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Article

Modification of SiO2, ZnO, Fe2O3 and TiN Films by Electronic Excitation under High Energy Ion Impact

1
Faculty of Science and Technology, Meijo University, Nagoya 468-8502, Japan
2
Japan Atomic Energy Agency (JAEA), Tokai 319-1195, Japan
*
Author to whom correspondence should be addressed.
Quantum Beam Sci. 2021, 5(4), 30; https://doi.org/10.3390/qubs5040030
Received: 30 August 2021 / Revised: 6 October 2021 / Accepted: 17 October 2021 / Published: 27 October 2021

Abstract

:
It has been known that the modification of non-metallic solid materials (oxides, nitrides, etc.), e.g., the formation of tracks, sputtering representing atomic displacement near the surface and lattice disordering are induced by electronic excitation under high-energy ion impact. We have investigated lattice disordering by the X-ray diffraction (XRD) of SiO2, ZnO, Fe2O3 and TiN films and have also measured the sputtering yields of TiN for a comparison of lattice disordering with sputtering. We find that both the degradation of the XRD intensity per unit ion fluence and the sputtering yields follow the power-law of the electronic stopping power and that these exponents are larger than unity. The exponents for the XRD degradation and sputtering are found to be comparable. These results imply that similar mechanisms are responsible for the lattice disordering and electronic sputtering. A mechanism of electron–lattice coupling, i.e., the energy transfer from the electronic system into the lattice, is discussed based on a crude estimation of atomic displacement due to Coulomb repulsion during the short neutralization time (~fs) in the ionized region. The bandgap scheme or exciton model is examined.

1. Introduction

Material modification induced by electronic excitation under high-energy (> 0.1 MeV/u) ion impact has been observed for many non-metallic solids since the late 1950’s; for example, the formation of tracks (each track is characterized by a long cylindrical disordered region or amorphous phase in crystalline solids) in LiF crystal (photographic observation after chemical etching) by Young [1], in mica (a direct observation using transmission electron microscopy, TEM, without chemical etching, and often called a latent track) by Silk et al. [2], in SiO2-quartz, crystalline mica, amorphous P-doped V2O5, etc. (TEM) by Fleischer et al. [3,4], in oxides (SiO2-quartz, Al2O3, ZrSi2O4, Y3Fe5O12, high-Tc superconducting copper oxides, etc.) (TEM) by Meftah et al. [5] and Toulemonde et al. [6], in Al2O3 crystal (atomic force microscopy, AFM) by Ramos et al. [7], in Al2O3 and MgO crystals (TEM and AFM) by Skuratov et al. [8], in Al2O3 crystal (AFM) by Khalfaoui et al. [9], in Al2O3 crystal (high resolution TEM) by O’Connell et al. [10], in amorphous SiO2 (small angle X-ray scattering (SAXS)) by Kluth et al. [11], in amorphous SiO2 (TEM) by Benyagoub et al. [12], in polycrystalline Si3N4 (TEM) by Zinkle et al. [13] and by Vuuren et al. [14], in amorphous Si3.55N4 (TEM) by Kitayama et al. [15], in amorphous SiN0.95:H and SiO1.85:H (SAXS) by Mota-Santiago et al. [16], in epilayer GaN (TEM) by Kucheyev et al. [17], in epilayer GaN (AFM) by Mansouri et al. [18], in epilayer GaN and InP (TEM) by Sall et al. [19], in epilayer GaN (TEM) by Moisy et al. [20], in InN single crystal (TEM) by Kamarou et al. [21], in SiC crystal (AFM) by Ochedowski et al. [22] and in crystalline mica (AFM) by Alencar et al. [23]. Amorphization has been observed for crystalline SiO2 [5] and the Al2O3 surface at a high ion fluence (though the XRD peak remains) by Ohkubo et al. [24] and Grygiel et al. [25]. The counter process, i.e., the recrystallization of the amorphous or disordered regions, has been reported for SiO2 by Dhar et al. [26], Al2O3 by Rymzhanov [27] and InP, etc., by Williams [28]. Density modification, i.e., a lower density in the track core surrounded by a shell with a higher density, has been observed for Al2O3 [10], amorphous SiO2 [11], Si3N4 [14] and amorphous SiN0.95:H and SiO1.85:H [16]. Interestingly, an electrically conducting track formation in tetrahedral-amorphous carbon (sp3 into sp2 bond transformation) has been observed by Gupta et al. [29]. The track radius, hillock height and diameter characterizing the surface morphology modification associated with the track are well described in terms of the electronic stopping power Se (defined as the energy loss due to electronic excitation and ionization per unit path length), and the velocity effect has been noticed [12]. The threshold of Se for the track formation has been reported [3,6,8,9,12,13,30] and the data appear to scatter, and it seems that the threshold Se depends on the observation method of the track [12]. No track formation by monatomic ions has been observed in AlN [19].
Moreover, electronic sputtering (the erosion of solid materials caused by electronic energy deposition) has been observed for various compound solids: UO2 by thermal-neutron-induced 235U fission fragments by Rogers [31,32] and by Nilsson [33], UO2 by energetic ions by Meins et al. [34], Bouffard et al. [35] and Schlutig [36], H2O ice by Brown et al. [37,38], Bottiger et al. [39], Baragiola et al. [40], Dartois et al. [41] and Galli et al. [42], frozen gas films of Xe, CO2 and SF6 [39], those of CO, Ar and N2 by Brown et al. [43], CO2 ice by Mejia et al. [44], SiO2 by Qui et al. [45], Sugden et. al. [46], Matsunami et al. [47,48], Arnoldbik et al. [49] and Toulemonde et al. [50,51], MgAl2O4 [48], UF4 ([34], by Griffith et al. [52] and Toulemonde et al. [53]), LiNbO3 [45], Al2O3 ([45] and by Matsunami et al. [54]), various oxides by Matsunami et al. (SrTiO3 and SrCeO3 [47,54], CeO2, MgO, TiO2 and ZnO [54], Y2O3 and ZrO2 [55], Cu2O [56,57], WO3 [58], CuO [59], Fe2O3 [60]), Si3N4 [45], Si3N4 and AlN by Matsunami et al. [55], Cu3N by Matsunami et al. [56,61], LiF ([50], by Assmann et al. [62] and Toulemonde et al. [63]), KBr [56], NaCl [63], CaF2 [53] and SiC [56]. The sputtering of frozen Xe films has been observed for low energy electron impact, against the anticipation of no atomic displacement [39], and the result confirms that the sputtering is caused by electronic excitation. Mechanisms of electronic excitation leading to atomic displacement will be discussed in Section 4.
As mentioned above, electronic sputtering has been observed for a variety of non-metallic materials, indicating that it seems to be a general phenomenon for non-metallic solids by high-energy ion impact. In many cases, ions with an equilibrium charge have been employed, which is usually attained by inserting thin foils, such as carbon, metals, etc., before impact on samples, and sputtered atoms are collected in carbon, metals, etc., followed by neutron activation and ion beam analysis to obtain sputtering yields. This article concerns the equilibrium charge incidence, though charge-state effects for non-equilibrium charge incidence have been observed and discussed ([23,29,34,49,52,58,62,64]). The electronic energy deposition or electronic stopping power (Se) at the equilibrium charge can be calculated using a TRIM or SRIM code by Ziegler et al. [65,66], a CasP code by Grande et al. [67] and the nuclear stopping power (Sn, defined as the energy loss due to elastic collisions per unit path length) [65,66]. Characteristic features of the electronic sputtering by high-energy ions are as follows:
(a).
Electronic sputtering yields (YSP) are found to be larger by 10–103 than nuclear sputtering yields due to elastic collision cascades, which can be estimated assuming linear dependence on Sn;
(b).
YSP super linearly depends on Se and is approximated by the power-law fit: YSP = (BSPSe)Nsp with 1 ≦ NSP ≦ 4 for most cases, with BSP being a material dependent constant.
Stoichiometric sputtering has been observed for many materials, whereas a considerable deviation from the stoichiometric sputtering has been reported for YBa2Cu3O7 [68], Gd3Ga5O12 and Y3Fe5O12 [69] and CaF2 and UF4 [53]. Only the heavy element of U [31,32,33,34,35,36] and the light element of O [49] have been detected.
Besides track formation and electronic sputtering, lattice disordering (the degradation of X-ray diffraction (XRD) intensity) with lattice expansion (an increase in the lattice parameter) by high-energy ions has been observed for the polycrystalline films of SiO2 [70] and WO3 [58], and lattice disordering with lattice compaction for those of Cu2O [57], CuO [59], Fe2O3 [60], Cu3N [61] and Mn-doped ZnO [71]. Only lattice disordering has been observed for the ultra-thin films of WO3 [72]. It should be noted that a comparison between high-energy and low-energy ion impact effects is important. Lattice expansion has been observed for a few keV D ion irradiation on Fe2O3 [73], and this can be understood by the incorporation of D into non-substitutional sites (incorporation or implantation effect). Thus, lattice expansion by medium-energy ions (100 keV Ne) on Fe2O3 [60] could be understood by Ne incorporation and/or interstitial-type defects generated by ion impact, with a possible stabilization by incorporating Ne in the film, whereas lattice compaction has been observed for a 100 MeV Xe ion impact on Fe2O3 [60]. It should be noted that the incorporation of ions in thin films does not take place for high-energy ions, since the projected range of ions (Rp) is much larger than the film thickness (e.g., Rp of 14 μm for 100 MeV Xe in SiO2), unless the thickness is too large. The lattice expansion due to the incorporation effect has been observed for a few keV H and D irradiation at a low fluence on WNOx with x ≈ 0.4, whereas lattice compaction has been observed at a high fluence of D [74]. Peculiarly, lattice expansion at a low ion fluence and compaction at a high fluence, as well as disordering, have been reported for medium-energy (100 keV Ne and N) and high-energy (100 MeV Xe and 90 MeV Ni) ion impact on WNOx [75,76]. One speculation is that the lattice compaction is due to vacancy-type defects generated by ion impact, which is to be investigated. Furthermore, a drastic increase in electrical conductivity has been observed for Cu3N [61], Mn-doped ZnO [71] and WNOx with x ≈ 0.4 [75,76]. The conductivity increase is ascribed to the increase in the carrier concentration and mobility.
There are a few reports on the Se dependence of the XRD intensity degradation per unit fluence (YXD) for SiC and KBr [56] and WO3 [72]. YXD is found to follow the power-law fit: YXD = (BXDSe)NXD, BXD being a material-dependent constant and the exponent NXD being comparable with the Nsp of the electronic sputtering. The results imply that similar mechanisms operate for lattice disordering and electronic sputtering. It is of interest to compare the Se dependence of lattice disordering with that of electronic sputtering for materials other than those mentioned above. In this paper, we have measured the lattice disordering of SiO2, ZnO, Fe2O3 and TiN films, and the sputtering of TiN. The XRD results are compared with the sputtering. The exciton model is examined and scaling parameters are explored for representing electronic excitation effects.

2. Materials and Methods

XRD has been measured using Cu-kα radiation. Accuracy of the XRD intensity is estimated to be approximately 10%, based on the variation of repeated measurements. Rutherford backscattering (RBS) has been performed with MeV He ions for evaluation of film thickness and composition. Similarly, accuracy of the RBS is estimated considering the variation of the repeated measurements. High-energy ion irradiation has been performed at room temperature and normal incidence. Irradiation of high-energy ion with lower incident charge than the equilibrium charge without carbon foil is often employed for the samples of XRD measurement; however, the effect of non-equilibrium charge incidence does not come into play because the length for attaining the equilibrium charge is much smaller than the film thickness, as described for each material in Section 3.
SiO2 films have been grown by thermal oxidation of Si(001) at 1300 °C for 5 hr. According to XRD, the films are polycrystalline with diffraction peaks at ~21°, 22°, 31°, 33°, 36° and 69°, with very weak peak at 44° and 47°. The peaks at ~21°, 22°, 44° and 47° have been assigned to (100), (002), (004) and (202) diffraction of hexagonal-tridymite structure [70]. The strong peak at 69° is Si(004) and peak at 33° is possibly Si(002). Film thickness is ~1.5 μm and the composition is stoichiometric (O/Si = 2.0 ± 5%) by RBS of 1.8 MeV He. Film density is taken to be the same as that of amorphous-SiO2 (a-SiO2), since it has been derived to be 2.26 gcm−3 from XRD results, which is close to that of a-SiO2 (2.2 gcm−3)
Pure ZnO films have been prepared on MgO (001) substrate by using a radio frequency magnetron sputtering (RFMS) deposition method with ZnO target, and it has been reported that the dominant growth orientation is (001) and (100) of hexagonal-wurtzite structure depending on the substrate temperature of 350 °C and 500 °C during the film growth, respectively [71,77,78]. The composition is stoichiometric, i.e., O/Zn = 1.0 ± 0.05, and film thickness is ~100 nm by He RBS. Here, the density is taken to be 4.2 × 1022 Zn cm−3 (5.67 gcm−3).
Preparation and characterization methods of Fe2O3 films are described in [60]. Briefly, Fe2O3 films have been prepared by deposition of Fe layers on SiO2-glass and C-plane cut Al2O3 (C-Al2O3) substrates using a RFMS deposition method with Fe target (99.99%) and Ar gas, followed by oxidation at 500 °C for 2–5 hr in air. According to RBS of 1.4–1.8 MeV He ions, the composition is stoichiometric (O/Fe = 1.5 ± 0.1) and film thickness used in this study is ~100 nm. Here, the density of 3.96 × 1022 Fe cm−3 (5.25 gcm−3) is employed. Diffraction peaks have been observed at ~33° and 36°, and crystalline structure has been identified as hexagonal Fe2O3 (hematite or α-Fe2O3). These correspond to (104) and (110) diffraction planes.
TiN films have been prepared on SiO2-glass, C-plane cut Al2O3 (C-Al2O3) and R-plane cut Al2O3 (R-Al2O3) substrates at 600°C using a RFMS deposition method with Ti target (99.5%) and high purity N2 gas. RBS of 1.4–1.8 MeV He ions shows that the composition is stoichiometric (N/Ti = 1.0 ± 0.05) and that the film thickness used in this study is ~170 nm (deposition time of 1 hr). Here, the density of 5.25x1022 Ti cm−3 (5.4 gcm−3) is employed. Diffraction peaks have been observed at 36.6°, 42.6° and ~77° on SiO2 glass and C-Al2O3. Crystalline structure has been identified as a cubic structure and these correspond to (111), (200) and (222) diffractions [79]. Diffraction intensity of (111) is larger than that of (200) on SiO2 glass, and diffraction of (111) on C-Al2O3 is very intensive. TiN on R-Al2O3 has preferential growth orientation of (220) of a cubic structure (diffraction angle at ~61°). Sputtered atoms are collected in the carbon foil (100 nm) and the sputtered atoms are analyzed by RBS to obtain the sputtering yields [54] (carbon collector method).

3. Results and Discussion

3.1. SiO2

The XRD intensity at the diffraction angle of ~22° (the most intensive (002) diffraction of hexagonal-trydimite) normalized to that of as-grown SiO2 films on Si(001) is shown in Figure 1 as a function of the ion fluence for 90 MeV Ni+10, 100 MeV Xe+14 and 200 MeV Xe+14 ion impact. The XRD intensity of the irradiated sample normalized to that of the unirradiated sample is proportional to the ion fluence to a certain fluence. Deviation from the linear dependence for the high fluence could be due to the overlapping effect. As observed in latent track formation (e.g., [5,6]), electronic excitation effects extend to a region (approximately cylindrical) with a radius of several nm and a length of the projected range or film thickness, and thus ions may hit the ion-irradiated part for a high ion fluence (called the overlapping effect). As described below, the XRD degradation yield per unit ion fluence (YXD) is reduced at a high fluence, and this could be understood as thermal annealing and/or a reduction in the disordered regions via ion-induced defects (recrystallization [26]). The damage cross-sections (AD obtained by RBS-channeling (RBS-C) technique and TEM [5]) are compared with YXD in Figure 2, and it appears that both agree well for Se > 10 keV. A discrepancy between AD and YXD is seen for Se < 10 keV, and the reason for this is not understood. In addition, sputtering yields are often reduced, and this is unlikely to be explained by the annealing effect. Therefore, the reasons for the sputtering suppression at a high fluence remain in question. The XRD degradation yields (YXD) per unit ion fluence are obtained and given in Table 1. The film thickness has been obtained to be ~1.5 μm, using 1.8 MeV He RBS. The attenuation length (LXA) of Cu-kα (8.0 keV) is obtained to be 128 μm [80] and the attenuation depth (LXA·sin(22°/2)) = 24.3 μm. The film thickness (~1.5 μm) is much smaller than the attenuation depth and thus no correction is necessary for the XRD intensity. The lattice expansion or increase in the lattice parameter of 0.5% with an estimated error of 0.2% at 1 × 1012 cm2 is found to be nearly independent of the electronic stopping power.
The electronic stopping power (Se*) appropriate for XRD intensity degradation is calculated using SRIM 2013, using a half-way approximation that the ion loses its energy for half of the film thickness (~0.75 μm), i.e., Se* = Se(E*) with E* = E(incidence) − Se(E) × 0.75 μm (Table 1). The correction for the film thickness on Se appears to be a few percent. It is noticed that the incident charge (Ni+10, Xe+14) differs from the equilibrium charge (+19, +25 and +30 for 90 MeV Ni, 100 MeV Xe and 200 MeV Xe, respectively (Shima et al.) [81], and +18.2, +23.9 and +29.3 (Schiwietz et al.) [82]), both being in good agreement. Following [64], the characteristic length (LEQ = 1/(electron loss cross-section times N)) for attaining the equilibrium charge is estimated to be 8.7, 8.3 and 7.9 nm for 90 MeV Ni+10, 100 MeV Xe+14 and 200 MeV Xe+14, respectively, from the empirical formula of the single-electron loss cross-section σ1L (10−16 cm2) of 0.52 (90 MeV Ni+10), 0.55 (100 MeV Xe+14) and 0.57 (200 MeV Xe+14) [83,84], N (2.2 × 1022 Si cm−3) being the density, and (target atomic number)2/3 dependence being included. Here, σ1L = σ1L(Si) + 2σ1L(O), ionization potential IP = 321 eV [85,86] with the number of removable electrons Neff = 8 and IP = 343 eV with Neff = 12 are employed for Ni+10 and Xe+14. LEQ is much smaller than the film thickness and hence the charge-state effect is insignificant.
The sputtering yields Ysp of SiO2 (normal incidence) are summarized in Table 2 for the comparison of the Se dependence of the XRD degradation yields YXD. There are various versions of TRIM/SRIM starting in 1985, and in this occasion, the results used the latest version of SRIM2013 are compared with those earlier versions. Firstly, the correction on the stopping power and projected range for carbon foils (20–120 nm), which have been used to achieve the equilibrium charge incidence, is less than a few %, except for low-energy Cl ions (several %). Secondly, Se by CasP (version 5.2) differs ~30% from that by SRIM 2013. Figure 2 shows the Se dependence of the XRD degradation yields YXD and Ysp. Both YXD and Ysp fit to the power-law of Se, and the exponents of XRD degradation NXD = 2.9 and Nsp = 3 (sputtering) are almost identical, indicating that the same mechanism is responsible for lattice disordering and sputtering. Further plotted is the sputtering yields vs. Se calculated using earlier versions of TRIM/SRIM (TRIM1985 to SRIM2010) [45,46,47,48,49,51], and the plot using earlier versions give the same exponent (Nsp = 3) with a 6% smaller constant Bsp in the power-law fit (20% smaller in the sputtering yields). This means that the plot and discussion using SRIM2013 do not significantly differ from those using the earlier versions of TRIM/SRIM. One notices that no appreciable difference in sputtering yields is observed among a-SiO2, films and single-crystal-SiO2 (c-SiO2) [45,46,47,48], even though the density of c-SiO2 is larger by 20% than that of a-SiO2, whereas much smaller yields (by a factor of three) have been observed for c-SiO2 [51]. The discrepancy remains in question. Sputtering yields YEC, which are due to elastic collision cascades, is estimated assuming YEC is proportional to the nuclear stopping power, discarding the variation of the α-factor (order of unity) depending on the ratio of target mass over ion mass (Sigmund) [87]. The proportional constant is obtained to be 2.7 nm/keV using the sputtering yields by low-energy ions (Ar and Kr) (Betz et al.) [88]. YEC is given in Table 2 and it is shown that Ysp/YEC ranges from 44 (5 MeV Cl) to 3450 (210 MeV Au).
In order to obtain the stopping powers (S) for the non-metallic compounds, such as SiO2, described above, we apply the Bragg’s additive rule, e.g., S(SiO2) = S(Si) + 2S(O) and S of the constituting elements is calculated using TRIM/SRIM and CasP codes. Before moving to the discussion of the Bragg’s deviation, the accuracy of S is briefly mentioned. It is estimated to be 8% (Be through U ions in Ag) near the maximum of S (~0.8 MeV/u) [66], 17% (K to U ions in Au) (Paul) [89]. Besenbacher et al. have reported no difference between solid and gas phases for 0.5–3 MeV He ion stopping in Ar with an experimental accuracy of 3% [90], and this could be understood by the fact that the binding energy (cohesive energy) of solid Ar is too small (0.08 eV (Kittel) [91]), compared with the ionization potential (IBethe) of 188 eV [92], to affect the stopping. On the other hand, Arnau et al. have reported a large deviation (~50% near the stopping power maximum at ~50 keV) for proton stopping between solid and gas phases of Zn, and the deviation reduces ~10% at ~1 MeV [93]. The cohesive energy of 1.35 eV [91] is much smaller than the mean ionization potential IBethe of 330 eV [92], and hence the small increase in IBethe cannot explain the Bragg’s deviation of Zn. They have argued that the difference in the 4s into 4p transition probability and screening effect between solid and gas phases are responsible. Both TRIM/SRIM and CasP codes are based on the experiments of conveniently available solid targets and molecular gas (e.g., N2, O2), and thus it is anticipated that the binding effect is included to some or large extent and that the Bragg’s deviation is not serious for nitrides and oxides, and is roughly 10% or less at around 1 MeV/u.

3.2. ZnO

The XRD intensity at a diffraction angle of ~34° ((001) diffraction) and 32° ((100) diffraction) normalized to those of unirradiated ZnO films on the MgO substrate is shown in Figure 3 as a function of the ion fluence for 90 MeV Ni+10, 100 MeV Xe+14 and 200 MeV Xe+14 ion impact. It appears that the XRD intensity degradation is nearly independent of the diffraction planes. The XRD intensity degradation per unit fluence YXD is given in Table 3, together with sputtering yields [54], stopping powers and projected ranges (SRIM2013). The X-ray (Cu-kα) attenuation length LXA is obtained to be 36.6 μm [80] and the attenuation depth is 11 and 10 μm for the diffraction angle of ~34° and 32°, respectively; thus, the X-ray attenuation correction is unnecessary. It appears that the appropriate energy for the YXD vs. Se plot, E − Seℓ/2, where ℓ = a film thickness of ~100 nm, is nearly the same as E* for sputtering, in which the energy loss of a carbon foil of 100 nm is considered. Similarly to the case of SiO2, the characteristic length (LEQ) is estimated to be 4.6, 4.4 and 4.2 nm for 90 MeV Ni+10, 100 MeV Xe+14 and 200 MeV Xe+14, respectively, from the empirical formula of the single-electron loss cross-section σ1L(10−16 cm2) of 0.52 (90 MeV Ni+10), 0.54 (100 MeV Xe+14) and 0.57 (200 MeV Xe+14) [83,84]. Here, σ1L = σ1L(Zn) + σ1L(O), and the ionization potential IP and Neff are described in Section 3.1. Again, LEQ is much smaller than the film thickness and the charge-state effect is insignificant.
Figure 4 shows the XRD intensity degradation YXD vs. electronic stopping power (Se) (SRIM2013 and TRIM1997) together with the sputtering yields Ysp vs. Se. Both YXD and Ysp follow the power-law fit and the exponent for YXD using TRIM1997 gives a slightly larger value than that using SRIM2013. The exponent of lattice disordering is nearly the same as that of sputtering. The change in the lattice parameter ∆ℓc appears to scatter, and roughly −0.2% and −0.1% with an estimated error of 0.1% are obtained for (100) and (002) diffractions by 100 MeV Xe at 10 × 1012 cm−2, assuming that ∆ℓc is proportional to the ion fluence. ∆ℓc is obtained at −0.3% for (002) diffraction by 200 MeV Xe at 5 × 1012 cm−2, and no appreciable change in the lattice parameter is observed by 90 MeV Ni ions at 40 × 1012 cm−2; more data are desired.

3.3. Fe2O3

The XRD intensity at a diffraction angle of ~33° and 36° (corresponding to diffraction planes of (104) and (110)) normalized to those of unirradiated Fe2O3 films on C-Al2O3 and SiO2 glass substrates as a function of the ion fluence is shown in Figure 5 for 90 MeV Ni+10, 100 MeV Xe+14 and 200 MeV Xe+14 ion impact. It appears that the XRD intensity degradation is nearly independent of the diffraction planes and substrates. The XRD intensity degradation per unit fluence YXD is given in Table 4, together with the sputtering yields [60] and stopping powers (SRIM2013). The X-ray (Cu-kα) attenuation length LXA is obtained to be 8.8 μm [80] and the attenuation depth is 2.5 and 2.7 μm for the diffraction angle of ~33° and 36°, respectively, which are much larger the film thickness of ~100 nm and thus the X-ray attenuation correction is unnecessary. The appropriate energy for the XRD vs. Se plot, using half-way approximation (E − Seℓ/2) with the film thickness ℓ of ~100 nm, again gives nearly the same as E* for sputtering, in which the energy loss of the carbon foil of 100 nm is taken into account.
Similarly to SiO2 and ZnO, the characteristic length (LEQ) is estimated to be 4.5, 4.3 and 4.1 nm for 90 MeV Ni+10, 100 MeV Xe+14 and 200 MeV Xe+14, respectively, from the empirical formula of the single-electron loss cross-section σ1L(10−16 cm2) of 0.56 (90 MeV Ni+10), 0.59 (100 MeV Xe+14) and 0.61 (200 MeV Xe+14) [83,84]. Here, σ1L = σ1L(Fe) + 1.5σ1L(O). LEQ is much smaller than the film thickness and the charge-state effect does not come into play.
Figure 6 shows XRD intensity degradation YXD vs. electronic stopping power (Se) (SRIM2013 and TRIM1997) together with the sputtering yields Ysp vs. Se. Both YXD and Ysp follow the power-law fit and the exponent using TRIM1997 gives a slightly larger fit than those using SRIM2013. The exponent of lattice disordering is two times larger than that of sputtering (Nsp is exceptionally close to unity, in contrast to the SiO2 and ZnO cases). The change in the lattice parameter appears to scatter depending on the substrate and diffraction planes, and is not proportional to the ion fluence. The average of the lattice parameter change in the (104) and (110) diffractions of Fe2O3 on C-Al2O3 is −0.2, −0.3% (an estimated error of 0.1%) and nearly zero at ~1 × 1012 cm−2 for 200 MeV Xe, 100 MeV Xe and 90 MeV Ni ion impact. The dependence of the lattice parameter change on the ion fluence and Se is complicated, and is to be investigated.

3.4. TiN

The XRD patterns are shown in Figure 7 for unirradiated and irradiated TiN films on the SiO2 glass substrate. As already mentioned in Section 2, (111) and (200) diffraction peaks are observed and the XRD intensity decreases due to ion impact. Figure 8 shows XRD intensities normalized to those of unirradiated TiN films on SiO2 glass, C-Al2O3 and R-Al2O3 substrates as a function of the ion fluence. It is seen that the XRD intensity degradation is nearly the same for the diffraction planes of (111) and (200) on SiO2, and for (111) on C-Al2O3. The XRD intensity degradation is less sensitive to the ion impact for the diffraction plane (220) on the R-Al2O3 substrate (~30% smaller than that for (111) and (200) on SiO2, and for (111) on C-Al2O3). The XRD intensity degradation per unit fluence YXD for (111) and (200) diffractions is given in Table 5, together with sputtering yields and stopping powers (TRIM1997 and SRIM2013). No appreciable change in the lattice parameter is observed, as shown in Figure 7. Similarly to the SiO2, ZnO and Fe2O3 cases, the appropriate energy, E − Seℓ/2, ℓ = film thickness of ~170 nm is taken into account, and the energy is close to that for sputtering, in which the energy loss of the carbon foil of 100 nm is considered. The X-ray (Cu-kα) attenuation length LXA is obtained to be 11.8 μm [80], and the attenuation depth is 3.7, 4.3 and 6.0 μm for diffraction angles of 36.6°, ~43° and 61°, respectively; thus, the X-ray attenuation correction is insignificant.
The characteristic length (LEQ) is estimated to be 4.5, 4.4, 4.2 and 4.0 nm for 60 MeV Ar+7, 90 MeV Ni+10, 100 MeV Xe+14 and 200 MeV Xe+14, respectively, from the empirical formula of the single-electron loss cross-section σ1L(10−16 cm2) of 0.43 (60 MeV Ar+7), 0.44 (90 MeV Ni+10), 0.46 (100 MeV Xe+14) and 0.48 (200 MeV Xe+14) [83,84]. Here, σ1L = σ1L(Ti) + σ1L(N), and the ionization potential IP and Neff are (IP = 143 eV and Neff = 1) for Ar+7, with those described in Section 3.1 for Ni+10 and Xe+14. LEQ is much smaller than the film thickness, and hence the charge-state effect is insignificant.
It is found that sputtered Ti collected in the carbon foil is proportional to the ion fluence, as shown in Figure 9 for 60 MeV Ar, 90 MeV Ni, 100 MeV Xe and 200 MeV Xe ions. The sputtering yield of Ti is obtained using the collection efficiency of 0.34 in the carbon foil collector [47] and the results are given in Table 5. Sputtered N collected in the carbon foil is obtained to be 0.4 × 1014 and 0.44 × 1014 cm−2 with an estimated error of 20% for 200 MeV Xe at 0.22 × 1012 cm−2 and 60 MeV Ar at 2.8 × 1012 cm−2, respectively, and this is comparable with the Ti areal density of 0.4 × 1014 cm−2 (200 MeV Xe) and 0.475 × 1014 cm−2 (60 MeV Ar). The results imply stoichiometric sputtering, due to the collection efficiency of N in the carbon foil collector of 0.35 [55], which is close to that of Ti. Thus, the total sputtering yield (Ti + N) is obtained by doubling Ysp (Ti) in Table 5. The sputtering yields of TiN (YEC) due to elastic collisions can be estimated assuming that YEC is proportional to the nuclear stopping power. Here, the proportional constant is obtained to be ~1.6 nm/keV using the experimental yields of 0.527 (0.6 keV Ar) and 0.427 (0.6 keV N) [94] and 0.7 (0.5 keV Cd) [88]. Ysp(TiN)/YEC ranges from 2.5 × 103 to 6 × 103. The XRD intensity degradations YXD and Ysp(Ti + N) are plotted as a function of the electronic stopping power Se in Figure 10. It appears that both fit to the power-law: YXD = (0.0224Se)1.26 and Ysp = (1.17Se)1.95. The exponents are comparable for XRD intensity degradation and sputtering.

4. Discussion

4.1. Comparison of Lattice Disordering with Sputtering

The electronic stopping power (Se) dependence of lattice disordering YXD, together with electronic sputtering, is summarized in Table 6, recognizing that most of the data have used TRIM1997. Results using SRIM2013 and TRIM1997 are compared in Section 3. Both exponents of the power-law fits are similar for SiO2, ZnO, Fe2O3, TiN and WO3 films, as well as for KBr and SiC. As mentioned in Section 3, it can be seen that the exponent of the lattice disordering NXD is comparable with that of sputtering Nsp, except for Fe2O3, in which Nsp is exceptionally close to unity, as in the case of Cu2O (Nsp = 1.0) [56] and CuO (Nsp = 1.08) [59]. The similarity of the exponent of lattice disordering and sputtering for SiO2, ZnO, Fe2O3, TiN, WO3, KBr and SiC imply that both phenomena originate from similar mechanisms, despite the fact that small displacements and annealing and/or the reduction in disordering via ion-induced defects are involved in the lattice disordering, whereas large displacements are involved in sputtering. The result of Fe2O3 indicates that the electronic excitation is more effective for lattice disordering. In the case of CuO, NXD is nearly zero [59]. In Table 6, YXD (10−12 cm2) at Se = 10 keV/nm and YXD/Ysp (×10−15 cm2) are listed. It is found that the ratio YXD/Ysp is an order of 10−15 cm2, except for ZnO, where the sputtering yields are exceptionally small. More data of lattice disordering would be desired for further discussion.

4.2. Electron–Lattice Coupling

Three models have been suggested for atomic displacement induced by electronic excitation: Coulomb explosion (CE) [3,4], thermal spike (TS) [50] and exciton model [30,95,96,97]. The neutralization time of the ionized region along the ion path is generally too short, and the fraction of the charged sputtered ions is small, e.g., 100 MeV Xe ions on SiO2 glass [48]. Hence, the CE model is unsound. However, a small atomic separation during the short time might be enough for electron–lattice coupling (a key for electronic excitation effects), which will be discussed later. A crude estimation of the evaporation yield for SiO2 based on the TS model appears to be far smaller than the experimental sputtering yield [55] and thus the TS model is also unsound. Moreover, the electron–lattice coupling or transfer mechanism of electronic energy into the lattice is not clear in the model. In the exciton model, the non-radiative decay of self-trapped excitons (STX, i.e., localized excited-state of electronic system coupled with lattice) leads to atomic displacement. According to the exciton model (or bandgap scheme), it is anticipated that the energy of the atoms in motion from the non-radiative decay of STX is comparable with the bandgap, leading to a larger sputtering yield with a larger bandgap, discarding the argument for the efficiency of STX generation from the electron–hole pairs, which is inversely proportional to the bandgap. This bandgap scheme is examined below. The effective depth contributing to the electronic sputtering of WO3 has been obtained to be 40 nm, which is nearly independent of Se [98], which would shed light on understanding the electronic sputtering; therefore, more data are desired.
The electronic sputtering yield Ysp super-linearly depends on the electronic stopping power (Se), and Ysp at Se = 10 keV/nm is taken to be a representative value, which is plotted as a function of the bandgap (Eg) in Figure 11 from [56], including the present TiN result. The optical absorbance (defined as log10(Io/I), Io and I being the incident and transmitted photon intensities) of TiN films are measured, and the direct bandgap Eg is obtained to be 4.5 eV for a film thickness of 25–50 nm, which decreases to 2.8 eV for a film thickness of ~180 nm by using the relation: (absorbance ● photon energy)2 is proportional to photon energy—Eg. The thickness dependence of Eg is under investigation by considering the influence of the reflectivity, film growth conditions and experimental problems, such as stray light, etc. A large variation has been reported for Eg, 4.0 eV (film thickness of 260 nm on Si substrate) (Popovic et al.) [99], 3.4 eV (thickness of ~100 nm on glass substrate) (Solovan et al.) [100] and 2.8–3.2 eV (film thickness of 460 nm on glass substrate) (Kavitha et al.) [101]. In this study, Eg is taken to be 4 eV and this choice is tolerable in the following discussion. It has been reported that the bandgap is reduced by 0.06 eV under a 400 KeV Xe ion implantation at 1016 cm−2 [99]. High-energy ion impact effects on optical properties are under way. It can be observed that the bandgap scheme seems to work for Eg > 3 eV [56]. A large deviation (two orders of magnitude) from the upper limit (dashed line indicated in Figure 11) is observed for ZrO2, MgO, MgAl2O4 and Al2O3. The existence of STX is known for limited materials, rare gas solids, SiO2 and alkali halides [30,95,96]. The STX does not exist for MgO and probably does not exist for Al2O3 [102]. The deviation for MgO and Al2O3 could be explained by the non-existence of STX. The numbers of electron–hole pairs leading to STX are inversely proportional to Eg, which could be a reason for the dependence of the sputtering yields for Eg < 3 eV. In any case, the single parameter of the band gap is insufficient for the explanation of the bandgap dependence of the sputtering yields.
Martin et al. [102] argued that STX exists for materials with small elastic constants. Following this suggestion, sputtering yields are plotted as a function of the elastic constant (C11) in Figure 12. Here, C11 (GPa) is taken to be 87 (SiO2), 348 (SrTiO3), 497 (Al2O3), 294 (MgO), 270 (TiO2), 210 (ZnO), 299 (MgAl2O4), 242 (Fe2O3), 35 (KBr) and 114 (LiF) [85], and, for other materials, 403 (CeO2) [103], 224 (Y2O3) [104], 400 (ZrO2) [105], 13 (WO3) [106], 126 (Cu2O) [107], 135 (CuO) [108], 388 (Si3N4) as an average of the values [109,110], 345 (polycrystalline-AlN) [111], which is smaller by 16% than 410 (AlN single crystal) [112], 234 (Cu3N) [113], 500 (SiC) [114] and 625 (TiN) [115]. It can be observed for oxides (the most abundant data are available at present) that Ysp decreases exponentially with an increase in the elastic constant for C11 < 300 GPa, except for MgO and ZrO2. Ysp for nitrides and SiC is larger than that for oxides at a given C11, and these are to be separately treated. It can be understood that the elastic constant represent the resistance of lattice deformation by electronic energy deposition. However, a single parameter, either the bandgap or elastic constant, is not adequate, and at least one more parameter is necessary. Furthermore, parameters other than those mentioned above are to be explored. More data for nitrides, alkali halides and especially carbides are desired.
Finally, a mechanism for the electron–lattice coupling is discussed. In an ionized region along the ion path, Coulomb repulsion leads to atomic motion, which is not adequate to cause sputtering because of its short neutralization time. Nevertheless, displacement comparable with the lattice vibration amplitude (one tenth of the average atomic separation, dav of ~0.25 nm for a-SiO2) is highly achievable during the neutralization time. As a first step, the time required for the Si+–O+ displacement of 0.025 nm (one tenth of dav) from dav is estimated to be ~15 fs using a formula [116]. Also, the time is estimated to be ~ 15 fs and ~ 12 fs for the Zn+–O+ displacement of 0.02 nm from dav of 0.2 nm in ZnO and for the Ti+–N+ displacement of 0.02 nm from dav of 0.2 nm in TiN, respectively. A similar situation has been reported for the Fe+–O+ displacement of 0.01 nm in Fe2O3 (~7 fs) [60], the K+–Br+ displacement of 0.01 nm in KBr (~9 fs) and the Si+–C+ displacement of 0.01 nm in SiC (~6 fs). These suggest a possibility that a small displacement comparable with the lattice vibration amplitude caused by Coulomb repulsion during the short neutralization time leads to the generation of a highly excited-state coupled with the lattice (h-ESCL), and h-ESCL is considered to be equivalent to STX or multi STX. The non-radiative decay of h-ESCL leads to atomic displacement (a larger displacement results in sputtering and smaller displacement results in phonon generation or lattice distortion).

5. Conclusions

We have measured the lattice disordering of polycrystalline SiO2, ZnO, Fe2O3 and TiN films, as well as the sputtering yield of TiN, by high-energy ion impact. It is found that lattice disordering is caused by electronic excitation and the degradation of the XRD intensity fits to the power-law on the electronic stopping power. The exponent in the fit of the XRD degradation is comparable with that of the electronic sputtering yield for these films, as well as the published results of WO3, KBr and SiC, implying that both lattice disordering and sputtering originate from similar mechanisms. In the case of Fe2O3, on the other hand, the exponent of the lattice disordering is larger by twice than that of the sputtering (the exponent for the sputtering is close to unity). The exciton mechanism seems to work for Eg > 3 eV, with some exceptions, and the elastic constant is examined as another scaling parameter for the electronic sputtering yields. A possibility of electron–lattice coupling is discussed based on a crude estimation that an atomic displacement comparable with the vibration amplitude due to Coulomb repulsion during the short neutralization time in the ionized region along the ion path can be attainable and, thus, the generation of a highly excited state coupled with the lattice is highly achievable, resulting in atomic displacement.

Author Contributions

Conceptualization and writing, N.M.; investigation, M.S., S.O. and B.T. All authors have read and agreed to the published version of the manuscript.

Funding

This research received no external funding.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Not applicable.

Acknowledgments

One of the authors (N.M.) thanks for utilization of XRD equipments of RIGAKU ULTIMA IV and RAD IIc at Radioisotope Research Center, Nagoya University and of AN Van de Graaff accelerator, Nagoya University for Rutherford backscattering spectroscopy (RBS) measurement. High-energy ion irradiation has been performed using a TANDEM accelerator at Tokai Research Center, JAEA.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. XRD intensity from (002) diffraction plane at ~22° normalized to as-grown films of SiO2 as a function of ion fluence for 90 MeV Ni (), 100 MeV Xe (o, ∆) and 200 MeV Xe (x) ions. Data of 90 MeV Ni () and 100 MeV Xe (∆) are from [70]. Linear fit is indicated by dashed lines. An estimated error of XRD intensity is 10%.
Figure 1. XRD intensity from (002) diffraction plane at ~22° normalized to as-grown films of SiO2 as a function of ion fluence for 90 MeV Ni (), 100 MeV Xe (o, ∆) and 200 MeV Xe (x) ions. Data of 90 MeV Ni () and 100 MeV Xe (∆) are from [70]. Linear fit is indicated by dashed lines. An estimated error of XRD intensity is 10%.
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Figure 2. XRD degradation per unit fluence YXD of polycrystalline SiO2 film (o, present result) and sputtering yield Ysp (x) of amorphous (or vitreous)-SiO2 (□, x)) and film of SiO2 (▲, , x, ◊, +) as a function of electronic stopping power (Se). Data (□, ▲) from (Qiu et al.) [45], () from (Sugden et al.) [46], (x) from (Matsnami et al.) [47,48], (◊) from (Arnoldbik et al.) [49] and (+) from (Toulemonde et al.) [51]. Se is calculated using SRIM2013, and power-law fits of YXD ((0.0545Se)2.9) and Ysp ((0.62Se)3.0) are indicated by blue and black dotted lines, respectively. Power-law fit (●) YXD ((0.055Se)3.4, TRIM1997) and Ysp ((0.58Se)3.0, TRIM1985 through SRIM2010) from [47,48,51] are indicated by black and green dashed lines. Damage cross sections (♦) are obtained by RBS-C and () by TEM from [5].
Figure 2. XRD degradation per unit fluence YXD of polycrystalline SiO2 film (o, present result) and sputtering yield Ysp (x) of amorphous (or vitreous)-SiO2 (□, x)) and film of SiO2 (▲, , x, ◊, +) as a function of electronic stopping power (Se). Data (□, ▲) from (Qiu et al.) [45], () from (Sugden et al.) [46], (x) from (Matsnami et al.) [47,48], (◊) from (Arnoldbik et al.) [49] and (+) from (Toulemonde et al.) [51]. Se is calculated using SRIM2013, and power-law fits of YXD ((0.0545Se)2.9) and Ysp ((0.62Se)3.0) are indicated by blue and black dotted lines, respectively. Power-law fit (●) YXD ((0.055Se)3.4, TRIM1997) and Ysp ((0.58Se)3.0, TRIM1985 through SRIM2010) from [47,48,51] are indicated by black and green dashed lines. Damage cross sections (♦) are obtained by RBS-C and () by TEM from [5].
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Figure 3. XRD intensity normalized to as-deposited films of ZnO as a function of ion fluence for 90 MeV Ni (∆, +), 100 MeV Xe (o, x) and 200 MeV Xe (□) ions. Diffraction planes are (002) at ~34° (∆, o, □) and (100) at ~32° (+, x). Linear fit is indicated by dashed lines. An estimated error of XRD intensity is 10%.
Figure 3. XRD intensity normalized to as-deposited films of ZnO as a function of ion fluence for 90 MeV Ni (∆, +), 100 MeV Xe (o, x) and 200 MeV Xe (□) ions. Diffraction planes are (002) at ~34° (∆, o, □) and (100) at ~32° (+, x). Linear fit is indicated by dashed lines. An estimated error of XRD intensity is 10%.
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Figure 4. XRD degradation per unit fluence YXD of polycrystalline ZnO film vs. electronic stopping power Se (TRIM1997 and SR2013). Power-law fit to YXD = (0.057Se)1.32 (TRIM1997) (o, blue dotted line) and (0.0585 Se)1.16 (SRIM2013) (+, black dotted line). Sputtering yield Ysp vs. Se (TRIM1997, □) and Se (SR2013, x) is also shown. Sputtering yield from [54]. Power-law fits to Ysp: (0.175 Se)1.57 for both Se from TRIM1997 and SR2013 is indicated by green dotted line.
Figure 4. XRD degradation per unit fluence YXD of polycrystalline ZnO film vs. electronic stopping power Se (TRIM1997 and SR2013). Power-law fit to YXD = (0.057Se)1.32 (TRIM1997) (o, blue dotted line) and (0.0585 Se)1.16 (SRIM2013) (+, black dotted line). Sputtering yield Ysp vs. Se (TRIM1997, □) and Se (SR2013, x) is also shown. Sputtering yield from [54]. Power-law fits to Ysp: (0.175 Se)1.57 for both Se from TRIM1997 and SR2013 is indicated by green dotted line.
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Figure 5. XRD intensity normalized to unirradiated films of Fe2O3 as a function of ion fluence for 90 MeV Ni (, , , ), 100 MeV Xe (o, , +, x) and 200 MeV Xe (●, ◊, ▼, ▷) ions. Diffraction peaks at ~33° of Fe2O3 films on C-Al2O3 substrate ((90 MeV Ni), o (100 MeV Xe), ● (200 MeV Xe)), ~36° of films on C-Al2O3 ( (90 MeV Ni), + (100 MeV Xe), ◊ (200 MeV Xe)), ~33° of films on SiO2 glass substrate ( (90 MeV Ni), (100 MeV Xe), ▼ (200 MeV Xe)) and ~36° of films on SiO2 glass substrate ( (90 MeV Ni), x (100 MeV Xe), ▷ (200 MeV Xe)). Data of 100 MeV Xe are from [60]. Linear fit is indicated by dotted lines. An estimated error of XRD intensity is 10%.
Figure 5. XRD intensity normalized to unirradiated films of Fe2O3 as a function of ion fluence for 90 MeV Ni (, , , ), 100 MeV Xe (o, , +, x) and 200 MeV Xe (●, ◊, ▼, ▷) ions. Diffraction peaks at ~33° of Fe2O3 films on C-Al2O3 substrate ((90 MeV Ni), o (100 MeV Xe), ● (200 MeV Xe)), ~36° of films on C-Al2O3 ( (90 MeV Ni), + (100 MeV Xe), ◊ (200 MeV Xe)), ~33° of films on SiO2 glass substrate ( (90 MeV Ni), (100 MeV Xe), ▼ (200 MeV Xe)) and ~36° of films on SiO2 glass substrate ( (90 MeV Ni), x (100 MeV Xe), ▷ (200 MeV Xe)). Data of 100 MeV Xe are from [60]. Linear fit is indicated by dotted lines. An estimated error of XRD intensity is 10%.
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Figure 6. XRD dgradation per unit fluence YXD of polycrystalline Fe2O3 films (, o) and sputtering yield Ysp (, x) as a function of the electronic stopping power (Se) in keV/nm. Power-law fits are indicated by dashed lines and Se is calculated using SRIM2013 (, ) and TRIM1997 (o, x): () YXD = (0.028 Se)2.28 (SRIM2013), (o) YXD = (0.029Se)2.54 (TRIM1997), () Ysp = (2.2 Se)1.05 (SRIM2013) and (x) Ysp = (1.16 Se)1.25 (TRIM1997). Sputtering data and power-law fit to the sputtering yields (TRIM1997) from [60].
Figure 6. XRD dgradation per unit fluence YXD of polycrystalline Fe2O3 films (, o) and sputtering yield Ysp (, x) as a function of the electronic stopping power (Se) in keV/nm. Power-law fits are indicated by dashed lines and Se is calculated using SRIM2013 (, ) and TRIM1997 (o, x): () YXD = (0.028 Se)2.28 (SRIM2013), (o) YXD = (0.029Se)2.54 (TRIM1997), () Ysp = (2.2 Se)1.05 (SRIM2013) and (x) Ysp = (1.16 Se)1.25 (TRIM1997). Sputtering data and power-law fit to the sputtering yields (TRIM1997) from [60].
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Figure 7. XRD patterns of TiN film on SiO2 glass substrate: unirradiated () and irradiated by 100 MeV Xe at 0.72 × 1012 cm−2 (+).
Figure 7. XRD patterns of TiN film on SiO2 glass substrate: unirradiated () and irradiated by 100 MeV Xe at 0.72 × 1012 cm−2 (+).
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Figure 8. XRD intensity normalized to unirradiated films of TiN as a function of ion fluence for 60 MeV Ar (, ), 90 MeV Ni (▼, +, , ▷), 100 MeV Xe (o, x, ◁) and 200 MeV Xe (▲, ●) ions. Diffraction plane (111) at diffraction angle of ~36.6° is indicated by , ▼, o and ▲ for SiO2 substrate, (200) at ~43° by , +, x and ● for SiO2 substrate, (111) by for C-Al2O3 substrate and (220) at ~61° by ▷ and ◁ for R-Al2O3 substrate. Linear fit is indicated by dotted lines. An estimated error of XRD intensity is 10%.
Figure 8. XRD intensity normalized to unirradiated films of TiN as a function of ion fluence for 60 MeV Ar (, ), 90 MeV Ni (▼, +, , ▷), 100 MeV Xe (o, x, ◁) and 200 MeV Xe (▲, ●) ions. Diffraction plane (111) at diffraction angle of ~36.6° is indicated by , ▼, o and ▲ for SiO2 substrate, (200) at ~43° by , +, x and ● for SiO2 substrate, (111) by for C-Al2O3 substrate and (220) at ~61° by ▷ and ◁ for R-Al2O3 substrate. Linear fit is indicated by dotted lines. An estimated error of XRD intensity is 10%.
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Figure 9. Areal density of sputtered Ti from TiN on SiO2 substrate collected in carbon foil vs. ion fluence for 60 MeV Ar (), 89 MeV Ni (∇), 99 MeV Xe (o) and 198 MeV Xe (∆) ions. An estimated error of areal density is 20%.
Figure 9. Areal density of sputtered Ti from TiN on SiO2 substrate collected in carbon foil vs. ion fluence for 60 MeV Ar (), 89 MeV Ni (∇), 99 MeV Xe (o) and 198 MeV Xe (∆) ions. An estimated error of areal density is 20%.
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Figure 10. XRD intensity degradation YXD (10−12 cm2) (o, +) and sputtering yields Ysp (Ti + N) (, x) vs. electronic stopping power Se (keV/nm). Se is calculated by TRIM1997 (o, ) and by SRIM2013 (+, x). Power-law fits are indicated by dotted lines: YXD = (0.0224Se)1.26 and Ysp = (1.17Se)1.95.
Figure 10. XRD intensity degradation YXD (10−12 cm2) (o, +) and sputtering yields Ysp (Ti + N) (, x) vs. electronic stopping power Se (keV/nm). Se is calculated by TRIM1997 (o, ) and by SRIM2013 (+, x). Power-law fits are indicated by dotted lines: YXD = (0.0224Se)1.26 and Ysp = (1.17Se)1.95.
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Figure 11. Sputtering yield at Se = 10 keV/nm vs. bandgap. Data from [56], TiN (present result) and LiF data from [62]. Dotted line is a guide for eyes (Eg > 3 eV).
Figure 11. Sputtering yield at Se = 10 keV/nm vs. bandgap. Data from [56], TiN (present result) and LiF data from [62]. Dotted line is a guide for eyes (Eg > 3 eV).
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Figure 12. Sputtering yield at Se = 10 keV/nm vs. elastic constant. Sputtering yield of TiN (present result), LiF from [62] and others from [56]. Dotted line is a guide for eyes for the most abundant available data of oxides (o, ).
Figure 12. Sputtering yield at Se = 10 keV/nm vs. elastic constant. Sputtering yield of TiN (present result), LiF from [62] and others from [56]. Dotted line is a guide for eyes for the most abundant available data of oxides (o, ).
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Table 1. XRD data of SiO2 films. Ion, incident energy (E in MeV), XRD intensity degradation (YXD), appropriate E* (MeV) considering the energy loss in the film and electronic stopping power in keV/nm (Se*) appropriate for YXD (see text). Se from SRIM2013. The deviation ∆Se* = (Se*/Se(E) − 1) × 100 is also given.
Table 1. XRD data of SiO2 films. Ion, incident energy (E in MeV), XRD intensity degradation (YXD), appropriate E* (MeV) considering the energy loss in the film and electronic stopping power in keV/nm (Se*) appropriate for YXD (see text). Se from SRIM2013. The deviation ∆Se* = (Se*/Se(E) − 1) × 100 is also given.
IonEnergyYXDE*Se*∆Se*
(MeV)(10−12 cm2)(MeV)(keV/nm)(%)
58Ni900.06684.57.246−0.32
136Xe1000.2791.011.56−3.2
136Xe2000.47518914.22−1.3
Table 2. Sputtering data of SiO2 (normal incidence). Ion, incident energy (E in MeV), energy (E* in MeV) corrected for the energy loss in carbon foils (see footnote), electronic stopping power (Se), nuclear stopping power (Sn), projected range (Rp) and sputtering yield (YSP). Se, Sn and Rp are calculated using SRIM2013. (Se(E*)/Se(E) − 1), (Sn(E*)/Sn(E) − 1) and (Rp(E*)/Rp(E) − 1) in % are given in the parentheses after Se(E*), Sn(E*) and Rp(E*), respectively. YSP in the parenthesis is for SiO2 films. Se(E) by CasP is also listed. YEC is the calculated sputtering yield due to elastic collisions.
Table 2. Sputtering data of SiO2 (normal incidence). Ion, incident energy (E in MeV), energy (E* in MeV) corrected for the energy loss in carbon foils (see footnote), electronic stopping power (Se), nuclear stopping power (Sn), projected range (Rp) and sputtering yield (YSP). Se, Sn and Rp are calculated using SRIM2013. (Se(E*)/Se(E) − 1), (Sn(E*)/Sn(E) − 1) and (Rp(E*)/Rp(E) − 1) in % are given in the parentheses after Se(E*), Sn(E*) and Rp(E*), respectively. YSP in the parenthesis is for SiO2 films. Se(E) by CasP is also listed. YEC is the calculated sputtering yield due to elastic collisions.
IonE(E*)Se(E*)Sn(E*)Rp(E*)YSPSe(CasP)YEC
(MeV)(keV/nm)(keV/nm)(μm)(keV/nm)
Qiu et al. [45]
35Cl5 (4.6)2.59 (−4.26)0.0426 (6.8)3.0 (−4.6)5.1 (4.4)1.870.12
35Cl20 (19.4)4.15 (−0.35)0.0134 (2.7)7.0 (−2.0)8.77 (8.22)3.830.036
Sugden et al. [46]
35Cl30 (29.9)4.24 (−2.4 × 10−3)9.3 × 10−3 (0.3)9.5 (−0.25)113.990.025
Matsunami et al. [47,48]
58Ni90 (89)7.265 (−0.055)0.0145 (0.84)18.3 (−0.60)1207.660.039
136Xe100 (99)11.88 (−0.49)0.091 (1.2)14.4 (−0.83)36214.00.246
136Xe200 (198)14.37 (−0.19)0.051 (0.73)21.9 (−0.51)40416.00.138
40Ar60 (60)4.406.5 × 10−316.3324.190.018
32S80 (80)3.493.3 × 10−3239.73.230.0089
Arnoldbik et al. [49]
63Cu50 (50)7.170.02711.6807.550.073
Toulemonde et al. [51]
197Au190 (190)16.90.14320.5142520.90.39
197Au190 (190) 1320
197Au197 (197)17.10.1420.9111021.20.38
197Au210 (210)17.40.1321.7123021.70.36
127I148 (148)12.90.062052515.30.16
58Ni69 (69)7.150.01815.51357.750.049
Equilibrium charge has been obtained by using carbon foils of 120 nm [45], 25 nm [46], 100 nm [47,48], 20 nm [49] and 50 nm [51].
Table 3. XRD data of ZnO films. Ion, incident energy (E in MeV), XRD intensity degradation (YXD), E* = E − ∆E (energy loss in carbon foil of 100 nm) (MeV) and electronic (Se*) and nuclear (Sn*) stopping powers in keV/nm and projected range Rp* (μm) at E* calculated using SRIM2013. Sputtering yield Ysp from [54]. Sputtering yield by 100 keV Ne ion is also given.
Table 3. XRD data of ZnO films. Ion, incident energy (E in MeV), XRD intensity degradation (YXD), E* = E − ∆E (energy loss in carbon foil of 100 nm) (MeV) and electronic (Se*) and nuclear (Sn*) stopping powers in keV/nm and projected range Rp* (μm) at E* calculated using SRIM2013. Sputtering yield Ysp from [54]. Sputtering yield by 100 keV Ne ion is also given.
IonEnergyYXDE*Se*Sn*Rp*Ysp
(MeV)(1014 cm2)(MeV)(keV/nm)(keV/nm)(μm)
32S80 806.620.007131.09
40Ar60 608.30.0149.52.08
58Ni900.7888913.720.031114.0
127I85 8418.920.2058.87.0
136Xe1001.39921.600.208.8
136Xe2001.719827.140.1121311
20Ne0.100.10 0.290.240.120.9
Table 4. XRD data of Fe2O3 films. Ion, energy (E in MeV), XRD intensity degradation (YXD), E* = E-∆E (energy loss in carbon foil of 100 nm) (MeV) and electronic (Se*) and nuclear (Sn*) stopping powers in keV/nm and projected range Rp* (μm) calculated using SRIM2013. Sputtering yield Ysp from [60]. Results by low energy (100 keV Ne) ion are also given.
Table 4. XRD data of Fe2O3 films. Ion, energy (E in MeV), XRD intensity degradation (YXD), E* = E-∆E (energy loss in carbon foil of 100 nm) (MeV) and electronic (Se*) and nuclear (Sn*) stopping powers in keV/nm and projected range Rp* (μm) calculated using SRIM2013. Sputtering yield Ysp from [60]. Results by low energy (100 keV Ne) ion are also given.
IonEnergyYXDE*Se*Sn*Rp*Ysp
(MeV)(10−12 cm2)(MeV)(keV/nm)(keV/nm)(μm)
58Ni900.128914.280.0309.838.3
136Xe1000.389923.250.197.957.9
136Xe2000.6019828.270.1111.781.7
20Ne0.10 0.100.3540.2580.122.3
Table 5. XRD data of TiN films. Ion, energy (E in MeV), XRD intensity degradation (YXD) for (111) and (200) diffraction on SiO2 and C-Al2O3, substrates, YXD for (220) diffraction on R-Al2O3 in the parenthesis, E* = E − ∆E (energy loss in carbon foil of 100 nm) (MeV) and electronic (Se*) and nuclear (Sn*) stopping powers in keV/nm and projected range Rp* (μm) calculated using SRIM2013 and sputtering yield Ysp of Ti. Se* (TRIM1997) is given in parenthesis.
Table 5. XRD data of TiN films. Ion, energy (E in MeV), XRD intensity degradation (YXD) for (111) and (200) diffraction on SiO2 and C-Al2O3, substrates, YXD for (220) diffraction on R-Al2O3 in the parenthesis, E* = E − ∆E (energy loss in carbon foil of 100 nm) (MeV) and electronic (Se*) and nuclear (Sn*) stopping powers in keV/nm and projected range Rp* (μm) calculated using SRIM2013 and sputtering yield Ysp of Ti. Se* (TRIM1997) is given in parenthesis.
IonEnergyYXDE*Se*Sn*Rp*Ysp(Ti)
(MeV)(10−12 cm2)(MeV)(keV/nm)(keV/nm)(μm)
40Ar600.14609.41 (9.33)0.01357.651.8
58Ni900.27 (0.2)8915.5 (16.5)0.03058.6147
136Xe1000.50 (0.35)9926.7 (25.5)0.196.9 380
136Xe2000.6019830.85 (30.25)0.1110529
Table 6. Summary of electronic stopping power (Se in keV/nm) dependence of lattice disordering YXD = (BXDSe)NXD for the present results of SiO2, ZnO, Fe2O3 and TiN films, and sputtering yields Ysp = (BspSe)Nsp of the present result for TiN. Lattice disordering and sputtering yields of WO3 film from [58,72], those of KBr and SiC from [56] and sputtering yields of SiO2, ZnO and Fe2O3 (see Section 3). Constant BXD and Bsp and the exponent NXD and Nsp are obtained using TRIM1997 and those using SRIM2013 are in parentheses. YXD at Se = 10 keV and YXD/Ysp (10−15 cm2) are given.
Table 6. Summary of electronic stopping power (Se in keV/nm) dependence of lattice disordering YXD = (BXDSe)NXD for the present results of SiO2, ZnO, Fe2O3 and TiN films, and sputtering yields Ysp = (BspSe)Nsp of the present result for TiN. Lattice disordering and sputtering yields of WO3 film from [58,72], those of KBr and SiC from [56] and sputtering yields of SiO2, ZnO and Fe2O3 (see Section 3). Constant BXD and Bsp and the exponent NXD and Nsp are obtained using TRIM1997 and those using SRIM2013 are in parentheses. YXD at Se = 10 keV and YXD/Ysp (10−15 cm2) are given.
SampleBXDNXDBspNspYXDYXD/Ysp
(nm/keV)(nm/keV)(10−12 cm2)(10−15 cm2)
(Se = 10 keV/nm)
SiO20.055 (0.0545)3.4 (2.9)0.58 (0.62)3.0 (3.0)0.130.67
ZnO0.057 (0.0585)1.32 (1.16)0.175 1.570.476198
Fe2O30.029 (0.028) 2.54 (2.28)1.16 (2.2)1.25 (1.05)0.0432.0
TiN0.02241.261.171.950.151.26
WO30.073552.650.653.60.440.53
KBr0.1272.40.773.01.783.9
SiC0.03771.971.861.530.151.7
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Matsunami, N.; Sataka, M.; Okayasu, S.; Tsuchiya, B. Modification of SiO2, ZnO, Fe2O3 and TiN Films by Electronic Excitation under High Energy Ion Impact. Quantum Beam Sci. 2021, 5, 30. https://doi.org/10.3390/qubs5040030

AMA Style

Matsunami N, Sataka M, Okayasu S, Tsuchiya B. Modification of SiO2, ZnO, Fe2O3 and TiN Films by Electronic Excitation under High Energy Ion Impact. Quantum Beam Science. 2021; 5(4):30. https://doi.org/10.3390/qubs5040030

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Matsunami, Noriaki, Masao Sataka, Satoru Okayasu, and Bun Tsuchiya. 2021. "Modification of SiO2, ZnO, Fe2O3 and TiN Films by Electronic Excitation under High Energy Ion Impact" Quantum Beam Science 5, no. 4: 30. https://doi.org/10.3390/qubs5040030

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