Abstract
Stacking fault free and planar defects (twin plane) free catalyzed Si nanowires (Si NWs) is essential for the carrier transport in the nanoscale devices applications. In this chapter, In-catalyzed, vertically aligned and cone-shaped Si NWs arrays were grown by using vapor–liquid–solid (VLS) mode on Si (111) substrates. We have successfully controlled the verticality and (111)-orientation of Si NWs as well as scaled down the diameter to 18 nm. The density of Si NWs was also enhanced from 2.5 μm−2 to 70 μm−2. Such vertically aligned, (111)-oriented p-type Si NWs are very important for the nanoscale device applications including Si NWs/c-Si tandem solar cells and p-Si NWs/n-InGaZnO Heterojunction LEDs. Next, the influence of substrate growth temperature (TS), cooling rate (∆TS/∆𝑡) on the formation of planar defects, twining along [112] direction and stacking fault in Si NWs perpendicular to (111)-orientation were deeply investigated. Finally, one simple model was proposed to explain the formation of stacking fault, twining of planar defects in perpendicular direction to the axial growth direction of Si NWs. When the TS was decreased from 600°C with the cooling rate of 100°C/240 sec to room temperature (RT) after Si NWs growth then the twin planar defects perpendicular to the substrate and along different segments of (111)-oriented Si NWs were observed.
Keywords
- silicon nanowires (Si NWs)
- VLS growth mode
- contact angle
- vertically aligned
- In-catalyst
- twining plane defects
- stacking fault
- nanoscale devices
- solar cells
1. Introduction
Recently much interest has been developed to control the band gap as a function of diameter of Si nanowires (Si NWs) to exploit the quantum size effect for photovoltaic applications [1, 2, 3, 4, 5], and its extension according to Moore’s law in view of the ongoing downscaling of integrated circuits (ICs) technologies as well as nano devices. Specially Si NWs are remarkably important for the fabrication of nanoscale devices such as transistors [6], sensors [7, 8], and thermoelectric devices [9]. Vapor–liquid–solid (VLS) mechanism provides a unique opportunity to investigate the crystalline quality and structures of single NW, where density, orientation, and periodicity of Si NWs system can be influenced by growth parameters such as temperature, pressure, plasma treatment, dopants [10, 11, 12], and the type of catalyst [13, 14], surface condition of substrate as well as size [15] of the metal Nano-droplets (NDs), shown in Figure 1. Si NWs growth by VLS mode using various material catalysts, such as Au, Al, Ga, In, Pb, Sn and Zn have been reported [16, 17, 18, 19, 20, 21, 22, 23, 24, 25].
Many researchers already grown vertically aligned Si NWs using Au-catalyst, which is not useful candidate for the application of nanoscale devices including solar cell and LEDs because it creates deep acceptor energy level at 0.54 eV in the Si band gap, whereas In-catalyst creates shallow accepter energy level at 0.16 eV in the Si band gap. Au-catalyst particles are strongly degrading the minority carrier life time, while In-catalyst particles are boosting to the carrier life time [26]. Previously, randomly oriented Si NWs were grown by Jeon and Kamisako
The defects study like twin planar defects as well as stacking fault were not rigorously investigated in the case of vertically aligned In-catalyzed (111)-oriented Si NWs. However, relatively very few investigations have been made about the stacking fault and twin planar defects in In-catalyzed Si NWs grown by VLS growth [26]. Zhan
2. Experimental methods
The basic mechanism behind the VLS growth mode is the transformation of the solid metal catalyst nanoparticle into a liquid alloy of the catalyst and compound of the semiconductor. In this case the liquid particle acts as a privileged site for Si deposition (precipitation via liquid catalyst), and has higher sticking coefficient as compared to the solid surfaces, shown in the schematic flow mechanism of Figure 1 [37]. Two different type of experiments were conducted to grow Si NWs. First, before the air-breaking condition p-type 300 μm-thick Cz-Si (111) substrate having resistivity of 1–10 Ω-cm, was cleaned by RCA washing. Next, In-NDs were grown on Si (111) substrate, using a conventional thermal evaporation system by evaporating pure In wire with base pressure (
In the next experiment everything was grown in the same sputtering chamber (without air-breaking condition). First, 300 μm-thick p-type Cz-Si (111) substrate, having resistivity ~1–10 Ω-cm was washed by RCA washing. Soon after RCA washing and drying the wafer was transferred to the plasma assisted and high vacuum sputtering chamber having background pressure,
The interface scenario between In-NDs and Si (111), surface morphologies, shape, density, and contact angle (
3. Results and discussion
Indium metal has a low melting point, and the eutectic temperature of the In–Si binary system coincides with the melting temperature of indium at 157°C [20, 38]. It is also worthy to note that the In–Si eutectic alloy exhibits a steep liquidus line, such that the liquid alloy can promptly be supersaturated with Si in a wide range of temperatures (at least up to 800°C) and In-Si eutectic alloy has an extremely low Si solubility approximately ∼10−4 at.% Si [26].
In this work the In-NDs were grown with the optimized growth condition, where the In-deposition time were increased from 20 min to 30 min in the plasma assisted crystal growth reactor as shown in Figure 2a and b. By decreasing the
In the second experiment (without air-breaking), we optimized the growth conditions to obtain good interface between In-NDs and Si substrate by eliminating the oxide layer to enhance the wettability. All the in-situ growth steps starting from In-NDs growth, plasma treatment, and up to the crystal growth of Si NWs were performed in a relatively clean and high vacuum chamber without air-breaking (sample-Nw). The θC of In on Si (111) was reported to be approximately 125° at 350°C by Mattila
Alet
We can see the planar defects as well as twining defects appearing in many segments of the Si NWs (sample-Nw) as shown in the inset of Figure 5b by green color arrow as well as red color arrow along the 〈112〉 direction. The (111)-oriented Si NWs at segments “P1–4” were also confirmed by HR-TEM image, as shown in Figure 4b. The corresponding selected area diffraction (SAED) shown in the in-set of Figure 4b has been recorded along the [110] zone axis from the Si NWs and which confirm the single crystal nature and its axial direction. But you can see twining of planes along 〈112〉 direction, which are distributed along the axial direction marked as S-I and S-II, as shown in Figure 4b for single (111)-oriented Si NW.
It can be established that the HR-TEM image of single Si NWs along with the SAED at “P1” and “P3” of Si NWs (taken from Figure 5b), marked by green color arrow, follow the same crystallinity of Si-substrate orientation marked at “P6” in Figure 5b. The diamond structure of the Si NWs with a lattice constant = 0.543 nm was confirmed by HR-TEM image and the SAED pattern at segments “P1” and “P3”, which corresponds to the space group Fd3m grown along the 〈111〉 direction. Unfortunately, the SAED at “P2” and “P4” (taken from Figure 4b) of Si NW shows small tilt with respect to orientation of Si-substrate, and the same have been marked by red color arrow, as shown in HR-TEM image given in Figure 5b. Spotty pattern both at “P2” and “P4” segments of the NW were confirmed by SAED observation, as shown in the inset of Figure 5b and same value of small angular tilt of both “P2” and “P4” segments were observed. We concluded that the Si NWs are vertically aligned except the twining of planar defects, which might be caused by the faster cooling rate 100°C/6 min, as well as due to the longer exposer time of downside wall to the reactive radicals as compared to the upper side of Si NWs [46]. We have to consider about the thermal conductivity mechanism, which is different for Si NWs as compared to the bulk Si-wafer [35].
Here we focus to discuss about the Si NWs crystallinity at the interface of emanating Si NWs from the Si-substrate, as shown in Figure 5a. Figure 5b gives the SAED pattern at point “P5”, which confirm the stacking fault in the grown Si NWs. The red arrow intensity originates from 1/2{111} spots is related to the 2H-polytype (stacking ABAB…) and the green arrow intensity originates from 1/3{111} spots is related to 9R (stacking ABCBCACAB…) [47].
Such structures are either attributed to scattering phenomena from two overlapping crystals with a stepped {111} twin boundary (parallel to the electron beam) [48, 49], or it might be the direct evidence of a 9R-polytype [23]. Furthermore, it has been established, when the crystallographic direction of the lattice abruptly changes in the In-Si material system then stacking fault may generated. Especially, when two crystals parts begins to grow separately and then meet at certain point, where the crystallographic direction remains the same, but each side of the boundary has an opposite phase. These kinds of stacking fault can be formed due to the complex dynamics of the In-NPs migration as well as mixing of the Si-atoms from the top of the Si NWs as well as precipitated Si atoms too. Such complex scenario will be explained later with the help of model given in Figure 7.
We realized that the In-NPs migration from the top of Si NWs toward the unused In-NDs and Si-substrate interface, where Si NWs are emanating from Si-substrate may cause to the stacking fault. We also know that an isolated defect like {111} faults have been observed to trap Au-atoms [34]. Therefore one cannot negate the possibility of In-atoms trapping by the planar defects of Si NWs during the VLS mode growth. Figure 6a gives the dark field-STEM image of the Si NWs, where one can clearly see the white spherical contrast of In-NPs around the side wall of the Si NWs. The compositional investigation by EDX taken at top of the Si NWs as well as taken from the pure In-NPs around the side wall of the Si NWs has been shown in the Figure 6b. The Kα X-ray energies for the In is 24.21 keV, and Lα X-ray energies for the same elements is 3.287 keV. Lα lines of the In-atoms can be separated, and this technique can be quantitatively used in a SEM. As shown in Figure 6c and d, we observed a reasonable mapping of Si and In sources, which originate from the Si NWs top and from the pure spherical In-NPs on the side wall of the Si NWs structure, respectively. The same phenomenon for Au NPs on the side of the Si NWs was observed by Krylyuk
The In-catalyzed Si NWs grown by VLS mechanism confronted with planar defects, twining and stacking fault, which were observed by HR-STEM, as shown in Figures 4a,b and 5a,b. To explain about the In-NPs migration to Si-substrate, Si NWs tapering, stacking fault, as well as limiting of NWs length by In-NPs migration, one simple model was anticipated. In this model, as a first step (Step-I), the In-NPs on Si(111), investigated, where most of the In-NDs got good wettability after Ar/H2-plasma treatment at 600°C with semispherical shape and few In-NDs were got spherical shape as shown in Figure 2a as well as in Figure 3a and the same is depicted schematically in Figure 7, Step-I. The spherical shape In-NDs may be confronted with thin oxide problems at the interface, shown in Figure 7 (Step-I). During the Si NWs growth as a second step (Step-II) we observed that the In-NPs migration from the top of Si NWs toward the Si-substrate are taking place and mixing with already grown In-NDs having thin oxide interface layer between Si-substrate and In-NDs, as shown in Figure 6a–d and the same is depicted schematically in Figure 7. Due to the migration of In-NPs toward the Si-substrate some new contact points on Si-substrate were created and the already In-NDs size were enhanced or elongated.
The supersaturation of the droplet, induced by the continuous gas phase supply of Si species, leads to the precipitation of Si nanowhiskers at the interface between the particle and the substrate, shown in Figure 1. Growth is obtained, as shown in Figure 7, Step-II, when a steady-state condition is reached between the flux of the Si through the In-NDs and the precipitation of Si on the Si-substrate [26, 37]. At the same time Si-atoms from sputtering source are also adsorbed on the same surface of In-NDs already deposited In-NDs on substrate having weak oxide interface layer, where the super saturation limit may be exceeded and now the precipitation of In-NPs were initiated as depicted in Figure 7, Step-II and the same has been confirmed in Figure 6a–d. Previously, the boron precipitation limit in BaSi2 were studied by in-plane and out of plane XRD, HR-STEM and TEM measurement and then successfully overcome the boron precipitation issue in bulk thin film
Suppression of the In-NPs migration from the top of the Si NWs are essential to grow longer Si NWs as well as to avoid NWs tapering and also to fix the stacking fault in Si NWs. We need to find new growth condition to suppress the In-NPs migration to find the lower optimal substrate temperature without compromising on the verticality control of Si NWs. The ultra-clean interface between Si-substrate and In-NDs is essential to get smaller
4. Summary
Using In-catalyzed VLS mode growth, we have successfully controlled the verticality and (111)-orientation of Si NWs and ultimately scaled down the diameter of NWs to 18 nm. The density of vertically aligned Si NWs was enhanced from 2.5 μm−2 to 70 μm−2. During the
Acknowledgments
This work was partially supported by the MEXT, FUTURE-PV Innovation (FUkushima Top-level United Center for Renewable Energy Research–Photovoltaics Innovation) Project.
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