Effect of high electronic energy deposition in semiconductors

Nuclear Instruments and Methods in Physics Research B 225 (2004) 111–128 www.elsevier.com/locate/nimb

Effect of high electronic energy deposition in semiconductors W. Wesch *, A. Kamarou, E. Wendler Institut f€ ur Festk€orperphysik, Friedrich–Schiller-Universit€at Jena, Max-Wien-Platz 1, D-07743 Jena, Germany Received 13 April 2004; received in revised form 27 April 2004

Abstract Track formation due to high electronic energy deposition during swift heavy ion irradiation is well known in insulating materials as well as in some intermetallic compounds and metals. In semiconductors the physical situation was much less clear, and only during the last few years several experimental results on the effect of high electronic energy deposition were published, which are summarised and discussed in the present paper. Like in insulators, swift heavy ion irradiation may cause amorphous tracks in some semiconductors, such as InP, InSb, InAs, GaSb and Ge, if a certain value of the electronic energy deposition ee per ion and unit length characteristic for the material and the ion is exceeded. The critical values of ee are much higher than in the other materials, and the corresponding ion energies are close to the maximum ion energies available with the existing high energy accelerators. On the other hand, with cluster ions, as e.g. C60 , tracks are easily formed in Si, Ge and GaAs with ion energies of several tens of MeV. Beside damage and track formation, annealing of damage was observed in the semiconductors, and it can be concluded that the effect of high electronic energy deposition represents itself as a competition between damage formation and annealing. Which of the two processes dominates is mainly determined by the electronic energy deposition. Qualitatively the observed behaviour can be explained in the framework of the thermal spike model. However, a quantitative description of all data available is not successful in the framework of this concept. This is obviously due to the fact that the effects are influenced not only by the electronic energy deposition, but also by a variety of other parameters, such as the ion velocity, the ion mass, the charge state of the impinging ions and the radial distribution of the electronic energy around the ions’ path. All these parameters have to be taken into account within a theoretical description. However, the present situation is characterised by a lack of sufficient experimental data, and further systematic work is required to make progress in this interesting field.
2004 Elsevier B.V. All rights reserved. PACS: 61.80.Jh; 61.82.Fk; 81.05.Ea Keywords: Semiconductors; Swift heavy ions; Electronic energy deposition; Damage formation and annealing

1. Introduction

*

Corresponding author. Tel.: +49-3641-947330; fax: +493641-947302. E-mail address: [email protected] (W. Wesch).

The growing interest in studying the interaction of swift heavy ions with energies of several hundreds of MeV up to GeV with solids is due to several aspects. The first one is a consequence of

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2004 Elsevier B.V. All rights reserved. doi:10.1016/j.nimb.2004.04.188

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future technological requirements: the tendency to move to higher implantation energies in order to form buried layers with modified properties requires to study the influence of the increasing electronic energy deposition on the structural modifications connected with the irradiation. A way to investigate this influence is the use of swift heavy ions because at these high energies the depth regions of dominating electronic and nuclear interaction are clearly separated. The second point is the damage caused in devices of future superhigh energy particle accelerators or in equipment used for space applications, which can also be studied using swift heavy ions. And last, but not least, the use of swift heavy ions to produce nanostructures in various materials is a further stimulating aspect. During the last years numerous results have been published concerning the effect of swift heavy ion irradiation in insulators and metals and, to a less extent, in semiconductors. While in many dielectrics and some metals the high electronic energy deposition leads to the formation of tracks around ion trajectories (see [1–3] and references therein), the situation is much more complicated in semiconductors. In thin amorphous Si and Ge layers crystalline tracks were found after swift heavy ion irradiation [4]. Clear evidence for the formation of amorphous tracks in crystalline material by irradiation with monoatomic ion beams was recently found in some compound semiconductors, as GeS [5], InP [6–10], InAs, InSb and GaSb [10] and in crystalline SiGe alloys [11]. Some evidence on track formation in Ge after irradiation with heavy ions in the GeV energy range was reported very recently [12]. The results published until now show that the formation of visible tracks in crystalline semiconductors, as in metals and insulators, requires a threshold value of the electronic energy deposition characteristic for the corresponding material which obviously cannot be reached with individual ions in various semiconductors, as Si and GaAs. On the other hand, amorphous tracks are easily formed in Si, Ge and GaAs by irradiation with MeV fullerenes with its higher electronic energy deposition [13– 16]. It should be mentioned that the data available do not show any correlation between track for-

mation and the macroscopic physical properties of the semiconductors, as e.g. band gap, melting point and electrical properties. The present paper summarises the experimental data published up to now on the effect of high electronic energy deposition in semiconductors with some emphasis on results obtained in our group.

2. Energy deposition An ion penetrating a solid suffers a successive loss of its energy mainly due to elastic collisions with the target nuclei and inelastic interactions with the electron system. These processes are described by the depth dependent energy depositions per ion and unit length en ðzÞ ¼ ðdE=dzÞn (nuclear energy deposition) and ee ðzÞ ¼ ðdE=dzÞe (electronic energy deposition) and can be calculated using the TRIM Monte Carlo code [17]. Which of the processes dominates in a certain depth is, for a given material and ion mass, determined by the ion energy. Generally, the ratio ee =en increases when moving from keV to MeV energies. This is illustrated in Fig. 1 for the example of Xe irradiation of InP. For 400 keV irradiation the depth distributions of electronic as well as nuclear energy deposition are located close to the surface, and en exceeds ee by a factor of approximately 10. In the case of 15 MeV the electronic energy deposition

Fig. 1. TRIM 97 calculations of the energy deposition in ionization processes (electronic energy deposition ee ) and in displacements (nuclear energy deposition en ) versus depth z for InP irradiated with Xe ions of different energies.

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near the surface lies already remarkably above the nuclear one, and en reaches its maximum at around 3.8 lm. When irradiating with 390 MeV, ee reaches a value of about 19 keV/nm which remains almost constant down to a depth of  5 lm. In this layer the electronic energy deposition exceeds the nuclear one by a factor of approximately 103 . At depths z larger than 5 lm the electronic energy deposition decreases and the nuclear energy deposition increases reaching its maximum at z ¼ 27 lm with en  0:9 keV/nm. This clearly illustrates the dominating effect of the electronic energy deposition in near surface layers of samples irradiated with ions in the several hundred MeV energy range, which allows an almost separate study of the influence of ionisation processes on damage formation and transformation.

3. Damage formation 3.1. Point defect production Since the early 1990s several papers have been published dealing with swift heavy ion irradiation of Si, Ge and GaAs [18–29]. Main goal of these studies was to investigate the modification of the electrical properties of the materials under swift heavy ion irradiation. For this purpose n- and pdoped substrates were irradiated with various ions (oxygen to uranium) with energies per nucleon ranging from 3 to 60 MeV/u corresponding to electronic energy depositions ee between about 1 and 28 keV/nm. In most cases in situ resistance and Hall mobility measurements were performed. Additionally, in some studies the samples were characterised ex situ by means of deep level transient spectroscopy (DLTS), photo luminescence (PL), optical absorption, electron paramagnetic resonance (EPR) and X-ray diffraction (XRD) techniques. The results can be summarised as follows. In n-Si and n-GaAs an ion fluence-dependent decrease of the mobility and an increase of the resistance, until saturation values are reached at certain ion fluences, were observed ([18,19,21,27] and references therein). This is illustrated for the example of 5.2 GeV Kr irradiation of n-Si in Fig. 2 [27]. In n-Ge the resistance

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Fig. 2. Resistance q and Hall mobility lH of n-Si (n ¼ 7  1011 P cm2 ) irradiated with 5.2 GeV Kr ions versus the ion fluence (data from [27]).

decreases after having passed a maximum, and the initially negative Hall mobility becomes positive indicating a type conversion from n- to p-type. This is shown in Fig. 3 for the case of 880 MeV Pb irradiation [27]. In p-Ge the conductivity and the majority carrier density increase with the ion fluence [20]. The observed modifications of the electrical properties are obviously a consequence of the formation of point defects due to swift heavy ion irradiation. These lead to a compensation of free carriers in n-Si and n-GaAs and a type conversion of n-Ge to p-Ge. A comparison of DLTS and PL results of swift heavy ion and proton, electron and neutron irradiated samples shows that the defects are similar in all cases. In Si and

Fig. 3. Resistance q and Hall mobility lH of n-Ge (n ¼ 2:4  1014 Sb cm2 ) irradiated with 880 MeV Pb ions versus the ion fluence (data from [27]).

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Ge defect complexes as the A-centre (vacancy– oxygen complex), the E-centre (vacancy-doping impurity complex) and the divacancy are the most frequent defects detected [18,23,28–30]. The main defect produced by swift heavy ions in GaAs is the As vacancy [27]. That swift heavy ion irradiation with an electronic energy deposition below 28 keV/ nm in GaAs and Ge leads only to the formation of point defects is also confirmed by Rutherford backscattering spectrometry in channelling configuration (RBS/c). As an example, Fig. 4 shows channelling RBS spectra of GaAs samples irradiated with various fluences of 390 MeV Xe20þ ions, ee ¼ 22:6 keV/nm in this case. It is clearly to be seen that under these irradiation conditions only a slight increase of the yield of backscattered ions occurs with respect to the aligned spectrum of virgin material. At an ion fluence NI ¼ 3  1014 cm2 the relative concentration of displaced lattice atoms, nda , calculated from the spectra by means of the DICADA code [31] amounts only to  0.01 (not shown) indicating that the increase of the backscattering yield is due to de-channelling of the analysing ions at point defects. To what extent the electronic energy deposition contributes to the observed defect production cannot be completely answered. At lower values of ee the defect production rates scale with the nuclear energy deposition in Si [18,24] and Ge [20]. In the case of heavier ions (increasing ee ) the defect production

Fig. 4. Energy spectra of 1.4 MeV He ions backscattered from GaAs irradiated at room temperature with 390 MeV Xe ions to various ion fluences.

rate decreases indicating in situ annealing of defects due to the high electronic excitation [18,20, 24]. Such an annealing due to high ee is also found in GaAs for ee ranging from  15 to  20 keV/nm [22]. For epitaxial GaAs layers it is concluded in [25] that the electronic energy deposition has an influence neither on defect production nor on annealing if ee 6 12 keV/nm [25]. Consequently there exists no clear evidence of ionisation induced defect production in Si, Ge and GaAs by individual ions for an electronic energy deposition below  28 keV/nm. The detected point defects and point defect complexes can mainly be ascribed to the effect of nuclear interaction, and these defects also may anneal already during the irradiation due to the electronic energy loss of the impinging ions. 3.2. Formation of tracks 3.2.1. Mono-atomic ion beams Until 1998 no indication concerning the formation of significant damage and/or amorphous tracks due to high electronic excitation was found in conventional semiconductors as Si, Ge and the binary III–V compounds. Exceptions were GeS [5] and the wide band gap semiconductors BN3 and diamond [32,33]. In GeS by means of high resolution transmission electron microscopy (TEM) elliptical amorphous tracks were registered after swift heavy ion irradiation with energies ranging from 5.9 to 13 MeV/u. In diamond and BN3 high pressure local regions were observed to correlate well with the depth distribution of the electronic stopping power and were interpreted as to be due to tracks. In 1997 it was reported that in Ge after irradiation with 100–262 MeV Au ions at an incidence angle of 75 to the surface normal thin amorphous surface layers are formed [34]. At ion fluences above  1014 cm2 surface rippling occurs. These surface modifications are absent when Ge is irradiated at nearly normal incidence of the ions and when lighter ions, such as iodine or nickel, are used. The observed effects cannot be explained by the electronic energy deposition, but they are obviously due to the nuclear interaction in combination with the influence of the free surface on damage accumulation. The question why these effects do not occur for irradiation under normal

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incidence remains to be answered. Also the occurrence of a buried sponge-like layer around the maximum of the nuclear energy deposition is more due to nuclear than electronic interaction. The formation of heavily damaged surface layers and amorphous tracks in InP after swift heavy ion irradiation was for the first time demonstrated in experiments performed in our group in 1998 [6,7] and was confirmed in 2002 by the data published by Szenes et al. [10]. The parameters for the InP irradiations are given in Table 1. In these studies the electronic energy deposition ee was varied between 12.5 and 30.4 keV/nm. Fig 5 shows, as an example, the RBS spectra of InP irradiated with various fluences of 250 MeV Xe17þ ions at room temperature [6]. The yield of He ions backscattered in the near surface region of the samples increases with the ion fluence and reaches the random level for fluences NI > 2  1013 cm2 indicating the formation of an amorphous layer. The relative concentration of displaced lattice atoms, nda , calculated from the RBS spectra with the use of the DICADA code [31] shows that in the depth region z P 0:1 lm nda is almost independent of z and increases with the ion fluence until the maximum value nda ¼ 1 is reached (Fig. 6). It is also to be seen that a  40 nm thick layer immediately behind the surface is only slightly damaged. Similar results were obtained for 340 MeV Xe20þ irradiation [7]. In order to understand the significant damage evolution in the region of high electronic energy deposition it is helpful to introduce the number of displacements per lattice atom, ndpa , according to   ndpa ¼ Ndispl NI =N0 (Ndispl is the number of displacements per ion and unit depth calculated by TRIM97 [17], NI the ion fluence, and N0 the target atomic density) which is proportional to the nuclear energy deposition. In Fig. 7 the relative Table 1 Irradiation parameters Ion

E (MeV)

E/u (MeV/u)

ee (keV/nm)

Kr Xe Au Pb

150 240–390 79–600 385–2080

1.78 1.83–2.98 0.40–3.04 1.85–10.0

12.5 18.7–19.7 13.2–28.6 27.2–30.4

Data in italics from [10].

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Fig. 5. Energy spectra of 1.4 MeV He ions backscattered from InP irradiated at room temperature with various ion fluences of 250 MeV Xe ions (from [6]).

Fig. 6. Relative concentration of displaced lattice atoms, nda , calculated from the spectra in Fig. 5 as a function of depth z (from [6]).

concentrations of displaced atoms, nda , calculated from the RBS spectra, behind the surface (z ¼ 0:3 lm) are depicted versus ndpa for various swift heavy ion irradiations of InP. For comparison, data of 2.5 MeV Kr irradiation in the maximum of the damage profiles (z ¼ zmax  1 lm) are included. It is obvious that for both Kr energies nda is approximately described by the same dependence on the number of displacements per atom, ndpa . In the case of 150 MeV Kr10þ irradiation up to ndpa ¼ 0:2 dpa only weak damage is produced. At the ion energy of 2.5 MeV damage formation is caused by nuclear interactions, and amorphisation requires ndpa  0:6 dpa. Therefore it can be

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Fig. 7. Relative concentration of displaced lattice atoms, nda , versus the number of primary displacements per lattice atom, ndpa , for 340 and 250 MeV Xe and 150 and 2.5 MeV Kr irradiation in InP (data from [7]).

concluded that the damage evolution in samples irradiated with 150 MeV Kr10þ has to be ascribed to the nuclear energy deposition en . At Xe ion energies of 250 MeV and 340 MeV nda ¼ 1, i.e. amorphisation is already reached for ndpa  0.016 dpa. In the case of 150 MeV Kr10þ irradiation at this energy deposition the measured relative damage concentration amounts to nda ¼ 0:08 only. It is known that the nuclear energy deposition required to amorphise InP, depending on the ion mass, corresponds to 0.2 dpa for 300 keV and 600 keV Se [35] and 1 dpa for 100 keV B [36]. For a nuclear energy deposition of 0.016 dpa with conventional ion energies (several 100 keV to a few MeV) a value nda 0:05 follows, i.e. only a very small defect concentration is produced due to nuclear processes. Therefore, the formation of amorphous surface layers by swift Xe irradiation as described above is not a consequence of nuclear energy deposition but due to the high electronic energy loss starting at the surface and extending to depths of several micrometers. Furthermore, since Kr irradiation does not produce amorphous material close to the surface up to ndpa ¼ 0:2 dpa, it has to be concluded that there exists a threshold value of the electronic energy deposition to form remarkable damage. In order to get information concerning morphology and kind of defects in the irradiated re-

gions selected samples were investigated by means of cross-sectional transmission electron microscopy (XTEM). In samples irradiated with 250 MeV Xe17þ and 340 MeV Xe20þ at ion fluences NI 6 5  1012 cm2 tracks could not be detected. The layers in these cases contain high concentrations of point defects and point defect complexes which is in agreement with the RBS data (Figs. 5 and 6). At ion fluences of (6–8) · 1012 cm2 , discontinuous and continuous amorphous tracks were found, the concentration of which increases with further rising ion fluence until at NI  5  1013 cm2 an amorphous surface layer is formed [6,7]. The morphology of the damage produced along the ion trajectory significantly changes with the depth. This is illustrated in Fig. 8 showing bright-field XTEM images taken at various depths of a sample irradiated with 250 MeV Xe17þ ions to a fluence of NI ¼ 7  1012 cm2 (corresponding to ndpa ¼ 3  103 dpa and nda ¼ 0:37 at z ¼ 0:3 lm; see Fig. 7). After a slightly damaged layer of  35 nm (see also RBS data in Figs. 5 and 6) in the depth region from  40 to 100 nm a layer containing a great number of tracks of isolated spherical or elongated cylindrical defect clusters along the ion trajectory is detected (part (A) in Fig. 8). In depths between 100 nm and about 10 lm (parts (A)–(C) in Fig. 8) straight lines with dark contrast occur in the XTEM images. The shape of the cross-section of these damage regions is nearly circular, and the track diameters were estimated from both plane-view (not shown) and cross-section TEM images to be about 7–15 nm. The density of tracks of 2 · 1011 cm2 is remarkably smaller than the ion fluence indicating that under these conditions not each impinging ion produces a visible track [6,7]. From the microdiffraction pattern (insertion in part (B) of Fig. 8) it is concluded that the inner structure of the columnar defects consists of a mixture of amorphous InP and probably a small amount of fine polycrystalline grains. Behind the track region at depths between 10 and 17 lm a crystalline InP layer was registered containing only clusters of point defects (part (D) of Fig. 8), and between 19 and 21 lm a band of heavily damaged InP (part (E) of Fig. 8) occurs which consists of amorphous and crystalline regions and coincides with the po-

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Fig. 8. Bright-field XTEM images of 250 MeV Xe-irradiated InP for NI ¼ 7  1012 cm2 obtained at different depths (from [6]).

sition of the maximum of the nuclear energy deposition. 250 MeV Xe17þ irradiation to an ion fluence of 5 · 1013 cm2 results in the formation of a similar layer sequence as observed in the case of irradiation at 7 · 1012 cm2 . The surface layer extending down to about 12 lm is now amorphous with some crystallites embedded (see Fig. 9(A) as an example for this depth region). This layer is followed by a transition region with spherical or elongated cylindrical zones consisting of amorphous and defective crystalline InP (z ¼ 12 to  15 lm; see Fig. 9(B)), an intermediate damaged but still crystalline layer (z ¼ 15–19 lm), and the second amorphous layer extending from 19 to 21 lm

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which is due to the nuclear energy loss in this depth. It should be mentioned that the ion fluence of 5 · 1013 cm2 corresponds to a nuclear energy deposition ndpa ¼ 0:55 dpa in the maximum of the ndpa depth distribution which causes the formation of an amorphous layer in InP at conventional ion energies too [35,37]. From panoramic XTEM images (Fig. 8) the complete depth distribution of damage was deduced which is schematically illustrated in Fig. 10 for NI ¼ 7  1012 cm2 together with the depth distributions of electronic and nuclear energy deposition. The figure shows that defects in the near surface region are only produced by the electronic energy loss if it exceeds a critical value of about 13 keV/nm, which is reached at a depth of 10 lm corresponding to the maximum depth of damage detected by means of XTEM. Furthermore, the existence of the intermediate crystalline layer supports the conclusion that the near surface damage is caused by electronic excitation and not by the small amount of nuclear energy loss in this region, because otherwise this crystalline region where the nuclear energy deposition is still higher  (Ndispl  10 (ion nm)1 corresponding to ndpa ¼  0:017 dpa) than close to the surface (Ndispl  1:6 1 (ion nm) corresponding to ndpa ¼ 0:0028 dpa) should not exist. Another interesting feature is the appearance of an only slightly damaged surface layer up to high Xe fluences. This may perhaps be an effect of the surface which could act as a sink for defects. However, irradiation through a 20 nm thick amorphous surface layer showed an intermediate slightly damaged region which was not present after irradiation through a 40 nm thick amorphised surface layer. This indicates that ion charge state effects are a more probable explanation than surface effects. Statistical fluctuations of the ion charge state which can reduce ee periodically below its threshold value for track formation may also explain the discontinuity of the tracks [12]. Our results show that in the case of Xe irradiation of InP with ee values around 19 and 20 keV/nm, amorphous tracks are formed if the critical value of ee  13 keV/nm is exceeded. Besides, it is obvious that not each ion produces a track but point defects and point defect complexes.

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Fig. 9. Bright-field XTEM images of 250 MeV Xe-irradiated InP for NI ¼ 5  1013 cm2 obtained at different depths: depth regions 1–10 lm (A) and 12–15 lm (from [7]).

If a critical concentration of these defects is exceeded amorphous tracks are formed, i.e. a certain pre-damage level is necessary under these conditions. In order to prove the influence of pre-damage on track formation, pre-damaged InP (1 MeV Si at TI ¼ 190 C, NI ¼ 2:5  1013 cm2 ) was irradiated with 250 MeV Xe17þ ions to a fluence of 7 · 1011 cm2 . The Xe ion fluence used is far below the value where tracks are formed in InP single crystals. As can be seen in Fig. 11, tracks are observed around z  1 lm which corresponds to the maximum of the nuclear energy deposition for 1 MeV Si ions [7]. This confirms the existence of predamage as a condition for track formation by Xe

irradiation in this range of the electronic energy deposition. It should be mentioned that preliminary experiments at the irradiation temperature TI ¼ 77 K resulted in a lower damage concentration than that produced at the same ion fluence at room temperature. This indicates the important role of the irradiation temperature for the damage formation by high electronic excitation. Recently Szenes et al. [10] reported on track formation in InP, InAs, InSb and GaSb due to Pb irradiation with energies in the range between 1.85 and 10 MeV/u corresponding to a variation of the electronic energy deposition ee from 27.2 to 30.4

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 Fig. 10. Calculated number (TRIM 87) of primary displacements, Ndispl , and electronic energy deposition, ee , for 250 MeV Xe irradiation in InP (a) and schematic illustration of the damage structure over the whole irradiated depth as deduced from TEM investigations (from [6]).

keV/nm. Whereas in InP and GaSb the tracks were identified by high-resolution TEM (HRTEM) as to be amorphous, the tracks found in InSb were completely recrystallised. Fig. 12 shows an XTEM image of InP irradiated with 385 MeV Pb ions at room temperature [10]. The number of tracks visible in the figure is in the order of 1010 cm2 which is in reasonable agreement with the ion fluence of 1 · 1010 cm2 in this case [38]. From this the authors conclude that each Pb ion produces a visible track in InP, which is in certain contradiction to the results obtained for Xe irradiation with ee values around 20 keV/nm [6,7]. For InSb and GaSb the same conclusion is drawn. Track diameters were determined from the channelling RBS spectra by using the overlap damage model introduced by Gibbons [39]. For InP the authors

give a mean diameter of 7 nm which is in good agreement with the value obtained from the TEM images and in the same order of magnitude as that determined in the case of 250 MeV Xe17þ [6]. For InSb in [10] mean track diameters between around 10 and 19 nm are given depending on the energy per nucleon, and in GaSb and InAs the diameters amount to 3.6 and 4.4 nm, respectively. The question is which model can describe the experimental findings concerning damage formation in semiconductors due to high electronic excitation. An often used description of track formation in solids is based on the thermal spike model (see e.g. [2] and references therein) in which it is assumed that the material melts along the ion trajectory up to depths for which the electronic energy loss exceeds a certain critical value. This

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Fig. 11. Bright-field XTEM image of an InP sample pre-damaged by 1 MeV Si ions (fluence NI ¼ 2:5  1013 cm2 , implanted at )190 C) and post-irradiated with 250 MeV Xe ions to NI ¼ 7  1011 cm2 obtained at a depth z around 1 lm (corresponding to the maximum energy loss of 1 MeV Si) (from [7]).

Fig. 12. Plane-view TEM micrograph of InP irradiated with 385 MeV/u Pb ions. The sample was tilted by 45 with respect to the ion beam (from [10]).

melt phase is followed by fast cooling and resolidification so that an amorphous track is formed within a crystalline matrix [2,40]. The critical energy loss depends on the target material and reflects the efficiency of the mechanism, which

converts electronic excitation energy into atomic motion. Such a model may hold well also for semiconductors in cases where, at sufficiently high electronic energy depositions, each ion is assumed to produce one track, as it was reported in [10] for Pb irradiation of InP, InSb and GaSb. However, when applying this model to the results obtained for Xe irradiation of InP, we have to assume that during the quenching of the thermal spike epitaxial recrystallisation must occur which is obviously not perfect, but leaves behind many defects and defect clusters. For a second ion impinging nearby, the recrystallisation of the molten core may be hindered because of a less perfect surrounding. Then, with ongoing irradiation the recrystallisation speed may become smaller than the cooling rate, and a rather continuous track is frozen in [6]. Up to now, most of the experiments with respect to track formation by swift heavy ions were done on InP and, therefore, an attempt is made to represent the results in the framework of a thermal spike model. Fig. 13 shows the radius R of the damaged regions obtained from RBS versus the electronic energy deposition, ee (a), and versus the ion velocity v (b) for SHI irradiations of InP (data are taken from [10,41,42]). From the figure it can be seen that in both parts the experimental data do not fall on one line. The experimental uncertainty of R is about 10% and cannot account for the observed deviations. The scattering of the data points reflects the complexity of the processes occurring during swift heavy ion irradiation of InP. The decrease of R for Pb ions at higher beam energies and correspondingly higher ion velocities (Fig. 13(b)) suggests a velocity effect as it was observed during swift heavy ion irradiation of insulators, in which also a decreasing damage efficiency was found at very high ion velocities (see e.g. [43]). Different values of R for Pb and Au ions at ee  28 keV/nm and for Au and Xe ions at ee  20 keV/nm (Fig. 13(a)) indicate an influence of the other parameters, as e.g. the ion velocity, the ion charge state, etc. Furthermore, one has to be careful with the quantity R resulting from RBS. It follows from the damage cross-section which is obtained by fitting the experimental data nda versus the ion fluence with Gibbons’ model [39] assuming each ion to produce a measurable defect

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Xe from those for Au is caused by both an influence of other effects and by an underestimation of R as discussed above. Now the values of R versus ee (Fig. 13(a)) are compared with the thermal spike model of Szenes [44,45] represented by the implicit equation   gee  LqpR2 R2 exp  2 Tmelt  Ttarget ¼ : qcpa2 ð0Þ a ð0Þ

Fig. 13. Track radii determined from channelling RBS spectra as a function of the electronic energy deposition, ee (a), and as a function of the velocity v (b) for various ions and energies in InP.

concentration. More specifically, this damage cross-section is the product of the relative defect concentration and the area cross-section produced by one ion. In the case of Pb irradiation, where each ion produces an amorphous track (and we assume this to be true also for Au), the relative defect concentration produced by one ion is one and, therefore, the values of R are realistic (which is proved by the comparison with TEM results [10]). The situation is much more complicated for Xe irradiation, where – at least at the beginning of the irradiation – not each ion creates an amorphous track but only point defects and point defect complexes. Consequently, the mean relative defect concentration produced per ion is lower than one, thus yielding too low values of R. Therefore we believe that the strong deviation of the results for

The curve in Fig. 13(a) is calculated using the material parameters of InP: melting temperature Tmelt ¼ 1335 K, heat of fusion L ¼ 425 027 J/kg, specific heat c ¼ 363:83 J/(kg K), density q ¼ 4:81  103 kg/m3 and the target temperature during the irradiation Ttarget ¼ 295 K. Open parameters are the fraction g of ee contributing to the lattice heating and the standard deviation að0Þ of the radial distribution of the lattice temperature at the moment of melting. From Fig. 13(a) it is obvious that the general tendency of the experimental values Rðee Þ is well represented by the model curve calculated with g ¼ 0:06 and að0Þ ¼ 4:6 nm, which are reasonable values. For a more serious analysis more experimental data for each ion species are needed. It is expected that for each ion species different values of g and að0Þ will be obtained, which might be connected with differences in the radial distribution of the electronic energy. However, to what extent a thermal spike model describes the processes correctly cannot be answered, because until now there is no direct experimental evidence that melting really occurs in semiconductors under swift heavy ion irradiation. In the case of Xe irradiation in InP where a certain pre-damage is obviously necessary to form visible amorphous tracks, a qualitative discussion of the experimental findings is also possible without assuming a melting process. Part of the energy of the ions may be directly transferred to atomic motion (without melting), leading to the formation of point defects around the ion trajectory. If a certain threshold concentration of point defects and point defect complexes is reached, with ongoing irradiation a transformation to the amorphous phase occurs similar to the amorphisation via homogeneous defect nucleation observed in semiconductors during ion implantation

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Fig. 14. Bright-field plan-view TEM image of a Si0:5 Ge0:5 layer irradiated with U ions. The arrows indicate the tracks (from [11]).

with ions of conventional energies (several hundred keV to several MeV). However, this model cannot unequivocally describe cases where each ion forms an amorphous track immediately. The first proof of discontinuous track formation in group IV semiconductors was recently published by Gaiduk et al. [11]. Strain-relaxed epitaxial p-type Si0:5 Ge0:5 layers were grown by molecular-beam epitaxy (MBE) on (0 0 1) silicon wafers by means of the stepwise compositionally grading technique and irradiated with 1.3 GeV U ions to a fluence of 1010 cm2 at a constant flux around 2 · 10 cm2 s1 . The energy depositions calculated by TRIM are ee ¼ 33:8 keV/nm and en ¼ 0:024 keV/nm under these conditions. Fig. 14 shows a typical bright-field plan-view TEM image of the alloy layer. Along the ion trajectories dotlike and elongated dark spots with an average size of 3–10 nm are formed like a string of pearls running parallel to each other. The tracks typically consist of two to five separated dots with varying defect structure and diameters along the ion trajectory. The inner structure consists of a crystalline core surrounded by a damaged periphery. The density of the isolated dots together with the dis-

continuous tracks is in good agreement with the ion fluence. The results strongly indicate that the track formation is a consequence of quenching of a melt phase and subsequent imperfect recrystallisation. Evidence of track formation in Ge single crystals during 710 MeV Bi and 1.3 GeV U irradiation corresponding to energy depositions ee ¼ 37 and 42 keV/nm, respectively, was recently found by Komarov et al. [12]. As in the case of Si0:5 Ge0:5 single dot-like defects as well as discontinuous tracks are registered along the ion trajectories. The density of the discontinuous tracks is about two orders of magnitude lower than the ion fluence which is similar to the behaviour found for Xe irradiation in InP [6,7]. The periodic appearance of dot-like defects within the discontinuous tracks is explained by fluctuations of the ion charge state along the ion trajectories which periodically reduces the electronic energy deposition below the threshold value for track formation. The distance between the defects can be explained by a periodic simultaneous loss of three electrons increasing the energy deposition by about 4–6 keV/nm. From this it is concluded that the formation of continuous tracks requires an electronic energy deposition of not less than 46–49 keV/nm. The discontinuity of the tracks and the fact that not each impinging ion forms a track indicate that the ee values reached under the irradiation conditions in this case are only slightly above the threshold value for the formation of visible damage. The scenario of track formation in the group IV semiconductors Si0:5 Ge0:5 and Ge therefore corresponds to that assumed for Xe irradiation in InP. 3.2.2. Cluster ion irradiation Another opportunity to obtain a high ratio of ee =en is the use of cluster ions with energies of

Table 2 Parameters of cluster ion irradiation for various semiconductors [13–16] Substrate

Particle

Energy (MeV)

ee (keV/nm)

en (keV/nm)

Reference

Si Ge GaAs

C60 C60 C60

30–40 20–40 20–40

47.9–57.2 37.3–50.9 38.4–52.3

0.88–0.71 1.65–1.04 1.66–1.05

[13,14] [15] [16]

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several 10 MeV. Consequently, in the last few years French groups investigated the damage formation due to high electronic energy deposition in Si, Ge and GaAs using the fullerene C60 with energies between 20 and 40 MeV [13–16] (see Table 2). Dunlop et al. [13] showed that 30 MeV C60 irradiation (ee ¼ 46 keV/nm, en ¼ 0:88 keV/nm) to ion fluences of several 109 cm2 causes continuous amorphous tracks in Si single crystals. The track density almost agrees with the ion fluence. The initial diameter of the tracks at the ion entrance into the target amounts to about 10 nm. Down to  80 nm this diameter remains almost constant, and at larger depths it gradually decreases as a consequence of an increasing spatial separation of the cluster constituents due to Coulomb repulsion between the charged fragments and multiple scattering in elastic collision processes. Very often at the end of the tracks a series of aligned droplets of damaged material was registered. Canut et al. [14] confirmed these findings for 30 and 40 MeV C60 irradiation and registered track diameters at the sample surface of 8.4 and 10.5 nm, respectively, for the two energies. From the dependence of the damage cross-sections on the electronic energy deposition a threshold value for track formation in Si of eet  30 keV/nm is estimated, which is by a factor of ten larger than that found in LiNbO3 under identical irradiation conditions [46,47]. Amorphous tracks were also found in C60 -irradiated Ge and GaAs by Colder et al. [15,16]. The diameter of the tracks increases with the energy of the clusters, i.e. with the electronic energy loss. This is illustrated in the high resolution TEM images in Fig. 15 for the example of GaAs irradiated with various cluster ion energies [16]. From the damage cross-sections determined from the TEM investigations threshold values of the electronic energy deposition eet ¼ 36 keV/nm for GaAs, eet ¼ 33 keV/ nm for Ge and eet ¼ 28 keV/nm for Si were determined. Contrary, irradiation of Ge with individual swift Bi and U ions, yielding significantly higher ee values (ee P 37 keV/nm), produces only discontinuous tracks [12]. This again points to more complicated effects influencing the behaviour of semiconductors under high electronic energy deposition, especially the importance of the radial distribution of the electronic energy.

123

Fig. 15. High-resolution TEM image of GaAs single crystals irradiated with (a) 20 MeV, (b) 30 MeV and (c) 40 MeV C60 ions (from [16]).

4. Damage transformation and annealing Previously some evidence of defect annealing due to high electronic energy deposition was found in Si, Ge and GaAs (see Section 3.1) [18,20,22–24]. Huber et al. [34] reported on recrystallisation of thin amorphised Ge layers after 252 MeV Au irradiation (in this case the value

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ee  28 keV/nm exceeds the threshold value for damage annealing due to electronic energy deposition of 10 keV/nm [18]). Furthermore, a partial crystallisation in 5 nm thick evaporated Si and Ge layers due to swift heavy ion irradiation was reported by Furuno et al. [4]. In the Ge layers, tracks were observed for electronic energy depositions ee P 5:3 keV/nm, the track diameters varied between 8 and 23 nm depending on ee . In the Si layers tracks were found for ee P 17 keV/nm, the diameter amounts to 7 nm. The tracks observed consist of small crystallites or recrystallised regions, and their formation is explained in the framework of the thermal spike model. Contrary, in bulk amorphous Si according to the theory of Trinkaus and Ryazanov [48], the formation of crystalline tracks is hardly to be expected, but a plastic deformation should occur and was very recently also detected [49]. The occurrence of tracks in the thin layers mentioned above may only be explained assuming that the layers were not completely amorphous but fine polycrystalline with grain sizes of typically 1 nm as mentioned in [50]. That melting in Si and Ge occurs at those small values of ee is in contradiction with the results obtained in single crystals. A probable explanation might be a restriction of the dissipation of the deposited energy due to the close vicinity of grain boundaries and surfaces [13]. A similar behaviour was found more than 40 years ago in thin or discontinuous metal films [51,52]. In these films, MeV ion irradiation could form tracks, whereas no damage was produced in bulk samples with the same ions. A more detailed study concerning ionisation stimulated damage transformation and annealing processes was performed by irradiating pre-damaged GaAs and InP crystals [41,53]. For this purpose GaAs and InP single crystals were implanted with 600 keV Ge and 600 keV Se ions, respectively, at 77 K. In order to achieve different damage levels between crystalline virgin and amorphised material, the ion fluence was varied between 2 · 1012 and 2.5 · 1013 cm2 . Pre-damaged as well as virgin samples were then irradiated with 390 MeV Xe21þ ions at room temperature using ion fluences ranging from 3 · 1011 to 1 · 1014 cm2 at a constant ion flux of about 1010 cm2 s1 . As an example,

Fig. 16. Relative concentration of displaced lattice atoms, nda , versus the depth z for Ge pre-damaged and Xe post-irradiated GaAs (from [53]).

Fig. 16 shows the relative concentration of displaced lattice atoms, nda , as a function of depth z for pre-damaged GaAs irradiated with various Xe ion fluences [53]. The Ge implantation causes a buried heavily damaged layer extending over a depth from 0.1 to 0.18 lm. Post-irradiation with Xe leads to an ion fluence dependent symmetric reduction of nda ðzÞ over the whole damage distribution. It should be remembered that Xe irradiation of single crystalline GaAs results in some damaging too (see Section 3.1, Fig. 4). But this damage is very low (nda  0:005), constant over the depth and almost independent of the Xe fluence up to NI ¼ 1014 cm2 [41]. The symmetric reduction of nda ðzÞ in pre-damaged GaAs therefore results mainly from in situ damage annealing due to the constant high electronic energy deposition in the region of the damaged layer. The observed damage annealing depends on the defect concentration in the pre-damaged layers: the smaller the initial damage concentration is, the higher is the relative annealed fraction (RAF). This is illustrated in Fig. 17 which shows the remax lative annealed fraction RAF ¼ ðninit:  nmax da da Þ= init: max init: max init: max nda as a function of nda ðnda and nmax da are the concentrations of the displaced lattice atoms in the maximum of the distributions for the Ge pre-damaged and Xe post-irradiated samples, respectively). For a pre-damage concentration max ninit: 6 0:75 the annealed fraction lies between da 80% and 90% corresponding to a remaining

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125

Fig. 17. Relative annealed fraction of damage in GaAs as a function of the initial damage concentration for Xe post-irradiation with different fluences (from [53]).

damage concentration of nda  0:05 at the highest Xe fluence. A considerable decrease of the RAF with increasing initial damage concentration is max observed if ninit: P 0:75. In this case of thick da heavily damaged layers only between 6% and 18% of the initial damage are annealed. This behaviour indicates that the annealing efficiency is much higher for low pre-damage levels than for those which are close to complete amorphisation. Further, it is also possible that in the case of high predamage levels the Xe irradiation increases the defect concentration, which would also cause a reduction of the relative annealed fraction, RAF. Consequently, the damage concentrations measured after the Xe irradiation are the result of a complex interplay between generation and annealing of defects by which annealing obviously dominates generation. Fig. 18 shows XTEM images of a Se-implanted GaAs sample with a pre-damage concentration of max ninit: ¼ 0:97 before (Fig. 18(a)) and after (Fig. da 18(b)) Xe irradiation. In the as-implanted sample (Fig. 18(a)) the depth interval from  50 to  220 nm mainly consists of amorphous material (dominating grey ‘‘background’’) with some crystalline inclusions (dark contrast spots correspond to crystalline material misaligned with the original crystalline matrix). The Xe irradiation (Fig. 18(b)) decreases the amount of damage considerably, the diffraction pattern shows no indication of amor-

Fig. 18. XTEM images of a GaAs sample first implanted with 2 · 1013 Se cm2 at 77 K and then irradiated with 1 · 1013 Xe cm2 at 300 K: (a) as-implanted, (b) after Xe irradiation (from [53]).

phous material which is obviously almost recrystallised. The behaviour of pre-damaged InP layers with respect to high electronic energy depositions differs remarkably from that observed in GaAs. 390 MeV Xe21þ irradiation (ee ¼ 19:7 keV/nm) causes only a slight reduction of the damage concentration at the lowest Xe fluences, with increasing ion fluence additional damage is produced in agreement with the damaging behaviour observed in single crystalline InP. But if the irradiations are performed with an electronic energy deposition below the threshold for damage formation, e.g. with 140 MeV Kr10þ corresponding to ee ¼ 12:5 keV/nm, significant damage annealing occurs similar to that observed in GaAs (for details see [41]).

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Our results show that in pre-damaged GaAs and InP damage annealing due to electronic energy deposition occurs for irradiation conditions for which in the corresponding crystalline material no defect formation due to the electronic energy loss is observed, i.e. for irradiation with ee below the threshold value for damage formation in crystalline samples. The damage reduction is caused by the growth of already existing crystalline inclusions as well as by a local crystallisation of amorphous regions. If the irradiation of predamaged InP is performed under conditions with ee above the threshold value for damage formation in crystalline samples, then damage formation dominates and slight annealing effects are observed only for very small ion fluences.

5. Summary and discussion The results available on swift heavy ion irradiation in semiconductors show that the situation concerning the effect of high local electronic energy deposition is extremely complex. As in other materials, high electronic energy loss may cause damage including the formation of amorphous tracks in crystalline semiconductors and complete amorphisation. Furthermore, under certain conditions damage annealing in samples containing defects or recrystallisation of thin polycrystalline layers can occur. One important parameter influencing the response of the material is the electronic energy deposition, ee . In Si, Ge and GaAs swift heavy ion irradiation with electronic energy depositions ee 6 28 keV/nm is connected with the formation of electrically compensating point defects identical to those formed by electron, neutron and proton irradiation. However, there is sufficient evidence that these defects can be ascribed to the effect of the small contribution of nuclear energy loss in the corresponding depth regions. Until now neither in Si nor in GaAs single crystals amorphous tracks were found after swift heavy ion irradiation with ee below  30 keV/nm. In Ge single crystals discontinuous tracks were registered after single-ion irradiation with an electronic energy deposition above 37 keV/nm, which

consist of pearl like distributions of amorphous dots along the ion trajectories. The discontinuity of the tracks is explained by charge state fluctuations which may periodically reduce ee below the threshold energy deposition for track formation. The absence of tracks in Si, for example, is ascribed by Chadderton et al. [54,55] to an in situ epitaxial recovery process due to the so called ‘‘projectile-assisted prompt anneal’’ (PAPA) within 1011 s. This process may be assisted by defects being specific for the target, as in Si the three-dimensional divacancy [56]. Here the question is whether the concentration of intrinsic defects in the crystals is sufficient to stimulate such a process. The fact that in Si irradiated with C60 molecular cluster projectiles tracks are registered may be explained within this approach by the high electronic energy deposition ( 48 keV/nm 6 ee 6 57 keV/nm) leading to an incomplete epitaxial recovery [57]. At electronic energy deposition above  19 keV/nm tracks were observed in InP, InAs, InSb and GaSb after irradiation with swift Xe, Au and Pb ions. Whereas the tracks in InSb are almost recrystallised, they are amorphous in InP and GaSb. For InP a threshold energy deposition eet  13 keV/nm was determined. Furthermore, at energy depositions around 20 keV/nm the track evolution depends on the ion fluence, i.e. not each impinging ion produces a visible track and predamage is obviously necessary to form tracks. Only for larger ion masses and higher ee values of about 28 keV/nm tracks seem to be directly produced in InP by each ion. In pre-damaged GaAs and InP damage annealing due to electronic energy deposition occurs for irradiation conditions for which in the corresponding crystalline material no defect formation due to the electronic energy loss is observed. If damage formation occurs in crystalline material (as in InP), damage formation dominates the annealing also in predamaged samples. The experimental data indicate that the electronic energy deposition alone cannot explain whether visible tracks in semiconductors are formed or not. An important role seem to play the ion mass, the ion velocity, the charge state of the impinging particles and – in connec-

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tion with these – the radial distribution of the electronic energy around the ions’ path. Furthermore, in semiconductors the influence of the electrical conductivity obviously cannot be neglected, as also our preliminary experiments performed on InP indicate. It turns out that in the semiconductors under swift heavy ion irradiation a strong competition between damage evolution, transformation and annealing occurs. Which of these processes dominates and whether tracks are formed or not depends on a variety of material parameters the influence of which is not yet sufficiently studied. Qualitatively most of the experimental data can be understood in the framework of the thermal spike concept although there is no clear experimental evidence that melting in semiconductors really occurs in all cases investigated. In the framework of this concept, the threshold energy deposition eet for track formation corresponds to the energy deposition causing melting of a certain region around the ion trajectory. The subsequent cooling and fast re-solidification can cause freezing in of damaged or amorphous tracks. In some cases (Xe irradiation of InP) within this concept recrystallisation must be assumed, leading to the formation of point defects. With further irradiation at a certain ion fluence, respectively a certain pre-damage level, a transformation to the amorphous phase occurs. However, the explanation of this behaviour not necessarily needs the assumption of melting. There are not yet sufficient experimental data available at the moment for a more extended theoretical description and further systematic studies on swift heavy ion irradiation in semiconductors are necessary. Acknowledgements The authors wish to thank K. G€ artner and S. Klaum€ unzer (HMI Berlin) for helpful discussions. The support of S. Klaumunzer during our experiments at the Hahn Meitner Institute is highly acknowledged. Our work was supported by the Bundesministerium f€ ur Bildung und Forschung, contract number 05KK1SJA/2.

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