of GIXRD) NiSi2 films were finally obtained after 120 s of SiH4 exposure. The strong enhancement of the intensities of the NiSi2 reflections indicated an increasing degree of crystallinity. This shows that phase pure NiSi can also be obtained at higher temperature for short silicidation times. In keeping with the effects of temperature and SiH4 partial pressure on Rs and haze (Figures 2a and 3a), the phase transformation from NiSi to NiSi2 with increasing SiH4 exposure time (Figure 6b) also resulted in a strongly increased haze (increased surface roughness) but only in rather insignificant changes in Rs (9.1−10.7 Ω/□). The increased haze after prolonged SiH4 exposure was consistent with the SEM morphologies in Figures 6c−f. While NiSi films after 15 s of SiH4 exposure were rather smooth (Figure 6c), increasing SiH4 exposure times up to 30 s led to a structure with randomly scattered pinholes (Figure 6d). Finally, the formation of phase pure NiSi2 at the longest SiH4 exposure times (60 and 120 s) resulted in rougher morphologies as evidenced in Figures 6e and f.

Figure 7. AFM morphologies of NiSi2 films synthesized at 500 °C using different SiH4 exposure times: (a) 60 and (b) 120 s. The chamber pressure was kept constant at 1.0 Torr.

films grown by 60 s of SiH4 exposure (Figure 7a) consisted of a heterogeneous morphology with scattered islandlike regions distributed in the NiSi2 matrix.13 Since RBS (Figure 4c) and GIXRD (Figure 6a) indicate that the overall films consisted predominantly of NiSi2, these regions may be attributed to remaining NiSi crystallites that are en route to reaction into NiSi2. After 120 s of SiH4 exposure, the (NiSi) islands disappeared completely and a homogeneous film morphology was observed, albeit with large roughness (Figure 7b). More quantitatively, the conversion of NiSi to NiSi2 was accompanied by a strong roughening of the surface with an rms value increase from 5.6 to 7.9 nm. This correlates well with the NiSi− NiSi2 phase transformation reported by previous solid-state reaction thin film studies.3 It should be mentioned that the strong roughening due to NiSi2 formation has been observed commonly in previous solid-state routes also. In the solid-state process, the formation of Ni2Si and NiSi occurs by a diffusion-dominated mechanism, while NiSi2 grows via a nucleation-limited process. Once NiSi2 nuclei have been formed in isolated regions, the nuclei grow very rapidly both vertically and laterally, resulting in a rough F

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Figure 8. (a), (b) GIXRD patterns (offset for clarity) of nickel silicide thin films synthesized using different SiH4 reaction times at different temperatures, as indicated. Reaction temperature: (a) 400, (b) 350 °C. (c), (d) Sheet resistance (Rs) and haze of nickel silicide thin films as a function of SiH4 reaction times at (c) 400 and (d) 350 °C. (e)−(h) SEM morphologies of NiSi thin films synthesized using various SiH4 reaction times with Ni at different reaction temperatures, as indicated. In all cases, the chamber pressure was 1.0 Torr.

film surface.3,36 Due to the very rapid growth rate of the NiSi2 nuclei at the high silicidation temperature required to obtain NiSi2, coexistence of NiSi and NiSi2 is not commonly observed.38,41 By contrast, in the catalyzed chemical vapor process described here, the transition from NiSi to NiSi2 is much more gradual (in both temperature and time), and a large coexistence window exists. The presence of local NiSi regions indicates that the transition from NiSi to NiSi2 also occurs in a localized way, presumably limited by nucleation, rather than uniformly converting NiSi layer-by-layer. We speculate that the coexistence of NiSi and NiSi2 stems from a slow growth rate of the NiSi2 nuclei due to the lower silicidation temperatures employed in the catalyzed chemical vapor process. The islandlike NiSi2 formation in combination with the volume expansion leads to film NiSi2 roughening in a very similar way as for solid-state reactions. However, it should be mentioned that, under the studied conditions, we did not observe any agglomeration of NiSi or NiSi2 on SiO2 surfaces and films were always continuous, as shown by the TEM images in Figure 5. Figure 8 shows the impact of the SiH4 exposure time on the phase, the Rs, and the morphology of nickel silicide films synthesized at temperatures of 400 and 350 °C. GIXRD (Figures 8a and b) indicates that short exposures of 15−30 s were insufficient to observe any silicidation reactions at such temperatures. For all longer exposures (up to 300 s), only NiSi was found at both studied temperatures. No GIXRD reflections that can be attributed to any other (e.g., Ni- or Si-rich) NixSiy phase were observed.

This indicates that varying the SiH4 exposure time did not lead to the observation of additional phases (e.g., Ni-or Si-rich silicides). At lower temperatures, long SiH4 exposures were also not able to induce the phase transformation to NiSi2 that was observed at 500 °C. The formation of the monosilicide thus appeared self-limiting. This indicates that the solubility of Si in NixSiy (or more precisely its temperature dependence) was the main factor that governed the phase transition. Due to the limited solubility of Si in Ni at lower temperatures, an additional supply of SiH4 in the form of longer exposure times did not have much impact due to saturation of the Si content in the films at 350 or 400 °C.3,5,22 Figures 8c and d show the evolution of both Rs and haze with SiH4 exposure time both at 400 and 350 °C. The Rs initially remained high (∼17 Ω/□) up to 30 s of SiH4 exposure but subsequently saturated at lower values for both temperatures. By contrast, the haze showed the opposite trend. The decrease in Rs (and increase in haze) indicated the onset of the silicidation reaction. Once NiSi was formed, Rs and haze were hardly influenced by the SiH4 exposure time (up to 300 s) at temperatures of 350−400 °C. This shows again that the silicidation reaction was self-terminating under these conditions. Figures 8e−h show top-view SEM morphologies of the NiSi films grown using different SiH4 exposure times at temperatures of 400 and 350 °C. The morphology appeared to be rather independent of the SiH4 exposure time and temperature indicating a broad process window for the formation of NiSi films. This was further substantiated by the AFM and RBS G

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analysis in Figure 9. Both the rms surface roughness and the Ni/Si stoichiometric ratio were found to be independent of the SiH4 exposure time.

Figure 9. AFM rms roughness and RBS Ni/Si stoichiometric ratio as a function of SiH4 exposure time at 350 °C. The chamber pressure was 1.0 Torr.

Although varying the SiH4 reaction time (15−300 s) in the temperature range of 350−500 °C resulted either in NiSi or NiSi2 phases, similar studies at 325 °C (Figure 10a) showed the presence of Ni-rich silicide (Ni3Si2) as a transient minority phase that coexisted together with the dominant NiSi phase (60−180 s). The minority phase completely disappeared when the reaction time was increased to 400 s improving both the film uniformity and sheet resistance (Figure 10b). At higher temperatures (500 °C), the transformation of Ni3Si2 into NiSi was sufficiently fast that even 15 s of SiH4 exposure were sufficient for the direct formation of NiSi, while in the lower temperature range 325−400 °C at least 60 s was required. The phase sequence in solid-state routes has been studied in great detail and can be considered to be well understood. Initially, Ni2Si is formed at temperatures around 200 °C, which, after reaching a certain thickness, transforms into NiSi at temperatures of 350−400 °C. Both growth processes are limited by diffusion. The NiSi phase then remains stable up to 700−800 °C (or agglomerates depending on the nature of the surface) followed by the rapid (“explosive”) formation of NiSi2 driven by nucleation.3,36,42,43 By contrast, the catalyzed chemical vapor process described here leads to a different phase formation sequence (Ni3Si2 → NiSi → NiSi2) as well as to the coexistence of NiSi and NiSi2 at low temperature (∼450 °C). This clearly illustrates that the silicide formation is not only determined by (Ni) diffusion but also that the reaction rate of SiH4 drives the kinetics of the phase formation and therefore influences the reaction products. The above results are in marked contrast with previous reaction studies on Ni thin films with SiH4.17,21,22 In these studies, drastic morphological changesin particular the formation of nanowires with very large aspect ratioshave been observed for certain process conditions that fall within the parameter range investigated here.21,22 In our studies, we have not observed any traces of nanowire formation after the exposure of Ni films to SiH4. To further mimic the process conditions of ref 21, we have studied also very long reaction times up to 45 min in the temperature interval of 350−500 °C (Figures 11a−c). In all cases, rough and discontinuous films were obtained, however, without the formation of any clear

Figure 10. (a) GIXRD patterns (offset for clarity) as well as (b) sheet resistance (Rs) and haze values as a function of SiH4 reaction time (60−300 s) at 325 °C The chamber pressure was 1.0 Torr.

nanowires or whiskers. In addition, thicker Ni films (80 nm) and higher SiH4 partial pressures (9.0 Torr) also resulted in large-crystallite-like morphologies (Figures 11d and e) again without any signs of nanowire growth, even for 45 min of SiH4 exposure. Although we could not explore the higher SiH4 pressures used in ref 21 as they are incompatible with our reactor chamber, the lower SiH4 pressures (lower supersaturation) used in this study should even foster nanowire growth further.21,23 Moreover, as mentioned above, the oxidation of the surface of the Ni films did fully impede any silicidation reactions, as expected and in marked contrast to the report in ref 22. These results suggest that the formation of silicide nanowires is not a general feature of such vapor−solid silicidation processes and that the proposed formation mechanisms21−23 of Ni silicide nanowire growth have to be reconsidered. We note that we did observe some whiskerlike structures at low density (Figure 11f) by SiH4 exposure of Ni nanoparticles that were formed by annealing Ni films at 900 °C in forming gas, in keeping with previous reports.19 Hence, it is possible that excessive roughness of the initial Ni film is a key prerequisite for nanowire formation. This will require further study.

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CONCLUSIONS We have studied the formation of nickel silicides during the catalytic chemical vapor reaction of thin Ni films with SiH4. In H

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Figure 11. SEM morphologies of nickel silicide thin films synthesized by reacting 10 nm Ni with SiH4 at 1.0 Torr for 45 min at different temperatures: (a) 350, (b) 400, and (c) 500 °C. (d), (e) Nickel silicide films grown using 80 nm Ni at 9.0 Torr SiH4 pressure for 45 min at 400 and 450 °C, respectively. (f) Whiskerlike structures (“nanowires”) were obtained after exposing Ni droplets (nanoclusters) at 400 °C for 45 min to 1.0 Torr SiH4. The inset shows the SEM of Ni droplets before reaction. The droplets were obtained by annealing 10 nm Ni films at 900 °C for 5 min in forming gas.

contrast to solid-state salicidation processes, these solid−vapor silicidation reactions already occurred at temperatures as low as 350 °C. The phase and the morphology were characterized as a function of different parameters such as silicidation temperature, the SiH4 partial pressure, and the exposure time. The formed phases, NiSi and NiSi2, were found to depend on the reaction temperature with NiSi2 requiring higher temperatures. A dependence on the SiH4 exposure time was only observed at the highest studied temperature (500 °C). At lower temperatures, the process was self-limiting. The low-temperature synthesis of NiSi and NiSi2 films led to resistivity values that were comparable to thin films synthesized by high-temperature solid-state annealing routes, indicating the high quality of the silicide films. In contrast to earlier reports,21−23 we have not observed the formation of silicide nanowires during the silicidation of regular Ni films. This was the case despite reproducing the experimental conditions described in the references as accurately as possible. This indicates that the formation of nickel silicide nanowires is not necessarily intrinsic to the solid−vapor process and may have an extrinsic origin. In particular, we could not confirm the role of a thick NiO layer on the surface in the nanowire formation, as brought forward in ref 22. By contrast, we have observed a complete suppression of the silicidation process by a thick continuous surface oxide, by the suppression either of catalytic decomposition of SiH4 or of the Ni and Si diffusion across the surface layer. We note that we could find some signs of nanowire (whisker) formation at low densities by exposing Ni nanoclusters (“droplets”) formed by high temperature annealing of Ni films, in keeping with ref 19.

This indicates that the underlying formation mechanisms of silicide nanowire formation have to be reconsidered and more work is necessary to obtain silicide nanowires with good control over phase, properties, and morphology, as required by potential nanoelectronic applications.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ABBREVIATIONS AFM, atomic force microscopy; CVD, chemical vapor deposition; CMOS, complementary metal−oxide−semiconductor; EDS, energy dispersive spectroscopy; EELS, electron energy loss spectroscopy; GIXRD, grazing incidence X-ray diffraction; HAADF-STEM, high-angle annular dark field scanning transmission electron microscopy; PECVD, plasmaenhanced chemical vapor deposition; PVD, physical vapor deposition; RBS, Rutherford backscattering spectrometry; Rs, sheet resistance; SEM, scanning electron microscopy; TEM, transmission electron microscopy; XRR, X-ray reflectivity

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REFERENCES

(1) Deckers, C. A.; Solanki, R.; Freeouf, J. L.; Carrythers, J. R.; Evans, D. R. Appl. Phys. Lett. 2004, 84, 1389. (2) Kim, J.; Anderson, W. A. Thin Solid Films 2005, 483, 60. (3) Zhao, F. F.; Zhang, J. Z.; Shen, Z. X.; Ospowicz, T.; Gao, W. Z.; Chan, L. H. Microelectron. Eng. 2004, 71, 104.

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(4) Kittl, J. A.; Opsomer, K.; Torregiani, C.; Demeurisse, C.; Mertens, S.; Brunco, D. P.; Van Dal, M. J. H.; Lauwers, A. Mater. Sci. Eng., B 2008, 154−155, 144. (5) Lee, C. Y.; Lu, M. P.; Liao, K. F.; Lee, W. F.; Huang, C.; Chen, S. Y.; Chen, L. J. J. Phys. Chem. C 2009, 113, 2286. (6) Yan, X. Q.; Yuan, H. J.; Wang, J. X.; Liu, D. F.; Zhou, Z. P.; Gao, Y.; Song, L.; Liu, L. F.; Zhou, W. Y.; Wang, G.; Xie, S. S. Appl. Phys. A: Mater. Sci. Process. 2004, 79, 1853. (7) Chiu, W. L.; Chiu, C. H.; Chen, J. Y.; Huang, C. W.; Huang, Y. T.; Lu, K. C.; Hsin, C. L.; Yeh, P. H.; Wu, W. W. Nano Res. Lett. 2013, 8, 290. (8) Sullivan, J. P.; Tung, R. T.; Schrey, F. J. Appl. Phys. 1992, 72, 478. (9) Kamineni, V. K.; Raymond, M.; Bersch, E. J.; Doris, B. B.; Diebold, A. C. J. Appl. Phys. 2010, 107, 093525. (10) Gao, X.; Andersson, J.; Kubart, T.; Nyberg, T.; Smith, U.; Lu, J.; Hultman, L.; Kellock, A. J.; Zhang, Z.; Lavoie, C.; Zhang, S. L. Electrochem. Solid-State Lett. 2011, 104, H268. (11) Adusumilli, P.; Seidman, D. N.; Murray, C. E. J. Appl. Phys. 2012, 112, 064307. (12) Anisur, M. R.; Osipowicz, T.; Chi, D. Z.; Wang, W. D. J. Electron. Mater. 2005, 34, 1110. (13) Julies, B. A.; Knoessen, D.; Pretorius, R.; Adams, D. Thin Solid Films 1999, 347, 201. (14) Do, K. W.; Yang, C. M.; Kang, I. S.; Kim, K. M.; Back, K. H.; Cho, H. I.; Lee, H. B.; Kong, S. H.; Hahm, S. H.; Kwon, D. H.; Lee, J. H. Jpn. J. Appl. Phys. 2006, 45, 2975. (15) De Keyser, K.; Van Bockstael, C.; Meirhaeghe, R. L.; Detavernier, C.; Verleysen, E.; Bender, H.; Vandervorst, W.; Sweet, J. J.; Lavoie, C. Appl. Phys. Lett. 2010, 96, 173503. (16) Liu, Z.; Zhang, H.; Wang, L.; Yang, D. Nanotechnology 2008, 19, 375602. (17) Kim, J. Appl. Phys. Lett. 2012, 101, 233103. (18) Dong, L.; Bush, J.; Chirayos, V.; Solanki, R.; Jiao, J. Nano Lett. 2005, 5, 2112. (19) Iriarte, G. F. J. Non-Cryst. Solid. 2010, 356, 1135. (20) Kim, J.; Shin, D. H.; Lee, E. S.; Han, C. S.; Park, Y. C. Appl. Phys. Lett. 2007, 90, 253103. (21) Kim, C. J.; Kang, K.; Sung, Y.; Ryu, K. G.; Moon, H.; Kim, J. M.; Zang, D. S.; Jo, M. H. Adv. Mater. 2007, 19, 3637. (22) Kang, K.; Kim, S. K.; Kim, C. J.; Jo, M. H. Nano Lett. 2008, 8, 431. (23) Silicon and Silicide Nanowires; Huang, Y., Tu, K. N., Eds.; CRC Press, Taylor & Francis: Boca Raton, FL, 2013. (24) Fan, X.; Zhang, H.; Du, N.; Yang, D. Mater. Res. Bull. 2012, 47, 3797. (25) Arrazat, B.; Daniel, A.; Nolot, E.; Favier, S.; Carron, V.; Nemouchi, F.; Hartmann, J. M.; Rolland, G.; Vandelle, P.; Kermarrec, O.; Campidelli, Y.; Descombes, S. Microelectron. Eng. 2007, 84, 2558. (26) Dellas, N. S.; Abraham, M. J. Mater. Res. 2011, 17, 2282. (27) Lavoie, C.; Detavernier, C.; Beseeler, P. In Silicide Technology for Integrated Circuits; Chen, L. J., Ed.: The Institution of Electrical Engineers: London, 2004. (28) Properties of Metal Silicides;Maex, K., van Rossum, M., Eds.; The Institution of Engineering and Technology: Stevenage, England, 1995. (29) Dubois, L. H.; Nuzzo, R. G. Surf. Sci. 1985, 149, 133. (30) Deng, F.; Johnson, R. A.; Asbeck, P. M.; Lau, S. S.; Dubbelday, W. B.; Hsiao, T.; Woo, J. J. Appl. Phys. 1997, 81, 8047. (31) Nuzzo, R. G.; Dubois, L. H. Surf. Sci. 1984, 19, 407. (32) Deeb, C.; Heuer, A. H. Appl. Phys. Lett. 2004, 84, 1117. (33) Lee, P. S.; Mangelinck, D.; Pey, K. L.; Ding, J.; Dai, J. Y.; Ho, C. S.; See, A. Microelectron. Eng. 2000, 51−52, 583. (34) Dubois, L. H.; Nuzzo, R. G. J. Vac. Sci. Technol., A 1984, 2, 441. (35) d’Heurle, F. M.; Gas, P. J. Mater. Res. 1986, 1, 205. (36) Gambino, J. P.; Colgan, E. G. Mater. Chem. Phys. 1998, 52, 99. (37) Colgan, E. G.; Gambino, J. P.; Hong, Q. Z. Mater. Sci. Eng., R 1996, 16, 43. (38) Park, B. J.; Jeong, S. Y.; Kim, J. H.; Seong, T. Y.; Choi, C. J. J. Vac. Sci. Technol., A 2011, 29, 0315031. (39) Murarka, S. P. Intermetallics 1995, 3, 173.

(40) Jiang, Y.; Singh, N.; Tan, L. H.; Lo, G. Q.; Chan, D. S. H.; Kwong, D. L. IEEE Electron. Device Lett. 2009, 30, 195. (41) Lee, K. M.; Kim, C. Y.; Choi, C. K.; Yun, S. W.; Ha, J. B.; Lee, J. H.; Lee, J. Y. J. Korean Phys. Soc. 2009, 55, 1153. (42) Mangelinck, D.; Hoummada. Appl. Phys. Lett. 2008, 92, 254101. (43) Lee, P. S.; Pey, K. L.; Mangelinck, D.; Chi, D. Z.; Osipowicz, T. J. Electrochem. Soc. 2005, 152, G305.

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dx.doi.org/10.1021/cm503810p | Chem. Mater. XXXX, XXX, XXX−XXX