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Hydrogen in CuInSe2
Journal of Physics and Chemistry of Solids 64 (2003) 1641–1647 www.elsevier.com/locate/jpcs
Hydrogen in CuInSe2 K. Ottea,b,*, G. Lippoldb, H. Neumannc, A. Schindlera a
Institut fu¨r Oberfla¨chenmodifizierung, Permoserstr. 15, 04318 Leipzig, Germany b Solarion GmbH, Ostende 5, 04288 Leipzig, Germany c Fritz-Simon-Str. 26/111, 04347 Leipzig, Germany
Abstract The surface of vaccum-cleaved and oxidized CuInSe2 single crystals with different deviations from ideal stoichiometry and oxidized CuGaSe2 thin films were exposed to low energy hydrogen ions (300 eV). Besides the removal of surface contaminations within a short exposure time as studied by X-ray photoelectron spectroscopy, this process leads to hydrogen incorporation into the bulk of the material accompanied by copper depletion up to the same depth. In addition, the hydrogen ions influence the defect equilibrium in such a way, that type-conversion from initially p-type to n-type CuInSe2 is possible. In case of Cu-rich material, hydrogen is found to remove the binary phases existing at the surface. A defect chemical model is proposed to explain the experimental observations. Interstitial hydrogen acts mainly as a donor and is able to passivate copper-vacancies (VCu). The Cu-depleted surface layer is the result of electromigration due to the built-in electrical field caused by band bending. Exposure time and sample temperature influence not only the hydrogen diffusion profile, but also the thickness of the Cu-depleted layer and the position of the produced pn-homojunction. q 2003 Elsevier Ltd. All rights reserved. Keywords: D. Defects
1. Introduction Hydrogen in semiconductors is an area of tremendous practical importance and has been intensively studied for various semiconductors such as silicon and binary compounds [1,2]. Much progress has been made in understanding the role of hydrogen as an impurity passivating or compensating extrinsic (dopants) or intrincic defects and thus affecting the electrical and optical properties of these materials [2 –4]. The promising results obtained for elemental and binary semiconductors motivated to consider hydrogenation as a possibility to modify in a controlled manner the electrical properties of ternary compounds like CuInSe2 (CIS) or Cu(In,Ga)Se2 (CIGS) used in the development and production of polycrystalline thin film solar cells. The most important result of first experiments was, that conventional hydrogen implantation as well as annealing in a hydrogen * Corresponding author. Address: Solarion GmbH, Ostende 5, 04288 Leipzig, Germany. Tel.: þ49-342-97169832; fax: þ 493497-15031. E-mail address: [email protected] (K. Otte).
plasma allowed near-surface n-type regions to be produced in p-type CIS crystals [5]. In analogy to models developed for binary semiconductors it has been then tentatively argued that the type conversion observed could be due to two main effects: (1) hydrogen atoms are trapped in empty copper sites and thus reduce the concentration of electrically active VCu acceptors and (2) there is an energetically stable interstitial position for hydrogen resulting in the formation of an electrically active donor [5,6]. However, this simple model remained speculative because of the complex intrinsic defect chemistry of the compound which can be influenced in an uncontrollable way by the hydrogenation methods used. The main disadvantage of plasma annealing is the produced lattice damage and the severe surface erosion by hydrogen related chemical effects leading to loss of selenium [5– 8] or even the formation of a surface layer of metallic indium [9]. The conventional high-energy implantation method generates damage at the surface, along the implantation path and within the implanted layer [5,7,10]. Hence, in order to introduce hydrogen into Cu-chalcopyrites without creating additional defects by the process itself, a method is required which allows (1) a separation of the plasma from the sample surface, (2) an
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energy and momentum transfer below the displacement thresholds of the constituent elements of the compound and (3) a diffusion of implanted hydrogen into the bulk of the sample. The combination of low-energy implantation (, 500 eV) with increased sample temperature (100 – 300 8C) was found to fulfil these requirements [11]. This communication reports results of surface and bulk property changes of single crystalline CIS and polycrystalline CIGS observed after hydrogenation using this technique.
2. Experimental Setup The hydrogen implantation has been performed with a Kaufman-type ion source with a mass separating multi grid system (JENION RAH-20) [12]at a sample temperature of 300 8C in an ultra high vacuum (UHV) system. The ion beam energy and the current density of the hydrogen ion beam were adjusted to 300 eV and 25 mA/cm2, respectively. At a working pressure of 8 £ 1022 Pa the ion beam mass þ þ content Hþ 1 :H2 :H3 is about 2:3:5 [12]. For most of the experiments we used single crystalline CuInSe2, grown by the vertical Bridgman technique. The p-type material was cut into small pieces, polished with 0.05 mm Al2O3 powder and annealed at 350 8C in vacuum for 30 min prior to the experiment. For some investigations we used polycrystalline CuGaSe2 and Cu-rich CuInS2 thin films. The X-ray photoelectron spectroscopy (XPS) investigations were done with a VG ESCALAB 220iXL electron spectrometer and a monochromatized Al Ka source at an energy resolution of 450 meV attached to the UHV chamber. The Raman spectra were taken using a DILOR XY800 Raman spectrometer with 514.53 nm Arþ ion laser excitation in unpolarized back-scattering configuration.
3. Results and discussion 3.1. Surface effects 3.1.1. Removal of surface contaminations The low-energy hydrogen implantation of single crystalline CIS as well as polycrystalline CuGaSe2 (CGS) and CIGS results in a removal of native surface oxides (In2O3, SeO2, Ga2O3) as well as contaminations (carbon, water) as proven by XPS [13,14]. Utilizing an ion energy of Eion ¼ 300 eV and an ion current density of j ¼ 25 mA/cm2 all surface contaminations are removed within exposure times of less than t ¼ 10 min (ion dose: d < 1 £ 1017 ions/cm2) at a sample temperature of T ¼ 300 8C: In Fig. 1 the progression of the In 3d5/2 core-level spectra as a function of hydrogen ion beam exposure time for an In-rich p-type CIS crystal is shown as an example. The binding energy of the In 3d-peak for
Fig. 1. XPS core-level spectra of the In 3d5/2 peak of a CuInSe2 sample for various hydrogen ion beam exposure times as indicated in the figure and one measurement on an oxidized (reference) sample. The dashed lines reveal different components as a result of curve fitting.
the reference sample (oxidized) is EB ¼ 444:7 eV with respect to the Fermi-level EF ; which is in good agreement with literature values [15]. The asymmetry at higher binding energies is due to an additional component at a position of EB ¼ 445:2 eV which is visible by curve fitting (dashed line in Fig. 1). The observation of this component can be explained by the existence of In2O3. This component vanishes after ion beam exposure of 10 s. Except the total binding energy shift due to the band bending (discussed later in this paper) no further chemical shift is observed, even after an hydrogen exposure dose of d ¼ 7 £ 1017 ions=cm2 [16]. The chemical reaction of hydrogen ions with SeO2 is a reduction process and yields volatile H2O and H2Se [13]. A similar reduction process resulting in metallic In or Ga at the surface of the sample, however is not observed by XPS measurements [13]. The following chemical reactions are proposed: In2 O3 þ 4Hþ þ 4e2 ! In2 O þ 2H2 O Ga2 O3 þ 4Hþ þ 4e2 ! Ga2 O þ 2H2 O
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The sublimation of Ga 2O has been reported for a temperature between 200 and 350 8C [17,18]. Hence, we believe that the same holds for In2O. In addition, the sublimation is probably supported by a slight increase of the surface temperature due to the energy and momentum transfer of the incident ions. A lower sample temperature results in the formation of metallic In and Ga by a further reduction of In2O and Ga2 O, respectively. During contamination removal and beyond, no indications of crystal damage could be observed by Raman spectroscopy [16,19,20]. 3.1.2. Reduction of binary phases Low-energy hydrogen implantation has also been studied for Cu-rich chalcopyrites [13,19,20]. These materials reveal binary phases at the surfaces, such as CuxSe ðx < 1Þ or CuxS ðx < 1Þ in the case of CuGaSe2 and CuInS2, respectively. In Fig. 2, Raman measurements on Cu-rich CuInS2 samples are shown. Besides the peak at 292 cm21 which can be attributed to the A1 mode of CuInS2, an additional mode at 474 cm21 is observed. This additional mode clearly demonstrates the existence of the binary phase CuS [21] at the surface of the reference sample. After hydrogen exposure (Eion ¼ 300 eV; j ¼ 2 mA=cm2 ; T ¼ 250 8C; t ¼ 60 min) the mode attributed to the binary phase disappears. This can be related to a removal of the binary
Fig. 2. Unpolarized Raman spectra of Cu-rich polycrystalline CuInS2 prior to (reference) and after low energy hydrogen implantation.
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phase. Similar results have been observed for the removal of CuSe on the surface of Cu-rich CuGaSe2 samples [13,19]. It should be noted, that annealing step in a hydrogen atmosphere under identical conditions does not result in a removal of the binary phase. This implies that the removal is indeed an effect caused by atomic hydrogen. In order to identify a chemical reaction path for the removal process, Raman measurements and in situ XPS measurements of the core-levels and the valence-band (VB) on a thin layer of CuSe on the surface of CIS single crystal after step-wise exposure to the hydrogen ion beam (Eion ¼ 300 eV; j ¼ 25 mA=cm2 ; T ¼ 300 8C) have been performed [22,23]. These experiments demonstrate that the chemical reaction of the hydrogen ions with the binary phase CuSe is not a pure removal process but rather a step-wise reduction process to metallic Cu with an intermediate reaction step to Cu2Se. One of the reaction products is H2Se which is volatile at the process conditions used. 3.1.3. Changes in surface composition Besides the removal of surface contaminations and binary phases, the hydrogen exposure of CuInSe2 results also in changes of the surface composition as revealed by step-wise in situ XPS measurements [14]. The strongest changes are observed during the removal of surface oxides. The calculation of the surface composition from the Cu 3p-, In 3d- and Se 3d- core-levels shows the existence of an apparent equilibrium in elemental composition after removal of the oxidized layer close to Cu2In4Se7 whereby the required hydrogen exposure dose strongly depends on the thickness of the contamination layer at the used exposure parameters (Eion ¼ 300 eV; j ¼ 25 mA=cm2 ; T ¼ 300 8C) [14]. Since neither the In- and Se-oxide removal leaves elemental In or Se at the surface nor the hydrogen ion beam exposure gives rise to a significant sputter effect at this energy [23], one can assume that the observed Cu-poor CIS-surface must have its origin in a Cu-diffusion process probably induced by hydrogen. This assumption is strongly supported by XPS measurements on CIS samples cleaved in ultra-high vacuum, where a similar Cu-reduction is observed after hydrogen exposure [23]. 3.1.4. Conductivity type-conversion XPS measurements of the VB of all investigated oxidized CuInSe2 single crystals reveal a position of the valence-band maximum ðEVBM Þ with respect to the Fermilevel EF of approximately EF 2 EVBM ¼ 0:4 ^ 0:1 eV [16, 24]. After hydrogen ion beam exposure (Eion ¼ 300 eV; j ¼ 25 mA=cm2 ; T ¼ 300 8C) the VBM shifts to EF 2 EVBM ¼ 0:9 ^ 0:1 eV: The observed shift of the VBM corresponds to a band bending at the surface due to the conductivity type-conversion from initially p-type to n-type CIS and is the reason for the In 3d- core-level shift depicted in Fig. 1. One might assume that the observed band bending is due to the oxide removal by the hydrogen ions. However, XPS measurements on oxid-free
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CuInSe2 has been described as a trap and release transport process rather than as intrinsic diffusion of interstitial hydrogen [26]. Hence, the maximum hydrogen concentration which can be incorporated into the material depends not only on the solubility but also on the concentration of intrinsic defects. The highest hydrogen concentration is detected at the surface of the sample, where, in addition, a Cu-deficiency is observed. Total-reflection X-ray fluorescence analysis on bevelled sections of a CuInSe2 single crystal, fabricated with a nitrogen ion beam etching process [27] prior and after hydrogen implantation, reveals a Cu-deficiency up to a depth of about 200 nm which is similar to the hydrogen diffusion profile under identical implantation conditions [25]. However, SIMS measurements on the same sample demonstrate, that the detected Cu-deficiency is less than 3 at.%. 3.3. Defect chemical explanation
Fig. 3. XPS valence band spectra of CuInSe2 single crystal cleaved in UHV prior to (reference) and after low energy hydrogen implantation.
CIS-samples—cleaved in UHV—reveal the same effect after low energy ion beam exposure as shown in Fig. 3. The VBM of the cleaved reference sample is shifted by about 500 meV after hydrogen exposure of 2 min (Eion ¼ 300 eV; j ¼ 25 mA=cm2 ; T ¼ 300 8C). In this figure the position of the VBM with respect to the Fermi-level is marked by arrows. Hence, the type-conversion is due to a hydrogen related change in defect properties rather than an oxide removal effect.
Summarizing the low-energy hydrogen implantation results, we have always observed a correlation between hydrogen incorporation, Cu-deficiency and type-conversion at the surface and, in addition, a hydrogen depth profile and Cu-depletion within the same depth range. In order to explain these correlations we have to take a closer look at the defect mechanism which might be responsible for these observations. It is known, that there exist VSe at the (112) surface of CIS resulting in an increased concentration of donors and hence leading to a band bending [28]. Due to the oxidation of the CIS-surface, these donors are passivated. A removal of surface oxides by hydrogen and consequently the reactivation of VSe cannot be the only possible explanation for the type-conversion since it has also been detected for oxid-free (UHV-cleaved) samples [16,24,25]. In previous publications [16,23– 25] we have proposed the following models to be responsible for the observed type-conversion of p-CIS: 1. interstitial hydrogen acting as a donor, 2. passivation of VCu by hydrogen: þ 0 V2 Cu þ H ! {VCu 2 H}
3.2. Bulk effects By exposing a CIS single crystal to the low-energy hydrogen ion beam (Eion ¼ 300 eV; j ¼ 25 mA=cm2 ; T ¼ 200 8C) for about 1 h, nuclear reaction analysis reveals an incorporation of hydrogen with concentrations from 1019 H/cm3 in a depth of about 250 nm up to some 10 21 H/cm3 next to the surface [25]. Additional secondary ion mass spectrometry (SIMS) investigations on CIS-samples implanted with hydrogen at different temperatures illustrate, that the hydrogen moves deeper into the sample at higher temperatures which clearly emphasizes a diffusion process. The movement of hydrogen within
or 3. step-wise interaction of hydrogen with the neutral defect complex resulting in a donor or an activation of InCu: þHþ
þHþ
2þ 0 2 2þ þ 2 þ 0 {2V2 Cu þ InCu } ! {VCu þ InCu } þ {VCu 2 H } ! 2 þ 0 In2þ Cu þ 2{VCu 2 H }
The proposed mechanism of the formation of InCu due to the removal of indium-oxide at the surface [24] can be ruled out, since the removal-process does not leave indium at the surface. Up to now, it was hard to distinguish which of the proposed models is the only
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reasonable explanation of the observed type-conversion. However, recent calculations by Kılıc¸ and Zunger [29] have shown that the models (1) and (2) are responsible for an increase of the net-donor concentration. Their theoretical predictions provide a clear picture, that model (1) is especially important for stoichiometric CIS. Model (3) is based on the assumption, that the VCu – H defect complex has a lower formation energy than that of the defect complex 2VCu – InCu itself. However, as shown by Kılıc¸ and Zunger [29], it seems more reasonable to consider the formation of a new complex, explicitly {2 VCu 2 InCu 2 H}; which acts as a donor, especially in non-stoichiometric CIS [29]. Finally, incorporated hydrogen gives rise to an increase of the donor concentration for either stoichiometric or defect-rich CIS and, hence, explains the observed type-conversion. The resulting band bending at the surface is shown schematically on the left hand side in Fig. 4A. Due to the shift of the Fermi-level close to the conduction band minimum (CBM), the defect formation
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energy of VCu decreases [30], resulting in a spontaneous generation of VCu and Cui. It seems justified to suppose, that electromigration [31] (the built-in electrical field of the band bending drives Cui at the increased sample temperature from the surface into the volume of the material) is the reason for the observed Cu-depleted surface. This is schematically shown on the right-hand side of Fig. 4A. In addition, the hydrogen implantation and diffusion profile after short-time exposure are added to this figure. The existence of a high concentration of VCu at the surface, acting as compensating acceptors, would hinder a type-conversion. However, hydrogen is able to eliminate the compensating defect by passivation. The question arises, how the correlation between hydrogen incorporation and Cu-deficiency within the same depth can be explained. If the above mentioned model is true, it should also hold for the description of the observation in the volume of the material. By increasing the exposure time (long-time exposure in Fig. 4B), the hydrogen filled layer and the n-type CIS surface layer
Fig. 4. Schematically drawing of the band bending (left side) and the Cu- as well as H-concentration profile (right side) after short-time exposure (A) and long-time exposure (B) to a hydrogen ion beam.
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increase in thickness, resulting in a pn-homo-junction. Sinces the built-in electric field of the junction is the reason for the electromigration of the Cui, the hydrogen filled layer is Cu-depleted. The position of the pn-junction is defined by hydrogen diffusion. The whole model is only applicable if hydrogen exists, in contrary to most III – V semiconductors, in charge state Hþ only independent of the position of the Fermi-level within the band-gap. This means that a possible transition energy level (þ/2 ) from charge state Hþ to charge state H2 should lie close to or above the CBM [23]. It could be shown [29] that this is the case at least for CuInSe 2. Experimental evidence supporting this result has been found in muon-experiments [32], where only the positive muon could be observed in CuInSe2 independent of the conductivity type of the material.
remarkable, that hydrogen does not only act as donor but has also the ability to passivate VCu. The latter fact might be especially interesting for CuGaSe2. There, the n-type doping is hindered by the creation of VCu [33]. However, if CuGaSe2 is implanted with a donor-like doping element, as for example with Ge [33], implanted hydrogen might be able to support the type-conversion by passivating the VCu. Further investigations are necessary to proof this approach.
Acknowledgements The authors are greatful to Th. Chasse´ for many useful hints and comments. The authors wish to express their sincere thanks to A. Zunger and C ¸ . Kılıc¸ for fruitful discussion and for providing unpublished results.
4. Conclusion From the results presented it is obvious that the method of low-energy hydrogen implantation can be used as a fast, gentle and controllable surface cleaning process. Whereas all contaminations are removed within a short time (less than 2 min), the utilized spectroscopic techniques do not give any indications of damage to the CuInSe2 crystal. The feasibility of the low-energy cleaning step for the fabrication of solar cells should be evaluated with regard to its capability of eliminating the detrimental effect of oxygen-related surface passivation whereas leaving the oxygen passivated grain boundaries unaffected. This might become especially important if the chemical bath deposition step for the buffer layer CdS will be replaced by a vacuum process. The influence of the incorporated hydrogen on the defect concentration results in a near surface conversion of initially p-type CuInSe2 to n-type with a band bending of about 0.8 meV. The suitability of this effect for a corresponding modification of the production of CIS-based solar cells depends strongly on the type-conversion stability. XPS measurements have shown [24], that the surface is reverted to p-type after long-time storage in air, presumably due to the re-oxidation of the surface. Hence, the utilization of the hydrogen implantation process requires an in situ continuation of all the following process steps. Hydrogen is able to remove the binary phases CuS and CuSe from the surface of Cu-rich material. However, this removal process results in small traces of metallic Cu at the surface. Hence, the influence of the additional Cu as well as the hydrogen incorporation on the solar cell properties has to be studied for a possible substitution of the KCN etch step in case of Cu-rich CuInS2. The property of hydrogen to dope p-CuInSe2 could be shown and explained by defect chemical models. It is
References [1] N.H. Nickel (Eds.), Hydrogen in Semiconductors. II. Semiconductor and Semimetals, vol. 61, Academic Press, San Diego, 1999. [2] S.J. Pearton, J.W. Corbett, M. Stavola (Eds.), Hydrogen in Crystalline Semiconductors, Springer, Berlin, 1992. [3] P. de Mierry, P. Etchegoin, M. Stutzmann, Phys. Rev. B 49 (1994) 5283. [4] J. Chevallier, M. Aucouturier, Ann. Rev. Mater. Sci. 18 (1988) 219. [5] R.D. Tomlinson, A.E. Hill, G.A. Stephens, M. Imanieh, P.A. Jones, R.D. Pilkington, P. Rimmer, M.V. Yakushev, H. Neumann, Proc. 11th. E. C. Photovolt. Solar Energy Conf., Montreux, Switzerland (1992) 791. [6] M.V. Yakushev, H. Neumann, R.D. Tomlinson, P. Rimmer, G. Lippold, Cryst. Res. Technol. 29 (1994) 417. [7] M.V. Yakushev, P.A. Jones, H. Neumann, G.A. Stephens, R.D. Tomlinson, Nucl. Instrum. Meth. B 84 (1994) 405. [8] M.V. Yakushev, A. Zegadi, H. Neumann, P.A. Jones, A.E. Hill, R.D. Pilkington, M.A. Slifkin, R.D. Tomlinson, Cryst. Res. Technol. 29 (1994) 427. [9] A.J. Nelson, S. Frigo, R. Rosenberg, J. Appl. Phys. 73 (1993) 8561. [10] M.V. Yakushev, R.D. Tomlinson, H. Neumann, Cryst. Res. Technol. 29 (1994) 125. [11] K. Otte, G. Lippold, A. Schindler, F. Bigl, H. Schlemm, M.V. Yakushev, R.D. Tomlinson, Inst. Phys. Conf. Ser. 152E (1998) 773. [12] K. Otte, A. Schindler, F. Bigl, H. Schlemm, Rev. Sci. Instrum. 69 (1998) 1499. [13] K. Otte, G. Lippold, D. Hirsch, A. Schindler, M.V. Yakushev, R.W. Martin, F. Bigl, Proc. 16th Eur. Photovolt. Solar Energy Conf., 1–5 May 2000, Glasgow, UK (2000) 915. [14] K. Otte, G. Lippold, D. Hirsch, T. Chasse´, A. Schindler, M.V. Yakushev, R.W. Martin, F. Bigl, Thin Solid Films 387 (2001) 185. [15] R. Scheer, H.J. Lewerenz, J. Vac. Sci. Technol. A 12 (1994) 56.
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[26] A. Weidinger, J. Krauser, Th. Riedle, R. Klenk, M.Ch. LuxSteiner, M.V. Yakushev, Mater. Res. Soc. Symp. Proc. 513 (1998) 177. [27] K. Otte, G. Lippold, F. Frost, A. Schindler, F. Bigl, M.V. Yakushev, R.D. Tomlinson, J. Vac. Sci. Technol. A 17 (1999) 19. [28] D. Cahen, R. Noufi, Appl. Phys. Lett. 54 (1989) 558. [29] C¸. Kılıc¸, A. Zunger, submitted to Phys. Rev B. [30] S.B. Zhang, S.H. Wei, A. Zunger, Phys. Rev. Lett. 78 (1997) 4059. [31] D. Cahen, L. Chernyak, Adv. Mater. 9 (1997) 861. [32] J.M. Gil, P.J. Mendes, L.P. Ferreira, H.V. Alberto, R.C. Vilao, N. Ayres de Campos, A. Weidinger, Y. Tomm, Ch. Niedermayer, M.V. Yakushev, R.D. Tomlinson, S.P. Cottrell, S.F.J. Cox, Phys. Rev. B. 59 (1999) 1912. [33] J.H. Scho¨n, N. Fabre, J. Oestreich, O. Schenker, H. RiaziNejad, M. Klenk, E. Arushanov, E. Bucher, Appl. Phys. Lett. 75 (1999) 2969.
Hydrogen in CuInSe2 K. Ottea,b,*, G. Lippoldb, H. Neumannc, A. Schindlera a
Institut fu¨r Oberfla¨chenmodifizierung, Permoserstr. 15, 04318 Leipzig, Germany b Solarion GmbH, Ostende 5, 04288 Leipzig, Germany c Fritz-Simon-Str. 26/111, 04347 Leipzig, Germany
Abstract The surface of vaccum-cleaved and oxidized CuInSe2 single crystals with different deviations from ideal stoichiometry and oxidized CuGaSe2 thin films were exposed to low energy hydrogen ions (300 eV). Besides the removal of surface contaminations within a short exposure time as studied by X-ray photoelectron spectroscopy, this process leads to hydrogen incorporation into the bulk of the material accompanied by copper depletion up to the same depth. In addition, the hydrogen ions influence the defect equilibrium in such a way, that type-conversion from initially p-type to n-type CuInSe2 is possible. In case of Cu-rich material, hydrogen is found to remove the binary phases existing at the surface. A defect chemical model is proposed to explain the experimental observations. Interstitial hydrogen acts mainly as a donor and is able to passivate copper-vacancies (VCu). The Cu-depleted surface layer is the result of electromigration due to the built-in electrical field caused by band bending. Exposure time and sample temperature influence not only the hydrogen diffusion profile, but also the thickness of the Cu-depleted layer and the position of the produced pn-homojunction. q 2003 Elsevier Ltd. All rights reserved. Keywords: D. Defects
1. Introduction Hydrogen in semiconductors is an area of tremendous practical importance and has been intensively studied for various semiconductors such as silicon and binary compounds [1,2]. Much progress has been made in understanding the role of hydrogen as an impurity passivating or compensating extrinsic (dopants) or intrincic defects and thus affecting the electrical and optical properties of these materials [2 –4]. The promising results obtained for elemental and binary semiconductors motivated to consider hydrogenation as a possibility to modify in a controlled manner the electrical properties of ternary compounds like CuInSe2 (CIS) or Cu(In,Ga)Se2 (CIGS) used in the development and production of polycrystalline thin film solar cells. The most important result of first experiments was, that conventional hydrogen implantation as well as annealing in a hydrogen * Corresponding author. Address: Solarion GmbH, Ostende 5, 04288 Leipzig, Germany. Tel.: þ49-342-97169832; fax: þ 493497-15031. E-mail address: [email protected] (K. Otte).
plasma allowed near-surface n-type regions to be produced in p-type CIS crystals [5]. In analogy to models developed for binary semiconductors it has been then tentatively argued that the type conversion observed could be due to two main effects: (1) hydrogen atoms are trapped in empty copper sites and thus reduce the concentration of electrically active VCu acceptors and (2) there is an energetically stable interstitial position for hydrogen resulting in the formation of an electrically active donor [5,6]. However, this simple model remained speculative because of the complex intrinsic defect chemistry of the compound which can be influenced in an uncontrollable way by the hydrogenation methods used. The main disadvantage of plasma annealing is the produced lattice damage and the severe surface erosion by hydrogen related chemical effects leading to loss of selenium [5– 8] or even the formation of a surface layer of metallic indium [9]. The conventional high-energy implantation method generates damage at the surface, along the implantation path and within the implanted layer [5,7,10]. Hence, in order to introduce hydrogen into Cu-chalcopyrites without creating additional defects by the process itself, a method is required which allows (1) a separation of the plasma from the sample surface, (2) an
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energy and momentum transfer below the displacement thresholds of the constituent elements of the compound and (3) a diffusion of implanted hydrogen into the bulk of the sample. The combination of low-energy implantation (, 500 eV) with increased sample temperature (100 – 300 8C) was found to fulfil these requirements [11]. This communication reports results of surface and bulk property changes of single crystalline CIS and polycrystalline CIGS observed after hydrogenation using this technique.
2. Experimental Setup The hydrogen implantation has been performed with a Kaufman-type ion source with a mass separating multi grid system (JENION RAH-20) [12]at a sample temperature of 300 8C in an ultra high vacuum (UHV) system. The ion beam energy and the current density of the hydrogen ion beam were adjusted to 300 eV and 25 mA/cm2, respectively. At a working pressure of 8 £ 1022 Pa the ion beam mass þ þ content Hþ 1 :H2 :H3 is about 2:3:5 [12]. For most of the experiments we used single crystalline CuInSe2, grown by the vertical Bridgman technique. The p-type material was cut into small pieces, polished with 0.05 mm Al2O3 powder and annealed at 350 8C in vacuum for 30 min prior to the experiment. For some investigations we used polycrystalline CuGaSe2 and Cu-rich CuInS2 thin films. The X-ray photoelectron spectroscopy (XPS) investigations were done with a VG ESCALAB 220iXL electron spectrometer and a monochromatized Al Ka source at an energy resolution of 450 meV attached to the UHV chamber. The Raman spectra were taken using a DILOR XY800 Raman spectrometer with 514.53 nm Arþ ion laser excitation in unpolarized back-scattering configuration.
3. Results and discussion 3.1. Surface effects 3.1.1. Removal of surface contaminations The low-energy hydrogen implantation of single crystalline CIS as well as polycrystalline CuGaSe2 (CGS) and CIGS results in a removal of native surface oxides (In2O3, SeO2, Ga2O3) as well as contaminations (carbon, water) as proven by XPS [13,14]. Utilizing an ion energy of Eion ¼ 300 eV and an ion current density of j ¼ 25 mA/cm2 all surface contaminations are removed within exposure times of less than t ¼ 10 min (ion dose: d < 1 £ 1017 ions/cm2) at a sample temperature of T ¼ 300 8C: In Fig. 1 the progression of the In 3d5/2 core-level spectra as a function of hydrogen ion beam exposure time for an In-rich p-type CIS crystal is shown as an example. The binding energy of the In 3d-peak for
Fig. 1. XPS core-level spectra of the In 3d5/2 peak of a CuInSe2 sample for various hydrogen ion beam exposure times as indicated in the figure and one measurement on an oxidized (reference) sample. The dashed lines reveal different components as a result of curve fitting.
the reference sample (oxidized) is EB ¼ 444:7 eV with respect to the Fermi-level EF ; which is in good agreement with literature values [15]. The asymmetry at higher binding energies is due to an additional component at a position of EB ¼ 445:2 eV which is visible by curve fitting (dashed line in Fig. 1). The observation of this component can be explained by the existence of In2O3. This component vanishes after ion beam exposure of 10 s. Except the total binding energy shift due to the band bending (discussed later in this paper) no further chemical shift is observed, even after an hydrogen exposure dose of d ¼ 7 £ 1017 ions=cm2 [16]. The chemical reaction of hydrogen ions with SeO2 is a reduction process and yields volatile H2O and H2Se [13]. A similar reduction process resulting in metallic In or Ga at the surface of the sample, however is not observed by XPS measurements [13]. The following chemical reactions are proposed: In2 O3 þ 4Hþ þ 4e2 ! In2 O þ 2H2 O Ga2 O3 þ 4Hþ þ 4e2 ! Ga2 O þ 2H2 O
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The sublimation of Ga 2O has been reported for a temperature between 200 and 350 8C [17,18]. Hence, we believe that the same holds for In2O. In addition, the sublimation is probably supported by a slight increase of the surface temperature due to the energy and momentum transfer of the incident ions. A lower sample temperature results in the formation of metallic In and Ga by a further reduction of In2O and Ga2 O, respectively. During contamination removal and beyond, no indications of crystal damage could be observed by Raman spectroscopy [16,19,20]. 3.1.2. Reduction of binary phases Low-energy hydrogen implantation has also been studied for Cu-rich chalcopyrites [13,19,20]. These materials reveal binary phases at the surfaces, such as CuxSe ðx < 1Þ or CuxS ðx < 1Þ in the case of CuGaSe2 and CuInS2, respectively. In Fig. 2, Raman measurements on Cu-rich CuInS2 samples are shown. Besides the peak at 292 cm21 which can be attributed to the A1 mode of CuInS2, an additional mode at 474 cm21 is observed. This additional mode clearly demonstrates the existence of the binary phase CuS [21] at the surface of the reference sample. After hydrogen exposure (Eion ¼ 300 eV; j ¼ 2 mA=cm2 ; T ¼ 250 8C; t ¼ 60 min) the mode attributed to the binary phase disappears. This can be related to a removal of the binary
Fig. 2. Unpolarized Raman spectra of Cu-rich polycrystalline CuInS2 prior to (reference) and after low energy hydrogen implantation.
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phase. Similar results have been observed for the removal of CuSe on the surface of Cu-rich CuGaSe2 samples [13,19]. It should be noted, that annealing step in a hydrogen atmosphere under identical conditions does not result in a removal of the binary phase. This implies that the removal is indeed an effect caused by atomic hydrogen. In order to identify a chemical reaction path for the removal process, Raman measurements and in situ XPS measurements of the core-levels and the valence-band (VB) on a thin layer of CuSe on the surface of CIS single crystal after step-wise exposure to the hydrogen ion beam (Eion ¼ 300 eV; j ¼ 25 mA=cm2 ; T ¼ 300 8C) have been performed [22,23]. These experiments demonstrate that the chemical reaction of the hydrogen ions with the binary phase CuSe is not a pure removal process but rather a step-wise reduction process to metallic Cu with an intermediate reaction step to Cu2Se. One of the reaction products is H2Se which is volatile at the process conditions used. 3.1.3. Changes in surface composition Besides the removal of surface contaminations and binary phases, the hydrogen exposure of CuInSe2 results also in changes of the surface composition as revealed by step-wise in situ XPS measurements [14]. The strongest changes are observed during the removal of surface oxides. The calculation of the surface composition from the Cu 3p-, In 3d- and Se 3d- core-levels shows the existence of an apparent equilibrium in elemental composition after removal of the oxidized layer close to Cu2In4Se7 whereby the required hydrogen exposure dose strongly depends on the thickness of the contamination layer at the used exposure parameters (Eion ¼ 300 eV; j ¼ 25 mA=cm2 ; T ¼ 300 8C) [14]. Since neither the In- and Se-oxide removal leaves elemental In or Se at the surface nor the hydrogen ion beam exposure gives rise to a significant sputter effect at this energy [23], one can assume that the observed Cu-poor CIS-surface must have its origin in a Cu-diffusion process probably induced by hydrogen. This assumption is strongly supported by XPS measurements on CIS samples cleaved in ultra-high vacuum, where a similar Cu-reduction is observed after hydrogen exposure [23]. 3.1.4. Conductivity type-conversion XPS measurements of the VB of all investigated oxidized CuInSe2 single crystals reveal a position of the valence-band maximum ðEVBM Þ with respect to the Fermilevel EF of approximately EF 2 EVBM ¼ 0:4 ^ 0:1 eV [16, 24]. After hydrogen ion beam exposure (Eion ¼ 300 eV; j ¼ 25 mA=cm2 ; T ¼ 300 8C) the VBM shifts to EF 2 EVBM ¼ 0:9 ^ 0:1 eV: The observed shift of the VBM corresponds to a band bending at the surface due to the conductivity type-conversion from initially p-type to n-type CIS and is the reason for the In 3d- core-level shift depicted in Fig. 1. One might assume that the observed band bending is due to the oxide removal by the hydrogen ions. However, XPS measurements on oxid-free
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CuInSe2 has been described as a trap and release transport process rather than as intrinsic diffusion of interstitial hydrogen [26]. Hence, the maximum hydrogen concentration which can be incorporated into the material depends not only on the solubility but also on the concentration of intrinsic defects. The highest hydrogen concentration is detected at the surface of the sample, where, in addition, a Cu-deficiency is observed. Total-reflection X-ray fluorescence analysis on bevelled sections of a CuInSe2 single crystal, fabricated with a nitrogen ion beam etching process [27] prior and after hydrogen implantation, reveals a Cu-deficiency up to a depth of about 200 nm which is similar to the hydrogen diffusion profile under identical implantation conditions [25]. However, SIMS measurements on the same sample demonstrate, that the detected Cu-deficiency is less than 3 at.%. 3.3. Defect chemical explanation
Fig. 3. XPS valence band spectra of CuInSe2 single crystal cleaved in UHV prior to (reference) and after low energy hydrogen implantation.
CIS-samples—cleaved in UHV—reveal the same effect after low energy ion beam exposure as shown in Fig. 3. The VBM of the cleaved reference sample is shifted by about 500 meV after hydrogen exposure of 2 min (Eion ¼ 300 eV; j ¼ 25 mA=cm2 ; T ¼ 300 8C). In this figure the position of the VBM with respect to the Fermi-level is marked by arrows. Hence, the type-conversion is due to a hydrogen related change in defect properties rather than an oxide removal effect.
Summarizing the low-energy hydrogen implantation results, we have always observed a correlation between hydrogen incorporation, Cu-deficiency and type-conversion at the surface and, in addition, a hydrogen depth profile and Cu-depletion within the same depth range. In order to explain these correlations we have to take a closer look at the defect mechanism which might be responsible for these observations. It is known, that there exist VSe at the (112) surface of CIS resulting in an increased concentration of donors and hence leading to a band bending [28]. Due to the oxidation of the CIS-surface, these donors are passivated. A removal of surface oxides by hydrogen and consequently the reactivation of VSe cannot be the only possible explanation for the type-conversion since it has also been detected for oxid-free (UHV-cleaved) samples [16,24,25]. In previous publications [16,23– 25] we have proposed the following models to be responsible for the observed type-conversion of p-CIS: 1. interstitial hydrogen acting as a donor, 2. passivation of VCu by hydrogen: þ 0 V2 Cu þ H ! {VCu 2 H}
3.2. Bulk effects By exposing a CIS single crystal to the low-energy hydrogen ion beam (Eion ¼ 300 eV; j ¼ 25 mA=cm2 ; T ¼ 200 8C) for about 1 h, nuclear reaction analysis reveals an incorporation of hydrogen with concentrations from 1019 H/cm3 in a depth of about 250 nm up to some 10 21 H/cm3 next to the surface [25]. Additional secondary ion mass spectrometry (SIMS) investigations on CIS-samples implanted with hydrogen at different temperatures illustrate, that the hydrogen moves deeper into the sample at higher temperatures which clearly emphasizes a diffusion process. The movement of hydrogen within
or 3. step-wise interaction of hydrogen with the neutral defect complex resulting in a donor or an activation of InCu: þHþ
þHþ
2þ 0 2 2þ þ 2 þ 0 {2V2 Cu þ InCu } ! {VCu þ InCu } þ {VCu 2 H } ! 2 þ 0 In2þ Cu þ 2{VCu 2 H }
The proposed mechanism of the formation of InCu due to the removal of indium-oxide at the surface [24] can be ruled out, since the removal-process does not leave indium at the surface. Up to now, it was hard to distinguish which of the proposed models is the only
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reasonable explanation of the observed type-conversion. However, recent calculations by Kılıc¸ and Zunger [29] have shown that the models (1) and (2) are responsible for an increase of the net-donor concentration. Their theoretical predictions provide a clear picture, that model (1) is especially important for stoichiometric CIS. Model (3) is based on the assumption, that the VCu – H defect complex has a lower formation energy than that of the defect complex 2VCu – InCu itself. However, as shown by Kılıc¸ and Zunger [29], it seems more reasonable to consider the formation of a new complex, explicitly {2 VCu 2 InCu 2 H}; which acts as a donor, especially in non-stoichiometric CIS [29]. Finally, incorporated hydrogen gives rise to an increase of the donor concentration for either stoichiometric or defect-rich CIS and, hence, explains the observed type-conversion. The resulting band bending at the surface is shown schematically on the left hand side in Fig. 4A. Due to the shift of the Fermi-level close to the conduction band minimum (CBM), the defect formation
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energy of VCu decreases [30], resulting in a spontaneous generation of VCu and Cui. It seems justified to suppose, that electromigration [31] (the built-in electrical field of the band bending drives Cui at the increased sample temperature from the surface into the volume of the material) is the reason for the observed Cu-depleted surface. This is schematically shown on the right-hand side of Fig. 4A. In addition, the hydrogen implantation and diffusion profile after short-time exposure are added to this figure. The existence of a high concentration of VCu at the surface, acting as compensating acceptors, would hinder a type-conversion. However, hydrogen is able to eliminate the compensating defect by passivation. The question arises, how the correlation between hydrogen incorporation and Cu-deficiency within the same depth can be explained. If the above mentioned model is true, it should also hold for the description of the observation in the volume of the material. By increasing the exposure time (long-time exposure in Fig. 4B), the hydrogen filled layer and the n-type CIS surface layer
Fig. 4. Schematically drawing of the band bending (left side) and the Cu- as well as H-concentration profile (right side) after short-time exposure (A) and long-time exposure (B) to a hydrogen ion beam.
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increase in thickness, resulting in a pn-homo-junction. Sinces the built-in electric field of the junction is the reason for the electromigration of the Cui, the hydrogen filled layer is Cu-depleted. The position of the pn-junction is defined by hydrogen diffusion. The whole model is only applicable if hydrogen exists, in contrary to most III – V semiconductors, in charge state Hþ only independent of the position of the Fermi-level within the band-gap. This means that a possible transition energy level (þ/2 ) from charge state Hþ to charge state H2 should lie close to or above the CBM [23]. It could be shown [29] that this is the case at least for CuInSe 2. Experimental evidence supporting this result has been found in muon-experiments [32], where only the positive muon could be observed in CuInSe2 independent of the conductivity type of the material.
remarkable, that hydrogen does not only act as donor but has also the ability to passivate VCu. The latter fact might be especially interesting for CuGaSe2. There, the n-type doping is hindered by the creation of VCu [33]. However, if CuGaSe2 is implanted with a donor-like doping element, as for example with Ge [33], implanted hydrogen might be able to support the type-conversion by passivating the VCu. Further investigations are necessary to proof this approach.
Acknowledgements The authors are greatful to Th. Chasse´ for many useful hints and comments. The authors wish to express their sincere thanks to A. Zunger and C ¸ . Kılıc¸ for fruitful discussion and for providing unpublished results.
4. Conclusion From the results presented it is obvious that the method of low-energy hydrogen implantation can be used as a fast, gentle and controllable surface cleaning process. Whereas all contaminations are removed within a short time (less than 2 min), the utilized spectroscopic techniques do not give any indications of damage to the CuInSe2 crystal. The feasibility of the low-energy cleaning step for the fabrication of solar cells should be evaluated with regard to its capability of eliminating the detrimental effect of oxygen-related surface passivation whereas leaving the oxygen passivated grain boundaries unaffected. This might become especially important if the chemical bath deposition step for the buffer layer CdS will be replaced by a vacuum process. The influence of the incorporated hydrogen on the defect concentration results in a near surface conversion of initially p-type CuInSe2 to n-type with a band bending of about 0.8 meV. The suitability of this effect for a corresponding modification of the production of CIS-based solar cells depends strongly on the type-conversion stability. XPS measurements have shown [24], that the surface is reverted to p-type after long-time storage in air, presumably due to the re-oxidation of the surface. Hence, the utilization of the hydrogen implantation process requires an in situ continuation of all the following process steps. Hydrogen is able to remove the binary phases CuS and CuSe from the surface of Cu-rich material. However, this removal process results in small traces of metallic Cu at the surface. Hence, the influence of the additional Cu as well as the hydrogen incorporation on the solar cell properties has to be studied for a possible substitution of the KCN etch step in case of Cu-rich CuInS2. The property of hydrogen to dope p-CuInSe2 could be shown and explained by defect chemical models. It is
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