Thin Solid Films 534 (2013) 90–99

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Growth of crystalline silicon nanowires on nickel-coated silicon wafer beneathsputtered amorphous carbon

Feng Ji Li a, Sam Zhang a,⁎, Jun Hua Kong b, Jun Guo b, Xue Bo Cao b, Bo Li c

a School of Mechanical and Aerospace Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore 639798, Singaporeb School of Materials Science and Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore 639798, Singaporec Central Iron and Steel Research institute, 76 South Xueyuanlu Rd, Haidan District, Beijing, 100081, China

⁎ Corresponding author. Tel.: +65 67904400.E-mail address: MSYZhang@ntu.edu.sg (S. Zhang).

0040-6090/$ – see front matter © 2013 Elsevier B.V. Allhttp://dx.doi.org/10.1016/j.tsf.2013.02.007

a b s t r a c t

a r t i c l e i n f o

Article history:Received 21 March 2012Received in revised form 1 February 2013Accepted 1 February 2013Available online 16 February 2013

Keywords:Rapid thermal annealingCrystalline siliconNanowiresNickel catalystOxidationSolid–liquid–solid growth

Growth of crystalline silicon nanowire of controllable diameter directly from Si wafer opens up another av-enue for its application in solar cells and optical sensing. Crystalline Si nanowire can be directly grownfrom Si wafer upon rapid thermal annealing in the presence of the catalyst such as nickel (Ni). However,the accompanying oxidation immediately changes the crystalline Si nanowire to amorphous SiOx. In thisstudy, amorphous carbon layer was sputtered to on top of the catalyst Ni layer to retard the oxidation. Scan-ning electron microscope, transmission electron microscope, Raman spectroscopy and X-ray photoelectronspectroscopy were employed to characterize the wires and oxidation process. A model was developed to ex-plain the growth and oxidation process of the crystalline Si nanowire.

© 2013 Elsevier B.V. All rights reserved.

1. Introduction

Following the discovery of carbon nanotubes [1], one-dimensional(1-D) anisotropic inorganic structures with nanometer diameters,such as nanowires, nanotubes and nanorods have been stimulatinggreat research interest and becoming the hottest research subjectsin recent years [2,3]. As 1-D nanostructure, Si-based nanostructures playcritical roles in interconnects and functional components in mesoscopicelectronic, optical devices and energy storage [4]. Their mechanical, elec-tronic and optical properties strongly rely on dimensionality and latticeorientation [2]. For instance, the ultimate mechanical strength of silicananowire is in excess of 12GPa and increases with the decrease of the di-ameter [5]. Diamond lattice Si only has an indirect 1.1-eV band gap show-ing weak infrared (IR) emission, whereas linear trans-polysilane chainsexhibit a 3.89-eV direct band gap showing strong ultraviolet (UV) emis-sion [6]. Silica nanowires could emit strong blue [7,8], red [7], green [6]and even UV light [7,8]. Moreover, it is capable of guiding light withinthe visible and near infrared spectral ranges with low optical losses [9].

However, the synthesis of Si-based nanostructures with desireddimensionality or morphologies is technically challenging [10]. Gener-ally speaking, top-down and bottom-up approaches are themain routesto synthesize Si nanowires. Top-down method prepares Si nanowiresvia dimensional reduction of bulk Si by lithography and etching,which is usually time-consuming, expensive and difficult to achieve

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uniform diameters below 10 nm [11,12]. Bottom-up approach is an as-sembly process joining Si atoms often via a vapor–liquid–solid (VLS)mechanism in generating single crystalline Si nanowires with diame-ters ranging from 5 nm to several hundreds of nanometers [3]. In themechanism [13,14], a liquid eutectic metal (such as Au, Ni, Fe, etc.) oralloy droplet composed of metal catalyst components and nanowirematerials (such as Si, C etc.) is formed under the reaction conditions [10].This liquid droplet serves as a preferred site for absorption of gas-phasereactant. Once supersaturated, it becomes the nucleation site for crystal-lization. During the growth, the catalyst droplet directs the nanowire'sgrowth orientation and defines the diameter of the crystalline nanowire.Ultimately, the growth terminates when the temperature is below theeutectic temperature of the alloy or the reactant is no longer available.The formation of eutectic metal/Si liquid droplets and the establishmentof the symmetry-breaking solid–liquid interface is the key for theone-dimensional crystal growth. However, it has relatively low yield ofproduction, and the use of flammable or toxic precursor gases such asSiH4/H2, SiCl4 or Si oxide vapor are the disadvantages. So far, chemicalvapor deposition [15], laser ablation [3], thermal evaporation [16–18],and thermal annealing have been applied to synthesize Si-based 1-Dnanostructures [10,19–21]. For instance, Goodey et al. synthesized Sinanowires from SiH4 at 500 °C by VLS mechanism using Au as the cata-lyst [22]. Alfredo et al. synthesized single-crystalline Si/SiOx nanowireswith crystalline cores vary from 6 to 20 nm, and length of 1 to 30 μmvia combining laser ablation nanometer-diameter iron catalyst clusterformation [3]. Justin et al. obtained defect-free Si nanowires with uni-form diameters ranging from 4 to 5 nm, length of several microns via a

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supercritical fluid solution-phase approach with Au nanocrystal as thecatalyst [2].

To avoid using flammable or toxic precursor gas as the Si source inVLS, solid–liquid–solid (SLS) growth of amorphous Si, amorphous SiOx

or crystalline-Si/amorphous-SiOx (i.e. c-Si/a-SiOx) nanowires directlyuses Si wafer as the source and adopts Ni [20], Au [23,24], NiO [18], orFe [25] as catalysts under nitrogen or argon atmosphere. Park et al. syn-thesized c-Si/a-SiOx nanowires by heating a NiO-catalyzed Si substrateunder a reductive environment of WO3, C and Ar atmosphere viaSLS growth [18]. A tentative SLS mechanism for the growth of thenanowire based on chemical reaction has been proposed: CO gasfrom the carbothermal reaction of WO3/C reacted with SiOx to pro-duce SiO and CO2, where SiO2 is formed from Si–Ni–O liquid dropletsin the NiO/Si interface. Si and SiOx are generated through the dispropor-tionation of SiO in the film [18]. The c-Si/a-SiOx nanowire is grown viathe phase separation of Si and SiOx from the super-saturated Si–Ni–Odroplets by the cooling effects of Ar flow. Taking consideration of theoxidation of the droplet at the beginning stage of the growth, Nie etal. developed a solid–liquid–solid growth mechanism catalyzed by thetemperature-dependent alloy, metallic and ionic Fe (Fe–FeOx) to de-scribe the growth of Si–SiOx core–shell nanowires during directlyheating Fe-coated Si wafers at three different temperatures. In themechanism, the formation of 1-D amorphous SiOx nanowires wasgrown from the fully oxidized Fe–Si alloying droplets as a result ofthe continuously feeding Si into the alloying droplets, whereas part-ly oxidized large ones form Fe–Si alloying cores and Fe–O–Si shells,which result in Si–SiOx core–shell nanowires.

Except Nie et al., almost all the other mechanisms overlooked theserious oxidation by residual oxygen in the reaction chamber at hightemperature, since the reaction is often under an inert gas atmosphererather than a high vacuum environment. Moreover, if these reactionstake place, Si crystal might be embedded in amorphous SiO2 matrix inthe form of film, rather than Si crystalline core encapsulated in SiOx

sheath in the form of the nanowire [26]. Therefore, the tentative mecha-nisms (reaction plus phase separation, or disproportionation reaction)are challenged to explain why the as-grown Si nanostructure is in theform of crystalline Si core wrapped in amorphous SiOx sheath. Amor-phous SiOx nanowires can be grown directly by heating a NiO-catalyzedsilicon substratewithout usingWO3/graphite powder in a reductive envi-ronment in a tube furnace [18], or by thermal heating Au or Ni thinfilm-coated Si substrates [27]. The residual oxygen is unavoidable in thechamber. The formation of amorphous all-SiOx nanowires is attributedto the serious oxidation of as-grown Si wires. Therefore, in order togrow crystalline Si nanowires, the oxidation of as-grown Si nanowiresmust be retarded. In summary, the main challenges in Si-related nano-structures are: (1) difficulty in controlling the dimensionality and mor-phologies. (2) Incredulity of the proposed SLS mechanisms based onphase separation of Si and SiOx or disproportionation reaction of SiOvapor in explaining the growth of c-Si/a-SiOx nanowires. (3) Inability inexplaining the formation of the amorphous SiOx nanowires by the dispro-portionation reaction mechanism.

As a promising alternative to standard furnace annealing, rapid ther-mal annealing presents advantages in precise control of annealing tem-perature and the annealing profile, ambient purity, short annealingtimes (from 1 s to 3 min), cycle time, and process flexibility. The funda-mental flexibility in creating different types of thermal processes arisesfrom the dynamic control of the heat source temperature, which permitsfast heating combined with dynamic optimization of temperature uni-formity [28,29]. However, the annealing chamber normally is under agas atmosphere instead of a vacuum environment due to the high cost.Alternatives of retarding the unavoidable oxidation must be workedout to grow crystalline Si nanowires. Our previous work reported thegrowth of amorphous all-SiOx nanowires via 3 min rapid thermalannealing of 345-nm a-C/Ni film using a-C layer as the oxidation re-tardation layer [30]. However, c-Si/a-SiOx nanowires were still notobserved due to the serious oxidation. This paper sputtered thicker

carbon film on a thin layer of Ni using Si wafer as substrate, followedby rapid thermal annealing for different duration of time to controlthe growth of the c-Si/a-SiOx nanowires. It is shown that the structureof c-Si/a-SiOx nanowires is controlled by the thickness of the a-C layerand the annealing duration. In the end, an oxidation-accompanyingSLS growth model is developed to explain the formation of the c-Sinanowires, amorphous all-SiOx nanowires and their transition.

2. Experimental details

2.1. Deposition of the film

The a-C/Ni film was prepared in an E303A magnetron sputteringsystem (Penta-Vacuum, Singapore). N-type Si (100) wafer was usedas the substrate (10 mm×10 mm in area, 475 μm in thickness and0.5 nm in Rq). Before loading into the sputtering chamber, the sub-strate was ultrasonically cleaned in acetone for 20 min, followed by10 min in alcohol. Once the base pressure reached 1.0×10−3Pa, Argas was introduced at a flow rate of 50 standard cubic centimeters perminute (sccm). High energy Ar+ ions were generated at a substratebias of−300 V for substrate cleaning for 15 min to further clean the sur-face contamination. Afterwards, a 5-min sputtering of Ni target (99.99%in purity) in 0.47 Pa pressure at room temperature resulted in about30-nm thickNi catalyst layer on the substrate. Si and Cwere subsequent-ly co-sputtered from Si and graphite target (99.999% in purity) on the Nilayer under the same pressure. Amorphous-C films of around 130 nmand 570 nm in thickness were respectively co-sputtered for 15 min and60 min. As a comparison study, a single Ni film was also deposited fromNi target for 5 min. The sputtering power densities for Ni, Si and C wererespectively RF 1.85 W/cm2, RF 2.47 W/cm2 and DC 12.22 W/cm2.

2.2. Rapid thermal annealing

The as-sputtered films underwent rapid thermal annealing (RTA,Jipelec Jetfirst 100 rapid thermal processor) in an Ar ambient at 1100 °Cfor 60 s and 180 s, respectively. The RTA chamber was purged 10 timeswith Ar (99.999% in purity) at 2000 sccm before ramping up to 600 °Cat 58 °C/s. After dwelling for 6 s at 600 °C, the chamber was furtherramped up to 1100 °C at 41.7 °C/s, and then held for 60 s and 180 s, re-spectively. After annealing, the chamber was firstly cooled down to500 °C at 43 °C/s, followed by natural cooling to room temperature. Dur-ing the whole annealing process (ramping, holding, and cooling), the in-flowof theAr gaswas kept at 2000 sccm tomaintain chamber pressure ofaround 1.08×105Pa.

2.3. Characterization

The surface chemical states of the films were studied by X-rayphotoelectron spectroscopy (XPS, Kratos) using a Kratos AXIS spec-trometer with monochromatic Al Kα (1486.71 eV) radiation (10 mAand 15 kV). The electron energy analyzer was operated at 0.1 eV scan-ning step size on a lowmagnification scan spot with 220 μm in diame-ter. The atomic concentration of each element is determined by theratio of the corresponding fitted core-level peak area over the sum ofthe fitted peak area of all the elements. The respective core-level XPScurve fitting was performed after a Shirley background subtraction bynonlinear least square fitting using a mixed Gauss/Lorentz functionafter charge-correcting by positioning C 1s peak to 284.8 eV. Field emis-sion scanning electron microscopy (JEOL, JBM-7600F, JEOL Ltd., Japan,5-kV operating voltage) showed that wire-like nanostructures wereformed on the surface of the as-annealed film surface. High resolu-tion transmission electron microscope (HRTEM, JEOL 2010 and 2100F,200-kV operating voltage) revealed that the wires were of a crystallinecore enveloped by an amorphous sheath. HRTEM samples were pre-pared by using a sharp stainless tweezers to scratch the as-grownwires off the substrate surface into a small plastic container of acetone,

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and subsequently followed by 1 min ultrasonic agitation for dispersion.The size of the wire was calculated by randomly taking the mean valueof 10 wires. Energy dispersive X-ray spectroscopy (EDX, EDAX Inc., US)cross-sectional elemental mapping and selected area electron diffraction(SAED) confirmed that the core was crystalline Si and the encapsulationwas amorphous SiOx. The carbon sp2 bonding structure was examinedwith a Renishaw Raman spectroscope at a wavelength of 633 nm radia-tion excited with a He–Ne laser.

3. Results and discussion

3.1. Surface chemical state of the film before and after annealing

Fig. 1(a) and (b) shows the XPS survey scan spectra of theas-sputtered and the 3-min annealed 570-nm a-C film. Note that nosignal for Ni 2p (binding energy ranging from 885 to 847.9 eV) isdetected in the spectrum except for Si 2p, Si 2s, O 1s, O 2s and C 1s,which means that Ni still locates at the bottom of the film. As seen,upon annealing, the signal of C 1s is dramatically trailed off, whilethe intensities of O 1s, O 2s, Si 2s and Si 2p are greatly boosted up. Be-fore annealing, the C 1s content (79.34 at.%) is ultrahigh in compari-son with Si 2p (4.16 at.%) and O 1s (16.50 at.%). The amount of Si inthe film is very small. After annealing, there is a tremendous decreasein C 1s content (5.2 at.%) and sharp increase in O 1s (67.10 at.%) andSi 2p (27.69 at.%).

To evaluate the chemical state of silicon in the as-sputtered filmand the oxidation of silicon after annealing, the Si 2p core-level peaksof the film before and after annealing are illustrated in Fig. 7(c) and(d). The five oxidation states of Sin+ (n=0, 1, 2, 3 and 4) correspondingto the five chemical structures of Si, Si2O, SiO, Si2O3 and SiO2, respective-ly [31], exist in nonstoichiometric SiOx (0bxb2) films. The position ofthe five synthetic peaks of Si 2p is determined thus: 99.6 eV in Si0,100.6 eV in Si1+, 101.6 eV in Si2+, 102.6eV in Si3+, and 103.6 eV inSi4+ [32,33]. The full width at half maximum of each peak is fixed tobe 1.90 eV during fitting. After charge correcting the Si 2p peak by posi-tioning the major component in C 1s peak to 284.8 eV, the above peakpositions are used to fit the Si 2p core-level lines. A variation in the po-sition of each peak of ±0.1 eV is considered acceptable. With the

Fig. 1. XPS spectra of the (a) as-sputtered film, (b) 3-min annealed films; deconvolution

possible existence of the five oxidation states in mind, after subtractingan integrated Shirley function background, the curve fitting is conductedby decomposing the spectrum into the Si 2p1/2 and Si 2p3/2 partner linesfor the 5 oxidization states without any pre-adjudication. Although thespin-orbit splitting and the Si 2p1/2 and Si 2p3/2 intensity ratio are notfixed in the fitting, the fitting yields reasonable results. The fitting yieldsthat the spin-orbit splitting is 0.6±0.05 eV and the intensity ratio isequal to the statistical value of 1:2 for all of the oxidation states.

As seen in Fig. 1(c) and (d), prolonging the annealing to 3 min, theatomic concentration of the Si4+ component is up-shifted and theatomic concentration of sub-oxide components is decreased. Note thatSi0 and Si1+ peaks don't appear in the whole process indicating Si0

does not exist in the a-C film. Huge difference of the atomic concentra-tions of the other four Si environments is that: for the as-sputtered film,the contents of Si3+ (38.73 at.%) and Si2+ (57.26 at.%) are high in com-parison with Si4+ (4.01 at.%); for the annealed one, there is a tremen-dous increase in the content of Si4+ (59.79 at.%) and sharp decreasein the concentration of Si3+ (29.53 at.%) and Si2+ (10.68 at.%). Uponannealing, the film is transformed from a randombondingmodel struc-ture into distinctmixtures of graphitic carbon and SiO2-like. The reasonsare summarized as follows. The annealing system is not a high vacuumdevice thus the presence of residual O2 is unavoidable. In addition, evenif the chamber pressurewith Ar (1.08×105Pa) is a little higher than theoutside atmosphere (1.01×105Pa), a small amount of O2 could still beintroduced into the chamber. Therefore, the additional O2 might havecome from the leakage of and the residual O2 in the annealing system.The presence of the additional O2 and the high annealing temperatureresult in the removal of carbon andcorresponding increase of oxidizationin the film. This, taken into consideration with an increase in the Si 2pbinding energy, intimates the formation of oxidized coatings which aremore SiOx>1.

In the previous work [30], the thickness of the a-C film is around345 nm deposited under the same target power density as in thepresent study. Since XPS can only detect the chemical compositionand bonding structures beneath the surface to around 10 nm. There-fore, the XPS survey spectra and high resolution spectra in [30] andthe present work look similar. However, there are still some differencesas listed below. a) The chemical composition shows difference: In [30],

of Si 2p core-level spectra of the (c) as-sputtered film, and (d) 3-min annealed film.

Fig. 2. Cross section of the (a) 570-nm a-C/30-nm Ni film, (b) 130-nm a-C/30-nm Ni film, and (c) 50-nm Ni film. Scale bar is 100 nm.

Table 1Details of the as-deposited a-C/Ni, Ni film and the as-grown nanowires.

Film structure Thickness ofthe a-C layer(nm)

Thickness ofthe Ni layer(nm)

Annealingduration(min)

Ratio of core diameterover sheath thickness

a-C/Ni 570 30 1 1.16±0.14a-C/Ni 570 30 3 0.25±0.09a-C/Ni 130 30 3 0Ni 0 50 3 0

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in as-sputtered film, C 1s (90.04±1.20 at.%), Si 2p (5.21±0.40 at.%) andO 1s (4.75±0.97 at.%), in as-annealed film, C1s (18.33±4.58 at.%), Si 2p(47.46±3.49 at.%) and O 1s (34.09±2.22 at.%). b) The sub-oxidesare different: In [30], the spectrum doesn't have Si0 and Si1+ forthe as-sputtered film, the Si sub-oxides in the as-annealed film wascompletely converted to Si4+.

To summarize, Ni catalyst particle doesn't move along with thewire. It remains at the bottom, rather than at the tip of the wire. Theamount of Si in the as-deposited film was very small. The as-annealedfilm has a close chemical composition to stoichiometric SiO2. Si0 doesn'texist in the as-sputtered film.

3.2. Morphology, chemistry and structure

3.2.1. Cross section of the filmThe typical cross section of the a-C film of 570 nm and 130 nm in

thickness are respectively shown in Fig. 2(a) and (b). The Ni-catalystlayer of around 30 nm is clearly observed between the a-C top layerand Si substrate in Fig. 2(b). However, it is not very clearly seen inFig. 2(a) due to the low contrast between the thick top layer and the Nilayer. But the white color of the Ni layer is also seen as highlighted bythe arrow at the interface. The thickness of the single Ni film is around50 nm as clear shown in Fig. 1(c). The details of the as-sputtereda-C/Ni and Ni film are summarized in Table 1.

3.2.2. Morphology of the wireFig. 3(a) shows the wires on the surface of the 600-nm a-C/Ni film

after 3-min annealing. The wires are of around 70 nm in diameter, upto 10 μm in length, entangling with each other. The wires become in-homogeneous with bright and transparent outer sheaths and a whiteinner core indicating a heterogeneous structure. Fig. 3(b) shows bun-dles of wires growing from the cracked film. Fig. 3(c) shows the wiresgrown after 3-min annealing of the 160-nm a-C/Ni film. The wires areof homogeneous structure and of round ends, lying on the film sur-face entangling into each other. Fig. 3(d) shows the wires grown via3-min annealing of the single Ni layer without the a-C layer on thetop. In summary, wire-like nanostructures grow during annealing ofthe a-C/Ni/Si-wafer and Ni/Si-wafer samples. Ni particle doesn't ap-pear at the tip of the wire.

Fig. 4(a) and (b) shows the structure of the nanowires after 1-minannealing of the 600-nm a-C/Ni film. The wires have amean outer diam-eter of 57.23±11.80 nm, amean core diameter of 20.97±5.07 nmand amean sheath thickness of 18.00±2.90 nm. The ratio of the core diameterover the sheath thickness is 1.16±0.14. After prolonging the annealing to3 min, the wire's outer diameter increases to 84.75±25.53 nm, and thecore diameter reduces a lot, down to 8.62±0.45 nm and the sheaththickness increased to 36.49±14.41 nm (cf. Fig. 4(c) and (d)). Theratio of the core diameter over the sheath thickness is 0.25±0.09,which is dramatically decreased compared to the 1-min annealing. Theend of the core is surrounded by a 12.1 nm thin sheath layer as shownin the inset of Fig. 4(d). Obviously, oxidation takes place inward to con-sume the core. Fig. 4(e) and (f) shows the nanowires via 3-min annealingof the 160-nm a-C/Ni film. Thewire in Fig. 4(f) is of 70.05 nm in diametershowing a homogeneous amorphous structure without an inside core.Fig. 4(g) and (h) shows the structure of the wires after 3-min annealingof the 50-nm Ni film without the a-C film. Homogeneous amorphouswires are also observed without the black core. The dark lines randomlyentangled with the Si nanowires do not belong to Si nanowires. Theyare the residual carbon dissolved in the solution during the preparationof the TEM samples. The details of the as-grown nanowires are summa-rized in Table 1.

In summary, with or without the a-C film on the top of the Nilayer, homogeneous or heterogeneous nanowires are able to form. Italso confirms that Si in the wire is from Si wafer. However, the coverlayer of a-C does provide oxidation reduction. As shown in Table 1,within the same annealing duration of 3 min, the ratio of the core diam-eter over the sheath thickness decreases from 0.25±0.09 to 0 as the a-C

Fig. 3. Wires grown via 3-min annealing of the (a)–(b) 600-nm a-C/Ni film, white arrow indicating the center of the wire, (c) 160-nm a-C/Ni film, and (d) 50-nm Ni film.

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layer thickness reduces from 570 nm to 130 nm (cf. Fig. 4(c)–(f)). Be-sides, for the same a-C layer thickness of 570 nm, prolonging theannealing from 1 min to 3 min gives rise to reduced ratio of the corediameter over the sheath thickness from 1.16±0.14 to 0.25±0.09(cf. Fig. 4(a)–(d)).

Fig. 4. TEM images of as-grown wires via (a)–(b) 1-min annealing of the 600-nm a-C/Ni fil160-nm a-C/Ni film, and (g)–(h) 3-min annealing of the 50-nm Ni film.

3.2.3. Chemistry and structure of the wireFig. 5(a) shows the cross-sectional elemental mapping of the

as-grown wire after 1-min annealing of the 600-nm a-C/Ni film. Asseen, it consists of Si and O only. The intensity of Si and O exhibits al-most the same parabolic function along the cross-sectional route, not

m, (c)–(d) 3-min annealing of the 600-nm a-C/Ni film, (e)–(f) 3-min annealing of the

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only indicating stable atomic concentration ratio (1:2) between Siand O along the route, but also implying that the cross section isvery round. Particularly, the intensity of O has a small plateau whilethat of the Si continues to increase while approaching the core ofthe nanowire, which could be regarded as an implication of the Sicore.

Fig. 5(b) shows a single Si nanowire consisting of a dark core witha diameter of 12.31 nm surrounded by a gray sheath of 21.17 nm inthickness. The ratio of the core diameter over the sheath thicknessis 0.58. The nanowire is straight with uniform smooth morphologyalong the growth direction. Inset in Fig. 5(b) shows the correspondingSAED patterns of the selected circle area of the wire. A unit cell ofcubic structure (lattice parameter=5.43 Å) was consistent with thediffraction pattern. Indexing of the SAED patterns shows the [111]zone axis of single-crystalline Si, which suggests that the nanowire isconstructed with crystalline Si growing along 110

h iorientation [34],

surrounded by amorphous SiOx sheath (i.e. c-Si/a-SiOx). Fig. 5(c) showsthe high resolution TEM image of the defect-free c-Si/a-SiOx nanowirevia 1-min annealing of the 570-nm-thick a-C film. The wire is consistedof a crystalline Si corewith diameter of 9.50 nm surrounded by an amor-phous SiOx sheath with thickness of 11.29 nm. The lattice spacing of3.84 Å agrees closely with that of the 110

� �lattice structure of silicon.

Fig. 5. (a) EDX cross sectional elemental mapping spectrum characterizes the chemicalconstructed with Si core and SiOx (x≈2) sheath, (b) TEM image of nanowire via 1-min annealolution TEM image of the nanowire grown via 1-min annealing of the 600-nm a-C/Ni film showFFT pattern of the core.

There is a dihedral angel of 120° between 110� �

and 011� �

(latticespacing, 3.84 Å) in the atom-resolved TEM image. The 2D fast Fou-rier transform (FFT) pattern of the lattice resolved core image isshown in the inset of Fig. 5(c). It is from the typical Si [111] zone axis.The FFT pattern in conjunction with the lattice-resolved core image re-veals the growth is along the 011

h iorientation.

3.2.4. Carbon bonding structureFig. 6 shows the transformation of carbon sp2 bonding structure

after 3-min annealing of the 570-nm film. The Raman spectra aredeconvoluted into D band (at ~1350 cm−1 A1g mode) and G band(at ~1580 cm−1, E2g mode) using amixture of Lorentzian and Gaussianfunctions [35,36]. D peak is the breathingmode of those sp2 sites only inrings [35]. Its intensity is directly dependent on the presence of thesix-fold rings. G peak is attributed to the in-plane stretching vibrationof any pair of sp2 bonded carbon atoms, either in C_C chains or in aro-matic rings [35]. ID/IG increases as the number of rings per cluster in-creases and the fraction of chain groups falls [37].

After annealing, G peak is up-shifted from 1533 cm−1 to 1588 cm−1.The corresponding ID/IG increases from 0.98 to 1.77. The second orderbands at approximate 2643 cm−1 and 2908 cm−1 are extruded. Theband at 2643 cm−1 is the most intensive one and attributed to the first

composition of the nanowire via 1-min annealing of the 600-nm a-C/Ni film to being of the 600-nm a-C/Ni film, inset shows the SAED pattern of the core, and (c) high res-ing the crystalline core and the amorphous sheath, the insets showing the corresponding

Fig. 6. Carbon sp2 bonding structure of the 600-nm a-C/Ni film (a) as-sputtered and (b)3-min annealed.

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overtone of the D band (around 2 times of 1350) [38,39]. An additionalhigher order band at 2908 cm−1 has been assigned to a combinationof the G and D modes characteristic for disturbed graphitic structures.The shape of the Raman spectra before and after annealing looks similarto the previous work in [30]. That's because they all underwent seriousgraphitization and oxidation after heating at 1100 °C. However, the dif-ference still exists after deconvoluting of the spectrum into G band andD band. The ID/IG ratio in [30] increases from to 0.94 to 1.29, while thatof the present study increases from 0.98 to 1.77, indicating the differentextent of graphitization. To summarize, once the temperature surpassesa certain point, the graphitization inclination and the crystallization isdramatically accelerated and strengthened. The tremendous increase ofID/IG, the huge up-shift of the G-band position, the appearance of thesecond-order bands at high Raman shift centers are attributed to grow-ing sp2 clusters and the conversion from sp3 to sp2 bonding configura-tion during the rapid thermal annealing.

3.3. Growth mechanism

3.3.1. Ni catalystFig. 7(a) shows a spherical Ni particle of around 38.1 nm in diam-

eter embedded at the end of a c-Si/a-SiOx nanowire. The hemisphericalparticle is enveloped by a thin SiOx layer (thickness of 5.7 nm) at the pe-riphery. The wire has an outer diameter of 32.8 nm, tapered inner di-ameter of 14.1 nm at the beginning (indicated by d1) and of 6.0 nm atthe top (indicated by d2). The tapered Si core indicates less oxidationat the wire end, implying the oxidation retardation effect of a-C layer.It has a clear interface with the Si core, suggesting that the Ni particlehas an orientation relationship or is partially coherent with the wire.The spherical shape suggests that the particlewas in liquid phase duringgrowth [13,14]. The typical lattice fringes of the Ni particle are clearlyobserved over the entire particle as magnified in Fig. 7(b). Fig. 7(c)shows a cucurbit-like Ni particle directing the growth of the amorphousSiOx nanowires. The particle has length of 119.0 nm and width of75.1 nm, surrounded by a layer of SiOx. The indexed SAED patternsshown in Fig. 7(d) are from Ni [113] zone axis [40]. These observationsconfirm that the Ni–Si phase does exist at the end of the wire and con-tributes to the growth of the wires.

3.3.2. Si sourceVLS growth of Si nanowires includes: a) absorption of Si atoms into

the liquid eutectic metal or alloy droplets from the vapor Si phase, the

liquid droplet serves as a preferential site; b) precipitation of Si nanowiresonce super-saturation of Si reaches, the liquid droplet serves as a nucle-ation site for crystallization; and c) termination of the growth when thetemperature is below the eutectic temperature of the catalyst metal oralloy, or the reactant is no longer available [20]. However, under thepresent conditions, Si concentration in the vapor phase is negligible atthe growth temperature because the specific surface/volume ratio ofbulk Si substrate is extremely low compared with that of the laserablation [3].

It is also different from the solid–liquid–solid growth of Park et al.[18], who grew c-Si/a-SiOx nanowires via 3 h heating a NiO-coated(solution method) Si substrate under a reductive environment of WO3,C andAr atmosphere. TheWO3 and Cwere in the formofmixed powdersbesides theNiO-coated Si substrate. However, in thepresent experiment,the sputtered carbon top layerwas separated by a layer of Ni. The longestannealing time is only 3 min. In addition, amorphous SiOx nanowiresdue to serious oxidation during annealing were also grown withoutsputtering the carbon layer on the top of nickel as shown in Fig. 3(d).Therefore, the disproportionation reaction of SiO vapor was excluded.The carbothermal reduction reaction between carbon and oxygen shouldbe consideredwhen the carbon layer is on the top. However, the functionof the reaction is only to retard the oxidation rather than generate Si andSiO2 through the disproportionation reaction of SiO vapor as in the workof Park et al. In addition, Si0 is not in the film as confirmed in Fig. 1(c).Therefore, the effect of small amount of Si in the as-deposited film canbe ignored. Si is neither from the vapor nor from the as-deposited a-Clayer. The only credible source is the bulk Si substrate.

3.3.3. Oxidation-accompanying solid–liquid–solid growth modelThe binary Ni–Si diagram shows that the eutectic temperature of

NiSi2 is around 993 °C [41]. Due to the melting effect of small-sizegrains [42], the eutectic compound NiSi2 can begin to form at a tem-perature lower than 964 °C. The sputtered Ni layer is able to melton the Si wafer substrate at the annealing temperature of 1100 °C,and form NiSi2 eutectic droplets. Considering the surface oxidation ofthe silicon wafer, the NiSi2 droplet is surrounded by Si–O. Because ofthe relatively high solubility of Si in NiSi2 eutectic alloy, more Si atomswould diffuse through the solid (substrate) liquid interface into the liq-uid phase. A second liquid–solid interface forms when the liquid phasebecomes supersaturatedwith Si atoms due to thermal or compositionalfluctuations, resulting in the growth of crystalline Si nanowire. The alloydroplets encapsulated with SiOx amorphous layer have been confirmed

Fig. 7. (a) Ni catalyst particle of 38.1 nm in diameter encapsulated at the end of the c-Si/a-SiOx wire, (b) high resolution TEM image of the Ni catalyst clearly showing the latticefringes, (c) cucurbit-like Ni particle encapsulated at the end of the amorphous SiOx nanowire and (d) the SAED patterns of the Ni particle are from the Ni [113] zone axis.

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in Fig. 7(a) and (c). The amorphous SiOx sheath is simultaneouslyformed due to oxidation. The growth involves not only the evolutionof solid–liquid–solid phases, but also relates to the simultaneous oxida-tion of the as-grown Si nanowire. Different from the Si-SiOx core–shellnanowires synthesized by Nie et al. without sputtering the top a-Clayer [25], the oxidation from the residual oxygen is effectively retardedin this work. The whole structure evolution responsible for the growthof c-Si/a-SiOx nanowires is thus developed taking place in the followingthree steps as schematically sketched in Fig. 8. To simplify the growthmechanism responsible for the evolution of Si nanowires, only theannealing duration is directly taken into consideration, whereas thea-C top layer is not drawn in the schematic diagram.

(I) Initiation of the Si nanowire1) Formation of the Ni–Si–O droplet. As sketched in Fig. 8(I), Siatoms are generated from the interface of Ni/Si wafer substrateat the annealing temperature of 1100 °C (which is higher thanthe eutectic temperature of the Ni–Si alloy), and then dissolveinto the liquid Ni to form Ni–Si droplet. The oxygen on the Siwafer surface partially oxidizes Ni–Si alloy to form Ni–Si–Oon the shell. 2) Precipitation of the c-Si/a-SiOx nanowire. Uponsuper-saturation of Si in the Ni–Si droplet, a concentration gradi-ent is formed due to excess Si at the interface between the Nilayer and Si substrate. Meanwhile, the gradual burning out of

carbon in the film at the Ar/O2 atmosphere gives rise to porousor cracked a-C film. As a result, the crystalline Si/amorphousSiOx (c-Si/a-SiOx) nanowire of dIC in core diameter precipitatesupward from the Si-supersaturated Ni–Si–O droplet. The crystal-line Si core is depicted with a preferred growth orientationb110> in the diagram.

(II) Oxidation and growth of the Si nanowireWith the precipitation of c-Si/a-SiOx wire, the residual oxygenin the reactor results in the inward oxidation of the c-Si core(dII

CbdIC). Note that the a-C film was partially cracked during

the annealing. A little bit of residual oxygen could further pen-etrate inside to oxidize the Ni–Si–O droplet.

(III) Transition from Si to SiOx nanowireAs time goes, the c-Si core is completely consumed and it com-pletes the transition from c-Si to a-SiOx nanowire. Fig. 8(III)shows the completion of the transition from c-Si to amorphousall-SiOx. SiOx layer envelopes the Ni particle inactivating it andending the growth of the wire, leaving a SiOx-rich surface.

High resolution TEM analysis reveals the presence of the Ni–Si–Ophase at the end of the wire, where the Ni crystal is embedded atthe end of the c-Si/a-SiOx nanowire (cf. Fig. 7(a) and (c)), and oxygenis resulted from the oxidation. Within the same time of annealing, asa result of the complete oxidation during annealing of the 160-nm

Fig. 8. Schematic growth and oxidation of the crystalline Si nanowire: (I) free Si atoms diffuse from Si substrate which dissolves into the Ni atoms until reaching a Ni–Si alloysuper-saturation (Ni–Si eutectic droplets), where Si is expelled from the Ni–Si core as a crystalline nanowire of dIC in diameter, the wire is depicted with a preferred b110> orien-tation. The oxygen on the Si wafer surface partially oxidizes Ni–Si alloy to form Ni–Si–O on the shell, thus the a-SiOx sheath surrounds the c-Si. The outer SiOx sheath is surroundedby the residual O2, (II) the c-Si core becomes smaller (dIICbdI

C) due to the inward oxidation, and (III) termination of the growth as the annealing time is up. The entire Si wire be-comes amorphous SiOx nanowire as a result of the complete oxidation. The Ni–Si–O droplet becomes a spherical Ni crystal surrounded by SiOx shell.

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a-C/Ni film, the entire c-Si core is oxidized into amorphous SiOx

nanowire (cf. Fig. 4(e)–(f)), whereas the 600-nm thick a-C/Ni layergives rise to c-Si/a-SiOx nanowires (cf. Fig. 4(c)–(d)). Furthermore,for the same a-C top layer, shorter annealing duration gives rise tothe c-Si/a-SiOx nanowire of bigger crystalline core diameter, thus higherratio of the core diameter over the sheath thickness (cf. Fig. 4(a)–(b)).

The driving force for such a continuous diffusion of Si atoms fromthe substrate into the liquid droplets, and then through the Ni–Si liquideutectic droplets into solid nanowires could be attributed to the kineticand environment aspects. From the kinetic point of view, the concentra-tion gradient and super-saturation due to the fluctuation in Ni–Si drop-lets are the driving force of the growth. From the environment point,the carrier Ar gas surrounds the semi-spherical shape Ni–Si dropletsand exchange energy andmomentumwith the atoms in theNi–Si drop-let, resulting in overcooling to the droplets. Such an overcooling iscritical to initiate the preferential unidirectional growth of c-Si/a-SiOx

nanowires. The resultant amorphous SiOx sheath is attributed to the ox-idation reaction between crystalline Si nanowire and residual O2. The a-Cfilm retards or attenuates the oxidation rate of the solid–liquid–solidgrowth of precipitated Si nanowire during rapid thermal annealing.Plus, the a-C gives rise to porous or cracked film to direct the growth ofSi nanowire during the annealing. Parameters such as the thickness ofthe a-C top layer, the annealing duration, Ni layer thickness, Ar flowrate, and annealing temperature might play important roles on themor-phology and structure of the nanowire.

In summary, at the annealing temperature of 1100 °C, eutectic Ni–Sidroplet forms at the Ni/Si wafer interface; here, Si atoms diffuse into the

Ni–Si droplet; upon super-saturation, Si precipitates and forms crystal-line Siwhich serves as the seed for Si nanowire.With time, crystalline Siwire forms. Meanwhile, residual oxygen in the annealing chamber oxi-dizes the crystalline Si wire to form amorphous SiOx sheath, giving riseto c-Si/SiOx structure. Prolonged annealing or reduced a-C layer thick-ness results in further inward oxidation of the crystalline Si core intoamorphous SiOx till the complete consumption of the c-Si core andthus the formation of amorphous all-SiOx nanowires.

4. Conclusions

Underneath amorphous carbon (a-C) layer, crystalline siliconnanowires can be grown directly from nickel-coated Si wafer sub-strate via rapid thermal annealing. The a-C layer has retarded theoxidation of the Si nanowire. Without the a-C layer, Si nanowire iscompletely oxidized to amorphous SiOx.Within the same annealing du-ration, thicker layer gives rise to bigger core diameter. For the samelayer thickness, prolonged annealing results in smaller core diame-ter. An oxidation-accompanying solid–liquid–solid model explainsthe growth of the Si nanowire and its oxidation process.

Acknowledgments

This work was supported by the Singapore Ministry of Education'sResearch Grant T208A1218 ARC4/08.

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