Swift heavy ion irradiation induced nanoparticle formation in CeO2 thin films

Nuclear Instruments and Methods in Physics Research B 269 (2011) 2786–2791

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Swift heavy ion irradiation induced nanoparticle formation in CeO2 thin films Mayora Varshney a,⇑, Aditya Sharma a,b, Ravi Kumar c,d, K.D. Verma a a

Material Science Research Laboratory, Department of Physics, S.V. College, Aligarh 202001, U.P., India Department of Applied Sciences & Humanities, Krishna Institute of Engineering and Technology, Ghaziabad 201206, U.P., India c Inter University Accelerator Center, Aruna Asaf Ali Marg, New Delhi 110067, India d Department of Material Science and Engineering, NIT, Hamirpur, H.P. 177005, India b

a r t i c l e

i n f o

Article history: Received 30 May 2011 Received in revised form 2 September 2011 Available online 8 September 2011 Keywords: Ion irradiation Thin films XRD AFM Raman

a b s t r a c t Nanoparticle formation in the, rf-sputtering grown, polycrystalline CeO2 thin films is achieved by the swift heavy ion (SHI) irradiation. Crystal structure and phases present in the as-grown and irradiated thin films are investigated by the X-ray diffraction (XRD) measurements. Irradiation induced formation of spherically shaped nanostructures, on the film surface, is confirmed by the atomic force microscopy (AFM). The Raman spectra of the irradiated samples show increased line-width and peak position shifting in the Raman active mode (F2g) of CeO2, indicative of the nanocrystallization in the irradiated CeO2 thin films. Formation of nanostructures in the irradiated samples is also briefly discussed in the light of ion energy and energy loss mechanisms. Ó 2011 Elsevier B.V. All rights reserved.

1. Introduction Nano-science/technology has become an attractive area of research as the nanodimensional systems/materials possess outstanding properties (i.e., optical, mechanical and electronic, etc.) and hence functionalities. The physicist finds nanomaterials as the quantum mechanically confined system where as the confinement in the dimensions is predictable by the size of the nanoparticle. The quality of altering all the physical and chemical properties by the small variation in size of the nanodimensional systems make them extremely suitable for a large variety of applications. There are various synthesis methods (i.e., sol–gel, co-precipitation, sputtering, laser ablation, and ion implantation/irradiation, etc.) by which doped and un-doped nanostructures can be synthesized [1–5]. In the field of nano-science/nanotechnology, the energetic ions are of use for, both, synthesis and modification [4,5] of the nanomaterials. The important issues in the nanomaterial synthesis are the (i) tailoring of shape of nanostructures, (ii) control of the size of particles and (iii) the size distribution. Ion beams play a significant role related to all these issues [5,6]. There are two ways of nanostructuring by ion beams namely; ion implantation and ion irradiation. In the ion implantation process the energetic ions have low energy (few tens of keV to few hundred keV). The desired ions can be implant in the suitable matrix by choosing the various implantation parameters like; type of ion, ion energy, range of ion, and ion fluence, etc. The synthesis of nanoparticles by ion implantation poses few problems. ⇑ Corresponding author. Tel.: +91 9319085147. E-mail address: [email protected] (M. Varshney). 0168-583X/$ - see front matter Ó 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.nimb.2011.09.001

First, when ions are implanted at low energies, their depth is very less and intersects the surface so that incoming ions may sputter previously implanted ions. When the ion fluence is increased (in low energy regime), the concentration of implanted element saturates. On the other hand, if the energy of the implanted ions is higher (to overcome problem of sputtering of implanted atoms), the duration of the implantation may be prohibitive because the concentration at the mean range decreases in inverse proportion to the profile width, which varies almost in proportion to the ion energy. Another major problem with ion implantation is the non-uniform size distribution of the so formed nanoparticles. This arises due to the concentration depth gradient of implanted atoms. Some efforts have been made to achieve a narrower size distribution of nanoparticles in the insulating and semiconductor materials via implantation followed by post annealing [4,6]. In contrast to nanostructuring by ion implantation the SHI irradiation induced nanocrystallization has emerged as more reliable and interesting technique. In case of SHI irradiation, the incident ions have velocity comparable to the orbital velocity of electrons of the target atoms. The energetic ions lose their energy by two distinct energy deposition processes; (i) electronic energy loss and (ii) nuclear energy loss. The electronic energy loss (Se or (dE/dx)e) is due to the inelastic collisions of the impinging ions with the atoms of the target material and this process dominant in the high energy regime (energy >1 MeV/nucleon) of the incident ion. On the other hand, the nuclear energy loss (Sn or (dE/ dx)n) is due to the elastic collisions of the incident ions with the target atoms and dominant in the low energy regime (energy of few keV/nucleon) of the incident ions. If material’s thickness is very less than the range of the projectile ions, the overall energy loss or the

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energy deposited by SHI beams is ascribed to the electronic energy loss process rather than the nuclear energy loss process. As a consequence of high energy deposition, within a very short span of time, track of molten zone is formed along the ions trajectory. This molten zone solidifies within the picoseconds. As a result of the rapid quenching of the material, nanocrystallization may take place within the matrix [6] or even near the surface of the sample [7]. Among the various known transparent conducting oxide materials (e.g., ZnO, SnO2, and In2O3, etc.) cerium oxide (CeO2) is one of the important candidate which have unique cubic crystal structure and high k-dielectric constant. CeO2 having a optical band gap of 3.3 eV and n-type carrier density which make it a promising candidate for various technological applications such as a buffer layer in silicon on insulator devices [8], basis in the field effect transistors [9], high kdielectric in the capacitors [10,11], catalysis [12], coating material [13], and electrolysis [14], etc. Attempts have, therefore, been made to improve its properties, and hence functionalities, by (i) creating nanostructures [15–17], (ii) transition metal ion doping [18–20] and (iii) bombarding with swift heavy ion (SHI) beams [21,22]. In case of ion irradiation effects on CeO2, this material is adopted for the simulation of radiation damage in oxide fuels (like; UO2, having same fluoride type crystal structure as of CeO2) in nuclear reactors which are subjected to high-energy particles interaction [23,24]. Although, some preliminary studies on the ion beam induced structural changes in crystalline films [25] and powder pellets [26] have been reported, but a detailed study on the SHI irradiation induced nano-crystallization and densification in the CeO2 thin films has not been carried out so far. In the present study, we report the synthesis of CeO2 thin films, on Si substrate, using RF-sputtering technique and their subsequent irradiation with 175 MeV Au ion beams at different irradiation fluences. We observed that irradiation at low irradiation fluence start to form narrow size distribution of particles. At the higher irradiation fluence spherical shaped and closely packed nanoparticles have been observed on the film surface. The formation of smaller sized nanocrystallites, in the irradiated thin films, is briefly discussed on the basis of energy loss mechanism, and experimental findings of XRD, AFM and Raman scattering measurements.

2. Experimental CeO2 ceramic targets, used in this work, were prepared from the commercial CeO2 powders (Sigma–Aldrich, 99.9% purity). The circular target of the diameter of 50 mm was formatted in a mould using a uni-axial pressure of 300 MPa. The target was sintered at 900 °C, in air, using a thermal cycle corresponding to a temperature upward and downward slope of 150 °C/h with a plateau of 2 h at maximum temperature (900 °C). The substrates (quartz) were cleaned with trichloro-ethylene and then acetone and methanol using an ultrasonic cleaner. The well cleaned substrates were loaded in the growth chamber having the base pressure of 6  104 Pa. The oxygen and argon gases were used (O2:Ar = 1:2) as the oxidizing reactant and the plasma generation gas, respectively. The sputtering power was 200 W and the substrates were kept at 600 °C temperature, during the deposition. The growth time was kept constant for all the samples. After the deposition the films were cooled down slowly to room temperature under the oxygen partial pressure of 240 Pa. After deposition, the film thickness was measured by a profilometer and found to 200 ± 10 nm. The well characterized CeO2 thin films were irradiated with 175 MeV Au13+ ions beams with two different fluences of 1  1012 ions/cm2 and 5  1012 ions/cm2 at room temperature using 15UD Tandem Accelerator at Inter-University Accelerator Center (IUAC) New Delhi, India. The ion beam was focused to a spot of 1 mm  1 mm and scanned over an area of 10 mm  10 mm using a magnetic scanner to achieve fluence uniformity across the

sample area which was typically 5 mm  5 mm. The fluence values were measured by collecting the charge falling on the sample mounted on an electrically insulated sample holder placed in secondary electron suppressed geometry. Ladder current was integrated with a digital current integrator and the charged pulses were counted using scalar counter. The structural study on un -irradiated and irradiated samples was done by glancing angle XRD using Brooker D8 advanced diffractometer with Cu Ka radiation (k = 1.540 Å). Surface morphology was studied by AFM using digital Nanoscope IIIa SPM, in tapping mode. The Raman scattering measurements were performed using In-Via Raman microscope. Excitation was provided by an argon-ion laser of wavelength 514 nm and a low incident power to avoid thermal effects.

3. Results and discussion Fig. 1(a)–(c) shows the glancing angle XRD pattern (plotted in log scale) of as deposited and irradiated thin films. The irradiation fluence, used to irradiate the sample, is mentioned in each panel of the figure. It is visible from the Fig. 1 that a high intense peak at 2h = 28.055° with two low intense peaks at 2h = 47.2° and 2h = 59.02° are observed in as-deposited and irradiated thin films. These peaks are assigned to (1 1 1), (2 2 0) and (2 2 2) plane of CeO2, having fluorite type structure (JCPDF#750390). It is visible from Fig. 1 that the intensity of the XRD peaks (e.g., peak (1 1 1)) is continuously decreased with increase of ion irradiation fluence. Beside this, the full width at half maximum (FHWM) of the XRD peaks is improved with increase in the irradiation fluence (see Fig. 1(d)). The average grain size is calculated by using the Scherrer relation; D = 0.9k/b cos h, where D is the average grain size, k is the wavelength of the used X-rays, and b is the FWHM of the diffraction peak. Thus calculated particle size is 13.82 nm, 11.86 nm and 9.72 nm for the as-deposited, 1  1012 ion/cm2, and 5  1012 ion/cm2 irradiated CeO2 thin films, respectively. It is clear that the crystalline size continuously decreased with increase in the irradiation fluence. The observed changes in the particle size, FWHM, and the XRD peak position may help us to understand the impact of SHI irradiation on the other structural properties like; lattice parameters values and strain present in the thin films etc. We calculated the lattice parameter (a) of as-deposited and irradiated thin films using the qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi 2 2 2 relation; a ¼ dðhklÞ ðh þ k þ l Þ, where d is the inter-planner distance and (hkl) is the miller index. The calculated unit cell parameters are found to 5.507 Å, 5.438 Å and 5.419 Å, respectively. This helps us to conclude that the cell parameter, and hence, unit cell volume shrink with increase the irradiation fluence. The percentage   a afilm strain is calculated using the relation r ¼ substrate  100 [27]. a substrate

Where, asubstrate , and afilm are the lattice parameters of substrate and film, respectively. The positive value of r corresponds to the tensile strain (the cell is elongated in the films plane and compressed along the out of plane growth direction). Whereas the negative value of r corresponds to the compressive stress (the cell is compressed in the films plane and elongated along the out of plane growth direction). Thus calculated strain comes to 1.77, 0.498 and 0.147, for the as-deposited 1  1012 ion/cm2, and 5  1012 ion/cm2 irradiated CeO2 thin films, respectively. It is clear from the calculations that the compressive strain is abridged with increasing the irradiation fluence. The net deduction in the lattice strain may cause the densification in the film material. Thus, our XRD results help us to conclude that ion irradiation, with different fluence, lead to (i) shrink in the unit cell parameter/volume, (ii) net deduction in the strain and, interestingly, (iii) generate smaller sized particles. To further examine the SHI irradiation induced formation of smaller sized particles in CeO2 thin films, systematic, AFM measurements were performed on the as-deposited and irradiated thin

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Fig. 1. (a–c) Glancing angle-XRD patterns of as-deposited and irradiated CeO2 thin films, (d) shows the evolution of the (1 1 1) peak, with different irradiation fluence.

films. Fig. 2 shows the two dimensional (2D) images of (a) asdeposited, (b) 1  1012 ion/cm2, and (c) 5  1012 ion/cm2 irradiated CeO2 thin films, respectively. It is clear from the Fig. 2(a) that, the RF-sputtering grown, un-irradiated CeO2 thin film, shows a few randomly distributed nanostructures, having different size and shape, on the film surface. From the morphology of the irradiated films (Fig. 2(b) and (c)), it is observed that irradiation with 175 MeV Au ions lead to modify the surface nanostructures. In case of 1  1012 ion/cm2 fluence irradiated sample (Fig. 2(b)), the nanostructures turn into the spherical shape and narrower in size. Beside this, the density of nanostructures is also increased. Further increase in the irradiation fluence (i.e., 5  1012 ion/cm2) leads to formation of smaller sized and more closely packed nanostructures on the film surface (Fig. 2(c)). However, some agglomerated lumps have also been observed in the higher fluence irradiated sample, but the morphology of the nanostructures remains spherical only. Thus our AFM results strengthened the formation of narrow sized particles in the irradiated CeO2 thin films and well support our XRD results. Thakurdesai et al. [28] have, also, examined the irradiation (at low fluence, i.e., order of 1012 ion/cm2) induced formation of smaller sized nano-phases of TiO2 by XRD and AFM measurements. However, at higher fluence (i.e., order of 1013 ion/cm2) the film show amorphous nature and nanocrystallization could not observe. In the present case, the irradiation fluence (at which

nanoparticle formation is observed) tallies the irradiation fluence used by Thakurdesai et al. It is evident that a complete characterization of structural and other properties of (i) single crystals, (ii) nanostructures and (iii) thin films requires the matching the results of different techniques. Here, our attention is addressed to Raman spectroscopy, which is widely used to study the various oxide materials [29–32]. Indeed, Raman scattering, even though not exhaustive, is able to give information on different phenomena like, determination of different phases of same material (e.g., anatase, and rutile phases of TiO2), [32], allowing the investigation of amorphous-to-crystalline phase transitions, oxygen defects, stress states and size effects etc. The latter fact derives from the possibility to manage the CeO2 structure at the nanoscale and/or in the radiation environment. A detailed analysis of Raman spectral features leading to the derivation of band parameters, such as energy peak, line-width, asymmetrical shape and relative intensity ratio, etc., and allow gaining important information to understand the functional behavior of the probed material. Fig. 3 shows the room temperature Raman spectra of as deposited and irradiated CeO2 thin films. For all the samples, spectra clearly shows six Raman lines and one kink positioned 301 cm1, 435 cm1, 464 cm1, 519 cm1, 610 cm1, 669 cm1 and 231 cm1, respectively. Peak at 464 cm1 is corresponding to the triply degenerate F2g Raman active mode of CeO2 thin film. The most intense peak at

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Fig. 2. AFM images of the as-deposited and irradiated CeO2 thin films.

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519 cm1 and other low intense peaks, including kink, are assigned to the various Raman active mode of Si [33]. The peak shape and intensity of the peak appeared at 519 cm1 is not affected by the irradiation. This indicates that the structural properties of Si substrate are not, significantly, affected by the irradiation and ions are expected to lose their energy within the film material. In case of a perfect crystal only the phonons near the center of the Brillouin zone (q0  0) contribute to the scattering of incident radiation due to the momentum conservation rule between phonons and incident light. As the size of the crystal is reduced, the vibration is limited to the size of the crystal, which gives rise to breakdown of the phonon momentum selection rule (q0  0), allowing phonons with q – 0 to contribute to the Raman spectrum [34,35] and leads to asymmetric broadening and dissimilarity in the Raman active modes. In the present case, the line-width of Raman active mode (F2g) of CeO2 is found to increase and the position of the F2g mode is also shifted toward the lower wave numbers with increasing the irradiation fluence (see the inset of Fig. 3). The observed Raman spectra is in accordance with the previously reported results on the evolution of F2g mode of CeO2 nanocrystals [36] which strongly depends on the confinement and/or strain effects due to the variation in the particle size and best described using spatial correlation model (also known as the phonon confinement model) [34,35]. Therefore, the present Raman spectra precisely validate the formation of nanoparticles in the irradiated samples. Hence, the Raman spectroscopy measurement supports the results of structural and microscopy measurements and strengthened the view of the irradiation induced formation of nanoparticles in the CeO2 thin films. Formation of smaller sized nanocrystallites, densification of the film material and changes in the structural properties of the CeO2 thin films can be understood by knowing the energy transfer processes of the energetic ions in the CeO2 matrix. In Fig. 4, we have plotted the depth versus Se and Sn of Au ions in the CeO2. The range of 175 MeV Au ions in CeO2 target is about to 12 lm, thus all the ions are expected to pass through the film thickness (200 nm) deep into the substrate. Therefore, the observed irradiation effects are to be analyzed as consequences of ion irradiation induced electronic energy transfer and not due to the implantation of ions in the film material. Two different models are often proposed to account for SHI irradiation induced material’s modification (i) The Coulomb explosion [37] and (ii) thermal spike [38]. In the Coulomb

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explosion model relaxation of localized excess of energy occurs via radial impulse of atoms lying in the surrounding area of the ion path, within the 1015 s. In this process a highly positive charged zone of target material explodes due to Coulomb repulsion before electrical neutrality is achieved. Therefore, this model predicts the ion irradiation induced amorphization and disorder in the target material. However, we have observed irradiation induced nanocrystallization in the CeO2 thin films, as judged form the evolution WFHM of XRD peaks, Raman spectra and AFM micrographs. Therefore, to understand such effects one has to consider the SHI irradiation induced formation of thermal spike. In the thermal spike model, energetic ions deposit their energy to the electronic subsystem of the target atoms and can lead to rising of the local temperature of the material via electron–phonon coupling [39]. The first step of energy deposition is the electronic excitation and ionization of the target atoms along the ion path. This step takes place in less than 1016 s. The local thermalization in the electronic sub-system is considered as the second step and completed in about 1014 s. The third and important step of the energy loss is the transfer of heat from the electronic to atomic subsystem, via electron–phonon coupling, and takes place between 1014 and 1012 s. The third step also depends on the magnitude of the coupling between the electron and phonon. As a result of electronic excitation and effective electron–phonon coupling, a cylindrical region around the track of the ion become fluid in the sense that any thermally induced shear stress relax in this region. After the time of more than few hundred pico-seconds the fluid region practically gets cooled down at ambient temperatures but usually in a transformed state [40]. The short term local melting of the material, followed by a rapid quenching of the liquid phase, leads to form observed nanocrystallites in the film material. The nanocrystallites are believed to nucleate and growth in size around the various point defects in the irradiated region [41]. Another possibility of irradiation induced formation of ordered nanostructures on the film surface or near the surface region is the lateral mass transport and dewetting [42]. Thermal spike model has been effectively used to understand the irradiation induced dewetting in the oxide thin films [43,44]. This model believes that when heavy ions, such as Au13+ in the present case, impinge on films a local melting occurs along the ion trajectory and persists for several picoseconds. Due to this melting, a radial temperature gradient is built up around the ion trajectory. The melting of material also establishes a gradient in the surface tension [43]. Due to this combination of the surface tension gradient and pressure gradient the molten zone forces the molecules onto the surface, away from the melt center. Therefore the viscous flow of molecules, driven by the capillary and elastic forces which originate from an imbalance of the surface and interface energy, leads to growth of nanostructures on the film surface or in the near surface region of the films [43,44]. Therefore, we believe that the formation of nanoparticles, on the CeO2 film surface, is the result of irradiation induced thermal spike and not due to the Coulomb explosion. In the present case, we have observed the growth of spherically shaped nanoparticles and densification of the film material. This may arise due to the fact that the spherical morphology presents minimum surface energy over the other morphologies; therefore, the CeO2 granulizes acquire spherical shaped nanoparticles during the re-growth of material under the ion irradiation. Besides the above discussed facts, the ion irradiation is also known to cause ultra fast solid–liquid–solid phase transitions [39,40]. Such anisotropic deformation by the rapid solid–liquid–solid transformation may lead to the densification of the film material, as predicted in the XRD results. However, studies on the ion beam irradiation induced nanoparticle formation may lead to further insight into the relationship between the ion’s energy loss processes, nature of film material, irradiation fluence and other important properties.

4. Conclusions High quality thin films of CeO2 were successfully deposited on Si substrates by RF-sputtering technique. The films were irradiated with 175 MeV Au beams at different fluence. It has been observed that FWHM of the XRD peaks is continuously increased with increasing irradiation fluence, an indicative of the formation of smaller sized nanocrystallites in the irradiated films. The AFM micrographs have clearly indicated the formation of spherically shaped nanoparticles on the film surface. The increased line-width and peak position shifting, towards to lower wave number, is likely to the breakdown of the phonon momentum selection rule (q0  0) and allowing phonons with q – 0 to contribute in the observed Raman spectrum. This is in accordance with the phonon confinement model and strengthened our view of the formation of narrower sized nanoparticles in the irradiated thin films. Our experimental results suggest that the irradiation induced thermal spike plays an important role in the formation of nanostructures in the irradiated films. Acknowledgements Authors (Aditya Sharma, Mayora Varshney and K.D. Verma) are thankful to Inter University Accelerator Centre, New Delhi, India for providing financial assistance under the UFUP research projects (codes; 41304 and 46302). References [1] W. Chen, J. Li, J. Appl. Phys. 109 (2011) 083930. [2] S. Colis, A. Bouaine, R. Moubah, G. Schmerber, C. Ulhaq-Bouillet, A. Dinia, J. Appl. Phys. 108 (2010) 053910. [3] A. Sharma, A.P. Singh, P. Thakur, N.B. Brookes, S. Kumar, C.G. Lee, R.J. Choudhary, K.D. Verma, R. Kumar, J. Appl. Phys. 107 (2010) 093918. [4] I. Szafraniak, I. Radu, R. Scholz, M. Alexe, U. Go Sele, Integr. Ferroelect. 55 (2003) 983. [5] U. Valbusa, C. Boragno, F. Buatier de Mongeot, J. Phys.: Condens. Matter 14 (2002) 8153. [6] D.K. Avasthi, Defence Sci. J. 59 (2009) 401. [7] S. Dhara, Crit. Rev. Solid State Mater. Sci. 32 (2007) 1. [8] T. Inoue, T. Ohsuna, L. Luo, X.D. Wu, C.J. Maggiore, Y. Yamamoto, Y. Sakurai, J.H. Chang, Appl. Phys. Lett. 59 (1991) 3604. [9] Y. Nishikawa, T. Yamaguchi, M. Yoshiki, H. Satake, N. Fukushima, Appl. Phys. Lett. 81 (2002) 4386. [10] A. Tiwari, V.M. Bosle, S. Ramachandran, N. Sudhakar, J. Narayan, S. Budak, A. Gupta, Appl. Phys. Lett. 88 (2006) 142511. [11] Y. Nishikawa, N. Fukushima, N. Yasuda, K. Nakayama, S. Ikegawa, Jpn. J. Appl. Phys., Part 1 41 (2002) 2480. [12] A.S. Deshpande, N. Pinna, P. Beato, M. Antonietti, M. Niederberger, Chem. Mater. 16 (2004) 2599. [13] V. Petrovsky, B.P. Gorman, H.U. Anderson, T. Petrovsky, J. Appl. Phys. 90 (2001) 2517. [14] I. Kosacki, T. Suzuki, V. Petrovsky, H. Anderson, Solid State Ion. 136–13 (2000) 1225. [15] V. Matolin, I. Matolinova, L. Sedlacek, K.C. Prince, T. Skala, Nanotechnology 20 (2009) 215707. [16] M. Li, S. Ge, W. Qiao, L. Zhang, Y. Zuo, S. Yan, Appl. Phys. Lett. 94 (2009) 152511. [17] L. Truffault, Q.W. Yao, D. Wexler, I.P. Nevirkovets, K. Konstantinov, T. Devers, S. Nightingale, J. Nanosci. Nanotechnol. 11 (2011) 4019. [18] Y.Q. Song, H.W. Zhang, Q.H. Yang, Y.L. Liu, Y.X. Li, L.R. Shah, H. Zhu, J.Q. Xiao, J. Phys.: Condens. Matter. 21 (2009) 125504. [19] Q.Y. Wen, H.W. Zhang, Q.H. Yang, Y.Q. Song, J.Q. Xiao, J. Appl. Phys. 107 (2010) 09C307. [20] P. Slusser, D. Kumar, A. Tiwari, Appl. Phys. Lett. 96 (2010) 142506. [21] N. Ishikawa, Y. Chimi, O. Michikami, Y. Ohta, K. Ohhara, M. Lang, R. Neumann, Nucl. Instrum. Methods Phys. Res. B 266 (2008) 3033. [22] M. Satoh, Y. Yamamoto, T. Inoue, Nucl. Instrum. Methods B 127/128 (1997) 166. [23] M.E. Cunningham, M.D. Freshley, D.D. Lanning, J. Nucl. Mater. 188 (1992) 19. [24] C.T. Walker, T. Kameyama, S. Kitajima, M. Kinoshita, J. Nucl. Mater. 188 (1992) 93. [25] M. Song, W. Weiguo, F. Chu, X. Yang, K. Mitsuishi, K. Furuya, H. Yasuda, Nucl. Instrum. Methods B 191 (2002) 586. [26] T. Sonoda, M. Kinoshita, N. Ishikawa, M. Sataka, Y. Chimi, N. Okubo, A. Iwase, K. Yasunag, Nucl. Instrum. Methods B 266 (2008) 2882.

M. Varshney et al. / Nuclear Instruments and Methods in Physics Research B 269 (2011) 2786–2791 [27] A. Sharma, D. Prakah, K.D. Verma, J. Optomechatronics Adv. Mater.: Rap. Commun. 1 (2007) 683. [28] M. Thakurdesai, A. Mahadkar, P.K. Kulriya, D. Kanjilal, V. Bhattacharyya, Nucl. Instrum. Methods B 266 (2008) 1343. [29] A. Dieguez, A. Romano-Rodriguez, A. Vila, J.R. Morante, J. Appl. Phys. 90 (2001) 1550. [30] A.K. Arora, M. Rajalakshmi, T.R. Ravindran, V. Sivasubramanian, J. Raman Spectrosc. 38 (2007) 604. [31] Y. Guhel, M.T. Ta, J. Bernard, B. Boudart, J.C. Pesant, J. Raman Spectrosc. 40 (2009) 401. [32] H. Rath, P. Dash, T. Som, P.V. Satyam, U.P. Singh, P.K. Kulriya, D. Kanjilal, D.K. Avasthi, N.C. Mishra, J. Appl. Phys. 105 (2009) 074311. [33] K.J. Krucnra, R.J. Hnvrrny, Am. Mineral. 79 (1994) 269. [34] A.K. Arora, M. Rajalakshmi, T.R. Ravindran, in: H.S. Nalwa (Ed.), Encyclopedia of Nanoscience and Nanotechnology, vol. 8, 2004, p. 499. ISBN:1-58883-064-0. [35] A. Dieguez, A. Romano-Rodrıguez, A. Vila, J.R. Morante, J. Appl. Phys. 90 (1550) 2001.

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[36] Z. Dohcevic-Mitrovic, Z.V. Popovic, M. Scepanovi, Acta Phys. Pol., A 116 (2009) 36. [37] E.M. Bringa, R.E. Johnson, Phys. Rev. Lett. 88 (2002) 165501. [38] M. Toulemonde, C. Dufour, E. Paumier, Phys. Rev. B 46 (1992) 14362. [39] D. Kaoumi, A.T. Motta, R.C. Birtcher, J. Appl. Phys. 104 (2008) 073525. [40] S.B. Ogale, Y.H. Li, M. Rajeswari, L. Salamanca Riba, R. Ramesh, T. Venkatesan, A.J. Millis, R. Kumar, G.K. Mehta, R. Bathe, S.I. Patil, J. Appl. Phys. 87 (2000) 4210. [41] T. Mohanty, P.V. Satyam, D. Kanjilal, J. Nanosci. Nanotechnol. 6 (2006) 2554. [42] J. Bischof, D. Scherer, S. Herminghaus, P. Leiderer, Phys. Rev. Lett. 77 (1996) 1536. [43] S. Kumar, S.K. Sharma, Allimuddin, R.J. Choudhary, D.M. Phase, R. Kumar, Nucl. Instrum. Methods B 266 (2008) 1741. [44] A. Sharma, K.D. Verma, M. Varshney, A.P. Singh, Y. Kumar, S. Srivastava, Y.K. Vijay, K. Asokan, R.J. Choudhary, R. Kumar, Adv. Sci. Lett. 4 (2011) 501.