Si(100) thin film induced by ion beam bombardment

Vacuum 89 (2013) 157e162

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Nanostructure formation of Cu/Si(100) thin film induced by ion beam bombardment G.S. Tang, H.Y. Liu, F. Zeng*, F. Pan* Key Laboratory of Advanced Materials (MOE), Department of Materials Science and Engineering, Tsinghua University, Qinghuayuan, Beijing 100084, PR China

a r t i c l e i n f o

a b s t r a c t

Article history: Received 22 September 2011 Received in revised form 11 March 2012 Accepted 13 March 2012

A simple method for fabricating self-organized Cu nano-dots on Si(100) substrate by low energy Arþ ion beam bombardment of a Cu thin film at room temperature over a large area is demonstrated. The morphological evolution has been investigated using scanning electron microscopy and atomic force microscopy. It was found that nano-ripple patterns formed on a Cu grain surface on a 110 nm thick polycrystalline Cu thin film under normal ion incidence. Uniformly distributed Cu nano-dots were obtained by bombardment of 55 nm thick nano-crystalline Cu thin films. The formation mechanism of the Cu nanostructures was discussed with the aid of numerical simulations using a modified damped KuramotoeSivashinsky equation. Ó 2012 Elsevier Ltd. All rights reserved.

Keywords: Nanostructure Self-organized Ion beam Bombardment Damped KuramotoeSivashinsky equation

1. Introduction Ion beam bombardment (IBB) induced self-organized nanostructures on solid surfaces have attracted great interest in scientific and industrial fields in recent years [1,2]. Under certain bombardment conditions, usually well-ordered ripple or dot nanostructures can be formed on the surface. Due to its simplicity and easy control of different process parameters, this method offers a simple and highly cost-efficient route for fabricating large area patterned surfaces. This self-organized process is also considered as a bottomeup method which has potential to overcome the limitations of a topedown approach such as lithographic techniques and ion etching. For normal ion incidence or oblique ion incidence with sample rotation, nanostructures with hexagonally arranged dots, rectangularly arranged dots, disordered dots and holes on different materials have been reported depending on experimental settings and process parameters [1e4]. The mechanism of nanostructure formation is a complex interplay between surface roughening induced by sputtering and a smoothing process on the surface. The theory accounting for evolution of the surface topography has also been investigated intensively in recent years [5e8]. However, the vast majority of published works focuses on

patterning the surfaces of bulk samples, few studies have been carried out on thin metal films [9,10]. Furthermore, metal nanodots on non-metallic substrates have wide applications in catalysis [11], high density data storage [12], and as nano-sensors based on surface-enhanced Raman scattering (SERS) or localized surface plasmon resonance (LSPR) [13]. Direct patterning of metal patterns on heterogeneous substrates by IBB is a challenging task, because the existence of EhrlicheSchwoebel (ES) barriers or preferential diffusion paths plays an important role in determining the final pattern characteristics [14]. ErhlicheSchwoebel (ES) barriers are the barriers to diffusion of adatoms and surface vacancies across step edges [14] so the presence of ES barriers can increase the roughness on the surface. It will be interesting to pattern metal nanostructures on non-metallic substrates by IBB. In this paper, we report Cu nanostructures formation on Si(100) semiconductor surface by IBB under normal incidence and discuss our observations of nano patterns by IBB using continuum model simulations. By modulating the microstructures of the Cu layer, we investigated the mechanism of morphology change between nanoripple and nano-dot patterns, which would guide us to get the desired pattern in thin metal film for further application. 2. Experimental process

* Corresponding authors. E-mail addresses: [email protected] mail.tsinghua.edu.cn (F. Pan).

(F.

Zeng),

0042-207X/$ e see front matter Ó 2012 Elsevier Ltd. All rights reserved. doi:10.1016/j.vacuum.2012.03.039

panf@

Ion bombardment was performed in a custom built ion beam sputtering system, equipped with a Kaufman-type broad-beam ion gun. The sputtering gas was 99.99% argon (Ar) at the working

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Fig. 1. XRD patterns of Cu/Si(100) thin films with thickness of 110 nm and 55 nm.

2

pressure of 2  10 Pa, and the base vacuum was better than 2  104 Pa. The bombardment was carried out under normal incidence to the sample surface at room temperature. Cu films with thickness of 55 nm and 110 nm (labeled as Cu55 and Cu110 respectively, hereafter) were deposited on Si(100) substrates using a DC magnetron sputtering system with 99.995% purity Cu sputter targets. Prior to deposition, commercially

available EPI-polished p-type Si(100) wafers with a root-meansquare (RMS) roughness of w0.2 nm and resistivity of w3 U cm were soaked in 20% HF solution for 1 min to remove the native oxide of silica, and then degreased by successive rinses in ultrasonic baths of acetone, ethanol, and de-ionized water. The base vacuum in the chamber during deposition was better than 5  104 Pa and the working pressure was fixed at 0.4 Pa, which was 99.99% argon. Then the Cu/Si(100) thin films were bombarded by Arþ ion beam under normal incidence to the sample surface with an ion flux Jion  1.1  1015 cm2 s1 and an ion energy Eion  3 keV. After deposition, the microstructure of the films was characterized by X-ray diffraction (XRD) using monochromatic Cu Ka radiation with l ¼ 1.5408 Å. The surface topography was investigated using an LEO 1530 field emission scanning electron microscopy (FE-SEM) and an SPM 9600 atomic force microscopy (AFM) operating in tapping mode. All AFM measurements were performed in air at room temperature with a resolution of 512  512 pixels. The RMS surface roughness w was calculated using the qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ! ! relation w ¼ h½hð r ; tÞ  hhð r ; tÞi2 i, where the term in brackets h/i denotes the spatial average. The two-dimensional (2D) autocorrelation function was calculated using the relation ! ! Cð r ; tÞ ¼ hhð r ; tÞhð0; tÞi. 3. Results and discussions

q  2q XRD patterns of the Cu films deposited on Si(100) substrates are shown in Fig. 1. The Cu films have a polycrystalline

Fig. 2. SEM images of morphological evolution of the Cu110 sample bombarded by Arþ ions (Jion ¼ 1.1  1015 cm2 s1, Eion ¼ 1.2 keV). (a) As deposited. (b) t ¼ 60 s. White arrows indicate ripple orientation. (c) t ¼ 180 s (d) t ¼ 240 s.

G.S. Tang et al. / Vacuum 89 (2013) 157e162

feature with grains randomly oriented. The full-width at half maximum (FWHM) values of the Cu(111) and Cu(200) spectrum peaks is increased by 240% and 170% respectively when the thickness of the Cu layer is changed from 110 nm to 55 nm. According to the Scherrer formula (DwK l=bcos q, where D is the average grain size, Kw0:89 is the shape factor, l is the wavelength of the X-rays, b is the FWHM and q is the Bragg angle) [15], the longitudinal grain size of the Cu110 and Cu55 sample was estimated to be w60 nm and w16 nm respectively, indicating that Cu grains grow rapidly as the film gets thicker. Moreover, the longitudinal grain size of the samples was estimated from the SEM images after the initial layer was sputtered, as described below. The electron microscopy experiments have been carried out to study the surface morphology of the Cu110 sample bombarded by Arþ ions with Jion ¼ 1.1  1015 cm2 s1 and Eion ¼ 1.2 keV. The initial surface topography shows a typical bump-like structure (Fig. 2(a)). After bombardment for 60 s, Cu grains with lateral size in, or exceeding, the sub-micrometer range can be clearly seen (Fig. 2(b)). Nano-ripples were formed on a single grain surface with ripple orientation randomly distributed as indicated by the white arrows. After further bombardment, the Cu structures turned to isolated nano-clusters, nano-dots and even nano-wires which are the remains of the nano-ripples formed on a single grain (Fig. 2(c)). Light Cu features on a darker background that represents the substrate can clearly be seen. After almost the entire Cu layer was removed, sparsely distributed Cu nano-dots were left on the substrate (Fig. 2(d)). The morphological evolution of the Cu55 sample bombarded by Arþ ions for a given ion flux and energy is shown in Fig. 3(a)e(f), and the time evolution of the surface roughness w is represented in Fig. 3(g). The lateral grain size of the Cu55 sample could not be distinguished clearly due to the resolution limit of the SEM, indicating the film had nano-crystalline grains. The initially continuous Cu film changed to a network of Cu clusters, isolated islands, nanodots combined with nano-rods and, finally, to nearly uniformly distributed nano-dots. A dark substrate feature was seen after with slightly longer exposure, which confirmed that the previous Cu

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pattern was observed on the Si substrate. The surface roughness w increased again with further bombardment and the coarsening process of the Si substrate took place to form finally hillock structures on the surface. Fig. 4 shows the AFM morphology of the Cu nano-dots on the Si substrate with variation of ion flux Jion and ion energy Eion. The shape, size and distribution of Cu nano-dots in AFM images agree with those shown in the SEM images (top-left insets in Fig. 4(a)e(c)). For instance, under Arþ ion etching with Jion ¼ 1.1  1015 cm2 s1 and Eion ¼ 2.5 keV, the Cu nano-dots with mean diameter of 12  4.3 nm, height of w5.3 nm and dot-density of 1.8  1010 cm2 were obtained. By reducing Jion to 0.6  1015 cm2 s1, the dotdensity increased to 2.5  1010 cm2, and the mean diameter was reduced to 8.4  5.5 nm and height to w4.6 nm. Dot-density of 1.2  1010 cm2, mean diameter of 30  13 nm and height of w7.8 nm were produced with Jion ¼ 1.1  1015 cm2 s1 and Eion ¼ 3.0 keV bombardment. The 2D-autocorrelation functions of the AFM images are represented in the top-right insets of Fig. 4(a)e(c), revealing the regular order of the dots. The autocorrelation functions exhibit a sign of hexagonal structure, demonstrating that a self-ordering process with a specific spatial frequency is in process. The Cu nano-dot pattern was found to be almost uniformly distributed over the whole sample surface (w2 cm2). This technology offers us a simple, high-efficiency and low-cost routine to pattern large area metal nano-dots on heterogeneous substrates, with an easy method to tailor the dot-size and dot-density by varying bombardment parameters such as ion flux and ion energy. To understand the mechanism of the pattern formation induced by ion bombardment, we have conducted a numerical simulation based on a continuum model. Under ion bombardment, four main mechanisms have been considered for surface morphological evolution: sputtering, surface diffusion, redeposition, and viscous flow [5,6,16,17]. The evolution of the surface during the ion bombardment is described by a surface ! height function h ¼ hð r ; tÞ to model the self-organized processes which lead to various patterns. A continuum model based on the damped KuramotoeSivashinsky (KS) equation has been applied for

Fig. 3. SEM images of morphological evolution of the Cu55 sample bombarded by Arþ ions (Jion ¼ 1.1  1015 cm2 s1, Eion ¼ 2.5 keV). (a) To (f) t ¼ 0, 30, 60, 90, 105 and 120 s respectively. (g) The surface roughness w, analyzed from AFM images, as a function of exposure time t.

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Fig. 4. Surface morphology of the Cu55 sample bombarded by Arþ ions. (a) Jion ¼ 1.1  1015 cm2 s1, Eion ¼ 2.5 keV, t ¼ 120 s. (b) Jion ¼ 0.6  1015 cm2 s1, ¼ 2.5 keV, t ¼ 180 s. (c) Jion ¼ 1.1  1015 cm2 s1, Eion ¼ 3.0 keV, t ¼ 90 s. Top-left: SEM image. Bottom: AFM image. Top-right: 2D-autocorrelation function calculated from corresponding AFM image. All of the images have the same scale. (d)e(f) histograms of particle size distribution for (a)e(c) respectively.

semiconductor or amorphous surface morphological evolution during erosion by an ion beam under normal incidence [7,10,18,19]:

l

vt h ¼ y0  gh  nV2 h  DV4 h þ ðVhÞ2 þh: 2

(1)

Here, y0 is the constant erosion velocity of the plane surface, n is the effective surface tension, D is the effective diffusion coefficient for the combined thermal diffusion [5], erosion induced diffusion [20], and viscous flow [16], l denotes the tilt-dependent sputtering yield [17], g is the damping term, accounting for the deposition of sputtered species on the surface [7], and h is Gaussian white noise resulting from the stochastic nature of the bombardment process. The time, where the surface roughness starts to saturate, is referred to as the crossover time tc. In the following, times larger than tc are called late time regime. The case of metals should be treated using a different approach. Due to the higher diffusivity with respect to semiconductors and amorphous materials and to the non-directional character of the metallic bond, metals do not amorphize upon ion bombardment, at least for low fluence [14]. Considering ES barriers or preferential diffusion paths on a metal surface, a simple correction has been proposed in [21] by introducing a different diffusion term Eq. (1) then becomes:

! X v2 h S! 2 ! n v! n n ! X l v4 h D0 V4 h þ þ ðVhÞ2 þh: D! 4 ! n 2 ! vn n

vt h ¼  y0  gh  

n V2 h þ

(2)

v2 h The term S! takes into account the diffusion arising from the 2 n v! n v4 h the diffusion along the crystalloES barrier and the term D! 4 n v! n ! graphic orientation n . By rotating the coordinate system (assuming ! ! 2 perpendicular n and setting one n parallel to the x axis), Eq. (2) becomes a modified anisotropic damped KS equation (adKSE).

Moreover, by redefining of h ¼ h  y0t, and by merging and rescaling the coefficients [19], Eq. (2) becomes dimensionless and reads:

    vt h ¼ gh  v2x þ av2y h  v4x þ bv2x v2y þ dv4y h þ ðVhÞ2 þh:

(3)

Here, a indicates the strength of the anisotropy of the BradleyeHarper effect [5] and the contribution to the diffusion arising from ES barriers. b marks the anisotropy of the ion-induced effective surface diffusion. d denotes the anisotropy of the diffusion along the preferential path. The limit a ¼ d ¼ 1, b ¼ 2 yields an isotropic damped KS equation (idKSE). Therefore, the variation of a, d, b and g determine the morphological evolution. Numerical simulations have been performed by integration of Eq. (3) using a standard discretization method with periodic boundary conditions, a mesh size of 256  256, spatial step dx ¼ 1.0, time step dt ¼ 0.01, an initially flat surface and an amplitude of the Gaussian white noise A ¼ 0.01. The value of g is fixed to 0.24 which has also been chosen in [7] to produce stable patterns. Simulated surface morphologies under normal incidence are shown in Fig. 5 with the coordinate axis indicated. The following cases are considered: (i) when no ES barriers and preferential diffusion paths exist, for which a hexagonally arranged dot pattern appears (Fig. 5(a)). The scenario is suited for the situation where the substrates are semiconductors or amorphous materials. For the Cu55 sample with nano-crystalline grains, the existence of ES barriers becomes improbable due to the lack of well-defined atomic steps on the surface [22]. These facts support the isotropic surface diffusivity rather than an anisotropic one. So the hexagonal structure is expected, as observed in the experimental results of the Cu55 sample (Fig. 4). (ii) When preferential diffusion paths are in both x and y axes, the value of b becomes anisotropic (e.g., b ¼ 8), and then a rectangularly arranged dot and hole pattern is formed (Fig. 5(b)). This case has occurred on a singlecrystal Cu(110) surface which had a similar pattern bombarded by Arþ ions at 320 K [21]. (iii) When the ES barrier takes place along the x axis, the value of a becomes anisotropic (a < 1, e.g., a ¼ 0.5), period ripples are produced with orientation parallel to the y axis (Fig. 5(c)). (iv) when only one preferential diffusion path

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Fig. 5. Surface morphology calculated from Eq. (3) in the late time regime with g ¼ 0.24. (a) a ¼ d ¼ 1, b ¼ 2. (b) a ¼ d ¼ 1, b ¼ 8. (c) a ¼ 0.5, d ¼ 1,,b ¼ 2. (d) a ¼ 1, d ¼ b ¼ 4. Insets: Corresponding 2D-autocorrelation function indicates the regular order of the image.

exists, the value of d and b become anisotropic (e.g., d ¼ b ¼ 4), highly ordered ripples are formed along the y axis which is the preferential diffusion path. For the Cu110 sample with lateral grain size at or exceeding the sub-micrometer range, surface diffusivity should be considered anisotropic. If local anisotropic diffusion occurs on a single grain surface due to ES barriers or preferential diffusion paths, a ripple pattern will be expected by simulation as demonstrated by case (iii) and/or (iv). This can explain the experimental observation in Fig. 2, where ripples formed on an individual grain. The variation of the ripple orientation indicated the change of azimuthal angle and/or crystallographic direction [23]. It can be inferred that the existence of ES barriers and preferential diffusion paths plays an important role on the surface morphology (nano-ripples) in sample Cu110. A reduced film thickness would lead to the formation of nano-crystalline grains, and the subsequent absence of ES barriers and preferential diffusion paths. In this scenario, uniformly distributed dots are formed in the Cu55 sample. It should be noted that, very recently, Bradly et al. developed a new method to model pattern formation on the solid surface by normal-incidence IBB [24e27]. In their theory, for a binary compound the coupling between the surface height and a surface layer of altered composition is the key to the regular nano-dots pattern formation [25,26], while for an elemental material concurrently deposited impurities can play a crucial role in determining the type of nano-scale pattern [27]. This certainly deserves further study.

4. Conclusions In conclusion, a simple, high-efficiency and low-cost method to fabricate self-organized Cu nano-dots by ion bombardment of Cu/Si(100) thin films at room temperature over a large area has been demonstrated. The microstructures of the Cu film were modulated by varying the film thickness. A uniformly distributed metal nano-dots pattern was obtained by IBB of a thin Cu film with nano-crystalline grains (Cu55), when the ES barriers or preferential diffusion paths lost efficacy. Moreover, dot-size and dot-density could be tailored by varying bombardment parameters such as ion flux and ion energy. On the other hand, a nanoripples pattern was formed on a polycrystalline Cu surface with lateral grain size at or exceeding the sub-micrometer range (Cu110), where ES barriers or preferential diffusion paths played an important role. The experimental results agree with simulations using a modified damped KuramotoeSivashinsky equation. Understanding these mechanisms will help us to pattern other metals such as Al, Fe, Co, Ni, Ag and Au, also using other substrates for practical application. Acknowledgments The authors are grateful for financial supports from the National Natural Science Foundation of China (Grant No. 50871066), National Hi-tech (Research and Development) project of China (Grant No. 2009AA034001), and National Basic Research Program

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