Effect of the Ga-content on the defect properties of CuIn1−xGaxSe2 single crystals

Thin Solid Films 361±362 (2000) 411±414 www.elsevier.com/locate/tsf

Effect of the Ga-content on the defect properties of CuIn12xGaxSe2 single crystals J.H. SchoÈn a, b,*, Ch. Kloc a, E. Bucher a, b a

Bell Laboratories, Lucent Technologies, 600 Mountain Avenue, Murray Hill, NJ07974, USA b UniversitaÈt Konstanz, FakultaÈt fuÈr Physik, P.O. Box X 916, D-78457 Konstanz, Germany

Abstract The defect properties of as-grown and annealed CuInSe2, CuIn0.5Ga0.5Se2, and CuGaSe2 single crystals grown by chemical vapor transport have been studied by photoluminescence, Hall effect, and resistivity measurements. The various observed defect levels are ascribed to intrinsic defect states taking into account stoichiometry, annealing conditions of the samples, and the formation enthalpy of these intrinsic defects. Furthermore, the obtained activation energies and the concentrations of these defects are compared for all three different compounds. The properties of the as-grown, p-type samples are dominated by copper and selenium vacancies (VCu, VSe). The activation energies of these defects decrease slightly with the increase of In-content, which is in line with the change of the band-gap. In contrast, a vast difference is observed for the properties of a donor level induced by annealing in the presence of the corresponding group III element. This defect level is much shallower and more easily formed in In-containing compounds. It is tentatively ascribed to a VCu±IIICu defect pair. These differences in the defect physics can account for the lack of p-type conduction and the limited solar cell performance of Ga-rich materials. q 2000 Elsevier Science S.A. All rights reserved. Keywords: CuIn12xGaxSe2; Annealing; Photoluminescence; Defects; Hall effect

1. Introduction

2. Experimental

Ternary Cu±III±VI2 chalcopyrites are very promising materials for thin ®lm solar cell applications [1]. The band gap in the CuIn12xGaxSe2 system can be varied from 1.04 eV for CuInSe2 (CIS) to 1.72 eV for CuGaSe2 (CGS). Therefore the addition of Ga should lead to a better match to the solar spectrum and a better performance of photovoltaic devices. However, the photovoltaic performance of CuIn12xGaxSe2 devices deteriorates for x . 0:3 [1±4]. Furthermore, Garich compounds lack n-type conduction [5]. In order to understand the differences in the defect chemistry of these compounds, we studied single crystals by means of Hall effect and photoluminescence measurements. Various defect levels are compared for different Ga-content and ascribed to intrinsic defects taking into account stoichiometry, electrical and optical properties, and the formation enthalpy of the defect levels.

CuInSe2 (CIS), CuIn0.5Ga0.5Se2 (CIGS), and CuGaSe2 (CGS) single crystals were grown by chemical vapor transport (CVT) using iodine as transport agent. All as-grown samples exhibit p-type conduction. Annealing experiments in group III atmosphere (elemental In, In/Ga, or Ga, respectively) were carried out in an open system horizontal furnace at temperatures ranging from 400 to 7008C. Argon ¯ow was used to prevent the formation of additional Sevacancies. Resistivity and Hall effect measurements were carried out using a conventional DC set-up in the temperature range between 30 and 400 K. Ohmic contacts were prepared by evaporation of Au and In for p- and n-type material, respectively. A Kr±ion laser was used as excitation source for steady-state photoluminescence (PL) measurements between 2 K and room temperature. The luminescent light was analyzed and detected by a grating monochromator and a liquid nitrogen cooled germanium detector or a S1photomultiplier tube, respectively. 3. Results and discussion

* Corresponding author. Tel.: 11-908-582-3052; fax: 11-908-5823260. E-mail address: [email protected] (J.H. SchoÈn)

All CVT-grown samples exhibit p-type conduction. Fig. 1 shows the carrier concentration of typical CIS, CIGS, and

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J.H. SchoÈn et al. / Thin Solid Films 361±362 (2000) 411±414

412

Fig. 1. Carrier concentration as function of reciprocal temperature for all three Cu±III±Se2 compounds. The dotted lines correspond to a ®t according to a two acceptor, one donor model.

CGS samples as function of temperature. The electrical properties in these materials can be described using a model consisting of two acceptor and compensating donor levels [6±8]. The second acceptor state is very shallow (about 15 meV [9]) and will not be discussed in this article. The dominant acceptor level (A1) in these compounds is generally ascribed to Cu-vacancies (VCu) [10]. The activation energy of A1 varies with acceptor concentration due to screening effects. The thermal ionization energy in the in®nite dilution limit Ea10 can be estimated using the following formula 1=3 Ea ˆ Ea10 2 aNa1

…1†

Values of 45, 50 and 60 meV are obtained for Ea10 in CIS, CIGS, and CGS, respectively. The product of the slope a and the relative dielectric constant 1 r has a constant value …1r a ˆ …35 ^ 5† £ 105 meV cm) for all three materials, which is in good agreement with an empirical value for many II±VI and III±V compounds found by PoÈdoÈr [11]. Annealing in In- or Ga-atmosphere leads to the formation of compensating donor states, but only CIS can be made ntype. CIGS and CGS remain compensated p-type semiconductors. The activation energy of the donor level D2 in CIS is estimated to be 20 meV. Fig. 2 shows the PL-spectra of as-grown and III-annealed CIS and CGS at 2 K. PL-spectra of CIGS samples exhibit the same features as both other compounds. The PL emission of the as-grown crystals of both materials is very similar. Four peaks are clearly observable. Peak 1 is ascribed to the excitonic recombination and peak 2 is identi®ed as freeto-bound transition. Excitation intensity dependent measurements revealed the donor±acceptor pair (DAP) character of peak 3. Because of the similar temperature and excitation intensity dependence peak 4 is ascribed to a phonon replica of the DAP transition peak 3. Hence, the energy of the longitudinal optical phonon can be estimated to 27, 32, and 36 meV in the three compounds, which is in good agreement with values reported in literature [12,13]. An activation energy of 45, 50, and 60 meV is obtained from the energetic position and the thermal quenching of peak 2, which corresponds to the A1 defect. Therefore this emission is ascribed to the transition between VCu and the conduction band. The activation energies of both defects of the donor±acceptor-pair transition (peak 3) are obtained

Fig. 2. Photoluminescence emission of as-grown (left) and annealed (right) CuInSe2 (top) and CuGaSe2 (bottom) single crystals at 2 K.

J.H. SchoÈn et al. / Thin Solid Films 361±362 (2000) 411±414

from the thermal quenching (45, 55, and 60 meV) and the excitation intensity dependent measurements of the peak position, which is given by ÿ




EDAP ˆ Eg 2 Ea 1 Ed 1 e2 =4p1r 10 r

…2†

where the last term corresponds to the Coulomb interaction of the donor±acceptor-pair at a distance r. Values of 60, 70, and 80 meV are obtained for the second level (D1) in CIS, CIGS, and CGS, respectively. This defect state D1 is generally ascribed to the Se-vacancy (VSe) [10,14±16], therefore peak 3 is interpreted as VCu±VSe transition. After annealing in III-atmosphere the PL is dominated by rather broad features. Furthermore the energetic region of the PL emission differs signi®cantly for CIS and CGS (see Fig. 3). The PL spectrum of CIGS is very similar to that of CGS, which is dominated by a donor±acceptor-pair transition (peak 6). The activation energy of the thermal quenching is approximately 60 meV for CGS and 50 meV for CIGS and therefore ascribed to the acceptor level A1. A donor (D2) activation energy of 110 and 130 meV is estimated for CIGS and CGS, respectively. The PL of annealed, n-type CIS consists of a shallow free-to-bound (peak 5) and a donor±acceptor-pair transition (peak 6). Since the PL-intensities of peak 5 and 6 exhibit the same temperature dependence and the material shows n-type conduction, we assume that one donor level (D2) determines the thermal quenching of both peaks. A value of 45 meV is estimated for the second defect level of the donor±acceptor-pair transition (peak 6), which is in very good agreement with the values obtained for A1 by Hall and PL measurements. However,

Fig. 3. Thermal quenching of the donor±acceptor-pair (DAP) transition for as-grown (peak 3) and annealed (peak 6) CuInSe2 and CuGaSe2 single crystals. In contrast to CuGaSe2 the activation energy changes for CuInSe2.

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Table 1 Defect levels in CuInSe2, CuIn0.5Ga0.5Se2, and CuGaSe2 single crystals

CuInSe2 CuIn0.5Ga0.5Se2 CuGaSe2

Ea1 (meV) VCu

Ed1 (meV) VSe

Ed2 (meV) VCu±IIICu

Eg (eV)

45 50 60

60 70 80

25 110 130

1.044 1.490 1.735

due to the strong shift of the peak position (6) with excitation intensity (9±12 mV/decade) the DAP model has to be seen as a rough estimate and spatial potential ¯uctuations [17±19] have to be taken into account for a more quantitative evaluation. Hence peak 5 is ascribed to a D2-valence band and peak 6 to a VCu±D2 transition. Hence, Hall and PL measurements reveal in good agreement, that the properties of the studied as-grown and annealed Cu III±Se2 materials are dominated by three defect levels (A1, D1, and D2). The activation energies of these defects are summarized in Table 1 and Fig. 4. The results of the electrical and optical characterization clearly reveal the formation of an additional donor level upon annealing in group-III atmosphere. Due to the annealing conditions we suggest that this state D2 is related to the IIICu antisite defect. Since theoretical calculations showed that this point defect should create much deeper levels in Cu III±Se2 compounds [5,20], we tentatively ascribe D2 to a IIICu±VCu defect complex. Pairing of these point defects pushes the deep IIICu level closer to the conduction band [5]. The annealing experiments indicate that this effect is formed much more easily in In-rich material, leading to ntype conduction in CIS. The smaller formation energy in CIS can be explained by the smaller band-gap and the smaller cohesive energy of In metal compared to Ga [20]. Furthermore, the activation energy in the Ga-containing compounds is signi®cantly higher as in pure CIS, which is in accordance with theoretical calculation. Two reasons can

Fig. 4. Comparison of the activation energies of the dominant defect levels in CuIn12xGaxSe2 as a function of the Ga-content of the sample. The dotted lines are a guide to the eye.

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be given for this increased activation energy upon Ga-alloying [5,20], (i) Since Ga-atoms are signi®cantly smaller than In-atoms, the lattice compression pushes the conduction band minimum upward, resulting in a higher defect activation energy in CIGS and CGS, (ii) Due to the s-wave character of the antisite defect the wave function is localized at the group-III atom. Since the Ga 4s level is about 0.7 eV lower compared to the In 5s level [21], the IIICu antisite defect is deeper in CIGS and CGS than in CIS. These differences in defect physics seem to be related to the lack of ntype conduction in Ga-rich compounds. In general, high ef®ciency Cu III±Se2 solar cells are prepared from Cupoor materials [22]. This leads to a type inversion at the top of the thin ®lm (n-type ordered vacancy compound) resulting in a `buried' p±n-junction, which is bene®cial for the suppression of interface-recombination [23,24]. Due to the lack of this type conversion for CIGS (x . 0:3), the p±njunction of photovoltaic devices moves to the `real' CdS/ CIGS interface causing higher recombination losses.

4. Summary and conclusion The electrical and optical properties of as-grown and IIIannealed Cu III±Se2 single crystals were examined by photoluminescence, Hall effect, and resistivity measurements. The dominant defect levels in these compounds are compared and ascribed to intrinsic defects taking into account stoichiometry and formation enthalpy of the defects. Cu-vacancies, which are the dominant acceptor states, exhibit very similar activation energies in all three materials. In as-grown samples they are partly compensated by Se-vacancies, which also show very comparable properties in the investigated Cu±III±Se2 chalcopyrites. The most striking difference between In-rich and Ga-rich compounds is found for IIICu-related defect levels formed after annealing in the presence of In or Ga, respectively. In both cases donor levels are formed, but the activation energy in In-rich materials is signi®cantly smaller. Furthermore, the formation of these defects is easier in CuInSe2 resulting in n-type conduction. These defect levels are tentatively ascribed to IIICu±VCu defect pairs, which act as shallow donors. The observed differences in the defect physics of Cu III±Se2 compounds seem to be responsible for the lack of n-type conduction and might limit the solar cell performance of Ga-rich chalcopyrites.

Acknowledgements We would like to thank V. Alberts, E. Arushanov, F.P. Baumgartner, K. Fess, K. Friemelt, M. Klenk, Ch. Kloc, L.L Kulyuk, H. Riazi-Nejad, O. Schenker and M. Steiner for technical assistance and fruitful discussions. References [1] J. HedstroÈm, H. Ohlsen, M. Bodegard, et al., Proc. 23rd IEEE Photovoltaic Specialists' Conf., IEEE, New York, 1993, p. 364. [2] A.M. Gabor, J.R. Tuttle, D.S. Albin, M.A. Contreras, R. Nou®, A.M. Hermann, Appl. Phys. Lett. 65 (1994) 198. [3] W.N. Shafarman, R. Klenk, B.E. McCandless, Proc. 25th IEEE Photovoltaic Spec. Conf., IEEE, New York, 1996, p. 763. [4] W.N. Shafarman, R. Klenk, B.E. McCandless, J. Appl. Phys. 79 (1996) 7324. [5] S.-H. Wei, S.B. Zhang, A. Zunger, Appl. Phys. Lett. 72 (1998) 3192. [6] J.H. SchoÈn, F.P. Baumgartner, H. Riazi-Nejad, Ch. Kloc, E. Bucher, J. Appl. Phys. 79 (1996) 6961. [7] D.J. Schroeder, A.A. Rockett, J. Appl. Phys. 82 (1997) 4982. [8] D.J. Schroeder, J.L. Hernandez, G.D. Berry, A.A. Rockett, J. Appl. Phys. 83 (1998) 1519. [9] J.H. SchoÈn, O. Schenker, H. Riazi-Nejad, K. Friemelt, Ch. Kloc, E. Bucher, Phys. Status Solidi A 161 (1997) 301. [10] S.B. Zhang, S.-H. Wei, A. Zunger, H. Katayama-Yoshida, Phys. Rev. B 57 (1998) 9642. [11] B. PoÈdoÈr, Semicond. Sci. Technol. 2 (1987) 177. [12] F.J. Ramirez, C. Rincon, Solid State Commun. 84 (1992) 551. [13] H. Tanino, T. Maeda, H. Fujikake, H. Nakanishi, S. Endo, T. Irie, Phys. Rev. B 45 (1992) 13323. [14] G. Masse, J. Phys. Chem. Solids 45 (1984) 1091. [15] S. Zott, K. Leo, M. Ruckh, H.W. Schock, Appl. Phys. Lett. 68 (1996) 1144. [16] J.H. SchoÈn, O. Schenker, H. Riazi-Nejad, K. Friemelt, Ch. Kloc, E. Bucher, Phys. Status Solidi A 161 (1997) 301. [17] B.I. Shklovskii, A.L. Efros, Electronic Properties of Doped Semiconductors, Springer, Berlin, 1984. [18] I. Dirnstorfer, Mt. Wagner, D.M. Wagner, D.M. Hofmann, M.D. Lampert, F. Karg, B.K. Meyer, Phys. Status Solidi A 168 (1998) 163. [19] J. Krustok, H. Collan, M. Yakushev, K. Hjelt, Physica Scr. 79 (1999) 179. [20] A. Zunger, S.B. Zhang, S.-H. Wei, Proc. 26th IEEE Photovoltaic Spec. Conf., IEEE, New York, 1997, p. 32. [21] A. Franceschetti, S.-H. Wei, A. Zunger, Phys. Rev. B 50 (1994) 8094. [22] V. Nadenau, D. Braunger, D. Hariskos, M. Kaiser, C. KoÈble, A. Oberacker, M. Ruckh, U. RuÈhle, R. SchaÈf¯er, et al., Prog. Photovoltaics 3 (1995) 363. [23] M.A. Contreras, H. Wiesner, D. Niles, K. Ramanathan, R. Matson, J. Tuttle, J. Keane, R. Nou®, Proc. 25th IEEE Photovoltaic Spec. Conf., IEEE, New York, 1996, p. 809. [24] D. Schmid, M. Ruckh, F. Grunwald, H.W. Schock, J. Appl. Phys. 79 (1996) 7324.