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Formation of Ni silicide from atomic layer deposited Ni
Current Applied Physics 16 (2016) 720e725
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Formation of Ni silicide from atomic layer deposited Ni Jaehong Yoon a, 1, Soo Hyeon Kim a, 1, Hangil Kim b, Soo-Hyun Kim b, Hyungjun Kim a, **, Han-Bo-Ram Lee c, * a
School of Electrical and Electronics Engineering, Yonsei University, Seoul, Republic of Korea School of Materials Science and Engineering, Yeungnam University, Gyeongsan, Republic of Korea c Department of Material Science Engineering, Incheon National University, Incheon, Republic of Korea b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 28 October 2015 Received in revised form 11 March 2016 Accepted 4 April 2016 Available online 6 April 2016
The silicidation of Ni deposited by plasma-enhanced atomic layer deposition with NH3 plasma and thermal atomic layer deposition using NH3 gas were comparatively studied. A SiNx interlayer was formed between the Ni deposited by plasma-enhanced atomic layer deposition and the Si substrate due to the direct exposure of the substrate to plasma, while no interlayer was observed when using thermal atomic layer deposition. In the plasma-enhanced atomic layer deposition, the diffusion of Ni was suppressed by the SiNx interlayer, so no Ni2Si phase was formed and its formation temperature increased. Ni formed by thermal atomic layer deposition showed sequential phase transformations to Ni2Si, NiSi, and NiSi2 with increased annealing temperatures. In the nanosized contact holes, a large amount of NiSi2 was formed due to the limited supply of Ni. These results provide important information for the fabrication of silicide in nanoscale 3D devices. © 2016 Elsevier B.V. All rights reserved.
Keywords: Vapor deposition Diffusion Phase transitions Nanostructures Electron microscopy
1. Introduction Metal silicides are used as important contact materials for complementary metal-oxide-semiconductor (CMOS) devices due to their low contact resistances and their compatibility with Si [1]. However, TiSi2 [2] and CoSi2 [3], which have been widely used as contacts, have technological limitations as such devices are continuously scaled down in size. The sheet resistance of TiSi2 steeply increases when the line width reaches 500 nm (known as the narrow line effect) since the low-resistivity TiSi2 phase (C54) cannot be formed within this narrow line during the annealing process [4]. Although CoSi2 is not significantly affected by the narrow line effect down to a line width of 100 nm, the greater consumption of Si during the formation of CoSi2 is another obstacle in these shallow junctions [5]. Thus, NiSi may be an alternative contact material since it has superior properties related to such applications; it is not affected much by the narrow line effect, requires low silicon consumption, and has a lower resistivity than
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (H.-B.-R. Lee). 1 These authors equally contributed this work. http://dx.doi.org/10.1016/j.cap.2016.04.005 1567-1739/© 2016 Elsevier B.V. All rights reserved.
(H.
Kim),
[email protected]
TiSi2 and CoSi2. In the typical fabrication of NiSi, metallic Ni films are deposited on Si and reacted with the Si by annealing through a solid-phase reaction. In general, three phases of Ni silicide (Ni2Si, NiSi, and NiSi2) are formed through solid phase reactions, and the NiSi phase is preferred as the contact material as it has low resistivity [3]. Physical vapor deposition (PVD) methods, such as sputtering, are typically used for Ni metal deposition. As devices are scaled down on the order of 101 nm, the structures of Si devices become more complex as they transition from 2D to 3D. This limits the applicability of conventional PVD techniques owing to their poor step coverage when used for nanoscale 3D structures [6]. Atomic layer deposition (ALD) is a promising deposition method for this purpose because it has excellent conformality in 3D structures, and it enables thickness control at the atomic level. Several studies regarding Ni ALD have been previously published. In an early study, Ni films were formed through multiple steps involving ALD NiO with a bis-(cyclopentadienyl)nickel precursor and water reactant, followed by H2 plasma reduction due to the lack of a suitable Ni precursor and reducing agent [7]. In another study, Ni films were obtained using a nickel-bis(1-dimethylamino2-methyl-2-butanolate) [Ni(dmamb)2] metal organic precursor and H2 gas reactant, and the resulting Ni films contained a large amount of carbon impurities and had a high resistivity [8]. Reactive NH3 gas was utilized as a counter reactant for the ALD Ni with a Ni(dmamb)2
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precursor, and the ALD Ni process was applied for area-selective deposition [9]. In addition, the same Ni(dmamb)2 precursor and NH3 plasma were used to obtain plasma-enhanced ALD (PE-ALD) Ni [10]. In general, because a plasma reactant has a higher reactivity than a gas reactant in many cases, PE-ALD shows some side effects from plasma exposure, as well as higher growth rates and a higher film density, compared to thermal ALD [11,12]. In our previous studies, we developed PE-ALD Co films using NH3 plasma and investigated silicide formation on these films [13e15]. In these studies, an amorphous SiNx interlayer was formed between the PE-ALD Co and the Si substrate due to the direct exposure of the Si substrate to NH3 plasma during the initial growth. During the annealing process, the SiNx interlayer interfered with the diffusion of Co into the Si substrate, resulting in the formation of epitaxial CoSi2. In the following report, a NH3 gas reactant was used instead of plasma. A SiNx interlayer was not observed, likely due to the lower reactivity of NH3 gas [9]. Similarly in our previous work, a SiNx interface layer was observed in the PE-ALD Ni film using a Ni(dmamb)2 precursor and NH3 plasma [10]. However, the effects of the SiNx interlayer on the silicide formation of Ni were not studied. Although the SiNx interlayer has some positive effects, such as acting as a diffusion control layer for the easy formation of epitaxial silicide [13,16], an interlayer-free ALD process is more desirable in order to provide a reliable silicide process. In addition, the silicide formation of ALD films inside nanoscale contact holes has not yet been reported. Since the process of silicide formation accompanies phase transformations and volume changes, an understanding of the formation of silicide that occurs inside confined 3D structures is vital. In this work, we conducted a comparative study on interface structures and the formation of Ni silicide in a PE-ALD Ni film prepared using NH3 plasma and Ni deposited by thermal ALD using NH3 gas. The interlayer formations of Ni on the 2D planar Si substrates were investigated in both deposition schemes, and the results were correlated with the phase transformation that occurred during the annealing process. Based on the results using a 2D planar structure, Ni was then deposited by thermal ALD in high aspect ratio Si hole patterns and annealed at various temperatures to investigate the formation of Ni silicide inside confined 3D structures. 2. Materials and methods A commercial ALD chamber (CN1 ATOMIC PREMIUM) with a loadlock was used. Reaction chamber has double shower head for supplying chemicals separately and uniformly. Nickel bis(1dimethylamino-2-methyl-2-butanolate) (Ni(dmamb)2) was used as the precursor, and NH3 gas and NH3 plasma were used as the reactants. For PE-ALD, a direct RF plasma system was equipped to the ALD chamber. The shower head was connected to RF plasma generator and wafer stage was grounded, so plasma was generated between the shower head and substrate. The precursor was kept in a glass bubbler maintained at 70 C to obtain a suitable vapor pressure. One ALD cycle was composed of four steps: precursor exposure (carried out for a duration ts), purging (carried out for a duration tp), reactant exposure (carried out for a duration tr), and purging. The flow of the Ar purging gas was 50 sccm (standard cubic centimeters per minute). When NH3 plasma was used as a reactant, ts, tr, and tp were 5 s, 5 s, and 8 s, respectively. NH3 gas was introduced into the chamber at a flow rate of 200 sccm and plasma was generated between the substrate and the showerhead using a plasma power of 300 W. For the case of NH3 gas, ts, tr, and tp were 8 s, 5 s, and 8 s, respectively. The deposition temperatures for the PE-ALD and thermal ALD were set to 200 C and 300 C, respectively. More detailed information on the experimental setup for the
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Ni ALD can be found in our previous reports [9,10]. Si(100) was used as the substrate for the silicide formation. Substrates were cleaned by immersion in a commercial buffered oxide etchant (6:1 volume ratio of 40% NH4F in water: 49% HF in water) for 10 s to remove native oxides. This was followed by rinsing with deionized (DI) water and blowing with N2, resulting in the formation of H-terminated Si. After cleaning, the Si(100) substrates were immediately loaded into the test chamber to prevent additional oxide formation. In addition to the planar Si substrate, Si substrates patterned with nanosized holes were also used, with holes 30 nm in diameter with aspect ratios of 5:1. The same cleaning steps were repeated to remove native oxide formed around the bottom region of the hole patterns. For the annealing experiments, a 30-nm-thick Ti capping layer was deposited ex situ by sputtering on the 20 nm Ni ALD films. The Ni films capped with Ti layers were annealed by rapid thermal annealing (RTA) at various annealing temperatures (Ta) from 400 to 700 C for 1 min under N2. RTA experiments were performed under vacuum conditions to prevent oxygen contamination, which negatively impacts silicide formation. The contact resistance of the NiSi due to the thermal ALD Ni was evaluated using the transmission line method (TLM). TLM patterns were fabricated by general photolithography and etching processes. Schematics of these processes and a top view optical microscopy image of the NiSi lines are presented in Fig. 1(a) and (b), respectively. A 10 nm Al2O3 film was deposited by ALD using trimethylaluminum (TMA) and H2O on a Si substrate as a diffusion mask. A photoresist (PR) film was coated on the Al2O3 and patterned by photolithography, after which the substrate was immersed in a diluted HF solution to open the line patterns. The distances between each PR line pattern were different, with distances of 50, 80, and 120 mm. Ni-thermal ALD and sputtering of the Ti cap were performed on the patterns. After RTA, the unreacted Ni and Ti capping layers were removed by selective etching and the Al2O3 diffusion mask was also removed using a HF solution. The resistance between the NiSi lines were measured using a probe station. Field emission scanning electron microscopy (FE-SEM, JEOL7001F) was used to analyze the thickness and surface morphology of the Ni films. The resistances of the ALD Ni films and the Ni silicide were measured using a 4-point probe, and the film resistivity was calculated from the analyzed thickness. To measure the resistance of only the Ni silicides, the unreacted Ni and Ti capping layers were removed by selective etching in a piranha solution (1:1 ¼ 30% H2O2 in DI water:97% H2SO4 in DI water) for 10 min. The phase transformation from metallic Ni to Ni silicides was analyzed by XRD (RIGAKU, Ultima IV). For the nanoscale structural and chemical composition analysis, high-resolution TEM (HRTEM) and scanning TEM (STEM) images were acquired using a TECNAI G2 F30 transmission electron microscope (300 kV accelerating voltage and a Schottky-type field emission electron gun) equipped with energy dispersive spectroscopy (EDS) analysis. EDS analysis was done in the STEM mode and the electron probe size used in the STEM-EDS mapping experiments was approximately 0.3 nm. 3. Results and discussion Fig. 2(a) and 2(b) show the cross-sectional view HRTEM images of the PE-ALD Ni formed using NH3 plasma and the thermal ALD Ni using NH3 gas, respectively. In Fig. 2(a), an interlayer having a brighter contrast than the Ni film and the Si substrate is observed in the PE-ALD Ni film, with a thickness of approximately 5 nm. The EDS line profile of the PE-ALD Ni film in Fig. 2(c) shows that N and O are localized in the interlayer, indicating the formation of SiOxNy. Consistent with our previous results, a nitride interlayer is formed
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Fig. 1. (a) Schematic of the fabrication process for the NiSi TLM and (b) a top-view optical image of the NiSi TLM patterns.
Fig. 2. Cross-sectional view HRTEM images of the (a) PE-ALD and (b) thermal ALD Ni films on Si(100). EDS line profiles of the (c) PE-ALD and (d) thermal ALD Ni films from region a’ and region b’, respectively.
due to the direct exposure of the Si substrate to the NH3 plasma reactant [10,13e15]. N was also detected in the region of the PE-ALD Ni film, consistent with X-ray photoelectron spectroscopy (XPS) results (see Fig. S1 in the Supplemental Information), likely due to
the higher reactivity of the NH3 plasma compared to the NH3 gas. The O signal most likely arises from residual oxygen present in native oxides. In our early report on PE-ALD Co, similar oxygen peak was detected in the interlayer [15]. By contrast, although a white
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contrast interlayer was formed during the thermal ALD Ni, it is very thin (<1 nm) as shown in Fig. 2(b). O and N signals are detected across the sample at background levels in the Ni deposited by thermal ALD, as shown in Fig. 2(d). The reactivity of the NH3 gas is lower than that of the NH3 plasma. For instance, SiNx is formed even at 300 C when NH3 plasma is used. SiNx is also formed by using NH3 gas at 300 C, but it has a negligible thickness (<1 nm) [17]. Therefore, in contrast to the PE-ALD Ni film, no nitride interlayer was formed in thermal ALD Ni, even though the Si substrate was exposed to the NH3 gas reactant during the initial growth. In addition, it is likely that the very thin interlayer between thermal ALD Ni and the Si substrate is composed of the native silicon oxide layer. Although native oxide etching is performed prior to Ni thermal ALD, a thin oxide layer could be formed again during the sample transfer process [18]. However, since we used 0.3 nm ebeam size for TEM analysis, oxygen in a very narrow interlayer region could not be detected. To investigate the formation of Ni silicide, the PE-ALD and thermal ALD Ni films on Si substrates were annealed at various temperatures and analyzed by X-ray diffraction (XRD). Because the SiOxNy interlayer is amorphous (Fig. 2(a)), no XRD peaks were observed. For the PE-ALD Ni film, nickel mono-silicide phases were observed at 44.8 for NiSi(210) [19] below 700 C as shown in Fig. 3(a). When the annealing temperature (Ta) was increased to 700 C, a strong XRD peak was detected at 33.27, corresponding to the NiSi2(200) [20] phase, while the NiSi(210) peak was not observed. The high intensity peak at 33 is a forbidden diffraction peak of the Si substrate, which is often observed in single crystal Si substrates [21]. In contrast to the PE-ALD Ni film, a different phase transformation was detected in the XRD spectra of the annealed thermal ALD Ni films. At Ta ¼ 400 C, a strong XRD peak was observed near 45 as shown in Fig. 3(b). The peak was deconvoluted to two peaks at 44.8 and 45.1, corresponding to NiSi(210) and Ni2Si(202) [22], respectively. At Ta ¼ 500 C, this strong peak shifted to a slightly lower diffraction angle, because the Ni2Si(202) peak at 45.1 was diminished while the NiSi(210) peak at 44.8 was not. Related to this, the small Ni2Si(112) [22] peak observed at 39.6 in the 400 C annealed sample was not observed at 500 C. The intensity of the NiSi(210) peak decreased, and eventually disappeared at 700 C. In addition, a small peak was detected at 29.2 at 500 C, assigned to NiSi2(111) [23], and the intensity of the peak increased slightly until 700 C. Compared to PE-ALD Ni, thermal ALD Ni resulted in different phase transformations as the annealing temperature increased. Ni deposited by PE-ALD was transformed to NiSi at 400 C without the formation of metal-rich phases such as Ni2Si, and NiSi2 was formed above 700 C. By contrast, NiSi and Ni2Si phases co-existed in the Ni deposited by thermal ALD that was annealed at 400 C, and only the Ni2Si phase disappeared above that temperature. The NiSi disappeared above 600 C, and the NiSi2
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that was detected at 500 C remained until 700 C. Consistently, the sheet resistance of Ni silicide from thermal ALD Ni showed the lowest value at 500 C. After annealing at 500 C, the sheet resistance of Ni silicides from thermal ALD Ni was 610 U/sq. and the resistivity of Ni silicide calculated from the resistance and thickness of Ni silicide was 580 mUcm. Since the NiSi phase has the lowest resistivity in three Ni silicide phases, the lowest sheet resistance of Ni silicide from ALD Ni at 500 C indicates the formation of the NiSi phase. It is known that Ni is a dominant diffusion species during the formation of Ni2Si, NiSi, and NiSi2 [24], and these three silicide phases are usually formed sequentially at 250 C, 350e500 C, and 700 C, respectively [1,25e27]. In the case of the PE-ALD Ni, the Ni atoms cannot easily diffuse into Si through the SiNx interlayer since the diffusion coefficient of Ni in SiNx is 2.28 1014 cm2/s at 650 C, much smaller than that of the Si substrate (6 104 cm2/s at 220e540 C) [28,29]. Moreover, the diffusivity of the Ni in the SiO2 was 1010e1014 cm2/s at 500e700 C, comparable to that of the SiNx layer. Thus, the oxygen species in the SiNx interlayer also interfered with the diffusion of Ni [30]. In addition, according to the NieN phase diagram, the formation of nickel nitride under these conditions is difficult and was not detected using XRD and EDS analyses. The SiNx interlayer is only a barrier for Ni diffusion but does not react with Ni, and as such the diffusion of Ni into the Si substrate is suppressed by the SiNx interlayer. As a result, a metalrich phase is not formed at 400 C and the transformation temperature of NiSi2 is higher than that from sputtered Ni, which does not have a SiNx interlayer. No SiNx interlayer is formed in the thermal ALD Ni, allowing the Ni atoms to more easily diffuse into the Si substrate which results in the sequential formation of three silicide phases as the annealing temperature is increased. The formation of the interfacial layer and phase transformation of the thermal and PE-ALD Ni are summarized in Table S1 in the Supplemental Information. The contact resistance of the Ni silicide obtained from the thermal ALD Ni was evaluated on a planar Si substrate using TLM (fabrication of the TLM line patterns was described previously in Fig. 1(a)). The patterned Ni silicide lines were used for transmission lines. After TLM pattern fabrication, the formation of NiSi was clearly observed in the optical microscopy image shown in Fig. 1(b). The resistance between the NiSi lines increased as the distance between lines increased, as shown in Fig. 4. The intercept on the yaxis that was obtained by linear fitting with three resistance points was 0.48 U and the calculated contact resistance was 0.24 U. This contact resistance is comparable to the typical value of the contact resistance for NiSi that is obtained from PVD Ni [31], indicating that thermal ALD Ni can be used as an alternative contact fabrication process to PVD Ni. We also investigated the silicide formation of thermal ALD Ni in
Fig. 3. XRD results of the (a) PE-ALD and (b) thermal ALD Ni films annealed from 400 C to 700 C.
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Fig. 4. Resistance versus distance between the NiSi lines for TLM using silicide formation of the Ni thermal ALD.
nanosized contact holes. As mentioned in Section 1, understanding the Ni silicide formation inside confined nanostructures is important because it involves both a phase transformation and a volume change. Ni was deposited for 230 cycles to form a 12-nm-thick Ni film, compared to the 25 nm diameter of the contact holes. After thermal ALD Ni, a Ti capping layer was deposited to prevent oxidation. After annealing at 500 C, silicide formation of the thermal ALD Ni in the nanosized contact holes was analyzed by HRTEM as shown in Fig. 5. In Fig. 5(a), the TEM image clearly shows that the Ni is conformally deposited by thermal ALD inside the nanosized contact holes. Fig. 5(b) shows that a void is formed at the bottom of the holes, while the deposited Ni in the middle and top regions remains intact. In addition, a dark contrast crystal facet is formed in the Si substrate region below the bottom of the hole.
Fig. 5(c) shows the EDS line profile along the white arrow. In the facet, the Ni signal gradually decreased along this direction to the Si substrate while the Si signal increased. Based on the compositional ratio of the nickel silicide phases, the region near the bottom of the hole (region a’) is occupied by the NiSi phase, and the region closer to the Si substrate (region b’) is occupied by the NiSi2 phase. The two regions in the facet were mathematically converted to diffraction patterns using fast Fourier transform (FFT) and the results are shown in Fig. 5(d). The FFT patterns from region a’ (close to the hole) show the existence of NiSi(210) and NiSi(200). In region b’ (close to the Si substrate), diffraction patterns of the NiSi2(111) phase were observed. NiSi(210) and NiSi2(111) crystal planes were also consistently observed in our XRD results after annealing at 500 C (Fig. 3(b)). In addition, the STEM high-angle annular darkfield (HAADF) image in Fig. 5(e) clearly shows the contrast differences within the facet region. Because the contrast in the STEM HAADF image is proportional to the atomic number [32], region a’ and region b’ are composed of NiSi and NiSi2, respectively. The XRD results shown in Fig. 3(b) indicate that the NiSi peak is much stronger than the NiSi2 peak in the case of thermal ALD followed by annealing at 500 C. It was reported that Ni films are mainly transformed to the NiSi phase rather than the NiSi2 phase after annealing at 500 C, and the transformation to the NiSi2 phase requires higher annealing temperatures [26]. As such, the Ni silicide from the Ni deposited by thermal ALD and then annealed at 500 C is primarily composed of NiSi with a small amount of NiSi2 also present. However, the volume of the NiSi2 phase is larger than that of the NiSi phase in the Ni silicide formed underneath the contact hole as shown in Fig. 5(d), which is inconsistent with the XRD results. In the planar sample, the whole Ni film is in contact with the Si substrate, while in the contact hole, Ni is only supplied from the bottom of each hole. As a result, it is likely that the supply of Ni atoms is more limited in the contact hole than in the planar substrate. At 500 C, the Ni atoms have enough thermal energy to diffuse into the Si substrate and form NiSi in the planar system.
Fig. 5. HRTEM images of the Ni deposited by thermal ALD in nanosized contact holes (a) before and (b) after annealing at 500 C. (c) EDS line profile along the arrow indicated in (b). (d) Magnified TEM image from (b) and FFT patterns collected from regions a’ and b’. (e) HAADF image of (d).
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Therefore, the formation of the NiSi2 phase in region b’ can be attributed to the geometrical limitations of Ni. 4. Conclusions We investigated the formation of silicide from Ni films prepared using PE- and thermal ALD using NH3 plasma and NH3 gas, respectively. A SiNx interlayer was formed in the PE-ALD Ni film due to direct exposure of the Si substrate to NH3 plasma, while no nitride interlayer was observed in the case of thermal ALD. Because the SiNx interlayer suppressed Ni diffusion, the Ni2Si phase was not formed and the formation temperature of NiSi2 increased. In contrast to PE-ALD, the Ni deposited by thermal ALD sequentially transformed to Ni2Si, NiSi, and NiSi2 as the annealing temperature increased. Using TLM, the contact resistance of the NiSi obtained from the thermal ALD Ni film was 0.24 U. The silicide formation of the Ni deposited by thermal ALD in the nanosized contact holes was different from that of the planar substrates. Inconsistent with the XRD results, a larger amount of NiSi2 was observed compared to NiSi. The Si-rich silicide phase (NiSi2) was formed at the bottom of the contact holes due to the limited supply of Ni atoms inside the nanostructures. The results from this study can provide important information for applications that require Ni deposited by thermal ALD for the fabrication of Ni silicide inside nanostructures. Acknowledgements This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education (2014R1A1A2059845), and by the MOTIE (Ministry of Trade, Industry & Energy) (10053098) and KSRC (Korea Semiconductor Research Consortium) support program for the development of the future semiconductor device. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.cap.2016.04.005. References [1] C. Lavoie, F.M. d'Heurle, C. Detavernier, C. Cabral Jr., Towards implementation of a nickel silicide process for CMOS technologies, Microelectron. Eng. 70 (2003) 144e157. [2] S.L. Zhang, U. Smith, Self-aligned silicides for Ohmic contacts in complementary metaleoxideesemiconductor technology: TiSi2, CoSi2, and NiSi, J. Vac. Sci. Technol. A 22 (2004) 1361e1370. € [3] S.L. Zhang, M. Ostling, Metal silicides in CMOS technology: past, present, and future trends, Crit. Rev. Solid State Mater. Sci. 28 (2003) 1e129. [4] J. Chen, J.-P. Colinge, D. Flandre, R. Gillon, J.P. Raskin, D. Vanhoenacker, Comparison of TiSi2, CoSi2, and NiSi for thin-film silicon-on-insulator applications, J. Electrochem. Soc. 144 (1997) 2437e2442. [5] S.P. Murarka, D.B. Fraser, A.K. Sinha, H.J. Levinstein, E.J. Lloyd, R. Liu, D.S. Williams, S.J. Hillenius, Self-aligned cobalt disilicide for gate and interconnection and contacts to shallow junctions, IEEE Trans. Electron Devices 34 (1987) 2108e2115.
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Formation of Ni silicide from atomic layer deposited Ni Jaehong Yoon a, 1, Soo Hyeon Kim a, 1, Hangil Kim b, Soo-Hyun Kim b, Hyungjun Kim a, **, Han-Bo-Ram Lee c, * a
School of Electrical and Electronics Engineering, Yonsei University, Seoul, Republic of Korea School of Materials Science and Engineering, Yeungnam University, Gyeongsan, Republic of Korea c Department of Material Science Engineering, Incheon National University, Incheon, Republic of Korea b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 28 October 2015 Received in revised form 11 March 2016 Accepted 4 April 2016 Available online 6 April 2016
The silicidation of Ni deposited by plasma-enhanced atomic layer deposition with NH3 plasma and thermal atomic layer deposition using NH3 gas were comparatively studied. A SiNx interlayer was formed between the Ni deposited by plasma-enhanced atomic layer deposition and the Si substrate due to the direct exposure of the substrate to plasma, while no interlayer was observed when using thermal atomic layer deposition. In the plasma-enhanced atomic layer deposition, the diffusion of Ni was suppressed by the SiNx interlayer, so no Ni2Si phase was formed and its formation temperature increased. Ni formed by thermal atomic layer deposition showed sequential phase transformations to Ni2Si, NiSi, and NiSi2 with increased annealing temperatures. In the nanosized contact holes, a large amount of NiSi2 was formed due to the limited supply of Ni. These results provide important information for the fabrication of silicide in nanoscale 3D devices. © 2016 Elsevier B.V. All rights reserved.
Keywords: Vapor deposition Diffusion Phase transitions Nanostructures Electron microscopy
1. Introduction Metal silicides are used as important contact materials for complementary metal-oxide-semiconductor (CMOS) devices due to their low contact resistances and their compatibility with Si [1]. However, TiSi2 [2] and CoSi2 [3], which have been widely used as contacts, have technological limitations as such devices are continuously scaled down in size. The sheet resistance of TiSi2 steeply increases when the line width reaches 500 nm (known as the narrow line effect) since the low-resistivity TiSi2 phase (C54) cannot be formed within this narrow line during the annealing process [4]. Although CoSi2 is not significantly affected by the narrow line effect down to a line width of 100 nm, the greater consumption of Si during the formation of CoSi2 is another obstacle in these shallow junctions [5]. Thus, NiSi may be an alternative contact material since it has superior properties related to such applications; it is not affected much by the narrow line effect, requires low silicon consumption, and has a lower resistivity than
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (H.-B.-R. Lee). 1 These authors equally contributed this work. http://dx.doi.org/10.1016/j.cap.2016.04.005 1567-1739/© 2016 Elsevier B.V. All rights reserved.
(H.
Kim),
[email protected]
TiSi2 and CoSi2. In the typical fabrication of NiSi, metallic Ni films are deposited on Si and reacted with the Si by annealing through a solid-phase reaction. In general, three phases of Ni silicide (Ni2Si, NiSi, and NiSi2) are formed through solid phase reactions, and the NiSi phase is preferred as the contact material as it has low resistivity [3]. Physical vapor deposition (PVD) methods, such as sputtering, are typically used for Ni metal deposition. As devices are scaled down on the order of 101 nm, the structures of Si devices become more complex as they transition from 2D to 3D. This limits the applicability of conventional PVD techniques owing to their poor step coverage when used for nanoscale 3D structures [6]. Atomic layer deposition (ALD) is a promising deposition method for this purpose because it has excellent conformality in 3D structures, and it enables thickness control at the atomic level. Several studies regarding Ni ALD have been previously published. In an early study, Ni films were formed through multiple steps involving ALD NiO with a bis-(cyclopentadienyl)nickel precursor and water reactant, followed by H2 plasma reduction due to the lack of a suitable Ni precursor and reducing agent [7]. In another study, Ni films were obtained using a nickel-bis(1-dimethylamino2-methyl-2-butanolate) [Ni(dmamb)2] metal organic precursor and H2 gas reactant, and the resulting Ni films contained a large amount of carbon impurities and had a high resistivity [8]. Reactive NH3 gas was utilized as a counter reactant for the ALD Ni with a Ni(dmamb)2
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precursor, and the ALD Ni process was applied for area-selective deposition [9]. In addition, the same Ni(dmamb)2 precursor and NH3 plasma were used to obtain plasma-enhanced ALD (PE-ALD) Ni [10]. In general, because a plasma reactant has a higher reactivity than a gas reactant in many cases, PE-ALD shows some side effects from plasma exposure, as well as higher growth rates and a higher film density, compared to thermal ALD [11,12]. In our previous studies, we developed PE-ALD Co films using NH3 plasma and investigated silicide formation on these films [13e15]. In these studies, an amorphous SiNx interlayer was formed between the PE-ALD Co and the Si substrate due to the direct exposure of the Si substrate to NH3 plasma during the initial growth. During the annealing process, the SiNx interlayer interfered with the diffusion of Co into the Si substrate, resulting in the formation of epitaxial CoSi2. In the following report, a NH3 gas reactant was used instead of plasma. A SiNx interlayer was not observed, likely due to the lower reactivity of NH3 gas [9]. Similarly in our previous work, a SiNx interface layer was observed in the PE-ALD Ni film using a Ni(dmamb)2 precursor and NH3 plasma [10]. However, the effects of the SiNx interlayer on the silicide formation of Ni were not studied. Although the SiNx interlayer has some positive effects, such as acting as a diffusion control layer for the easy formation of epitaxial silicide [13,16], an interlayer-free ALD process is more desirable in order to provide a reliable silicide process. In addition, the silicide formation of ALD films inside nanoscale contact holes has not yet been reported. Since the process of silicide formation accompanies phase transformations and volume changes, an understanding of the formation of silicide that occurs inside confined 3D structures is vital. In this work, we conducted a comparative study on interface structures and the formation of Ni silicide in a PE-ALD Ni film prepared using NH3 plasma and Ni deposited by thermal ALD using NH3 gas. The interlayer formations of Ni on the 2D planar Si substrates were investigated in both deposition schemes, and the results were correlated with the phase transformation that occurred during the annealing process. Based on the results using a 2D planar structure, Ni was then deposited by thermal ALD in high aspect ratio Si hole patterns and annealed at various temperatures to investigate the formation of Ni silicide inside confined 3D structures. 2. Materials and methods A commercial ALD chamber (CN1 ATOMIC PREMIUM) with a loadlock was used. Reaction chamber has double shower head for supplying chemicals separately and uniformly. Nickel bis(1dimethylamino-2-methyl-2-butanolate) (Ni(dmamb)2) was used as the precursor, and NH3 gas and NH3 plasma were used as the reactants. For PE-ALD, a direct RF plasma system was equipped to the ALD chamber. The shower head was connected to RF plasma generator and wafer stage was grounded, so plasma was generated between the shower head and substrate. The precursor was kept in a glass bubbler maintained at 70 C to obtain a suitable vapor pressure. One ALD cycle was composed of four steps: precursor exposure (carried out for a duration ts), purging (carried out for a duration tp), reactant exposure (carried out for a duration tr), and purging. The flow of the Ar purging gas was 50 sccm (standard cubic centimeters per minute). When NH3 plasma was used as a reactant, ts, tr, and tp were 5 s, 5 s, and 8 s, respectively. NH3 gas was introduced into the chamber at a flow rate of 200 sccm and plasma was generated between the substrate and the showerhead using a plasma power of 300 W. For the case of NH3 gas, ts, tr, and tp were 8 s, 5 s, and 8 s, respectively. The deposition temperatures for the PE-ALD and thermal ALD were set to 200 C and 300 C, respectively. More detailed information on the experimental setup for the
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Ni ALD can be found in our previous reports [9,10]. Si(100) was used as the substrate for the silicide formation. Substrates were cleaned by immersion in a commercial buffered oxide etchant (6:1 volume ratio of 40% NH4F in water: 49% HF in water) for 10 s to remove native oxides. This was followed by rinsing with deionized (DI) water and blowing with N2, resulting in the formation of H-terminated Si. After cleaning, the Si(100) substrates were immediately loaded into the test chamber to prevent additional oxide formation. In addition to the planar Si substrate, Si substrates patterned with nanosized holes were also used, with holes 30 nm in diameter with aspect ratios of 5:1. The same cleaning steps were repeated to remove native oxide formed around the bottom region of the hole patterns. For the annealing experiments, a 30-nm-thick Ti capping layer was deposited ex situ by sputtering on the 20 nm Ni ALD films. The Ni films capped with Ti layers were annealed by rapid thermal annealing (RTA) at various annealing temperatures (Ta) from 400 to 700 C for 1 min under N2. RTA experiments were performed under vacuum conditions to prevent oxygen contamination, which negatively impacts silicide formation. The contact resistance of the NiSi due to the thermal ALD Ni was evaluated using the transmission line method (TLM). TLM patterns were fabricated by general photolithography and etching processes. Schematics of these processes and a top view optical microscopy image of the NiSi lines are presented in Fig. 1(a) and (b), respectively. A 10 nm Al2O3 film was deposited by ALD using trimethylaluminum (TMA) and H2O on a Si substrate as a diffusion mask. A photoresist (PR) film was coated on the Al2O3 and patterned by photolithography, after which the substrate was immersed in a diluted HF solution to open the line patterns. The distances between each PR line pattern were different, with distances of 50, 80, and 120 mm. Ni-thermal ALD and sputtering of the Ti cap were performed on the patterns. After RTA, the unreacted Ni and Ti capping layers were removed by selective etching and the Al2O3 diffusion mask was also removed using a HF solution. The resistance between the NiSi lines were measured using a probe station. Field emission scanning electron microscopy (FE-SEM, JEOL7001F) was used to analyze the thickness and surface morphology of the Ni films. The resistances of the ALD Ni films and the Ni silicide were measured using a 4-point probe, and the film resistivity was calculated from the analyzed thickness. To measure the resistance of only the Ni silicides, the unreacted Ni and Ti capping layers were removed by selective etching in a piranha solution (1:1 ¼ 30% H2O2 in DI water:97% H2SO4 in DI water) for 10 min. The phase transformation from metallic Ni to Ni silicides was analyzed by XRD (RIGAKU, Ultima IV). For the nanoscale structural and chemical composition analysis, high-resolution TEM (HRTEM) and scanning TEM (STEM) images were acquired using a TECNAI G2 F30 transmission electron microscope (300 kV accelerating voltage and a Schottky-type field emission electron gun) equipped with energy dispersive spectroscopy (EDS) analysis. EDS analysis was done in the STEM mode and the electron probe size used in the STEM-EDS mapping experiments was approximately 0.3 nm. 3. Results and discussion Fig. 2(a) and 2(b) show the cross-sectional view HRTEM images of the PE-ALD Ni formed using NH3 plasma and the thermal ALD Ni using NH3 gas, respectively. In Fig. 2(a), an interlayer having a brighter contrast than the Ni film and the Si substrate is observed in the PE-ALD Ni film, with a thickness of approximately 5 nm. The EDS line profile of the PE-ALD Ni film in Fig. 2(c) shows that N and O are localized in the interlayer, indicating the formation of SiOxNy. Consistent with our previous results, a nitride interlayer is formed
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Fig. 1. (a) Schematic of the fabrication process for the NiSi TLM and (b) a top-view optical image of the NiSi TLM patterns.
Fig. 2. Cross-sectional view HRTEM images of the (a) PE-ALD and (b) thermal ALD Ni films on Si(100). EDS line profiles of the (c) PE-ALD and (d) thermal ALD Ni films from region a’ and region b’, respectively.
due to the direct exposure of the Si substrate to the NH3 plasma reactant [10,13e15]. N was also detected in the region of the PE-ALD Ni film, consistent with X-ray photoelectron spectroscopy (XPS) results (see Fig. S1 in the Supplemental Information), likely due to
the higher reactivity of the NH3 plasma compared to the NH3 gas. The O signal most likely arises from residual oxygen present in native oxides. In our early report on PE-ALD Co, similar oxygen peak was detected in the interlayer [15]. By contrast, although a white
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contrast interlayer was formed during the thermal ALD Ni, it is very thin (<1 nm) as shown in Fig. 2(b). O and N signals are detected across the sample at background levels in the Ni deposited by thermal ALD, as shown in Fig. 2(d). The reactivity of the NH3 gas is lower than that of the NH3 plasma. For instance, SiNx is formed even at 300 C when NH3 plasma is used. SiNx is also formed by using NH3 gas at 300 C, but it has a negligible thickness (<1 nm) [17]. Therefore, in contrast to the PE-ALD Ni film, no nitride interlayer was formed in thermal ALD Ni, even though the Si substrate was exposed to the NH3 gas reactant during the initial growth. In addition, it is likely that the very thin interlayer between thermal ALD Ni and the Si substrate is composed of the native silicon oxide layer. Although native oxide etching is performed prior to Ni thermal ALD, a thin oxide layer could be formed again during the sample transfer process [18]. However, since we used 0.3 nm ebeam size for TEM analysis, oxygen in a very narrow interlayer region could not be detected. To investigate the formation of Ni silicide, the PE-ALD and thermal ALD Ni films on Si substrates were annealed at various temperatures and analyzed by X-ray diffraction (XRD). Because the SiOxNy interlayer is amorphous (Fig. 2(a)), no XRD peaks were observed. For the PE-ALD Ni film, nickel mono-silicide phases were observed at 44.8 for NiSi(210) [19] below 700 C as shown in Fig. 3(a). When the annealing temperature (Ta) was increased to 700 C, a strong XRD peak was detected at 33.27, corresponding to the NiSi2(200) [20] phase, while the NiSi(210) peak was not observed. The high intensity peak at 33 is a forbidden diffraction peak of the Si substrate, which is often observed in single crystal Si substrates [21]. In contrast to the PE-ALD Ni film, a different phase transformation was detected in the XRD spectra of the annealed thermal ALD Ni films. At Ta ¼ 400 C, a strong XRD peak was observed near 45 as shown in Fig. 3(b). The peak was deconvoluted to two peaks at 44.8 and 45.1, corresponding to NiSi(210) and Ni2Si(202) [22], respectively. At Ta ¼ 500 C, this strong peak shifted to a slightly lower diffraction angle, because the Ni2Si(202) peak at 45.1 was diminished while the NiSi(210) peak at 44.8 was not. Related to this, the small Ni2Si(112) [22] peak observed at 39.6 in the 400 C annealed sample was not observed at 500 C. The intensity of the NiSi(210) peak decreased, and eventually disappeared at 700 C. In addition, a small peak was detected at 29.2 at 500 C, assigned to NiSi2(111) [23], and the intensity of the peak increased slightly until 700 C. Compared to PE-ALD Ni, thermal ALD Ni resulted in different phase transformations as the annealing temperature increased. Ni deposited by PE-ALD was transformed to NiSi at 400 C without the formation of metal-rich phases such as Ni2Si, and NiSi2 was formed above 700 C. By contrast, NiSi and Ni2Si phases co-existed in the Ni deposited by thermal ALD that was annealed at 400 C, and only the Ni2Si phase disappeared above that temperature. The NiSi disappeared above 600 C, and the NiSi2
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that was detected at 500 C remained until 700 C. Consistently, the sheet resistance of Ni silicide from thermal ALD Ni showed the lowest value at 500 C. After annealing at 500 C, the sheet resistance of Ni silicides from thermal ALD Ni was 610 U/sq. and the resistivity of Ni silicide calculated from the resistance and thickness of Ni silicide was 580 mUcm. Since the NiSi phase has the lowest resistivity in three Ni silicide phases, the lowest sheet resistance of Ni silicide from ALD Ni at 500 C indicates the formation of the NiSi phase. It is known that Ni is a dominant diffusion species during the formation of Ni2Si, NiSi, and NiSi2 [24], and these three silicide phases are usually formed sequentially at 250 C, 350e500 C, and 700 C, respectively [1,25e27]. In the case of the PE-ALD Ni, the Ni atoms cannot easily diffuse into Si through the SiNx interlayer since the diffusion coefficient of Ni in SiNx is 2.28 1014 cm2/s at 650 C, much smaller than that of the Si substrate (6 104 cm2/s at 220e540 C) [28,29]. Moreover, the diffusivity of the Ni in the SiO2 was 1010e1014 cm2/s at 500e700 C, comparable to that of the SiNx layer. Thus, the oxygen species in the SiNx interlayer also interfered with the diffusion of Ni [30]. In addition, according to the NieN phase diagram, the formation of nickel nitride under these conditions is difficult and was not detected using XRD and EDS analyses. The SiNx interlayer is only a barrier for Ni diffusion but does not react with Ni, and as such the diffusion of Ni into the Si substrate is suppressed by the SiNx interlayer. As a result, a metalrich phase is not formed at 400 C and the transformation temperature of NiSi2 is higher than that from sputtered Ni, which does not have a SiNx interlayer. No SiNx interlayer is formed in the thermal ALD Ni, allowing the Ni atoms to more easily diffuse into the Si substrate which results in the sequential formation of three silicide phases as the annealing temperature is increased. The formation of the interfacial layer and phase transformation of the thermal and PE-ALD Ni are summarized in Table S1 in the Supplemental Information. The contact resistance of the Ni silicide obtained from the thermal ALD Ni was evaluated on a planar Si substrate using TLM (fabrication of the TLM line patterns was described previously in Fig. 1(a)). The patterned Ni silicide lines were used for transmission lines. After TLM pattern fabrication, the formation of NiSi was clearly observed in the optical microscopy image shown in Fig. 1(b). The resistance between the NiSi lines increased as the distance between lines increased, as shown in Fig. 4. The intercept on the yaxis that was obtained by linear fitting with three resistance points was 0.48 U and the calculated contact resistance was 0.24 U. This contact resistance is comparable to the typical value of the contact resistance for NiSi that is obtained from PVD Ni [31], indicating that thermal ALD Ni can be used as an alternative contact fabrication process to PVD Ni. We also investigated the silicide formation of thermal ALD Ni in
Fig. 3. XRD results of the (a) PE-ALD and (b) thermal ALD Ni films annealed from 400 C to 700 C.
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Fig. 4. Resistance versus distance between the NiSi lines for TLM using silicide formation of the Ni thermal ALD.
nanosized contact holes. As mentioned in Section 1, understanding the Ni silicide formation inside confined nanostructures is important because it involves both a phase transformation and a volume change. Ni was deposited for 230 cycles to form a 12-nm-thick Ni film, compared to the 25 nm diameter of the contact holes. After thermal ALD Ni, a Ti capping layer was deposited to prevent oxidation. After annealing at 500 C, silicide formation of the thermal ALD Ni in the nanosized contact holes was analyzed by HRTEM as shown in Fig. 5. In Fig. 5(a), the TEM image clearly shows that the Ni is conformally deposited by thermal ALD inside the nanosized contact holes. Fig. 5(b) shows that a void is formed at the bottom of the holes, while the deposited Ni in the middle and top regions remains intact. In addition, a dark contrast crystal facet is formed in the Si substrate region below the bottom of the hole.
Fig. 5(c) shows the EDS line profile along the white arrow. In the facet, the Ni signal gradually decreased along this direction to the Si substrate while the Si signal increased. Based on the compositional ratio of the nickel silicide phases, the region near the bottom of the hole (region a’) is occupied by the NiSi phase, and the region closer to the Si substrate (region b’) is occupied by the NiSi2 phase. The two regions in the facet were mathematically converted to diffraction patterns using fast Fourier transform (FFT) and the results are shown in Fig. 5(d). The FFT patterns from region a’ (close to the hole) show the existence of NiSi(210) and NiSi(200). In region b’ (close to the Si substrate), diffraction patterns of the NiSi2(111) phase were observed. NiSi(210) and NiSi2(111) crystal planes were also consistently observed in our XRD results after annealing at 500 C (Fig. 3(b)). In addition, the STEM high-angle annular darkfield (HAADF) image in Fig. 5(e) clearly shows the contrast differences within the facet region. Because the contrast in the STEM HAADF image is proportional to the atomic number [32], region a’ and region b’ are composed of NiSi and NiSi2, respectively. The XRD results shown in Fig. 3(b) indicate that the NiSi peak is much stronger than the NiSi2 peak in the case of thermal ALD followed by annealing at 500 C. It was reported that Ni films are mainly transformed to the NiSi phase rather than the NiSi2 phase after annealing at 500 C, and the transformation to the NiSi2 phase requires higher annealing temperatures [26]. As such, the Ni silicide from the Ni deposited by thermal ALD and then annealed at 500 C is primarily composed of NiSi with a small amount of NiSi2 also present. However, the volume of the NiSi2 phase is larger than that of the NiSi phase in the Ni silicide formed underneath the contact hole as shown in Fig. 5(d), which is inconsistent with the XRD results. In the planar sample, the whole Ni film is in contact with the Si substrate, while in the contact hole, Ni is only supplied from the bottom of each hole. As a result, it is likely that the supply of Ni atoms is more limited in the contact hole than in the planar substrate. At 500 C, the Ni atoms have enough thermal energy to diffuse into the Si substrate and form NiSi in the planar system.
Fig. 5. HRTEM images of the Ni deposited by thermal ALD in nanosized contact holes (a) before and (b) after annealing at 500 C. (c) EDS line profile along the arrow indicated in (b). (d) Magnified TEM image from (b) and FFT patterns collected from regions a’ and b’. (e) HAADF image of (d).
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Therefore, the formation of the NiSi2 phase in region b’ can be attributed to the geometrical limitations of Ni. 4. Conclusions We investigated the formation of silicide from Ni films prepared using PE- and thermal ALD using NH3 plasma and NH3 gas, respectively. A SiNx interlayer was formed in the PE-ALD Ni film due to direct exposure of the Si substrate to NH3 plasma, while no nitride interlayer was observed in the case of thermal ALD. Because the SiNx interlayer suppressed Ni diffusion, the Ni2Si phase was not formed and the formation temperature of NiSi2 increased. In contrast to PE-ALD, the Ni deposited by thermal ALD sequentially transformed to Ni2Si, NiSi, and NiSi2 as the annealing temperature increased. Using TLM, the contact resistance of the NiSi obtained from the thermal ALD Ni film was 0.24 U. The silicide formation of the Ni deposited by thermal ALD in the nanosized contact holes was different from that of the planar substrates. Inconsistent with the XRD results, a larger amount of NiSi2 was observed compared to NiSi. The Si-rich silicide phase (NiSi2) was formed at the bottom of the contact holes due to the limited supply of Ni atoms inside the nanostructures. The results from this study can provide important information for applications that require Ni deposited by thermal ALD for the fabrication of Ni silicide inside nanostructures. Acknowledgements This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education (2014R1A1A2059845), and by the MOTIE (Ministry of Trade, Industry & Energy) (10053098) and KSRC (Korea Semiconductor Research Consortium) support program for the development of the future semiconductor device. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.cap.2016.04.005. References [1] C. Lavoie, F.M. d'Heurle, C. Detavernier, C. Cabral Jr., Towards implementation of a nickel silicide process for CMOS technologies, Microelectron. Eng. 70 (2003) 144e157. [2] S.L. Zhang, U. Smith, Self-aligned silicides for Ohmic contacts in complementary metaleoxideesemiconductor technology: TiSi2, CoSi2, and NiSi, J. Vac. Sci. Technol. A 22 (2004) 1361e1370. € [3] S.L. Zhang, M. Ostling, Metal silicides in CMOS technology: past, present, and future trends, Crit. Rev. Solid State Mater. Sci. 28 (2003) 1e129. [4] J. Chen, J.-P. Colinge, D. Flandre, R. Gillon, J.P. Raskin, D. Vanhoenacker, Comparison of TiSi2, CoSi2, and NiSi for thin-film silicon-on-insulator applications, J. Electrochem. Soc. 144 (1997) 2437e2442. [5] S.P. Murarka, D.B. Fraser, A.K. Sinha, H.J. Levinstein, E.J. Lloyd, R. Liu, D.S. Williams, S.J. Hillenius, Self-aligned cobalt disilicide for gate and interconnection and contacts to shallow junctions, IEEE Trans. Electron Devices 34 (1987) 2108e2115.
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