Medium-energy ion irradiation of Si and Ge wafers: studies of surface nanopatterning and

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1 Home Search Collections Journals About Contact us My IOPscience Medium-energy ion irradiation of Si and Ge wafers: studies of surface nanopatterning and signature of recrystallization in 100 kev Kr + bombarded a-si This content has been downloaded from IOPscience. Please scroll down to see the full text Semicond. Sci. Technol ( View the table of contents for this issue, or go to the journal homepage for more Download details: IP Address: This content was downloaded on 09/02/2016 at 03:41 Please note that terms and conditions apply.

2 Semiconductor Science and Technology Semicond. Sci. Technol. 31 (2016) (9pp) doi: / /31/3/ Medium-energy ion irradiation of Si and Ge wafers: studies of surface nanopatterning and signature of recrystallization in 100keV Kr + bombarded a-si Pravin Kumar Inter University Accelerator Centre (IUAC), New Delhi , India vishakhapk@gmail.com Received 19 October 2015, revised 4 January 2016 Accepted for publication 11 January 2016 Published 8 February 2016 Abstract We report new and exciting experimental results on ion-induced nanopatterning of a-si and a-ge surfaces. The crystalline Si (100) and Ge (100) wafers were amorphized and an a/c interface was developed by pre-irradiation with a 50 kev Ar + beam at normal incidence with an ion fluence of ions cm 2. These amorphized surfaces were post-irradiated with Ar + and Kr + beams at an angle of 60. The post irradiation was done with ion fluences of ions cm 2. For each beam, two energies (50 and 200 kev for Ar +, 100 and 250 kev for Kr + ) were chosen to ensure ion stopping in both sides of the a/c interface. Regular nanopatterning (in the form of ripples) is observed on the Ge surface only with the post irradiation of the Kr + beam. The Si surface showed regular nanopatterning with the irradiation of both beams with two energies. For the ion beams crossing the a/c interface, ripples of higher amplitude and longer wavelength were formed. Further, the irradiation with a heavy beam yielded surface ripples of relatively larger amplitudes. The Raman measurements confirm amorphization of the pre-irradiated surfaces. Surprisingly, the post-irradiated Si surface with the 100 kev Kr + beam showed evidence of recrystallization. In the paper we discuss the physics at the interface and explain the experimental findings. Keywords: Ion beams, surface nanopatterning, interface, ion irradiation, recrystallization, Raman measurements (Some figures may appear in colour only in the online journal) Introduction The alteration of electron behaviour in low-dimensional (1 to 100 nm) materials makes them interesting to physicists as well as technologists [1, 2]. The fascinating properties of these materials are being studied worldwide for developing new technology. The main challenge lies in fabricating the materials with excellent shape and size distributions [3]. For fabricating such materials, various chemical and physical routes are available, each with some advantages/disadvantages. Ion irradiation is a widely used physical process to fabricate nanostructures on solid surfaces [4, 5] with great control over size and size distributions. Surface amorphization of a variety of materials by low- and medium-energy ions has been practised since 1960 [6]. Diverse patterns such as ripples and dots were observed, comprising self-organized periodic nanostructures on the surfaces bombarded by the ion beams [7, 8]. Many theoretical models have been proposed and modified to explain the nanoscale patterning in correlation with the surface and the ion-beam parameters [9 17], however, a generalized model that fits with the experimental observations of the surface patterning is still lacking. The dependence of surface nanopatterning on the angle of ion incidence is most crucial. Recently, Kumar et al showed that incompressible solid flow through an amorphous Si layer yields the formation of ripples at the a/c interface and hence on the surface [18, 19]. Moreno-Barrado et al [20] showed that non-uniform generation of the stress across the damaged /16/ $ IOP Publishing Ltd Printed in the UK

3 Figure 1. Schematic of the experimental plan followed in the present study. amorphous layer induced by the irradiation is the key factor behind the range of experimental observations. This effect, in synergy with the non-trivial evolution of the a/c interface, can be applied to a variety of semiconductors and to ion irradiation conditions for explaining surface nanopatterning. In the present study, the amorphized surfaces of Ge and Si (achieved by pre-irradiation with a 50 kev Ar + beam at normal incidence with an ion fluence of ions cm 2 ) are irradiated with Ar + and Kr + beams at 60 with an ion fluence of ions cm 2. To understand the role of the interface, the irradiation was done with two energies for each ion beam (50 and 200 kev for Ar + ; 100 and 250 kev for Kr + ). These two energies were chosen to ensure ion stopping above (with low energy) and below (with high energy) the a/c interface. We present the experimental results of the surface nanopatterning and discuss the role of the a/c interface position (in correlation with ion energy) on the nanoscale surface patterning. The mechanism of recrystallization in the 100 kev Kr + bombarded a-si is also discussed. Experiment The experimental plan is illustrated in figure 1. The ion irradiation was completed in two steps using the upgraded version of the low-energy ion beam facility (LEIBF) [21]. In first step, a 50 kev Ar + beam was pre-irradiated on Si and Ge surfaces at normal incidence with an ion fluence of ions cm 2. The beam current during pre-irradiation was maintained at 1 μa. This pre-irradiation amorphized the Si and Ge surfaces leading to an a/c interface in both materials. The depths of the a/c interfaces, as calculated using the TRIM software, were 72 nm and 53 nm in Si and Ge, respectively. A total of 5 samples of each material were prepared by pre-irradiation. In the second step, 4 samples of each material were bombarded by the 4 beams (50 kev Ar +, 200 kev Ar +, 100 kev Kr + and 250 kev Kr + at an angle of 60 with ion fluences of ions cm 2 ). The beam current during post-irradiation was maintained at 15 μa. Figure 2. AFM images (scan size; 10 μm 10 μm) of the preirradiated (50 kev Ar + beam at a fluence of ions cm 2 ; normal incidence) Si and Ge surfaces. The fifth sample of each material, which was not processed by the beams, is referred to as pre-irradiated Si/Ge in the manuscript and used as the reference. With lower energies, the stopping range of Ar + and Kr + ions in both materials is less than the depth of the a/c interface. With higher energies, both ion beams pass through the interface in Si and Ge. The samples (pristine, pre- and post-irradiated) were analysed by a Nano Scope IIIa atomic force microscope (AFM; Bruker AXS Inc., USA) under ambient conditions in the tapping mode. Micro-Raman analysis of the samples was carried out to observe the structural changes in the surface layer. For the micro-raman measurements, the Ar + laser was used to excite the samples at nm with 50 mw of power. The cross-sectional SEM measurements of the pre-irradiated and 200 kev Ar + beam irradiated samples were carried out to determine the exact depth of the a/c interface. To see the exact depth of the a/c interface in pre-irradiated and 200 kev Ar + irradiated Si and Ge samples, the cross-sectional SEM images were captured using a fieldemissionscanningelectronmicroscope (FE-SEM): MIRA II LMH from TESCAN with a resolution of 1.5 nm at 30 kv. In this model, the imaging is done using a secondary electron (SE) and a backscattered electron (BSE) detector. The recoil and recoil energy distributions for the irradiation by the 4 beams on a-si and a-ge surfaces were calculated using the TRIM software. The densities of amorphous and crystalline Si and Ge as a layer material (two layer systems) were used as an input to calculate the distributions. Results and discussion The surface morphology of pre-irradiated Si and Ge surfaces is shown in figure 2. The nanopatterning on the surfaces is 2

4 Figure 3. The AFM images of post-irradiated Si surfaces at an angle of 60 with an ion fluence of ions cm 2. (1) 50 kev Ar + (2) 200 kev Ar + (3) 100 kev Kr + (4) 250 kev Kr +. completely missing. In comparison with Ge, the Si surface is smoother. Figure 3 shows AFM images of post-irradiated Si surfaces. The surface patterning (nanoripples) is clearly visible with the irradiation using both beams (Ar and Kr) with two energies each. However, more ordered nanoripples are formed with the irradiation of Ar ion beam. In comparison with the 50 kev Ar + beam, irradiation with the 200 kev Ar + beam (which yields a/c interface 3

5 Figure 4. The AFM images of post-irradiated Ge surfaces at an angle of 60 with an ion fluence of ions cm 2. (1) 50 kev Ar + (2) 200 kev Ar + (3) 100 kev Kr + (4) 250 kev Kr +. crossing of the ions) causes surface nanoripples of reduced amplitude and enhanced wavelength (almost double). Irradiation with the 100 kev Kr + beam results in the formation of short range nanoripples. Under irradiation with the 250 kev Kr + beam (which also yields a/c interface crossing of the ions) and in comparison to 100 kev Kr + beam, reduction in the amplitude of the ripples was observed. The ripple wavelength does not change much for this irradiation condition. The surface morphologies of the post-irradiated Ge surfaces are shown in figure 4. Regular nanopatterning on the surface is completely absent with irradiation by the 50 kev Ar + beam. With the 200 kev Ar + beam, the surface patterns start to evolve. The 100 kev Kr + beam-irradiated surface 4

6 Figure 5. The cross-sectional SEM measurements of: (1) and (2) preirradiated Ge and Si, respectively. (3) and (4) 200 kev Ar irradiated Ge and Si, respectively. Figure 7. (a) Raman measurements of pristine, pre- and postirradiated Si surfaces. (b) Raman measurements of pristine, pre- and post-irradiated Ge surfaces. Figure 6. The vacancy profile of the pre-irradiated Ge and Si. The pre-irradiation was done with a 50 kev Ar beam at normal incidence. The a/c interfaces, as seen from vacancy profile, appear to be at 175 nm and 225 nm for Ge and Si, respectively. shows the formation of ordered nanoripples. The uniformity and the regularity of the surface nanoripples are broken by irradiation with the 250 kev Kr + ion beam. The increase in wavelength and amplitude is confirmed by irradiation with high-energy Kr + ions. Owing to the fact that there is not a well-defined saturation value of the surface roughness, the amplitude of the nanoripple (as a function of ion beam parameters) is not measurable with the help of existing theoretical models. At low surface temperatures (as is the case in the present study), the wavelength of the nanoripples varies linearly with the ion energy [22]. The present experimental values of the ripple wavelength at only two energy points have some consistency (particularly for Ar + beam irradiation on Si and Kr + beam irradiation on Ge) with the theoretical predictions. Understanding the role of the a/c interface on surface patterning is one of the key issues in the present study. Therefore, the depth of a/c interface, as calculated using the computer code SRIM [23], has been verified in a few samples with cross-sectional SEM measurements, which are shown in figure 5. The thickness of the a/c interface in the pre-irradiated and 200 kev Ar irradiated samples is much higher than that calculated by SRIM. This difference could be due to the fact that a collision cascade propagates beyond the ion range with reduced events. Therefore, the region beyond the ion range is also modified and a sharp a/c interface is not achieved in the SEM images. The vacancy profile shown in figure 6, as calculated by TRIM in pre-irradiated Si and Ge, fits well with the cross-sectional SEM measurements. The Raman measurements of the pristine, pre- and postirradiated Si and Ge surfaces are shown in figures 7(a) and (b), respectively. The pre-irradiation induces complete amorphization of the two surfaces as evidenced by the broadening and shifting of Si-Si ( 520 cm 1 ) and Ge-Ge 5

7 Table1. The TRIM calculations for the chosen beam irradiation on Si and Ge surfaces. The yields are per ion irradiation, and the irradiation of 2000 ions was considered. Beam irradiation Sputtering yield (S) in Si Backscattering yield (B) in Si S/B insi Sputtering yield (S) in Ge Backscattering yield (B) in Si S/B in Ge 50 kev Ar kev Ar kev Kr kev Kr ( 300 cm 1 ) LO phonon peaks towards higher wavelengths. In the present study, the nanoporosity (if any) beyond the amorphous phases as seen by Impellizzeri et al [24] in thin films is difficult to check with Raman spectra. However, a very recent study [25] shows the formation of voids/porosity in the amorphous phase of the low-energy bombarded Ge wafer and the subsequent repair by thermal and athermal treatments. The nanoporosity in the crystalline structure is characterized by the shifting of sharp Raman peaks towards higher wavelengths [26]. The post-irradiation treatments, except for 200 kev Ar + on Si, lead to a higher degree of amorphization of both surfaces. Some improvement in the Raman intensity due to the irradiation of the 200 kev Ar + beam in Si might be due to statistical error (sample-to-sample variation). But the appearance of the Si-Si LO phonon peak at 500 cm 1 along with a broad shifted peak (like in the other pre- and post-irradiated samples) upon irradiation with the 100 kev Kr + beam is a clear signature of recrystallization in a-si. The sputtering and backscattering yields calculated by TRIM (as included in the SRIM software) for each postirradiation condition are shown in table 1. It can be seen that the backscattering yield in Si is less than 7% for all beam irradiations. However, it varies from 10 to 20% for ion irradiation of Ge. All samples for which the backscattering yield is less than 14% show regular patterning. To explain ion-induced surface nanopatterning, Bradley and Harper (B-H) [27] proposed an erosion-based theory that combines curvature-dependent sputtering with a temperature-dependent mass redistribution/diffusion resulting in surface instabilities and surface smoothing, respectively. The transition from smooth surfaces at normal ion incidence to rippled surfaces at some critical angle could not be explained by the B-H model. Carter and Vishnyakov [28] removed this discrepancy for small angles and showed the existence of atomic flux parallel to the surface, which is generated during ion bombardment, counteracting the BH instability. A critical angle as the starting point for the transition was also predicted by this model. Madi et al [29] suggested that the curvature coefficients depend on the erosive and the mass-redistribution components. Norris et al proved this dependence with the help of molecular dynamics simulations [30]. Castro et al and Norris et al showed that surface nanopatterning is driven by stress-induced solid flow [31, 32]. Teichmann et al [33] investigated low-energy ( a few kev) inert ion-induced nanopatterning on Ge surfaces and showed that for ions with masses lower than the mass of the Ge target atoms (Ne, Ar), no pattern formation occurs and surface smoothing is observed for all angles of ion incidence. In contrast, for erosion with higher-mass ions (Kr, Xe), the ripple formation starts at incidence angles of about 65 depending on the ion energy. For low-energy inert ion irradiation on Ge surfaces, the curvature-dependent coefficient changes its sign at 45 representing the transition from stability (smoothening) to instability (nanopatterning). The amplitude of the curvature-dependent coefficient increases with the mass and the energy of ions. The experimental results were explained with the help of the angular distribution of the sputtering yield and backscattering/reflection of ions. The smoothening of a rough surface by the backscattered/reflected ions is beautifully explained by Hauffe [34]. The results presented here for medium energies as explained by angle-dependent sputtering yield and the backscattering of ions (see table 1), particularly on Ge surfaces, are in perfect agreement with the low-energy experiments. Therefore, the ratio of the sputtering yield to the backscattering yield is an important factor. The position of the a/c interface in correlation with ion energy does not seem to hold for nanoscale surface patterning in general. Molecular dynamics and continuum viscous flow models along with experiments using the bombardment of lowenergy Xe + and Ar + ions on an Si(100) target showed that the generation of non-uniform stress across the damaged amorphous layer induced by the irradiation explains surface nanopatterning on a variety of semiconductors [20]. The densities of Ge and Si are 5.3 and 2.3 gm cm 3, respectively. At a 60 angle of incidence, the longitudinal ion straggling of 200 kev Ar + and 100 kev Kr + in Ge is 50 nm and 15 nm, respectively. The mass of a Kr atom is 2.1 times of the mass of an Ar atom. Therefore, the free volume to accommodate ions in Ge is less for 100 kev Kr + irradiation and more stress across the damaged amorphous layer is expected in this case. Similar qualitative arguments can be made for other medium-energy irradiation conditions to justify the role of non-uniform stress across the amorphous layer in nanoscale surface patterning. The recoil ion energy and target atom distributions in the post-irradiated Si samples are shown in figures 8(a) and (b), respectively. In calculations, the layer 1 (the depth of the sample processed by the beam; the projected range of ions in Ge) and layer 2 (the bulk region unaffected by the ion beam) are considered to be fully amorphized and crystalline, respectively. The material densities (1.7 gm cm 3 for a-si and 2.23 gm cm 3 for c-si) were used accordingly to deduce the 6

8 Figure 8. (a) The recoil ion energy distributions in Si for irradiation by: (1) 50 kev Ar + (2) 200 kev Ar + (3) 100 kev Kr + (4) 250 kev Kr +. The black vertical line represents the position of the a/c interface. (b): The recoil atom distributions in Si for irradiation by: (1) 50 kev Ar + (2) 200 kev Ar + (3) 100 kev Kr + (4) 250 kev Kr +. The black vertical line represents the position of the a/c interface. 7

9 distributions of target recoil atoms and their energies. In figure 8(a), the recoil ion energy is 20 ev when the 100 kev Kr + ion beam is used to bombard the sample (see a3). The number of recoil target atoms/ions that have energies of 20 ev in a-si is more than that in c-si (see figure 8(b3)). The molecular dynamic simulations by Marqués et al [35] showed that the introduction of low-energy recoils in the amorphous matrix and directed at crystalline grains can initiate the recrystallization process. It was further shown that by varying the recoil ion energy and the position of launching the recoils near the interface, we can either amorphize the system, recrystallize it or maintain a stationary state [36]. The measured efficiency of recrystallization was 67% for both recoil energies, 20 ev and 15 ev. Therefore, we believe that the recrystallization in 100 kev Kr + beam-irradiated a-si (in the present experimental study) is induced by the low-energy recoils generated near the interface. Conclusion The a/c interfaces of Si and Ge were prepared by the irradiation of 50 kev Ar + beam (at normal incidence) with an ion fluence of ions cm 2. The pre-irradiated Si and Ge samples were exposed to various ion beams (50 kev and 200 kev Ar + ; 100 kev and 250 kev Kr + ) at an angle of 60. The post irradiation was carried out with an ion fluence of ions cm 2. Two energies for each beam were chosen to ensure ion stopping either side of (above or below) an a/c interface in each material. The pre-irradiated samples undergo complete amorphization and do not show any nanoscale surface patterning. The Si surface shows regular nanopatterning with post irradiation of both beams with two energies. Regular nanopatterning on the Ge surface is not observed with post irradiation using Ar + beams. However, post irradiation of Ge surfaces with Kr + beams results in regular surface nanopatterning. The results, particularly on Ge surfaces, show perfect analogy with the low-energy experiments, the results of which are explained in terms of ioninduced sputtering and backscattering yields. Qualitatively, the generation of stress across the amorphous layer under various irradiation conditions also supports the experimental results. The position of the a/c interface in correlation with ion energy does not seem to be an important parameter for nanoscale surface patterning. The unexpected recrystallization of a-si under irradiation with a 100 kev Kr + beam (at RT) is induced by the low-energy recoils generated near the interface and directed towards the crystalline grains. Acknowledgments The help received from Kedar Mal for conducting the ion irradiation experiments, S A Khan for SEM measurements, F Singh for Raman measurements and I Sulania for AFM measurements is greatly appreciated. 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