Low-temperature damage formation in ion implanted InP

Nuclear Instruments and Methods in Physics Research B 307 (2013) 377–380

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Low-temperature damage formation in ion implanted InP E. Wendler a,⇑, A. Stonert b, A. Turos b,c, W. Wesch a a

Friedrich-Schiller-Universität Jena, Institut für Festkörperphysik, Max-Wien-Platz 1, 07743 Jena, Germany National Center of Nuclear Research, 05-400 Swierk/Otwock, Poland c Institute of Electronic Materials Technology, Wolczynska 133, 01-919 Warsaw, Poland b

a r t i c l e

i n f o

Article history: Received 18 September 2012 Received in revised form 14 December 2012 Accepted 16 December 2012 Available online 20 January 2013 Keywords: Ion implantation Damage formation III–V Semiconductors InP

a b s t r a c t Damage formation in ion implanted InP is studied by quasi–in situ Rutherford backscattering spectrometry (RBS) in channelling configuration. Subsequent implantation steps are performed at 15 K each followed by immediate RBS analysis without changing the environment or the temperature of the sample. 30 keV He, 150 keV N and 350 keV Ca ions were applied. The depth distribution of damage is in good agreement with that calculated with the SRIM code. The evolution of damage at the maximum of the distribution as a function of the ion fluence is described assuming damage formation within single ion impacts and stimulated growth of damage when the collision cascades start to overlap with cross sections rd and rg, respectively. These cross sections are found to depend on the primary energies deposited in the displacement of lattice atoms and in electronic interactions calculated with the SRIM code. The obtained empirical formulas are capable to represent the experimental results for different III–V compounds implanted at 15 K with various ion species. Ó 2013 Elsevier B.V. All rights reserved.

1. Introduction Indium phosphide (InP) is a semiconductor with superior carrier mobility and a direct band gap. It has been used for the fabrication of optoelectronic devices such as laser diodes, photodetectors and solar cells. And there is a continuous research to enhance the application of ion beams in InP device technology. This includes, for example, the formation of ternary compounds such as In1xGaxP by high-fluence Ga ion implantation (see e.g. [1,2]), the development of ion-cutting of InP-based field effect transistors via masked hydrogen implantation [3], the study of Mn implants in InP (see e.g. [4,5]) or the formation of nanodots under focused ion beam irradiation [6]. Any application of ion beams requires a sufficient understanding of the processes of ion–solid interaction and the formation of radiation damage. It was shown that damage formation in ion implanted InP at moderate and slightly elevated temperatures [7,8] is well described using the concept of critical temperatures (see e.g. [9,10]). In InP the critical temperatures (Tc) are between 340 and 420 K (see e.g. [11]). This range of temperatures coincides with an annealing stage which is attributed to the mobility of P interstitials [12,13]. However, a concentration-dependent relaxation and annealing of implantation-induced defects is also found in InP and related compounds between 77 K and room temperature [11,14,15] and even during storage at room temperature [16]. Therefore, the study of primary effects of ⇑ Corresponding author. E-mail address: [email protected] (E. Wendler). 0168-583X/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.nimb.2012.12.074

ion–solid interaction in InP requires implantation and subsequent damage analysis to be done at cryogenic temperatures. The aim of the present paper is to investigate the primary effects of damage formation in ion implanted InP. The experiments were carried out at 15 K without changing the temperature or the environment of the samples between implantation and measurement. Rutherford backscattering spectrometry (RBS) in channelling configuration is applied to register damage formation. From that almost no information is obtained regarding the microscopic effects of damage formation. Here cross sections of damage formation are extracted from the fluence dependencies basing on a common model of damage build-up. The main aim of this paper is to analyse these cross sections using calculated quantities which represent the energy deposition in electronic processes and in the displacement of lattice atoms. The results obtained are compared with those obtained for other III–V compound semiconductors ion implanted under similar conditions. 2. Experimental conditions h1 0 0i oriented single crystalline InP samples are implanted with 30 keV He, 150 keV N and 350 keV Ca ions. Ion energies are chosen to obtain comparable thicknesses of the implanted layers. Depending on the ion species used, the ion fluence NI varies between 1012 and 1015 cm2 (for details of the implantation conditions see Table 1). RBS with 1.4 MeV He ions and a scattering angle of 170° is used to collect RBS spectra in aligned (Yal) and random (Yra) direction. All experiments are performed in a special

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Table 1 Implantation conditions (ion fluence NI), depth of maximum damage (zmax), number of displacements per ion and unit depth (N displ ) at depth z = zmax, electronic energy loss per ion and unit depth (Sel) at depth z = zmax and cross sections rd and rg used to calculate the curves in Figs. 2 and 3 with Eq. (2). The errors given for rd and rg take into account the statistical error of the measurement and a 5% uncertainty of the ion fluence. The last column gives the cross section rd calculated with Eq. (3).

4  1013–7  1014 5  1012–1.5  1014 5  1011–4  1013 1.5  1012–3  1013

zmax (lm) 0.16 0.24 0.21 0.18

Sel at zmax (108eV/ion/cm)

rd (1014 cm2)

rg (1014 cm2)

(108/ion/cm)

rd (1014 cm2) calculated

0.112 0.803 3.390 3.760

9.46 30.1 59.6 61.6

0.055 ± 0.010 1.22 ± 0.25 6.07 ± 0.92 6.42 ± 0.80

0.42 ± 0.06 3.06 ± 0.96 12.1 ± 3.9 15.6 ± 3.1

0.051 0.557 3.76 4.38

N displ at zmax

two-beam target chamber at (15 ± 2) K, which allows to do subsequent ion implantations each followed immediately by the RBS measurement [17]. The minimum yield vmin ¼ Y al =Y ra versus depth z is used to calculate the depth distribution of the relative concentration of displaced lattice atoms nda(z) using the code DICADA [18]. For the calculation a random displacement of lattice atoms within the lattice cell is assumed. This corresponds to point defects, point defect complexes or amorphous zones which are to be expected in InP at such a low temperature. nda(z) is referred to as relative damage concentration or damage profile in the following text. For comparison of results for different ion species, the number of displacements per implanted ion and unit depth, N displ versus depth z is calculated with the code SRIM (version 2008_04) [19]. The density of InP was taken to be 4.787 g/cm3 corresponding to an atomic density of N0 = 3.9541022 at./cm3. The displacement energies of 6.6 eV for In and 8.8 eV for P are taken from Ref. [20]. Resulting values of N displ and the energy deposited in electronic interactions per ion and unit depth, Sel, are included in Table 1. The ion fluences NI are converted to the number of displacements per lattice atom ndpa by ndpa ¼ N I  N displ =N 0 with N displ taken in the maximum of the distribution (see Table 1). 3. Results and discussion

2500

(a)

+

backscattering yield

30 keV He 150 keV N 350 keV Ar 350 keV Ca

NI (cm2)

InP : N (150 keV)

2000

ion fluence 14 -2 in 10 cm

1500 1000 500

1.50 1.00 0.70 0.40 0.20 0.10 0.05

0 200

300

400

500

channel number

defect concentration nda

Ion, energy

(b)

1.0 0.8

*

Ndispl

0.6

in a.u.

0.4 0.2 0.0

dnmax max max da ¼ rd ð1  nmax da Þ þ rg nda ð1  nda Þ: dNI

ð1Þ

0.0

0.1

0.2

0.3

0.4

0.5

0.6

depth z (μm)

max

Fig. 1. Energy spectra of 1.4 MeV He ions backscattered on InP implanted with 150 keV N ions (a). The random spectrum (uppermost in the figure) and the aligned spectrum of the unimplanted InP (lowest in the figure) are given as solid lines. The calculated damage profiles are shown in part (b). The number of displacements per implanted ion and unit depth, N displ , as calculated with SRIM is included in arbitrary units for comparison.

rel. damage concentr. nda

Fig. 1a shows typical energy spectra of 1.4 MeV He ions backscattered on InP after implantation of various ion fluences at 15 K, here for implantation of 150 keV N ions. The evolution of a damage peak is clearly seen in the In part of the spectra (below channel number 412). The peak finally reaches the random level at about 1  1014 cm2, which is commonly taken as an indication for amorphisation. The damage profiles nda(z) resulting from the In part of the spectra in Fig. 1a are plotted in Fig. 1b. For comparison the distribution of primary displacements is included in arbitrary units. Despite a small deviation between the calculated and measured profiles at the trailing edge, a good agreement is observed with respect to the depth of maximum damage and the shape of the profile. This is also observed for the implantation of 30 keV He and 350 keV Ca ions (not shown). As can be seen from Fig. 1b, the position of maximum damage, zmax, does not change with ion fluence and the corresponding values for the implants considered here are given in Table 1. Fig. 2 summarises the damage evolution at the maximum of the distribution as a function of the ion fluence, nmax da ðN I Þ. For all three ion species studied a continuous transition towards amorphisation is observed. This can be well represented assuming two mechanisms contributing to damage formation: (i) damage formation within direct ion impacts with the cross section rd and (ii) stimulated growth of already existing damage with the cross section rg. Taking into account that the probabilities for the two processes to max max occur are given by ð1  nmax da Þ and nda ð1  nda Þ, respectively, one obtains [10,22]

1

InP

0.1 He N Ca

0.01 10

12

10

13

ion fluence NI

T = 15 K 14

10

10

15

(cm-2)

Fig. 2. Relative damage concentration at the maximum of the distribution, nmax da , versus the ion fluence NI for 30 keV He, 150 keV N and 350 keV Ca ion implantation into InP at 15 K. The lines are fitted to the experimental data using Eq. (2) and the resulting parameters are given in Table 1.

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An analytical solution of Eq. (1) is provided in Ref. [10] with

nmax da ðN I Þ

¼ 1  ðrd þ rg Þ=frg þ rd exp½ðrd þ rg ÞNI g:

ð2Þ

The curves shown in Fig. 2 are fitted to the experimental data by adjusting the parameters rd and rg. rd follows from the almost linear increase of the relative damage concentration nmax da with ion fluence at very low concentrations, whereas rg accelerates the transition towards the saturation value nmax da ¼ 1. The resulting parameters are included in Table 1. From Fig. 2 it can be seen that the curves yield a good representation of the experimental data. This indicates that the two mechanisms mentioned above give a reasonable description of what is observed. To compare the results for different ion species, in Fig. 3 the maximum relative damage concentration nmax is plotted against da the number of displacements per lattice atom ndpa. In this figure data for 350 keV Ar ion implantation are included which are taken from Ref. [21]. The figure shows that almost no difference occurs in nmax da ðndpa Þ for ion species between N and Ca. That is the damage concentration measured after implantation at 15 K is only determined by the energy deposited in the displacement of lattice atoms. About 0.2. . .0.25 dpa are necessary to amorphise InP. In difference to that a higher number of displacements per atom is required for the light element He (see Fig. 3). For comparison, in GaAs values of about 0.3. . .0.4 dpa are needed to achieve amorphisation using heavy elements and more than 0.8 dpa when implanting the light element He [11,23]. As found here for InP, the value for He implantation differs from that for ion species heavier than N. In general, the values for GaAs are clearly higher than those for InP. The differences between damage formation in InP and GaAs ion implanted at room temperature and at 80 K with damage analysis performed at room temperature [24] can be related to different ranges of the critical temperature [11]. However, at the temperature used here and doing quasi–in situ damage analysis, thermal effects can be widely excluded. Therefore, the difference between InP and GaAs observed here must be related to intrinsic properties of the materials being determined by their bond strength. The cross section rd is the area integral over the relative damage concentration produced per ion at the depth zmax. If one assumes a homogeneous damage distribution across the area, rd is given by the product of the area and the relative damage concentration per ion. If one further assumes that each ion produces an amorphous cluster corresponding to a relative damage concentration per ion equal to unity, then rd is the area damaged by one ion

at the corresponding depth. Under these assumptions the diameter of the clusters produced by individual ions is 0.24 nm for He, 1.28 nm for N, 2.74 nm for Ar and 2.88 nm for Ca implantation. The distance between nearest neighbours in InP is dnn = 0.254 nm. For He implantation the diameter of the assumed cluster is of the same size as dnn. This means the clusters would consist of very few atoms only for which an amorphous state cannot be defined. This suggests that at least for He ions the assumption of complete amorphisation within a single ion impact is not correct. For the other ion species the estimated diameters are not in conflict with the assumption that heavily damaged and/or amorphous material is produced within a single ion impact, which grows during further irradiation. High-resolution transmission-electron-microscopy studies on GaAs and GaP ion implanted at 30 K revealed that the visible damage produced within the core of isolated collision cascades created by heavy-ion irradiation (Ar and heavier ions) is amorphous [25]. Given the fact that InP is amorphised at lower ndpa values than GaAs (see above), it can be expected that the same is true in the case of InP ion implanted at 15 K. For the light element He the situation is different. As discussed above the formation of heavily damaged or amorphous clusters is not to be expected. In this case one may suspect that amorphisation does not occur heterogeneously by direct ion impacts but homogeneously on a finer scale by the coalescence of point defects and point defect complexes as suggested for ion implanted SiC [26]. This does not inhibit the use of Eq. (2) because point defects/point defect complexes and the lattice distortions caused by them contribute to the yield of backscattered ions (see e.g. [27]). As already mentioned, for our implantation temperature thermal effects can probably be excluded. Therefore, the damage formation is determined by primary ion-induced effects. This reinforces the attempt to compare the cross sections rd and rg with the energy deposited in the displacement of lattice atoms. The latter can be quantified by the cross section rSRIM ¼ N displ =N 0 . In Fig. 4 the cross sections rd and rg are depicted versus the calculated cross section rSRIM. The increase of rd with rSRIM is between a linear and a quadratic dependence. This is in agreement with findings for ion implanted GaAs [21]. The stronger-than-linear increase of the cross section of damage formation rd with rSRIM suggests non-linear processes to occur within the primary collision cascades. In contrast to that, a linear dependence on rSRIM is found for the cross section of stimulated growth rg. This means that the stimulated growth of damage is proportional to the density of the primary collision cascades. Such a linear dependence is also observed for other III–V compounds and the line in Fig. 4b repre-

1.0

0.6 0.4

He N

0.2

Ar Ca

T = 15 K 0.0 0.00

0.05

0.10

0.15

0.20

0.25

0.30

displacements per atom ndpa Fig. 3. Relative damage concentration at the maximum of the distribution, nmax da , versus the number of displacements per lattice atom, ndpa, for 30 keV He, 150 keV N, 350 keV Ca and 350 keV Ar ion implantation into InP at 15 K. The data for Ar ion implantation are taken from Ref. [21]. The lines are the same as in Fig. 2 with the ion fluences NI converted to ndpa as described in the text. Notice the linear scale in difference to Fig. 2.

-13

cross section σd (cm2)

10

cross section σd (cm2)

rel. damage concentr. nda

max

InP 0.8

InP 10

-14

10

-15

M1 N1 T = 15 K 10

-15

10

(a) -14

-13

10

cross section σSRIM (cm2)

10

-13

10

-14

He N Ar Ca

(b) 10

-15

10

-14

cross section σSRIM (cm2)

Fig. 4. Cross sections of damage formation rd (a) and of stimulated growth of damage rg (b) versus the calculated cross section rSRIM for InP ion implanted at 15 K. In part (a) the solid and dashed lines indicate a linear and quadratic dependence of rd on rSRIM, respectively. The line in part (b) is calculated with rg = 15.6 rSRIM which was obtained as a mean value for various III–V compounds exhibiting a stimulated growth of damage [21].

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E. Wendler et al. / Nuclear Instruments and Methods in Physics Research B 307 (2013) 377–380

sents a mean dependence for various ion species implanted into different III–V compounds at 15 K [21]. In a simple picture a collision cascade can be thought to consist of regions which are heavily damaged contributing significantly to the backscattering of analysing He ions, and of those which are only weakly damaged containing point defects or small clusters of them which contribute only weakly to the backscattering of He ions. Within such a picture, the stimulated growth of damage can be interpreted in two main ways. (i) The weakly damaged regions need to overlap for the point defects to agglomerate to larger complexes which are then more visible to the ion beam. Such a process may be driven by the strain introduced by the point defects. A possible mechanism of amorphisation is the transformation of material into the amorphous state when the strain introduced by point defects exceeds some critical value [28–30]. (ii) When the point defects are close to a heavily damaged cluster produced by previous ions, the latter cluster can grow by a process which is known as ion beam induced interfacial amorphisation [31]. Both scenarios rely on the availability of point defects. Therefore, our results imply that within single collision cascades point defects are produced in such an amount that the growth of damage (as detectable by RBS) increases linearly with the cross section rSRIM. As already mentioned, at 15 K this effect is independent of the chosen III–V compound and the ion species implanted [21]. This finding suggests some kind of universal dependence. In order to represent rd in terms of the energy deposited into the primary processes, an empirical formula was introduced in Ref. [21]. To account for the non-linear dependence of rd(rSRIM) (see Fig. 4a) the ansatz rd = rSRIM fenh/fred is used with fenh being a function of enhancement and fred a function of reduction of damage formation within the primary collision cascades of individual ions impinging on crystalline material. In analogy to the macroscopic dependence rg / rSRIM the function of enhancement is taken to be fenh / rSRIM. For the function of reduction of damage formation fred a useful dependence was found to be fred / rSRIM0.2 Sel (Sel energy deposited in electronic interactions per ion and unit depth). This considers that both nuclear and electronic energy deposition contribute to in-cascade annealing (for details see [21]). Finally these assumptions yield

rd ¼ n

r1:8 SRIM Sel

:

ð3Þ

The parameter n is obtained to be n = (4.68 ± 0.4)  1021 eVcm13/5 as an optimum value for various ion species implanted into different III–V compounds [21]. In Table 1 the calculated values of the cross section of damage formation per individual ion are given. In the case of He implantation a good agreement between the experimental and calculated value of the cross section is obtained. For the other ion species the calculated value underestimates the experimental one by (32. . .54)%. These deviations cannot solely be explained by the statistics of the measurement and assuming an uncertainty of the ion fluence of 5% (see Table 1). Therefore, further investigations are necessary to collect data with higher precision and to improve the empirical formula. 4. Summary Implantation of He, N, Ar and Ca ions into InP at 15 K is studied applying quasi–in situ RBS to measure the relative damage concen-

tration versus depth. The fluence dependence of damage concentration at the maximum of the distribution is analysed assuming the occurrence of damage formation by direct ion impacts and stimulated growth of damage. The corresponding cross sections rd and rg are shown to depend on the energy deposited in the displacement of lattice atoms represented by the cross section rSRIM and the electronic energy loss per ion and unit depth Se. This is meaningful because at 15 K thermal effects are widely excluded and primary ion-induced effects can be studied. The cross section rg of stimulated growth of damage increases linearly with rSRIM and a uniform dependence is observed for different III–V compounds ion implanted with various ion species. This means that in the case where stimulated growth of damage occurs, the related cross section is proportional to the density of primary collision cascades. The cross section rd of direct-impact damaging reveals a dependence on rSRIM which is between a linear and a quadratic one indicating the occurrence of non-linear processes. An empirical formula was deduced which is capable to represent the rd data obtained for different III–V compounds ion implanted at 15 K with various ion species as a function the calculated quantities rSRIM and Se. This allows the prediction of damage formation in this group of materials during low-temperature ion implantation. Further this result can be taken as input for the modelling of damage formation as a function of temperature and ion flux. References [1] Hsin-Chiao Fang, Chuan-Pu Liu, Sandip Dhara, Nucl. Instr. Meth. B 269 (2011) 324. [2] Hubert Gnasen, Surf. Interface Anal. 43 (2010) 28. [3] Wayne. Chen, T.F. Kuech, S.S. Lau, J. Electrochem. Soc. 158 (2011) H727. [4] I.G. Bucsa, R.W. Cochrane, S. Roorda, J. Appl. Phys. 107 (2010) 073912. [5] K. Bharuth-Ram, W.B. Dlamini, H. Masenda, D. Naidoo, H.P. Gunnlaugsson, G. Weyer, R. Mantovan, T.E. Mølholt, R. Sielemann, S. Ólafsson, G. Langouche, K. Johnston, The ISOLDE Collaboration, Nucl. Instr. Meth. B 272 (2012) 414. [6] K.A. Grossklaus, J.M. Millunchick, J. Appl. Phys. 109 (2011) 014319. [7] E. Wendler, T. Opfermann, P.I. Gaiduk, J. Appl. Phys. 82 (1997) 5965. [8] U.G. Akano, I.V. Mitchell, F.R. Shepherd, Can. J. Phys. 70 (1992) 789. [9] F.F. Morehead, B.L. Crowder, Radiat. Eff. 6 (1970) 27. [10] W.J. Weber, Nucl. Instr. Meth. B 166–167 (2000) 98. [11] E. Wendler, B. Breeger, C. Schubert, W. Wesch, Nucl. Instr.Meth. B 147 (1999) 155. [12] K. Karsten, P. Ehrhardt, Phys. Rev. B 51 (1995) 10508. [13] H. Hausmann, P. Ehrhart, Phys. Rev. B 51 (1995) 17542. [14] A. Turos, A. Stonert, L. Nowicki, R. Ratajczak, E. Wendler, W. Wesch, Nucl. Instr. Meth. B 240 (2005) 105. [15] R. Ratajczak, A. Turos, A. Stonert, L. Nowicki, W. Strupinski, Acta Phys. Pol. A 120 (2010) 136. [16] U.G. Akano, I.V. Mitchell, F.R. Shepheard, Appl. Phys. Lett. 59 (1991) 2570. [17] B. Breeger, E. Wendler, W. Trippensee, C. Schubert, W. Wesch, Nucl. Instr. Meth. B 174 (2001) 199. [18] K. Gärtner, Nucl. Instr. Meth. B 227 (2005) 522. [19] J.P. Biersack, J.F. Ziegler, The Stopping and Ranges of Ions in Matter, vol. 1, Pergamon Press, Oxford, 1985. [20] R. Bäuerlein, Z. Phys. 176 (1963) 498. [21] E. Wendler, L. Wendler, Appl. Phys. Lett. 100 (2012) 192108. [22] N. Hecking, K.F. Heidemann, E. TeKaat, Nucl. Instr. Meth. B 15 (1986) 760. [23] B. Breeger, E. Wendler, Ch. Schubert, W. Wesch, Nucl. Instr. Meth. B 148 (1999) 468. [24] E. Wendler, W. Wesch, G. Götz, Nucl. Instr. Meth. B 55 (1991) 789. [25] M.W. Bench, I.M. Robertson, M.A. Kirk, Nucl. Instr. Meth. B 59 (60) (1991) 372. [26] W.J. Weber, L.M. Wang, N. Yu, N.J. Hess, Mater. Science Eng. A 253 (1998) 62. [27] E. Wendler, W. Wesch, G. Götz, Phys. Status Solidi A 112 (1989) 289. [28] T. Egami, Y. Waseda, J. Non-Cryst. Solids 64 (1984) 113. [29] G. Linker, Solid State Commun. 57 (1986) 773. [30] P. Partyka, R.S. Averback, D.V. Frobes, J.J. Coleman, P. Ehrhart, W. Jäger, Appl. Phys. Lett. 65 (1994) 421. [31] J. Linnros, R.G. Elliman, W.L. Brown, J. Mater. Res. 3 (1988) 1208.