- Email: [email protected]
Gallium doping dependence of single-crystal n-type Zno grown by metal organic chemical vapor deposition
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Journal of Crystal Growth 283 (2005) 279–285 www.elsevier.com/locate/jcrysgro
Gallium doping dependence of single-crystal n-type Zno grown by metal organic chemical vapor deposition J.D. Yea, S.L. Gua,, S.M. Zhua, S.M. Liua, Y.D. Zhenga, R. Zhanga, Y. Shia, H.Q. Yub, Y.D. Yeb a
Key Laboratory of Advanced Photonic and Electronic Materials and Department of Physics, Nanjing University, Nanjing 210093, PR China b Materials Analysis Center of Nanjing University, Nanjing 210093, PR China Received 3 May 2005; received in revised form 30 May 2005; accepted 1 June 2005 Available online 26 July 2005 Communicated by R. Bhat
Abstract High-quality single-crystal Ga-doped ZnO films have been epitaxially deposited on (0 0 0 2) sapphire substrate by low-pressure metal organic chemical vapor deposition (MOCVD) technique. The dependence of structural, electrical and optical properties of films on Ga doping concentration was investigated. As grown at the Ga/Zn gas ratio of 3.2 at%, the film shows a narrow linewidth of 0.261 for ZnO (0 0 0 2) peak, high carrier concentration of 2.47 1019 cm3, and high optical transparency over 90%. The carrier concentration increased sharply and became saturated at higher doping level due to the onset of carrier compensation. The Burstein–Moss blueshift of the absorption edge energy increased as expected with the carrier concentration up to 2.47 1019 cm3. In addition, doping-induced photoluminescence (PL) emission linewidth broadening and bandgap renormalization (BGR) effects have also been observed. The intensity of PL emission decreased with increasing Ga dopant concentration, which was believed to be a direct consequence of the doping-enhanced nonradiative recombination rates. r 2005 Elsevier B.V. All rights reserved. Keywords: A3. Metal organic chemical vapor deposition; B1. Gallium doping; B1. ZnO
1. Introduction Zinc oxide, a wide-gap semiconductor with a high exciton binding energy (60 meV), has emerged Corresponding author.
E-mail addresses: [email protected] (J.D. Ye), [email protected] (S.L. Gu).
as a candidate for use in short-wavelength optoelectronic devices such as ultraviolet (UV) lasers and light-emitting diodes [1–5]. In addition, due to its high thermal and chemical stability, good electrical conductivity and high optical transparency, ZnO has also attracted more attention in the applications of display devices, solar cells and electronic transducers [4–8]. In general,
0022-0248/$ - see front matter r 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2005.06.030
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undoped ZnO films always exhibit n-type conduction with typical carrier concentration of 1017 cm3, which, however, is far from the demands of device applications. Thus, high conductivity controlled by intentional doping is necessary. Group III elements (Al, Ga and In), in comparison with group VII elements (Cl, Br) are more suitable for n-type doping of ZnO materials because of their lower vapor pressure [9]. ZnO: Al films have been studied widely; however, oxidation of the Al source during ZnO growth owing to the high reactivity of Al is a problem. The covalent bond length of Ga–O (0.192 nm) is slightly smaller than that of Zn–O (0.197 nm), which will make the deformation of the ZnO lattice small even in the case of high Ga doping concentration [10]. Thus, Ga is an attractive dopant for n-type ZnO. Polycrystalline ZnO:Ga thin films have been achieved by many methods [11–14]; however, few reports have so far focused on the control of carrier concentration in single-crystal Ga-doped ZnO epilayers. Single-crystal ZnO:Ga films with carrier concentration up to 1.13 1020 cm3 was first achieved on GaN substrate via molecular beam epitaxy (MBE) [10]. Recently, Zhong et al. [15] reported the MOCVD growth of Ga-doped ZnO single-crystal nanotips with a low resistivity of 4 103 O cm, where the Ga doping mechanism was not studied. In the present study, we report the successful epitaxy of single-crystal Ga-doped ZnO films with carrier concentration up to 2.47 1019 cm3 by low-pressure MOCVD technique. The dependence of structural, electrical and optical properties of ZnO on Ga doping concentration and the associated doping mechanisms have also been studied.
2. Experiments The undoped and Ga-doped ZnO films were uniformly deposited on a sapphire (0 0 0 1) substrate in a vertical, homemade, low-pressure MOCVD reactor. Prior to entering the reactor, 2 in sapphire substrates were degreased in acetone, ethanol and then chemically etched in H2SO4 at 200 1C for 10 min followed by deionized water
rinsing and nitrogen flow drying. During the deposition process, diethylzinc (DEZn), 5 N-purity oxygen, and triisopropylgallium (TIPGa) were used as Zn, O and Ga sources, respectively. Argon gas of 6 N-purity was used as metalorganic (MO) carrier gas. The Ga-doped ZnO thin films were deposited on an undoped ZnO buffer layer to improve the crystallinity of ZnO:Ga layers. The Ga/Zn mole ratio was modulated from 0.0 to 6.3 at%, by controlling the temperature of TIPGa MO source. The growth temperature was 450 1C and the thickness of ZnO:Ga epilayers is approximately 300 nm. In our study, no obvious variation of thickness was observed for different doping concentrations. The structure and crystallinity of samples were analyzed by X-ray diffraction (XRD). Carrier concentrations were evaluated by capacitance–voltage (C2V ) method at a frequency of 300 kHz using a mercury probe as Schottky contact. Transmittance spectra were recorded using an UV/visible spectrometer (U-3410, Hitachi) with a resolution of 1 nm. Photoluminescence (PL) spectra were recorded using He–Cd laser (325 nm) as the excitation source with an incident power density of 40 mW/cm2. All measurements are carried out at room temperature.
3. Results and discussion Fig. 1 shows the typical XRD curve of a ZnO:Ga thin film with a Ga/Zn mole ratio of 3.2 at% and the inset shows the comparison of the (0 0 0 2) peak for the undoped and Ga-doped ZnO thin films. The Ga-doped thin film exhibited (0 0 0 2) and (0 0 0 4) peaks at 2y ¼ 34.561 and 72.881 for wurtzite structure ZnO, indicative of preferential orientation with the c-axis perpendicular to the substrate. Besides these peaks, no other diffraction peaks were observed, that is to say, there was no second phase corresponding to gallium or its oxides even at a high Ga/Zn mole ratio of 3.2 at% during film growth. In addition, the small full-width at half-maximum (FWHM) of the (0 0 0 2) peak (0.261) confirms the single-crystal quality of Ga-doped ZnO film. However, as shown in the inset of Fig. 1, the (0 0 0 2) peak for
ARTICLE IN PRESS
undoped ZnO ZnO:Ga 3.2%
Al2O3 (0006)
ZnO (0002)
XRD Intensity (a.u.)
ZnO:Ga 3.2%
33.6
34.0
34.4
34.8
35.2
35.6
2θ (deg)
ZnO (0004)
30
40
50 2θ (deg)
60
70
Fig. 1. X-ray diffractions of ZnO:Ga thin film with Ga/Zn mole ratio of 3.2%. The inset figure shows the blown-up ZnO (0 0 0 2) peak of undoped and Ga-doped (3.2%) ZnO thin films.
Ga-doped ZnO layer broadened slightly and shifted to 34.561, higher than that of the undoped film (34.421), indicative of the decreasing lattice constant along the c-axis due to Ga incorporation. These phenomena generally exist in impuritydoped crystals and can be attributed to an increase in local strain around impurity atoms or point defect associated with the impurity atoms [16]. In fact, the covalent bond length of Ga–O is slightly smaller than that of Zn–O [10], and Ga substituted at Zn sites would reduce the lattice constant along the c-axis, and induce an uniaxial compressive strain in ZnO films, which is responsible for the shift of peak (0 0 0 2). It should also be noted that the mismatch of Ga–O and Zn–O, which is as small as 3%, is expected to make the deformation of the lattice small and cause only a slight degradation in crystal structure even in the case of high Ga doping concentrations. Fig. 2 shows room temperature carrier concentrations of Ga-doped ZnO films as a function of Ga/Zn gas mole ratio. C– V method is an alternative characterization technique that has been used extensively on wide bandgap semiconductors such as GaN [17], SiC [18] and ZnO [19,20]. The Hg probe has been used as good Schottky contact even in the case of high carrier concentration up to 1019 cm3 [18,19]. The carrier concentrations (n) can be obtained by n ¼
Carrier concentrations (cm-3)
J.D. Ye et al. / Journal of Crystal Growth 283 (2005) 279–285
281
Measured by C-V method 10
20
10
19
10
18
10
17
10
16
10
15
0
1
2
3
4
5
6
Ga/Zn mole ratio (%) Fig. 2. Room temperature carrier concentrations of Ga-doped ZnO films measured by the C– V method as a function of Ga/ Zn gas mole ratio
2=q ½dð1=C 2 Þ=dV [21], where q is the electron charge, e is the relative dielectric constant with the value of 8.5 for ZnO material, C is the barrier capacitance per unit area, and V is the applied voltage, respectively. In our experiments, C2V characteristics show that the capacitance increased under the positive bias voltage, indicative of n-type conduction and the curves of 1/C2 versus voltage are quasi-linear. The electron concentrations are derived from the gradient of 1/C2 versus voltage, as shown in Fig. 2. The undoped film shows high resistance with a low carrier concentration of 3.3 1015 cm3. The carrier concentration initially increased sharply, reached a maximum of 2.47 1019 cm3 with the Ga/Zn gas mole ratio of 3.2 at%. It can be explained that, due to the similar atomic radii, Ga atoms easily substituted the host zinc atoms and ionized into Ga3+, contributing a free electron, and also, the similarity of electronegativities results in the localization effect of conduction electrons due to gallium to be small [11]. However, it is surprising that the amplitude of variation in carrier concentration is far larger than that of the Ga/Zn mole ratios, which is also observed in the Ga-doped ZnO materials deposited by RF sputtering and spray pyrolysis methods [11,13,22]. Reddy et al. [11] found that the resistivity initially decreased exponentially at Ga doping concentrations smaller
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than 3.8 at%. It may be related to the doping efficiency during the growth process. The doping efficiency is determined by the growth conditions such as growth temperature or VI/II mole ratio [6,23]. Here, Ga/Zn mole ratio is the ratio in the gas phase, rather than the ratio of atomic concentrations in the films. When the doping level is small, the doping efficiency will increase with the Ga/Zn mole ratio, causing the sharp increase of carrier concentration. At higher Ga/Zn ratios, the carrier concentration saturates and decreases at even higher Ga/Zn ratios, which was also observed by Hu et al. [6] in ZnO:Ga layers, due to carrier compensation at high doping level. After a certain level of doping, Ga atoms have a tendency to occupy interstitial sites as neutral defects, and even substitute oxygen site as an acceptor. In addition, the acceptor-like complex defects such as (GaZn–Oi) and (GaZn–VZn) are easily formed in highly Gadoped ZnO. In the case of Ga-doped ZnSe, complexes such as (GaZn–VSe) and (Gai–VZn) have been found and are known to act as compensating native defects, as determined from the results of ion channeling measurements and positron beam analysis [24]. Roberts et al. [25] pointed out that 2 Ga3þ is formed by doubly charged oxygen 2 Oi with two Ga3+, serving as double–electron traps and thus reducing the carrier density. Transmittance measurements were carried out at room temperature, and we found that both undoped and Ga-doped films have a high transmittance of over 90%, indicative of good optical properties. The absorption coefficient is given by a ¼ ðlog ð1=TÞÞð1=dÞ aC [26], where T is the value of transmittance, d is the film thickness, and the quantity aC takes into account reflection losses at the sample and window surfaces. Fig. 3 shows the derived absorption curves for the undoped and Ga-doped films. It is found that the absorption edge of the undoped film is slightly steeper than that of the Ga-doped film, which may be due to the band tailing effect caused by gallium incorporation. In addition, the absorption edge shifts towards a higher energy level with the increase in the dopant concentration. This feature is commonly observed for highly doped films and is known as the Burstein–Moss (BM) shift [27,28]. At a carrier concentration greater than critical
Expt. Fit
Absorption (a.u.)
282
Eabs
3.387eV
3.2 at % n=2.47x1019
Eabs
3. 3 01eV
0.0 at % n=3.3x1015 3.0
3.2
3.4
3.6
3.8
Photon Energy (eV) Fig. 3. The dotted and solid lines are the experimental and fitted values, respectively, of the absorption coefficient a (E) in the region of the direct band gap derived from transmittance spectra through the relation a ¼ ðlogð1=TÞÞð1=dÞ aC . Indicated by arrows are the fundamental absorption edges obtained.
Mott density, i.e., n41019 cm3, the Fermi level shifts above the bottom of the conduction band, and in this case, the absorptive optical transition only occurs from the valence band to the state above the Fermi level. Taking the BM effect into account by the introduction of a Fermi level filling factor (FLFF), the absorption curve can be fitted by the following formula [29]. aðEÞ ¼
A f1 þ exp½ðE abs EÞ=kT gE " R0 G0 ðE 0 R0 EÞ2 þ G20 # Z þ1 1 G0 dE 0 þ 1 exp½2pzðE 0 Þ ðE E 0 Þ2 þ G20 E0
where E is the photon energy, the term f1 þ exp½ðE abs EÞ=kT g is the FLFF, Eabs is the absorption edge, G0 is the broadening parameter for both the discrete and continuum exciton terms, zðEÞ ¼ ½R0 =ðE E 0 Þ1=2 , R0 is the exciton binding energy (59 meV). The parameter kT is adjustable so as to take into account the inhomogeneous
ARTICLE IN PRESS J.D. Ye et al. / Journal of Crystal Growth 283 (2005) 279–285
broadening effect. The solid lines of Fig. 3 are the results of this fitting scheme, in surprisingly good agreement with the experimental data over a sufficiently wide energy region. The absorption edge energy is evaluated to be 3.301 and 3.387 eV for the undoped and Ga-doped films, respectively, demonstrating an expected BM blueshift. Considering the nonparabolic nature of conduction band and bandgap renormalization (BGR) effects, the absorption edge energy shift DE abs ðnÞ is found to follow an approximate functional dependence of n2=3 . The comparison of theoretical calculation and experimental data has been shown elsewhere in detail [30]. Fig. 4 shows room-temperature PL spectra of the undoped layer and a Ga-doped ZnO layer with Ga/Zn mole ratio of 6.3 at%. Both spectra were dominated by the near-band-edge (NBE) optical transition around 3.29 eV while deep-level emissions were negligibly weak. However, the intensity of NBE decreased with increasing Ga incorporation in the ZnO layer, which is also observed in Sidoped AlN layers [31]. These facts imply that Ga doping does not contribute in creating defects responsible for deep-level emission, but it produces local strain around Ga atoms or nonradiative channels associated with Ga impurities, annihilating the exciton radiative transition and thus limiting the radiative efficiency of NBE [6]. Meanwhile, the linewidth of NBE increased
283
gradually with the Ga doping concentration, as shown in Figs. 4 and 5(a). One of the potential mechanisms responsible for broadening linewidth of PL is the BM effect. Due to indirect transitions (violating the k selection rule) between filled states in the conduction band and valance band, the width of PL line was found empirically to increase with the energy shift [32]. The indirect character of such transitions was attributed to scattering of carriers by ionized donors and to Auger processes [33]. The other main mechanism is potential fluctuations caused by the random distribution of doping impurities [34]. Randomly distributed impurities lead to unavoidable fluctuations of the doping concentration on a microscopic scale, and thus give rise to potential fluctuation or tail states of band edges, resulting in broadening of the luminescence line. Also, higher compensation always leads to greater potential fluctuation for a fixed carrier concentration. This mechanism has been elucidated in detail for heavily Ga-doped ZnO [35] and Si-doped GaN layers [36]. An additional systematic study is necessary to elucidate which mechanism is mainly responsible for the broadening of the NBE. Apart from broadening due to deliberately doped Ga donors, the residual acceptors also contribute to the luminescence broadening. Fig. 5(a) shows the NBE transition of samples with different Ga dopant concentrations. The peak
Room Temp.
PL Intensity (a.u.)
PL Intensity (a.u.)
ZnO: Ga 6.3 at %
11 10 9 8 7 6 5 4 3 2 1 0
n-type ZnO
(a)
2.5
3.0
3.5
Photon Energy (eV) Fig. 4. Room temperature photoluminescence spectra of the undoped and Ga-doped (6.3 at%) films.
3.30 3.29 3.28
0.0 at % Exp. 1/3 Fit: ∆EPL =-Kn
0.9 at %
3.27
-5
3.2 at %
(K=1.3X10 meV cm)
3.26
6.3 at %
3.0 3.2 3.4 3.6
2.0
(b)
Photon Energy (eV)
E PL (eV)
Undoped ZnO
15
10
16
10
17
10
18
10
19
10
n (cm-3 )
Fig. 5. (a) Room temperature photoluminescence of NBE transition for ZnO-doped at different Ga doping concentration and (b) the dependence of peak energy positions (EPL) on the carrier concentration.
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position of the NBE transition (EPL) shifted from 3.298 to 3.260 eV as Ga/Zn ratio increases from 0.0 to 3.2 at%, and then increased to 3.288 eV at a higher ratio of 6.3 at%. It is found that EPL shifts to low energy monotonously with increasing carrier concentration rather than the dopant concentration, as shown in Fig. 5(b). The redshift of NBE with carrier concentration can be interpreted as doping induced BGR effect, which is also reported in Si-doped GaN and AlN materials [31,36]. This effect is usually due to free carrier screening and the bandgap shrinkage is given by DE PL ¼ Kn1=3 , where n is the carrier concentration and K is BGR coefficients depending on the material. As shown in Fig. 5(b), our experimental data are fitted well by the relation. The BGR coefficients (K) evaluated by the least-squares fitting was 1.3 105 meV cm. The present K value is in good agreement with the value of ZnO derived from GaN material by Reynolds et al. [37] and somewhat weaker than that for GaN (2.4 105 meV cm) and GaAs (1.6 105 meV cm) due to the difference in dielectric constant and conduction band effective mass.
4. Conclusions Single-crystal Ga-doped ZnO films with high structural quality, high conductivity and optical transmittancy have been epitaxially deposited by the MOCVD technique. High Ga doping level causes a small deformation of the lattice and only a slight degradation in crystal structure, indicating that Ga is a good dopant for ntype ZnO. The BM shift of the absorption edge energy increases with the carrier concentration up to 2.47 1019 cm3. The near-bandgap emission has a monotonic redshift with increasing free carrier concentration, which is attributed to the BGR effect, and can be fitted by an n1=3 power law with a BGR coefficient of 1.3 105 meV cm. The carrier concentration increased sharply and became saturated at a higher doping level due to the onset of carrier compensation commonly observed in heavily doped semiconductors.
Acknowledgments This research was supported by special funds from the Major State Basic Research Project of China (Project no. G001CB3095), National Natural Science Foundation of China (Project no. 60276011 and Project no. 60390073) and Project of High Technology Research & Development of China (Project no. 2002AA311060).
References [1] A. Tsukazaki, A. Ohtomo, T. Onuma, et al., Nat. Mater. 4 (2005) 42. [2] D.C. Look, Mater. Sci. Eng. B 80 (2001) 383. [3] D.C. Look, B. Claflin, Phys. Status Solidi B 241 (2004) 624. [4] S.J. Pearton, D.P. Norton, K. Ip, et al., Prog. Mater. Sci. 50 (2005) 293. [5] S.J. Pearton, Y.W. Heo, M. Ivill, et al., Semicond. Sci. Technol. 19 (2004) R59. [6] J. Hu, R.G. Gordon, J. Appl. Phys. 72 (1992) 5381. [7] A.E. Delahoy, M. Cherny, Mater. Res. Soc. Symp. Proc. 426 (1996) 467. [8] W.J. Jeong, G.C. Park, Sol. Energy Mater. Sol. Cell 65 (2001) 37. [9] H. Kato, M. Sano, K. Miyamoto, et al., J. Crystal Growth 237–239 (2002) 538. [10] H.J. Ko, Y.F. Chen, S.K. Hong, et al., Appl. Phys. Lett. 77 (2000) 3761. [11] K.T.R. Reddy, H. Gopalaswamy, P.J. Reddy, et al., J. Crystal Growth 210 (2000) 516. [12] Y. Li, G.S. Tompa, S. Liang, et al., J. Vac. Sci. Technol. A 15 (1997) 1063. [13] G.A. Hirata, J. Mckittrick, J. Siqueiros, et al., J. Vac. Sci. Technol. A 14 (1996) 791. [14] K.Y. Cheong, N. Muti, S.R. Ramanan, Thin Solid Films 410 (2000) 142. [15] J. Zhong, S. Muthukumar, Y. Chen, et al., Appl. Phys. Lett. 83 (2003) 3401. [16] K. Ohkawa, T. Mitsuyu, O. Yamazaki, J. Appl. Phys. 62 (1987) 3216. [17] C.E. Stutz, M. Mack, M.D. Bremser, et al., J. Electron. Mater. 27 (1998) L26. [18] C. Sartel, V. Souliere, M. Zielinski, et al., Mater. Sci. Forum 483 (2005) 121. [19] A. dadgar, N. Oleynik, J. Blasing, et al., J. Crystal Growth 272 (2004) 800. [20] X. Tang, A. Clauzonnier, H.I. Campbell, et al., Appl. Phys. Lett. 84 (2004) 3043. [21] P. Blood, Semicond. Sci. Technol. 1 (1996) 7. [22] M. Miyazaki, K. Sato, A. Mitsui, et al., J. Non-Cryst. Solids 218 (1997) 323.
ARTICLE IN PRESS J.D. Ye et al. / Journal of Crystal Growth 283 (2005) 279–285 [23] P. Kozodoy, S. Keller, S.P. DenBaars, et al., J. Crystal Growth 195 (1998) 265. [24] T. Hino, T. Haga, Y. Abe, et al., Appl. Phys. Lett. 82 (1997) 1196. [25] N. Roberts, R.P. Wang, A.W. Sleight, et al., Phys. Rev. B 57 (1998) 5734. [26] A.R. Goni, A. Cantarero, K. Syassen, et al., Phys. Rev. B 41 (1990) 10111. [27] E. Burstein, Phys. Rev. 93 (1954) 632. [28] T.S. Moss, Proc. Phys. Soc. London Sect. B 55 (1954) 775. [29] M. Munoz, F.H. Pollak, M. Kahn, et al., Phys. Rev. B 63 (2001) 233302. [30] J.D. Ye, S.L. Gu, S.M. Zhu, et al., Appl. Phys. Lett. 86 (2005) 192111.
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[31] K.B. Nam, M.L. Nakarmi, J. Li, et al., Appl. Phys. Lett. 83 (2003) 2787. [32] E. Lliopoulos, D. Doppalapudi, H.M. Ng, et al., Appl. Phys. Lett. 73 (1998) 375. [33] S.M. Ryvkin, Phys. Status Solidi 11 (1965) 285. [34] E.F. Schubert, I.D. Goepfert, W. Grieshaber, et al., Appl. Phys. Lett. 71 (1997) 921. [35] T. Makino, Y. Segawa, S. Yoshida, et al., Appl. Phys. Lett. 85 (2004) 759. [36] I. Lee, J.J. Lee, P. Kung, et al., Appl. Phys. Lett. 74 (1999) 102. [37] D.C. Reynolds, D.C. Look, B. Hogai, J. Appl. Phys. 88 (2000) 5760.
Journal of Crystal Growth 283 (2005) 279–285 www.elsevier.com/locate/jcrysgro
Gallium doping dependence of single-crystal n-type Zno grown by metal organic chemical vapor deposition J.D. Yea, S.L. Gua,, S.M. Zhua, S.M. Liua, Y.D. Zhenga, R. Zhanga, Y. Shia, H.Q. Yub, Y.D. Yeb a
Key Laboratory of Advanced Photonic and Electronic Materials and Department of Physics, Nanjing University, Nanjing 210093, PR China b Materials Analysis Center of Nanjing University, Nanjing 210093, PR China Received 3 May 2005; received in revised form 30 May 2005; accepted 1 June 2005 Available online 26 July 2005 Communicated by R. Bhat
Abstract High-quality single-crystal Ga-doped ZnO films have been epitaxially deposited on (0 0 0 2) sapphire substrate by low-pressure metal organic chemical vapor deposition (MOCVD) technique. The dependence of structural, electrical and optical properties of films on Ga doping concentration was investigated. As grown at the Ga/Zn gas ratio of 3.2 at%, the film shows a narrow linewidth of 0.261 for ZnO (0 0 0 2) peak, high carrier concentration of 2.47 1019 cm3, and high optical transparency over 90%. The carrier concentration increased sharply and became saturated at higher doping level due to the onset of carrier compensation. The Burstein–Moss blueshift of the absorption edge energy increased as expected with the carrier concentration up to 2.47 1019 cm3. In addition, doping-induced photoluminescence (PL) emission linewidth broadening and bandgap renormalization (BGR) effects have also been observed. The intensity of PL emission decreased with increasing Ga dopant concentration, which was believed to be a direct consequence of the doping-enhanced nonradiative recombination rates. r 2005 Elsevier B.V. All rights reserved. Keywords: A3. Metal organic chemical vapor deposition; B1. Gallium doping; B1. ZnO
1. Introduction Zinc oxide, a wide-gap semiconductor with a high exciton binding energy (60 meV), has emerged Corresponding author.
E-mail addresses: [email protected] (J.D. Ye), [email protected] (S.L. Gu).
as a candidate for use in short-wavelength optoelectronic devices such as ultraviolet (UV) lasers and light-emitting diodes [1–5]. In addition, due to its high thermal and chemical stability, good electrical conductivity and high optical transparency, ZnO has also attracted more attention in the applications of display devices, solar cells and electronic transducers [4–8]. In general,
0022-0248/$ - see front matter r 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2005.06.030
ARTICLE IN PRESS 280
J.D. Ye et al. / Journal of Crystal Growth 283 (2005) 279–285
undoped ZnO films always exhibit n-type conduction with typical carrier concentration of 1017 cm3, which, however, is far from the demands of device applications. Thus, high conductivity controlled by intentional doping is necessary. Group III elements (Al, Ga and In), in comparison with group VII elements (Cl, Br) are more suitable for n-type doping of ZnO materials because of their lower vapor pressure [9]. ZnO: Al films have been studied widely; however, oxidation of the Al source during ZnO growth owing to the high reactivity of Al is a problem. The covalent bond length of Ga–O (0.192 nm) is slightly smaller than that of Zn–O (0.197 nm), which will make the deformation of the ZnO lattice small even in the case of high Ga doping concentration [10]. Thus, Ga is an attractive dopant for n-type ZnO. Polycrystalline ZnO:Ga thin films have been achieved by many methods [11–14]; however, few reports have so far focused on the control of carrier concentration in single-crystal Ga-doped ZnO epilayers. Single-crystal ZnO:Ga films with carrier concentration up to 1.13 1020 cm3 was first achieved on GaN substrate via molecular beam epitaxy (MBE) [10]. Recently, Zhong et al. [15] reported the MOCVD growth of Ga-doped ZnO single-crystal nanotips with a low resistivity of 4 103 O cm, where the Ga doping mechanism was not studied. In the present study, we report the successful epitaxy of single-crystal Ga-doped ZnO films with carrier concentration up to 2.47 1019 cm3 by low-pressure MOCVD technique. The dependence of structural, electrical and optical properties of ZnO on Ga doping concentration and the associated doping mechanisms have also been studied.
2. Experiments The undoped and Ga-doped ZnO films were uniformly deposited on a sapphire (0 0 0 1) substrate in a vertical, homemade, low-pressure MOCVD reactor. Prior to entering the reactor, 2 in sapphire substrates were degreased in acetone, ethanol and then chemically etched in H2SO4 at 200 1C for 10 min followed by deionized water
rinsing and nitrogen flow drying. During the deposition process, diethylzinc (DEZn), 5 N-purity oxygen, and triisopropylgallium (TIPGa) were used as Zn, O and Ga sources, respectively. Argon gas of 6 N-purity was used as metalorganic (MO) carrier gas. The Ga-doped ZnO thin films were deposited on an undoped ZnO buffer layer to improve the crystallinity of ZnO:Ga layers. The Ga/Zn mole ratio was modulated from 0.0 to 6.3 at%, by controlling the temperature of TIPGa MO source. The growth temperature was 450 1C and the thickness of ZnO:Ga epilayers is approximately 300 nm. In our study, no obvious variation of thickness was observed for different doping concentrations. The structure and crystallinity of samples were analyzed by X-ray diffraction (XRD). Carrier concentrations were evaluated by capacitance–voltage (C2V ) method at a frequency of 300 kHz using a mercury probe as Schottky contact. Transmittance spectra were recorded using an UV/visible spectrometer (U-3410, Hitachi) with a resolution of 1 nm. Photoluminescence (PL) spectra were recorded using He–Cd laser (325 nm) as the excitation source with an incident power density of 40 mW/cm2. All measurements are carried out at room temperature.
3. Results and discussion Fig. 1 shows the typical XRD curve of a ZnO:Ga thin film with a Ga/Zn mole ratio of 3.2 at% and the inset shows the comparison of the (0 0 0 2) peak for the undoped and Ga-doped ZnO thin films. The Ga-doped thin film exhibited (0 0 0 2) and (0 0 0 4) peaks at 2y ¼ 34.561 and 72.881 for wurtzite structure ZnO, indicative of preferential orientation with the c-axis perpendicular to the substrate. Besides these peaks, no other diffraction peaks were observed, that is to say, there was no second phase corresponding to gallium or its oxides even at a high Ga/Zn mole ratio of 3.2 at% during film growth. In addition, the small full-width at half-maximum (FWHM) of the (0 0 0 2) peak (0.261) confirms the single-crystal quality of Ga-doped ZnO film. However, as shown in the inset of Fig. 1, the (0 0 0 2) peak for
ARTICLE IN PRESS
undoped ZnO ZnO:Ga 3.2%
Al2O3 (0006)
ZnO (0002)
XRD Intensity (a.u.)
ZnO:Ga 3.2%
33.6
34.0
34.4
34.8
35.2
35.6
2θ (deg)
ZnO (0004)
30
40
50 2θ (deg)
60
70
Fig. 1. X-ray diffractions of ZnO:Ga thin film with Ga/Zn mole ratio of 3.2%. The inset figure shows the blown-up ZnO (0 0 0 2) peak of undoped and Ga-doped (3.2%) ZnO thin films.
Ga-doped ZnO layer broadened slightly and shifted to 34.561, higher than that of the undoped film (34.421), indicative of the decreasing lattice constant along the c-axis due to Ga incorporation. These phenomena generally exist in impuritydoped crystals and can be attributed to an increase in local strain around impurity atoms or point defect associated with the impurity atoms [16]. In fact, the covalent bond length of Ga–O is slightly smaller than that of Zn–O [10], and Ga substituted at Zn sites would reduce the lattice constant along the c-axis, and induce an uniaxial compressive strain in ZnO films, which is responsible for the shift of peak (0 0 0 2). It should also be noted that the mismatch of Ga–O and Zn–O, which is as small as 3%, is expected to make the deformation of the lattice small and cause only a slight degradation in crystal structure even in the case of high Ga doping concentrations. Fig. 2 shows room temperature carrier concentrations of Ga-doped ZnO films as a function of Ga/Zn gas mole ratio. C– V method is an alternative characterization technique that has been used extensively on wide bandgap semiconductors such as GaN [17], SiC [18] and ZnO [19,20]. The Hg probe has been used as good Schottky contact even in the case of high carrier concentration up to 1019 cm3 [18,19]. The carrier concentrations (n) can be obtained by n ¼
Carrier concentrations (cm-3)
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281
Measured by C-V method 10
20
10
19
10
18
10
17
10
16
10
15
0
1
2
3
4
5
6
Ga/Zn mole ratio (%) Fig. 2. Room temperature carrier concentrations of Ga-doped ZnO films measured by the C– V method as a function of Ga/ Zn gas mole ratio
2=q ½dð1=C 2 Þ=dV [21], where q is the electron charge, e is the relative dielectric constant with the value of 8.5 for ZnO material, C is the barrier capacitance per unit area, and V is the applied voltage, respectively. In our experiments, C2V characteristics show that the capacitance increased under the positive bias voltage, indicative of n-type conduction and the curves of 1/C2 versus voltage are quasi-linear. The electron concentrations are derived from the gradient of 1/C2 versus voltage, as shown in Fig. 2. The undoped film shows high resistance with a low carrier concentration of 3.3 1015 cm3. The carrier concentration initially increased sharply, reached a maximum of 2.47 1019 cm3 with the Ga/Zn gas mole ratio of 3.2 at%. It can be explained that, due to the similar atomic radii, Ga atoms easily substituted the host zinc atoms and ionized into Ga3+, contributing a free electron, and also, the similarity of electronegativities results in the localization effect of conduction electrons due to gallium to be small [11]. However, it is surprising that the amplitude of variation in carrier concentration is far larger than that of the Ga/Zn mole ratios, which is also observed in the Ga-doped ZnO materials deposited by RF sputtering and spray pyrolysis methods [11,13,22]. Reddy et al. [11] found that the resistivity initially decreased exponentially at Ga doping concentrations smaller
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than 3.8 at%. It may be related to the doping efficiency during the growth process. The doping efficiency is determined by the growth conditions such as growth temperature or VI/II mole ratio [6,23]. Here, Ga/Zn mole ratio is the ratio in the gas phase, rather than the ratio of atomic concentrations in the films. When the doping level is small, the doping efficiency will increase with the Ga/Zn mole ratio, causing the sharp increase of carrier concentration. At higher Ga/Zn ratios, the carrier concentration saturates and decreases at even higher Ga/Zn ratios, which was also observed by Hu et al. [6] in ZnO:Ga layers, due to carrier compensation at high doping level. After a certain level of doping, Ga atoms have a tendency to occupy interstitial sites as neutral defects, and even substitute oxygen site as an acceptor. In addition, the acceptor-like complex defects such as (GaZn–Oi) and (GaZn–VZn) are easily formed in highly Gadoped ZnO. In the case of Ga-doped ZnSe, complexes such as (GaZn–VSe) and (Gai–VZn) have been found and are known to act as compensating native defects, as determined from the results of ion channeling measurements and positron beam analysis [24]. Roberts et al. [25] pointed out that 2 Ga3þ is formed by doubly charged oxygen 2 Oi with two Ga3+, serving as double–electron traps and thus reducing the carrier density. Transmittance measurements were carried out at room temperature, and we found that both undoped and Ga-doped films have a high transmittance of over 90%, indicative of good optical properties. The absorption coefficient is given by a ¼ ðlog ð1=TÞÞð1=dÞ aC [26], where T is the value of transmittance, d is the film thickness, and the quantity aC takes into account reflection losses at the sample and window surfaces. Fig. 3 shows the derived absorption curves for the undoped and Ga-doped films. It is found that the absorption edge of the undoped film is slightly steeper than that of the Ga-doped film, which may be due to the band tailing effect caused by gallium incorporation. In addition, the absorption edge shifts towards a higher energy level with the increase in the dopant concentration. This feature is commonly observed for highly doped films and is known as the Burstein–Moss (BM) shift [27,28]. At a carrier concentration greater than critical
Expt. Fit
Absorption (a.u.)
282
Eabs
3.387eV
3.2 at % n=2.47x1019
Eabs
3. 3 01eV
0.0 at % n=3.3x1015 3.0
3.2
3.4
3.6
3.8
Photon Energy (eV) Fig. 3. The dotted and solid lines are the experimental and fitted values, respectively, of the absorption coefficient a (E) in the region of the direct band gap derived from transmittance spectra through the relation a ¼ ðlogð1=TÞÞð1=dÞ aC . Indicated by arrows are the fundamental absorption edges obtained.
Mott density, i.e., n41019 cm3, the Fermi level shifts above the bottom of the conduction band, and in this case, the absorptive optical transition only occurs from the valence band to the state above the Fermi level. Taking the BM effect into account by the introduction of a Fermi level filling factor (FLFF), the absorption curve can be fitted by the following formula [29]. aðEÞ ¼
A f1 þ exp½ðE abs EÞ=kT gE " R0 G0 ðE 0 R0 EÞ2 þ G20 # Z þ1 1 G0 dE 0 þ 1 exp½2pzðE 0 Þ ðE E 0 Þ2 þ G20 E0
where E is the photon energy, the term f1 þ exp½ðE abs EÞ=kT g is the FLFF, Eabs is the absorption edge, G0 is the broadening parameter for both the discrete and continuum exciton terms, zðEÞ ¼ ½R0 =ðE E 0 Þ1=2 , R0 is the exciton binding energy (59 meV). The parameter kT is adjustable so as to take into account the inhomogeneous
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broadening effect. The solid lines of Fig. 3 are the results of this fitting scheme, in surprisingly good agreement with the experimental data over a sufficiently wide energy region. The absorption edge energy is evaluated to be 3.301 and 3.387 eV for the undoped and Ga-doped films, respectively, demonstrating an expected BM blueshift. Considering the nonparabolic nature of conduction band and bandgap renormalization (BGR) effects, the absorption edge energy shift DE abs ðnÞ is found to follow an approximate functional dependence of n2=3 . The comparison of theoretical calculation and experimental data has been shown elsewhere in detail [30]. Fig. 4 shows room-temperature PL spectra of the undoped layer and a Ga-doped ZnO layer with Ga/Zn mole ratio of 6.3 at%. Both spectra were dominated by the near-band-edge (NBE) optical transition around 3.29 eV while deep-level emissions were negligibly weak. However, the intensity of NBE decreased with increasing Ga incorporation in the ZnO layer, which is also observed in Sidoped AlN layers [31]. These facts imply that Ga doping does not contribute in creating defects responsible for deep-level emission, but it produces local strain around Ga atoms or nonradiative channels associated with Ga impurities, annihilating the exciton radiative transition and thus limiting the radiative efficiency of NBE [6]. Meanwhile, the linewidth of NBE increased
283
gradually with the Ga doping concentration, as shown in Figs. 4 and 5(a). One of the potential mechanisms responsible for broadening linewidth of PL is the BM effect. Due to indirect transitions (violating the k selection rule) between filled states in the conduction band and valance band, the width of PL line was found empirically to increase with the energy shift [32]. The indirect character of such transitions was attributed to scattering of carriers by ionized donors and to Auger processes [33]. The other main mechanism is potential fluctuations caused by the random distribution of doping impurities [34]. Randomly distributed impurities lead to unavoidable fluctuations of the doping concentration on a microscopic scale, and thus give rise to potential fluctuation or tail states of band edges, resulting in broadening of the luminescence line. Also, higher compensation always leads to greater potential fluctuation for a fixed carrier concentration. This mechanism has been elucidated in detail for heavily Ga-doped ZnO [35] and Si-doped GaN layers [36]. An additional systematic study is necessary to elucidate which mechanism is mainly responsible for the broadening of the NBE. Apart from broadening due to deliberately doped Ga donors, the residual acceptors also contribute to the luminescence broadening. Fig. 5(a) shows the NBE transition of samples with different Ga dopant concentrations. The peak
Room Temp.
PL Intensity (a.u.)
PL Intensity (a.u.)
ZnO: Ga 6.3 at %
11 10 9 8 7 6 5 4 3 2 1 0
n-type ZnO
(a)
2.5
3.0
3.5
Photon Energy (eV) Fig. 4. Room temperature photoluminescence spectra of the undoped and Ga-doped (6.3 at%) films.
3.30 3.29 3.28
0.0 at % Exp. 1/3 Fit: ∆EPL =-Kn
0.9 at %
3.27
-5
3.2 at %
(K=1.3X10 meV cm)
3.26
6.3 at %
3.0 3.2 3.4 3.6
2.0
(b)
Photon Energy (eV)
E PL (eV)
Undoped ZnO
15
10
16
10
17
10
18
10
19
10
n (cm-3 )
Fig. 5. (a) Room temperature photoluminescence of NBE transition for ZnO-doped at different Ga doping concentration and (b) the dependence of peak energy positions (EPL) on the carrier concentration.
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position of the NBE transition (EPL) shifted from 3.298 to 3.260 eV as Ga/Zn ratio increases from 0.0 to 3.2 at%, and then increased to 3.288 eV at a higher ratio of 6.3 at%. It is found that EPL shifts to low energy monotonously with increasing carrier concentration rather than the dopant concentration, as shown in Fig. 5(b). The redshift of NBE with carrier concentration can be interpreted as doping induced BGR effect, which is also reported in Si-doped GaN and AlN materials [31,36]. This effect is usually due to free carrier screening and the bandgap shrinkage is given by DE PL ¼ Kn1=3 , where n is the carrier concentration and K is BGR coefficients depending on the material. As shown in Fig. 5(b), our experimental data are fitted well by the relation. The BGR coefficients (K) evaluated by the least-squares fitting was 1.3 105 meV cm. The present K value is in good agreement with the value of ZnO derived from GaN material by Reynolds et al. [37] and somewhat weaker than that for GaN (2.4 105 meV cm) and GaAs (1.6 105 meV cm) due to the difference in dielectric constant and conduction band effective mass.
4. Conclusions Single-crystal Ga-doped ZnO films with high structural quality, high conductivity and optical transmittancy have been epitaxially deposited by the MOCVD technique. High Ga doping level causes a small deformation of the lattice and only a slight degradation in crystal structure, indicating that Ga is a good dopant for ntype ZnO. The BM shift of the absorption edge energy increases with the carrier concentration up to 2.47 1019 cm3. The near-bandgap emission has a monotonic redshift with increasing free carrier concentration, which is attributed to the BGR effect, and can be fitted by an n1=3 power law with a BGR coefficient of 1.3 105 meV cm. The carrier concentration increased sharply and became saturated at a higher doping level due to the onset of carrier compensation commonly observed in heavily doped semiconductors.
Acknowledgments This research was supported by special funds from the Major State Basic Research Project of China (Project no. G001CB3095), National Natural Science Foundation of China (Project no. 60276011 and Project no. 60390073) and Project of High Technology Research & Development of China (Project no. 2002AA311060).
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