Micro-Raman and electronic structure study on kinetics of electronic excitations induced monoclinic-to-tetragonal phase transition in zirconium oxide films

Abstract

Monoclinic-to-tetragonal phase transformation (PT) in sputtering grown zirconium oxide (ZrO₂) films on silicon substrates by electronic excitation (EE) induced by swift heavy ion (SHI) irradiation is reported. The density of EEs and the fluences of irradiation were varied for the better insight of phase transformation kinetics. The phase transition is well evident from the investigations using grazing incidence X-ray diffraction (GIXRD) and micro-Raman spectroscopy (mRS). Studies reveal a PT from the monoclinic to tetragonal phase. It is noted that at high fluence of Ag ion irradiation partly PT to cubic phase is also observed. However, it is clear from this study that this PT is not only due to transient temperature induced by SHI, but also attributed to the strain in the lattice created under the influence of the induced density of defects in the lattice. Interestingly, it may be noted that strain is well evident by the stiffening of the characteristic Raman modes of monoclinic phase. The modifications in electronic and local structure revealed using soft X-ray absorption spectroscopy (XAS) and X-ray absorption fine structure (XAFS) and found after fitting of Zr K-edge XAFS that phase transformation from m-ZrO₂ to t-ZrO₂ and/or c-ZrO₂ upon Ni and Ag irradiation. Studies would elucidate a deeper understanding about the kinetics of PT under such non-equilibrium conditions.

---

Introduction

Zirconium oxide (ZrO₂) also known as zirconia, is a scientifically and technologically crucial material as it possesses high melting temperature, high refractive index, low thermal conductivity, hardness, and corrosion barrier properties. It can be used in a wide range of industrial applications such as an oxygen sensor in fuel cells,¹ catalytic support medium,² durable thermal barrier coatings and optical coatings absorbing harmful ultraviolet radiation. It is also proposed as a gate dielectric material in metal-oxide semiconductor devices.^3,4 Interestingly, zirconia is also known as one of the most radiation hard ceramics,^5,6 which leads to extensive importance in the nuclear industry as a passivating medium for hydrogen ingress in pressure tubes. The deeper understanding on the evolution of the structural and electronic properties of zirconia upon variation in temperature and pressure in the polymorphs of pure zirconia i.e. monoclinic, tetragonal or cubic phases is the subject of intense experimental and theoretical studies. It is known that at high-temperature it has the cubic structure. However, upon increasing in external pressure the monoclinic phase transforms into an orthorhombic phase. The tetragonal and cubic phases of zirconia could be stabilized by extrinsic doping of yttrium.

Interestingly, crystalline to crystalline phase transformations (PT) such as monoclinic to tetragonal phase by swift heavy ions (SHIs) have also been reported along with its dependence on irradiations parameters. SHIs of energy ≥1 MeV per nucleon with materials has enormous potential to modify material structure and thereby its properties. It is well established that these modification are the effect of energy loss by energetic ions via two independent processes, firstly electronic energy loss (S_e) dominates over second one which is nuclear energy loss (S_n) under high energy regime. Most of the previous work has shown that the ion irradiation produces damages, crystalline-to-amorphous and/or crystalline-to-crystalline phase transformation along their trajectories in many insulators/semiconductors where later processes have its dependence on certain threshold value of electronic energy loss. The kinetics of these transformations has shown that most of them are the result of one ion impact mechanism leading to formation of latent tracks. Two major models as Coulomb explosion^7,8 and thermal spike^9,10 have been proposed to understand the tracks formation mechanism and material mutation with phase transition along with them. Where the first one is based on electrostatic repulsion, while the second one consider the energy transfer to the lattice atoms in the form of heat reaching to melting temperature and then rapid quenching leading to track formation in very short time scale. Recent experiments on transition metal oxide have evidence that beside amorphization and damage creation, a crystalline to crystalline phase transformation could occur in these oxides.^11–16 In addition ion irradiation induce transformation has also been observed in more complex oxides.^17,18

Understanding of these crystalline phase transformation under ion irradiation has a great importance from both physics and technical point of view. In present paper our investigation is focused on crystalline phase transformation under ion irradiation in zirconium oxide (ZrO₂) thin film. The oxide exhibit three basic allotropic phases, at ambient pressure and temperature thermodynamically stable one is monoclinic phase (m-ZrO₂) with structure belong to (C⁵_2h, P2₁/c) space group having seven fold cation coordination. With increase in temperature zirconia transforms to tetragonal phase at 1273 K with space group (D¹⁵_4h, P4₂/nmc) with eightfold cation coordination, further increase in temperature up to 2573 K leads to cubic phase having space group as (O⁵_h, Fm3m) also having eightfold cation coordination.¹⁹ Both monoclinic and tetragonal phases are considered as distortive form of cubic fluorite structure. Pervious ion irradiation study performed over bulk ZrO₂ has shown partial phase transformation and found effective threshold electronic energy loss for these transformation at around ∼11.3 keV nm⁻¹.^12,20–22 Beside this recent experiment has revealed a complete transformation of phase for 0.6 GeV Pb (lead) ions irradiation with an electronic energy loss of ∼42 keV nm⁻¹.²³ It is also known that the phase transformations depend on Poisson behavior and some minimum number of ion impacts is needed for the transformations. However, the effects of strain under the influence of modifications in the electronic structure and phonon structure during PT are yet to be reported. Hence, these investigations are reported in the present manuscript. It may be noteworthy that partial transformation from tetragonal to cubic phase by energetic ions are investigated.

Therefore, present manuscript reports the kinetics of such PT in ZrO₂ films with two different energy ions, firstly with 120 MeV Ag (silver) ions and secondly with 130 MeV Ni (nickel) ions. The ion and energy are chosen in such a way that Ag ions have electronic energy loss well above the threshold value nearly ∼21 keV nm⁻¹ while the Ni ions are set aside to threshold value ∼13 keV nm⁻¹. The modifications in the electronic and local structure are investigated using O and Zr K-edges XANES and XAFS which are reported along with X-ray diffraction (XRD) and micro-Raman spectroscopy to identify the crystallographic phase of the irradiated film. More, many structural parameters are also estimated from the electronic structure analysis and envisage some critical understanding of the PT under such non-equilibrium conditions of ion irradiation.

Experimental

Highly crystalline ion beam sputtered ZrO₂ films with monoclinic structure on silicon with silica as buffer layer on silicon substrates. The film was then irradiated by 120 MeV Ag and 130 MeV Ni ions at different fluences at room temperature using 15UD Tandem Pelletron Accelerator at Inter-University Accelerator Centre (IUAC) New Delhi. The energy loss parameters of irradiation are calculated by using SRIM (2008) code and are summarized in Table 1. The structural study of pristine and irradiated sample was done by XRD measurements using Bruker D8 advance diffractometer with Cu Kα radiation (λ = 1.540 Å). The Raman scattering measurements were performed using Ranishaw InVia Raman microscope under the excitation by 514 nm argon-ion laser at 50 mW power.

Table 1 Irradiation parameters calculated using SRIM 2008

Irradiated material	Ion species	Energy (MeV)	Ion range (Rp) (μm)	Electronic energy loss (keV nm^−1)	Nuclear energy loss (eV nm^−1)
ZrO2 thin film	^107Ag	120	9.3	21.3	10.4
	^58Ni	130	13.1	13.3	21.1

---

The XAS measurements of pristine and irradiated thin films across O K-edges were carried out in total electron yield (TEY) mode at soft X-rays absorption spectroscopy (SXAS) beamline (BL-01) of the Indus-2 Synchrotron Source at Raja Ramanna Centre for Advanced Technology (RRCAT), Indore, India. Experiments were performed in an ultra high vacuum (UHV) chamber with a base pressure of ∼10⁻¹⁰ Torr. Energy resolution at O K edge energy was ∼200 meV. Zr K-edge X-rays absorption fine structure (XAFS) spectra were recorded at BL-9, scanning extended edge X-rays absorption fine structure (EXAFS) Beamline of Indus-2. XAFS measurements were done in fluorescence mode using Vortex energy dispersive detector. The beamline consists of Rh/Pt coated meridional cylindrical mirror for collimation and Si (111) based double crystal monochromator to select excitation energy. The energy range of XAFS was calibrated using Zr foil at 17 998 eV.

Results and discussion

X-ray diffraction study

X-ray spectrum of ZrO₂ film irradiated at ambient temperature and pressure with Ni (nickel) and Ag (silver) ions at a fluence of 5 × 10¹³ and 3 × 10¹³ ions per cm² along with pristine are displayed in Fig. 1. Pristine sample shows two dominating peaks for 2θ ∼28.2° and 31.5° indexed at (−111) and (111) reflection of monoclinic ZrO₂. The nickel ion irradiated film with increasing fluence from 1 × 10¹² to 5 × 10¹³ ions per cm² intensity of the tetragonal peak at 30.2° indexed (101) increases giving a clear evidence of phase transformation with considerable amount of monoclinic phase still present in it at maximum fluence. While for Ag ion irradiation the maximum fluence used is 3 × 10¹³ ions per cm² at which insignificant amount of monoclinic phase is present. With this the central line (101) or (111) belonging to tetragonal or cubic phase is not well resolved. A quantitative XRD analysis to get the fraction of tetragonal phase with increasing ion fluence is performed using the procedure purposed by Garvie and Nicholson:²⁴where I_t(101), I_m(1̄11) and I_m(111) are the integrated intensity of the indices (101), (1̄11) and (111) of tetragonal and monoclinic. It may be noted that the XRD pattern of the highest fluence is only shown in Fig. 1, and rest all is used for the calculations as cited in Table 2(a) and (b).

Fig. 1 XRD pattern of pristine and irradiated films. Evolution of tetragonal phase can be seen for both Ag (120 MeV) and Ni (130 MeV) ions at 3 × 10¹³ ions per cm² and 5 × 10¹³ ions per cm². A complete transformation observed for Ag ion while traces of monoclinic are present for Ni irradiated samples.

Table 2 (a) and (b) Calculated fraction of tetragonal phase with ion fluence for Ni and Ag ions, respectively

Ion fluence (ions cm^−2)	Ct (%)
(a)
Ni ∼ 1 × 10^12	6.9
Ni ∼ 1 × 10^13	44.8
Ni ∼ 5 × 10^13	81.6
(b)
Ag ∼ 5 × 10^11	9.2
Ag ∼ 1 × 10^12	17.4
Ag ∼ 1 × 10^13	96.4
Ag ∼ 3 × 10^13	98.0

---

Raman spectroscopy study

In addition to X-ray diffraction Raman spectroscopy is a fine technique to look around local structural changes. The space groups symmetry of m-phase and t-phase of ZrO₂ are P2₁/c (C⁵_2h) and P4₂/nmc (D¹⁵_4h) respectively. Group theory investigation has shown that there are 36 lattice vibration modes of monoclinic ZrO₂: G_mono = 9A_g + 9A_u + 9B_g + 9B_u out of which 9A_g and 9B_g a total of 18 modes are Raman active, 8A_u and 7B_u are IR-active remaining are acoustic mode. Eighteen degrees of freedom for lattice vibration in tetragonal ZrO₂ provides 18 phonon branch according to group theory analysis out of which 6 are Raman active modes (A_1g + 2B_1g + 3E_g), three are IR-active (A_2u + 2E_u) and remaining are acoustic phonons while there is single Raman active mode of cubic ZrO₂ which appears at 620 cm⁻¹.^25–27

Room temperature Raman spectra of irradiated films were collected to investigate the effect of ion irradiation on crystal structure and its symmetry (phonon modes). The study of affected phonon modes is useful to understand role of irradiation energy in stabilizing and transformation of the phases. Raman spectra for pristine and ion irradiated samples of ZrO₂ are shown in Fig. 2(a) and (b). It clearly show 8 Raman active modes of monoclinic phase for the pristine film identified at around 179 cm⁻¹ (A_g), 190 cm⁻¹ (A_g), 305 cm⁻¹ (A_g), 334 cm⁻¹ (B_g), 348 cm⁻¹ (A_g), 381 cm⁻¹ (B_g), 476 cm⁻¹ (A_g), 615 cm⁻¹ (B_g), and 637 cm⁻¹ (A_g).^26,28,29 Upon increase in Ag ion irradiation fluence at 1 × 10¹² ions per cm² for both ions Raman spectra remain unchanged as observed in previous studies.^30–33 At a fluence of 3 × 10¹² ions per cm² small modification in the structure become visible such as decrease in the intensity of bands at 179 and 190 cm⁻¹ corresponding to monoclinic phase and new broad bands start appearing at 148, 265 and ∼642 cm⁻¹ signifying the evolution of tetragonal phase with increasing fluence.³³ Moving further to higher fluences a complete transformation comes into sight for Ag ions as in XRD patterns whereas Ni ions were not able to do so and in this case a fraction of monoclinic phase still exist along with the evolved tetragonal phase as shown in Fig. 2(a) and (b).

Fig. 2 (a) and (b) shows the Raman spectra for different fluence of Ni and Ag ion showing the presence of tetragonal modes; while (c) and (d) shows comparative shifts of newly developed modes under Ag and Ni ion irradiation.

In order to have a deeper understating about the kinetics of PT, the zoomed spectra of region 130–210 cm⁻¹ and 575–725 cm⁻¹ is shown in Fig. 2(c) and (d), respectively. Here, Fig. 2(c) illustrates more profound spectral feature of pristine and irradiated samples primarily for the fluences at which phase transformation initiated. Firstly, the A_g mode centered at 179 cm⁻¹ exhibits the stiffening of about 2 cm⁻¹ during the partial transformation from monoclinic to tetragonal phase. After the complete PT to tetragonal phase this mode complete disappears. Secondly, the center of E_g mode signifying the tetragonal phase at 148 cm⁻¹ also stiffened by 1 cm⁻¹ upon high fluence irradiation by Ag ions as compared to Ni ion irradiation. Importantly, the second region as shown in Fig. 2(d) is marked by three different dashed line. The first dashed line is marked on 615 cm⁻¹ (B_g) mode of monoclinic phase and upon transformation to tetragonal phase this mode also stiffened to 621 cm⁻¹. Interestingly, the mode at 637 cm⁻¹ stiffened to 642 cm⁻¹ for highest fluence of Ni ion irradiations. On the other hand, this mode shifts to 647 cm⁻¹ for the Ag ion irradiation and becomes very pronounced. The pronounced nature of this mode confirms the dominant fraction of tetragonal phase in agreement with XRD results. However, the large stiffening for Ag ion irradiations reveals the presence of some fraction of strained cubic phase. The evolution of very weak and broad mode at around 470 cm⁻¹ for highest fluence of Ag ion irradiation is also confirms the presence of strained cubic phase with highly dominant tetragonal phase. Moreover, as mentioned above that the tetragonal mode stiffened by 1 cm⁻¹ can also be ascribed to the presence of strained cubic phase in agreement with literature.³⁴ This shift could be assigned to different type of stress fields and the difference of 1 cm⁻¹ to decrease in Zr–O bond length due to disorder of oxygen sub-lattice caused by oxygen vacancies. The roll of stress fields is well studied by B. Schuster et al. with external applied pressure finding that the required fluence has reduces by one order of magnitude for partial phase transformation.^30,35 However, it is our believe that stress field produced by ions along the trajectory is more intense and affect the crystal structure locally which is leading us toward the traces of strained cubic phase. Such features of mRS spectra are further supported by the XAFS spectral analysis along with the tetragonal/cubic phase detection discussed in next section.

X-Ray absorption spectroscopy study

Oxygen K-edge XAS characteristics. Fig. 3 shows the normalized O K-edge XAS spectra of pristine and irradiated films using 130 MeV Ni and 120 MeV Ag ions for 5 × 10¹³ ions per cm² and 3 × 10¹³ ions per cm², respectively. It clearly shows the change in the electronic structure under different condition of irradiation of the ion beam irradiated ZrO₂ films. Both, the wide spread and the large intensity of the absorption signal indicate important covalent contribution to the bonding. The spectra can be divided into two well defined regions in the energy range of 20 eV above threshold. The first region goes from absorption edge up to 5–7 eV above the threshold. This region is attributed to transition from transition metal oxide 1s core state to O 2p states hybridized with metal nd states (here Zr 4d) which are split by crystal field effects. Second region falls on higher energy side i.e. between 9 and 18 eV above the threshold shows more complex structure. This region is attributed to the O 2p states hybridized with the metal (n + 1) sp states (here Zr 5sp).^36,37 The spectral features of O K-edge spectra of three different phases (cubic, tetragonal and monoclinic) of ZrO₂ thin films are recently studied.³⁸ Such reports corroborate that shape of O K-edge photo-absorption spectra more significantly depends on the phase composition. The e_g and t_2g peaks of tetragonal and cubic phases exhibit same spectral features i.e., sharp e_g and t_2g peaks. However, the e_g peak is found to be split in case of monoclinic ZrO₂ a possible reason for this could be the exchange splitting. The energy splitting (Δd), between t_2g and e_g peaks, is found to be 3.5 eV, 3.0 eV and 2.9 eV for cubic, tetragonal, monoclinic phases of ZrO₂, respectively.³⁹ The intensity ratio of e_g and t_2g peaks is found to be 2 : 3 in case of monoclinic ZrO₂. This figure also indicates that there are no measurable changes in the second region of the spectra with increasing the irradiation fluence; however the spectral features are noticeably modified within the first region. It is evident from the first region of the spectra of pristine film that; (i) e_g peak shows broader spectral feature due to peak splitting (splitting of e_g peak is marked by dashed line) (ii) intensity ratio of e_g and t_2g peak is 2 : 2.7 and (iii) the energy difference between e_g and t_2g peaks (Δd) is 4.1 eV. The observed intensity ratio of e_g and t_2g peaks and splitting of e_g peak are in accordance with the spectra of monoclinic ZrO₂ samples. Therefore, O K-edge XAS spectrum of pristine film indicates the presence of monoclinic phase of ZrO₂ which are in agreement with the findings of our XRD and Raman investigations. After the ion irradiation the Δd value becomes to 3.4 eV. This value of Δd is close to the reported values for the cubic ZrO₂.^39,40 At the highest Ag ion irradiation fluence, the spectral features (mainly P and Q) indicate almost complete phase transformation from monoclinic to tetragonal phase.

Fig. 3 Comparison between normalized O K-edge XANES spectra of pristine and irradiated ZrO₂ thin films. The raw data have been shifted vertically for more clarity.

Zr K-edge XAFS characteristics. The comparison of the X-ray absorption near edge structure (XANES) spectra recorded for Zr K-edge of pristine and irradiated samples are shown in Fig. 4(a) and k²-weighted χ(k) XAFS spectra in Fig. 4(b). For pristine sample the main peak is around 18023 eV with no other features showing the presence of monoclinic symmetry whereas for Ni ion irradiation it has a shoulder marked as C and for Ag ion irradiation a peak is observed shown as B and the main peak shows a split feature marked as A. The split feature A of Ag irradiated sample suggests to have cubic phase of ZrO₂. The k²-weighted χ(k) XAFS data shows very different dramatic change in local structure of pristine and irradiated films. To reveal change in m-ZrO₂ phase to t-ZrO₂ and/or c-ZrO₂, we have done XAFS data fitting.

Fig. 4 (a) Comparison between the normalized Zr K-edge XANES spectra of Zr foil, pristine and irradiated ZrO₂ films, and (b) k²-weighted χ(k) spectra of pristine, Ni and Ag irradiated ZrO₂ films.

Fig. 5(a and b) represents the best fit of the magnitude and real Fourier transform (FT) of χ(R) Zr K-edge for pristine and irradiated films. The k-range of 3–8.7 Å⁻¹ was used for FT of Zr XAFS data. The fitting of pristine were performed using m-ZrO₂ structure in R-space range of 1–3.3 Å. For m-ZrO2 structure, the first shell Zr bond with 8 oxygen atoms (Zr–O) at mean value 2.16 Å distance and second next near-neighbor Zr surrounded by 12 Zr atoms (Zr–Zr) at 3.72 Å distance. The R-space fitting of pristine shows not much variation in Zr–O bond distance whereas Zr–Zr bond distance has large reduction for Ni irradiated films XAFS data fitting were done with mixed phase of m- and t-ZrO2 and for Ag irradiated t- and c-ZrO2 structure in R-range of 1–3.63 Å. The results of best fit values of the parameters are listed in Table 3. It has been noted from fitting that Ni irradiated film has 54% t-ZrO2 structure with large Zr–O bond distance and remaining 46% m-ZrO2 structure. Similarly, the XAFS fitting of Ag irradiated film shows higher % of t-ZrO2 phase mixed with c-ZrO2 phase. Thus, the combination of Zr K-edge XANES and XAFS fitting confirm phase transformation from m-ZrO₂ to t-ZrO₂ and/or c-ZrO₂.

Fig. 5 (a) Magnitude and (b) real component of the Fourier transform of χ(R) vs. R for pristine and irradiated ZrO₂ films. Dotted lines represent experimental data whereas the solid lines represent the best fit.

Table 3 XAFS results for the best fit values

Sample	% of ZrO2 structure	RZr–O (Å)	σ^2Zr–O (Å^2)	RZr–Zr (Å)	σ^2Zr–Zr (Å^2)
Pristine	m-ZrO2	2.147 (2)	0.0073 (2)	3.418 (3)	0.0132 (3)
513 Ni	m-ZrO2	3.158 (3)	0.0156 (3)	3.898 (2)	0.0207 (4)
	t-ZrO2	3.158 (3)	0.0156 (3)	3.898 (2)	0.0207 (4)
313 Ag	t-ZrO2 (66 ± 2)	3.103 (2)	0.0192 (3)	3.512 (3)	0.0207 (4)
	c-ZrO2 (34 ± 2)	2.095 (3)	0.0032 (8)	3.565 (5)	0.0034 (4)

---

Track radius a view under thermal spike simulation. Understanding of track formation and amorphization of materials due to huge electronic energy loss in the wake of SHI irradiation has lead us to two competing mechanisms “Coulomb explosion” and “Thermal spike” model. Where the earlier one consider an electrostatic repulsion between the charged ions created due to the kinetic energy transfer by heavy ions along their trajectory leading to track formation and amorphization when electrostatic potential overcome the chemical bond energy, while the later one has a completely different approach. It considers a radial distribution of energy density deposited by ions in the form of thermal spike. The thermal spike model itself has its two versions to explain the track formation and phase transformations along the ion trajectories. First one is inelastic thermal spike model^9,23,41 where energy loss by heavy ions is deposited on electronic system by electron–electron interaction and then to atomic system by electron–phonon coupling. A cylinder with very high temperature above the melting or even vaporization temperature is assumed to form around the ion trajectory. Due to small volume of the cylinder the cooling rate of the molten material is very high resulting in rapid quenching of molten phase to solid phase which eventually give rise to amorphous tracks. Mathematically inelastic thermal spike model can be expressed by two coupled equations governing the energy distribution on electronic and lattice subsystem. For cylindrical geometry a time dependent thermal transient process can be expressed as:where C_e, C(T_a), T_e, T_a and K_e, K(T_a) are the specific heat, the temperature and the thermal conductivity of the electronic and atomic systems, respectively. ρ is the specific mass of the lattice and g is the electron phonon coupling constant. B(r,t) is the energy density^42,43 supplied by the incident ion to the electronic system by the ballistic collision.

Numerical solutions of coupled equations using simulation codes yields the temperature of atomic subsystems as a function of radial distance from track core and time which is then converted into corresponding energy density (eV at⁻¹). Fig. 6(a) and (b) are representing the evaluation of temperature for ZrO₂ with the replacement of melting temperature T_m to phase transformation temperature T_m–t in simulation codes and taking the value of electron–phonon mean free path at around 8.5 nm obtained from the experimental fit of damage cross section as a function of electronic energy loss. The macroscopic thermodynamic parameters used for simulation are listed in Table 4. We found a reasonable agreement between the radial distribution of phase transformation temperature and experimental damage cross section as seen from figures and the correlation are discussed in next section.

Fig. 6 Evolution of atomic temperature versus time at different radial distance from the ion trajectory for ZrO₂ using thermal spike simulation codes where the melting temperature has been replaced by m–t phase transition temperature, (a) represents Ag ion effect having energy 120 MeV with an electronic energy loss of 21.3 keV nm⁻¹, (b) represents Ni ion effect having energy 130 MeV with an electronic energy loss of 13.3 keV nm⁻¹.

Table 4 Thermodynamic parameters used in thermal spike calculations

Parameters	Values
Band gap (eV)	5
Solid density (g cm^−3)	5.68
Liquid density (g cm^−3)	5.68
Mean diffusion length (nm)	8.5
Se (keV nm^−1), SRIM 2008	21.3(Ag)/13.3(Ni)
Melting temperature (K)	2973
Latent heat of fusion (J g^−1)	760

---

The analysis gave us understanding that the transformation of the phase relies mainly on damage cross section or on the radius of the track along the trajectory of the ions and radius is found to depend on the electronic energy loss either for the bulk sample^12,23,42 or for the thin films, as we observed in this study. The electronic energy loss for Ag ∼120 MeV ion is 21.3 keV nm⁻¹ and 13.3 keV nm⁻¹ for the Ni ∼130 MeV. With such extensive change in the electronic energy loss there is a significant difference in the concentration of the PT from m-to-t phase as viewed from XRD as well as Raman spectra analysis. It is well clear that at a fluence of 1 × 10¹³ ions per cm² of Ag ions almost all the fraction of monoclinic transformed to tetragonal. So, by taking the track size of 4.5 nm as estimated using the thermal spike calculations, it required more than 6 ion impacts for the complete PT. On the other hand for the highest dose of Ni ions only about 80% fraction transformed to tetragonal phase. It means by taking into account the track size of 1 nm as estimated by thermal spike calculations, it required about 2 ion impacts for the Ni ions.

Two major outcome of the present investigations are firstly the kinetic rate of phase transformation is much faster or one can say a sudden transformation of phase has taken place in comparison with previous studies on bulk samples^12,22,23 either for Ni ∼130 MeV ion or for Ag ∼120 MeV. Secondly only partial phase transformation has been reported so far for the bulk samples except a study which shows a complete phase transformation with a massive electronic energy loss of around ∼42 keV nm⁻¹¹⁹ while in present experiment a complete transformation of phase has occurred under Ag ion irradiation at 1 × 10¹⁴ ions per cm² with an electronic energy loss of ∼21 keV nm⁻¹ which is just half of the energy required for the bulk samples. A possible reason for such a fast kinetic rate and complete phase transformation can be the accumulation of high defect concentration at the interface of monoclinic–tetragonal phase due to fast production rate of defect as compare to recombination leading to a non-equilibrium defect population state under the first ion impact while second ion impact just after the first impact with a very short time delay increases the temperature of the lattice providing an easy path for defects to diffuse through the lattice. Therefore, for Ag ions large numbers of ion impacts are required as compared to Ni ions.

Conclusions

In summary, monoclinic ZrO₂ films on Si substrate with SiO₂ as buffer layer were investigated for phase transition studies at room temperature under swift heavy ion irradiation with varying density of electronic excitations (EE) at various ion fluences. Structural and Raman studies confirms the monoclinic-to-tetragonal phase transformation (PT). A quantitative fraction of tetragonal phase analysis is also carried out to understand the kinetics of PT. It is found that for thin film transformation follow multiple ion impact mechanism in agreement with previous studies on bulk material. A partly PT to cubic phase is also noticed at high fluence of irradiation at high density of EE. However, it is clear that this PT is not only due to transient temperature induced by SHI, but also attributed to the change in the high density of defects induced in the lattice. The modifications in electronic and local structure investigated using XAS and XAFS techniques, also confirms PT and various parameters are reported. Thus, the reported studies would shine light on a better understanding of the PT under such non-equilibrium conditions of SHI irradiation.

Acknowledgements

Authors are grateful to Director (IUAC, New Delhi), Director (UGC-CSR Indore) Director (RRCAT, Indore) for their moral encouragement and for extending experimental facilities including the synchrotron facility (BL-1 & BL-9) at RRCAT Indore. Authors also acknowledge Mr Rakesh for his help during the XAS measurements of BL-01 beam line, while Dr R. G. Singh (DU) and Mr Subodh K. Gautam for their support during irradiation of films and Pelletron group IUAC for providing stable beam during irradiation. One of the authors (M. R.) acknowledges fellowship from University Grants commission (UGC), New Delhi and UGC-CSR for providing project: CRS-128-2014-15.

References

1. M. O. Figueiredo, A. C. D. Santos, S. Meriani and C. Palmonari, "Zirconia' 88: Advances in Zirconia Science and Technology" 1989, p. 81.
2. E. J. Walter, S. P. Lewis and A. M. Rappe, Surf. Sci., 2001, 495, 44–50.
3. V. Fiorentini and G. Gulleri, Phys. Rev. Lett., 2002, 89, 266101.
4. R. Puthenkovilakam, E. A. Carter and J. P. Chang, Phys. Rev. B: Condens. Matter Mater. Phys., 2004, 69, 155329.
5. A. Meldrum, L. A. Boatner and R. C. Ewing, Phys. Rev. Lett., 2001, 88, 25503.
6. C. Morant, J. M. Sanz and L. Galan, Phys. Rev. B: Condens. Matter Mater. Phys., 1992, 45, 1391.
7. A. Audoard, E. Balanzat, S. Bouffard, J. C. Jousset, A. Chamberod, A. Dunlop, D. Lesueur, G. Fuchs, R. Spohr, J. Vetter and L. Thome, Phys. Rev. Lett., 1990, 65, 875.
8. A. Brbu, A. Dunlop, D. Lesueur and R. Averback, Europhys. Lett., 1991, 15, 37–42.
9. M. Toulemonde, C. Dufour, A. Meftah and E. Paumier, Nucl. Instrum. Methods Phys. Res., Sect. B, 2000, 166, 903–912.
10. G. Szenes, Phys. Rev. B: Condens. Matter Mater. Phys., 1995, 51, 8026–8029.
11. S. K. Gautam, F. Singh, I. Sulania, R. G. Singh, P. K. Kulriya and E. Pippel, J. Appl. Phys., 2014, 115, 143504.
12. A. Benyagoub, Phys. Rev. B: Condens. Matter Mater. Phys., 2005, 72, 094114.
13. F. Lu, et al., Phys. Chem. Chem. Phys., 2012, 14, 12295.
14. F. Lu, et al., J. Phys. Chem. C, 2011, 115, 7193.
15. Y. Zhang, et al., Phys. Rev. B: Condens. Matter Mater. Phys., 2010, 82, 184105.
16. J. Lian, J. Zhang, F. Namavar, Y. Zhang, F. Lu, H. Haider, K. Garvin, W. J. Weber and R. C. Ewing, Nanotechnology, 2009, 20, 245303.
17. J. Lian, L. M. Wang, S. X. Wang, J. Chen, L. A. Boatner and R. C. Ewing, Phys. Rev. Lett., 2001, 87, 145901.
18. L. Wang, W. Gong, S. Wang and R. C. Ewing, J. Am. Ceram. Soc., 1999, 82, 3321–3329.
19. A. Christensen and E. Carter, Phys. Rev. B: Condens. Matter Mater. Phys., 1998, 58, 8050–8064.
20. A. Benyagoub and F. Levesque, Europhys. Lett., 2002, 60, 580.
21. A. Benyagoub, Nucl. Instrum. Methods Phys. Res., Sect. B, 2003, 206, 132–138.
22. A. Benyagoub, Nucl. Instrum. Methods Phys. Res., Sect. B, 2006, 245, 225–230.
23. A. Benyagoub, Nucl. Instrum. Methods Phys. Res., Sect. B, 2010, 268, 2968–2971.
24. R. C. Garvie and P. S. Nicholson, J. Am. Ceram. Soc., 1972, 55, 303–305.
25. G. Siu, M. Stokes and Y. Liu, Phys. Rev. B: Condens. Matter Mater. Phys., 1999, 59, 3173–3179.
26. X. L. Tang and X. H. Zheng, J. Mater. Sci. Technol., 2004, 20, 485–489.
27. A. Kuwabara, T. Tohei, T. Yamamoto and I. Tanaka, Phys. Rev. B: Condens. Matter Mater. Phys., 2005, 71, 64301.
28. P. Bouvier, E. Djurado, G. Lucazeau and T. Le Bihan, Phys. Rev. B: Condens. Matter Mater. Phys., 2000, 62, 8731–8737.
29. P. Bouvier and G. Lucazeau, J. Phys. Chem. Solids, 2000, 61, 569–578.
30. B. Schuster, F. Fujara, B. Merk, R. Neumann, T. Seidl and C. Trautmann, Nucl. Instrum. Methods Phys. Res., Sect. B, 2012, 277, 45–52.
31. J. A. Valdez, Z. Chi and K. E. Sickafus, J. Nucl. Mater., 2008, 381, 259–266.
32. A. Benyagoub, F. Levesque, F. Couvreur, C. Gibert-Mougel, C. Dufour and E. Paumier, Appl. Phys. Lett., 2000, 77, 3197–3199.
33. C. Gibert-Mougel, F. Couvreur, J. M. Costantini, S. Bouffard, F. Levesque, S. Hémon, E. Paumier and C. Dufour, J. Nucl. Mater., 2001, 295, 121–125.
34. C. L. Tracy, M. Lang, F. Zhang, C. Trautmann and R. C. Ewing, Phys. Rev. B: Condens. Matter Mater. Phys., 2015, 92, 174101.
35. B. Schuster, M. Lang, R. Klein, C. Trautmann, R. Neumann and A. Benyagoub, Nucl. Instrum. Methods Phys. Res., Sect. B, 2009, 267, 964–968.
36. F. M. F. De Groot, J. C. Fuggle, B. T. Thole and G. A. Sawatzky, Phys. Rev. B: Condens. Matter Mater. Phys., 1990, 42, 5459–5468.
37. L. Soriano, M. Abbate, J. Faber, C. Morant and J. Sanz, Solid State Commun., 1995, 93, 659–665.
38. A. Kikas, J. Aarik, V. Kisand, K. Kooser, T. Käämbre, H. Mändar, T. Uustare, R. Rammula, V. Sammelselg and I. Martinson, J. Electron Spectrosc. Relat. Phenom., 2007, 156–158, 303–306.
39. D. McComb, Phys. Rev. B: Condens. Matter Mater. Phys., 1996, 54, 7094–7102.
40. S. Ostanin, A. J. Craven, D. W. McComb, D. Vlachos, A. Alavi, A. T. Paxton and M. W. Finnis, Phys. Rev. B: Condens. Matter Mater. Phys., 2002, 65, 224109.
41. A. Meftah, J. M. Costantini, N. Khalfaoui, S. Boudjadar, J. P. Stoquert, F. Studer and M. Toulemonde, Nucl. Instrum. Methods Phys. Res., Sect. B, 2005, 237, 563–574.
42. Z. G. Wang, C. Dufour, E. Paumier and M. Toulemonde, J. Phys.: Condens. Matter, 1995, 7, 2525.
43. M. Toulemonde, C. Dufour, Z. Wang and E. Paumier, Nucl. Instrum. Methods Phys. Res., Sect. B, 1996, 112, 26–29.

---