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Evidence and analysis of parallel growth mechanisms in Cu2O films prepared by Cu anodization
Electrochimica Acta 55 (2010) 4353–4358
Contents lists available at ScienceDirect
Electrochimica Acta journal homepage: www.elsevier.com/locate/electacta
Evidence and analysis of parallel growth mechanisms in Cu2 O films prepared by Cu anodization F. Caballero-Briones a,b,∗,1 , A. Palacios-Padrós a , O. Calzadilla c , Fausto Sanz d,a,b,∗∗ a
Department of Physical Chemistry, Universitat de Barcelona, Martí i Franquès 1, 08028 Barcelona, Spain CIBER-BBN, María de Luna 11, 50018 Zaragoza, Spain Facultad de Física, Universidad de La Habana, San Lázaro y L, Colina Universitaria, 10400 Vedado, La Habana, Cuba d Institute for Bioengineering of Catalonia (IBEC), Edifici Hèlix, Baldiri i Reixac 15-21, 08028 Barcelona, Spain b c
a r t i c l e
i n f o
Article history: Received 29 July 2009 Received in revised form 8 September 2009 Accepted 8 October 2009 Available online 20 October 2009 Keywords: Cuprous oxide Anodic films Reflectance Thickness Band gap Urbach tail parameter Dissolution Growth mechanism
a b s t r a c t We have studied the preparation of Cu2 O films by copper anodization in a 0.1 M NaOH electrolyte. We identified the potential range at which Cu+ dissolution takes place then we prepared films with different times of exposure to this potential. The morphology, crystalline structure, band gap, Urbach energy and thickness of the films were studied. Films prepared with the electrode unexposed to the dissolution potential have a pyramidal growth typical of potential driven processes, while samples prepared at increasing exposure times to dissolution potential present continuous nucleation, growth and grain coalescence. We observed a discrepancy in the respective film thicknesses calculated by coulometry, atomic force microscopy and optical reflectance. We propose that anodic Cu2 O film formation involves three parallel mechanisms (i) Cu2 O nucleation at the surface, (ii) Cu+ dissolution followed by heterogeneous nucleation and (iii) Cu+ and OH− diffusion through the forming oxide and subsequent reaction in the solid state. © 2009 Elsevier Ltd. All rights reserved.
1. Introduction Cu2 O is a p-type semiconductor with a band gap around 2 eV that has recently aroused interest because of its potential applications in random access memories [1], low-cost solar cells [2], and gas sensors [3] among other devices. The electrochemical fabrication of Cu2 O has been usually reported by electrodeposition onto transparent conductive oxides [2,4–7]. Copper electrochemistry and Cu2 O growth in alkaline media have been extensively studied [8–12]: the main anodic peak observed around U = −400 mV (vs silver/silver chloride electrode, henceforth SSC) in the potentiodynamic curves is associated to the formation of Cu2 O. Additionally, at potentials negative with respect to Cu2 O formation, the development of Cu+ hydrous species and OH− adsorption has been described. Copper (I) soluble species,
∗ Corresponding author at: Department of Physical Chemistry, Universitat de Barcelona, Martí i Franquès 1, 08028 Barcelona, Spain. Tel.: +34 934039239; fax: +34 934039239. ∗∗ Corresponding author at: Department of Physical Chemistry, Universitat de Barcelona, Martí i Franquès 1, 08028 Barcelona, Spain. E-mail addresses: [email protected] (F. Caballero-Briones), [email protected] (F. Sanz). 1 On leave from: CICATA-IPN Altamira, Km 14.5 Carr. Tampico-Pto Industrial, 89600 Altamira, Mexico. 0013-4686/$ – see front matter © 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.electacta.2009.10.031
namely (Cu2 O2 H)− aq can also be produced. The Cu2 O growth starts from the reconstructed OH-populated surface [11]. Nevertheless only few works focus on the preparation of thick Cu2 O anodic films and although soluble Cu+ formation has been repeatedly suggested, we did not found reports with experimental evidence of copper dissolution and its effects on film properties. In a recent contribution [13] we reported an anodization method of Cu0 in alkaline media that allows the preparation of Cu2 O films as thicker as 100 nm. We studied the electronic properties of the films and proposed an improved electronic band diagram of the Cu|Cu2 O|electrolyte interface. Herewith, a Cu+ dissolution potential range forming Cu+ species in solution before the Cu2 O growth potential range has been identified by current enhancement with the addition of a complexing agent for Cu+ and then, several films with different exposure times to the dissolution potential have been prepared. The contribution of this exposure to the film properties is analyzed in terms of optical, microstructural and thickness changes. The band gap and Urbach parameter were modulated as achieved. Moreover, the consideration of this dissolution effect brings new information on the growth mechanism of Cu2 O films. 2. Experimental details Cu2 O layers were prepared onto polycrystalline Cu disks in a 0.1 M NaOH electrolyte by an electrochemical routine described
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Fig. 1. Potentiodynamic response of the Cu electrode in the 0.1 M NaOH electrolyte. Inset depicts the enhanced current response in the dissolution peak at around U = −450 mV, when NH4 NO3 was added.
elsewhere [13]. Briefly, the electrochemical procedure involves a series of potential steps that lead to different processes, i.e. the reduction of the native oxide, the formation of an OH-adsorbed submonolayer, a dissolution stage and oxide formation. The growth was performed using a PGSTAT 12 Autolab potentiostat in the three electrode configuration, using a homemade glass cell with
the Cu disk as working electrode placed at the bottom of the cell, a platinum coil as auxiliary electrode and a true reference SSC electrode. Potentiodynamic curves were used to monitor the electrode state and to identify the working potentials. In Fig. 1 is shown the cyclic voltammetry of the Cu electrode in the 0.1 M NaOH electrolyte where the main oxidation peak at U = −400 mV is observed. At around U = −550 mV a shoulder that corresponds to the OH-adsorption processes is distinguished. To characterize the dissolution potential, 500 l of a 0.01 M ammonia nitrate complexing solution was added to the electrolyte solution in order to enhance the current response due to Cu+ dissolution, as observed in the inset of Fig. 1, featuring the Cu+ dissolution peak at around U = −450 mV. The Cu2 O films have been prepared by exposing the electrode to the dissolution potential, Ud = −450 mV, during selected times (td = 0 s, td = 50 s, td = 100 s, td = 300 s, td = 500 s and td = 1000 s), and then to the growth potential, Ug = −400 mV, for 60 min. Current versus time transients were also registered to assess the mechanisms of Cu2 O growth. The amount of dissolved Cu+ and film thickness were calculated by the transferred charge at the potentials Ud and Ug respectively. The morphology of the films was studied by atomic force microscopy (AFM) in the intermittent contact mode with a multimode microscope from Veeco with Si cantilevers of 35 N/m spring constants. The roughness (Rrms ) and the average height were obtained from 2 m × 2 m AFM images. Grazing incidence X-ray diffraction (GIXRD) patterns were obtained using a Siemens D500 X-ray diffractometer with the detector pivoting on the 2 circle, with unfiltered Cu K␣ radiation and an incidence angle of 0.5◦ .
Fig. 2. (a) 2 m × 2 m AFM images of the films prepared at different exposition times at the dissolution potential Ud = −450 mV; (b) 500 × 500 nm AFM images of the td = 0 s and td = 1000 s films in the contour representation to better appreciate the nature of the grain formation. Z scales were equalized for image comparison.
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Fig. 3. Roughness and average height of the films, calculated from 2 m × 2 m AFM images. Inset: standard deviation of the roughness and average height as function of time. Continuous lines are only for guideline.
Phases were identified using the PDF database [14]. The diffuse reflectance measurements, from which the thickness, band gap and Urbach tail parameter were calculated, had been performed in a Shimadzu UV-2101 PC Spectrometer with an integrating sphere, using a polished Cu substrate as a reference. The optical properties of the Cu2 O thin films in the spectral range of interest, 300–800 nm, were fitted by three types of electronic excitations: band gap transitions, interband transitions from the bulk of the valence band into the bulk of the conduction band, and intraband transitions of the electrons in the conduction band. These transitions were modeled with standard formulae available in the SCOUT 98 program [15]. For the band gap transitions, we used the Tauc-Lorentz model [16] that has been proposed to model the band gap. The transitions into the upper half of the conduction band are represented by an harmonic oscillator. The Drude Model was used to calculate the contribution of free electrons to the susceptibility. All the procedures used for the fitting are implemented in the SCOUT 98 program. The experimental data were fitted with a standard deviation better than 10−4 . 3. Results 3.1. Morphology Fig. 2a presents 2 m × 2 m AFM images of the different films prepared at several exposure times. In the absence of a dissolution pathway, grains have a pyramidal shape with a wide size distribution. When the electrode is largely exposed to the dissolution potential, the grains evolve to a more irregular shape and grains become more definite and their size increases with the exposure time. Fig. 2b shows in more detail the morphology of the films td = 0 s and td = 1000 s. In the image corresponding to td = 0 s film the pyramids evolving from the bottom of the layer are evident, constituting a typical morphology of a diffusion limited process where planar diffusion hinders some of the nuclei to grow [17,18]; for the td = 1000 s film, the coalescence and grain size distribution observed in the respective image suggest continuous formation of nucleus and grain growth. Fig. 3 shows the morphology parameters, Rrms and average height as function of exposure time at the dissolution potential. They are typical values from at least five images obtained in different zones of the sample. We report the roughness from the 2 m × 2 m AFM images because at this size Rrms does not depend on the length scale used to probe it and the surface can be described with a single number [17]. In Fig. 3 it is observed that both average height and roughness increase steadily from td = 0 s to td = 1000 s. Similar behavior of the roughness values has been observed in
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Fig. 4. Crystallite size in the (1 1 1) Cu2 O direction vs exposure time at the dissolution potential.
chemical bath deposited films at long deposition times [19] and is attributable to increasing film compactness. The values of Rrms of the films prepared with exposure to the dissolution potential, reach a similar value to that of film td = 0 s only after 300 s of exposure. Dispersions of the parameters, taken as the standard deviation, are presented in the inset. Roughness dispersion tends to reduce from film td = 300 s indicating that for longer exposure times, grain growth becomes the dominant mechanism not formation of new nuclei. Correspondingly, height dispersion does not reduce until growth contribution uniforms grain sizes. 3.2. Structure and optical properties The cuprite phase was identified in all the films by X-ray diffraction, without evidence of CuO. The main cuprite peak at a Bragg angle around 2 = 36.4◦ is always observed and only in the td = 0 s and the td = 1000 s films the (2 0 0) Cu2 O reflection becomes slightly noticeable, thus lattice parameters of the prepared samples were not calculated because the lack of other distinct diffraction peaks and due to the indetermination of the Bragg angle in the GIXRD measurements. Even so, the reported bulk Cu2 O lattice parameter is 0.427 nm and that of Cu is 0.361 nm, leading to a lattice mismatch of about 18%. The generated epitaxial stress between the Cu2 O and Cu lattices is released with a tilt of 5◦ of the Cu2 O lattice with respect to that of Cu and a lattice compression of 1% [11,12]. The crystallite size of the films can be estimated from the Scherrer formula [20]: =
K ˇ cos
(1)
Here, = mean crystallite dimension, K = shape factor ∼0.9, = wavelength (1.54 Å), ˇ = line broadening (FWHM), = Bragg angle, taken for this samples as the position of the powder pattern of the (1 1 1) reflection of Cu2 O (2 = 36.417◦ ) due to the indetermination of the peak position in the GIXRD measurement. Fig. 4 plots the crystallite size versus exposure time at the dissolution potential. The crystallite size of the films increases with the exposure time from 9 to 14 nm, and it is only at 300 s that it reaches a similar size (around 12 nm) than when the electrode was not exposed. The increase in the crystallite size would affect the stress in the films, but in the absence of more diffraction peaks it is not possible to apply the Williamson-Hall treatment to determine it. Nevertheless, as it is mentioned below, the stress effect will be included in the disorder analysis quantified by the Urbach tail parameter.
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Fig. 5. Diffuse reflectance spectra from the prepared samples. Inset shows a transformed spectra used for calculation of the band gap energy (detail in the text).
The UV–vis reflectance spectra of the samples are shown in Fig. 5, where it is observed a reduction in the reflectance in the nonabsorption zone, from 650 to 800 nm that contains the information from which the film thickness was extracted (shown in Fig. 8) [15]. The inflection zone where the band gap is located is around 500 nm and below it the intraband transitions can be observed as two peaks around 300 and 350 nm. The inset in Fig. 5 shows the calculation of the direct band gap from the reflectance spectra as proposed by Kumar et al. [21]. For a direct transition the band gap energy (Eg ) is obtained from the linear zone intercept with the energy axis. In the inset, after the absorption zone, the so-called Urbach region manifests itself as an exponential tail where the information of the disorder is enclosed [22]. From the absorption versus energy plot, the fitting of this exponential region below the gap gives the value of the Urbach energy (E0 ). In Fig. 6, Eg and E0 are plotted versus the exposure time at the dissolution potential. The reported bulk value for Cu2 O is Eg = 2.10 eV [23]. For the studied films Eg values are around 2.07–2.10 eV for the films between td = 0 s and td = 300 s and then increases to about 2.20 eV at td = 500 s. The Urbach energy shows a similar behavior for the samples exposed to the dissolution potential, i.e. E0 increases with time, but for t = 0 a high value (ca. 400 meV) is measured. The increase in Eg would be explained by the formation of nanoparticulated Cu2 O [23], but this seems to be in contradiction with the crystallite size determined by GIXRD. This blue shift has
Fig. 6. Eg and E0 plot of the prepared films vs the exposure time at dissolution potential.
Fig. 7. Semilogarithmic plot for the current transients recorded at the growth potential, Ug = −400 mV.
also been observed in electrodeposited Cu2 O films without quantum confinement possibilities and it has been explained by the formation of potential barriers at the surface due to the preparation of films by repeating deposition and oxidation [24]. This explanation is also valid in our case, because it supports the dual nucleation-growth mechanism suggested in the morphology analysis. The detailed analysis of E0 values in Cu2 O films related to the contribution of the different sources of disorder is subject of a more extended contribution (in preparation), nevertheless is interesting to point out some issues to explain the present observations. As mentioned before, the Urbach energy is a sum of several factors of disorder: crystalline structure defects, grain boundaries (directly related with grain size), stress and electronic interactions [22]. Then, film growth mechanisms that lead to a lack of uniformity, strong dispersion of particle sizes and structural stress would direct to increments in the Urbach energy depending, of course, on the deposition conditions. In our case, we observe strong correlation of the increase in roughness and size dispersion with the time of exposure at the dissolution potential. The suggestion of parallel mechanism in anodic Cu2 O film growth gains more evidences with this analysis. 3.3. Growth mechanism In Fig. 7 the current versus time curves registered at the growth potential, Ug = −450 mV, are shown. For the td = 0 s film curve the typical shape of a nucleation process is observed, with the spike minimum around 60 s. The current then increases due to nucleation and growth of Cu2 O to reach the maximum value after 1800 s of growth and then reduce slowly as the diffusion layer expands. Chronoamperometric curves obtained with increasing dissolution times have the spike minimum at the very beginning of the growth time and after the current maximum is reached it tends to decrease rapidly again. This reduction could be caused by electrode conduction drop due to the thicker Cu2 O film grown. It is interesting to observe that the curves of the films with td = 500 s and td = 1000 s present secondary current peaks near the end of the process confirming that concurrent processes are taking place. The film prepared without exposure to the dissolution potential, td = 0, reaches its maximum current value at 1900 s, compared with the one with a dissolution time of 1000 s that reaches it at 36 s. For the case of td = 0 s the growth depends on OH− ion diffusion to the electrode surface. The film td = 1000 represents a completely different situation, an important amount of Cu+ is now present in the electrode vicinity, thus additionally to a rapid and extensive nucle-
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Fig. 8. Film thickness calculated from reflectance measurements, from the charge transfer calculated during the growth potential application and estimated from the average heights from AFM.
ation process, growth caused by heterogeneous nucleation from the incoming Cu+ ions is expected. Fig. 8 presents the comparison between the thicknesses calculated using three different magnitudes: the transferred Faradaic charge, the average height calculated from the AFM images and the thickness calculated from the optical reflectance measurements. There are two striking features: the first one is the huge discrepancy between the thicknesses calculated optically and by coulometry. This is consequence of parallel mechanisms of potentiostatic growth and chemical deposition; as mentioned above, the first one depends only on OH− diffusion to the electrode, while dissolved Cu+ is readily available in the second. The second striking feature is that the sample prepared without exposition to dissolution potential, td = 0, results much thinner by the charge transfer calculation than optically. The calculated “charge transfer thickness” for td = 0 fairly corresponds to the AFM value, accounting for the “external” growth. This second feature evidences that an important amount of the Cu2 O film grows underneath the electrode surface by another concurrent mechanism. We propose that this mechanism arises from diffusion of oxygen and copper species
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through the growing oxide layer that subsequently react in the solid state. Following the work of Krishnamurthy et al. [25] we propose this mechanism to be potential driven and the copper and oxygen species could be copper vacancies and oxygen interstitials as they are reported to be the main defects in the bulk Cu2 O [26]. It is expected that such a mechanism would cause an important amount of structural stress and electronic defects. This would explain the abnormally high E0 observed in the td = 0 film. On the other side, the reduction in thickness observed at the higher times could be explained because the electric field reduction that would reduce ion migration through the oxide layer. Nevertheless, we believe that more experiments are still needed to clarify the identity of the diffusing species and their role in the charge transfer, to better understand the kinetics of film growth. With the aforementioned evidences of morphology, current transient analysis and thickness, we propose that additionally to the electrochemical process for Cu2 O formation and the diffusionsolid state reaction, there is another mechanism that can be described with the following considerations. • Dissolved Cu+ and OH− reach the supersaturation level at electrode vicinity. • Cu2 O nuclei form on the electrode surface and grow upon ion incorporation. • Homogeneous nucleation does not take place: any Cu2 O precipitation nor porous layer. An scheme of the proposed growth mechanisms is shown in Fig. 9a and b. Fig. 9a shows the scheme of the well-known chemical deposition process [27] that forms Cu2 O nuclei growing from the Cu+ and OH− ions present in solution. This mechanism accounts for the surface inhomogeneity in grain size and roughness increase. Fig. 9b depicts the electrochemical formation of Cu2 O and the solid state reaction mechanism, with Cu+ reacting with the OH− ions in the solution to form an hydroxylated Cu(I) surface [11,13], then after OH− dehydration, oxygen and Cu+ migrate in opposite directions to form Cu2 O that grows beneath the electrode surface [12]. This migration would cause a high density of defects and stress that would be responsible for the lower crystallinity observed. The “lost charge” deduced by the analysis of the td = 0 s film thickness is partially due to this mechanism, but as mentioned before, further studies are needed to verify the role of defects in charge transfer.
Fig. 9. Schemes of (a) chemical deposition of Cu2 O from dissolved Cu+ in alkaline media. (b) Place exchange mechanism of Cu2 O growth under applied potential; electrochemical reaction takes place at electrode surface.
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4. Conclusions We prepared Cu2 O films by Cu anodization using an electrochemical program that includes electrode exposition to Cu+ dissolution potential at selected times. The films were crystalline Cu2 O, with an energy band gap that was tailored from 2.07 to 2.21 eV. From the analysis of the film morphology and thickness and from the current transients during the film growth we propose that the Cu2 O film growth mechanism includes concurrent processes: a pure electrochemical step, an heterogeneous nucleation of Cu2 O from the dissolved Cu+ and OH− , and the solid state reaction between diffusing Cu+ and oxygen that causes Cu2 O formation underneath the electrode surface. Acknowledgements FCB recognizes IPN-México for COTEPABE License; APP acknowledges financial support through an UB-Collaboration Grant at the Physical Chemistry Department. OCA acknowledges a UBVisiting Professor Aid. The support of the Scientific Technical Services of the University of Barcelona is kindly recognized, particularly the Units of X-ray Diffraction, Molecular Spectroscopy and Nanometric Techniques. This work was partially financed by the Project CTQ-2007-68101-C02-01 from the Ministerio de Educación y Ciencia, Spain. References [1] R. Dong, D.S. Lee, W.F. Xiang, S.J. Oh, D.J. Seong, S.H. Heo, H.J. Choi, M.J. Kwon, S.N. Seo, M.B. Pyun, M. Hasan, H. Hwang, Appl. Phys. Lett. 90 (2007) 042107. [2] S.S. Jeong, A. Mittiga, E. Salza, A. Masci, S. Passerini, Electrochim. Acta 53 (2008) 2226.
[3] A. Mittiga, E. Salza, F. Sarto, M. Tucci, R. Vasanthi, Appl. Phys. Lett. 88 (2006) 163502-1. [4] T. Hsueh, C. Hsu, S. Chang, P. Guo, J. Hsieh, I. Chen, Scripta Mater. 57 (2007) 53. [5] M. Izaki, K. Mizuno, T. Shinagawa, M. Inaba, A. Tasaka, J. Electrochem. Soc. 153 (9) (2006) C668. [6] A.H. Jayatissa, P. Samarasekara, G. Kun, Phys. Status Solidi A 206 (2) (2009) 332. [7] A.-L. Daltin, F. Bohr, J.-P. Chopart, Electrochem. Acta 54 (2009) 5813. [8] J.M.M. Droog, C.A. Alderliesten, P.T. Alderliesten, G.A. Bootsma, J. Electroanal. Chem. 111 (1980) 61. [9] M.R. Gennero de Chialvo, J.O. Zerbino, S.L. Marchiano, A.J. Arvia, J. Appl. Electrochem. 16 (1986) 517. [10] L.D. Burke, T.G. Ryan, J. Electrochem. Soc. 137 (1990) 1358. [11] J. Kunze, V. Maurice, L.H. Klein, H. Strehblow, P. Marcus, J. Phys. Chem. B 105 (2001) 4263. [12] M. Kang, A.A. Gewirth, J. Phys. Chem. B 106 (2002) 12211. [13] F. Caballero-Briones, J.M. Artés, I. Díez-Pérez, P. Gorostiza, F. Sanz, J. Phys. Chem. C 113 (2009) 1028. [14] PDF Cards 050667 (Cu2 O) and 040836 (Cu), Powder Diffraction Files-2 Data Base, International Centre for Diffraction Data, Newtown Square, PA, USA, 2000. [15] M. Theiss, Hardware and Software for Optical Spectroscopy-SCOUT 98, Dr. Bernhard-Klein-Str. 110, 52078 Aachen, Germany, www.mtheiss.com. [16] G.E. Jellison Jr., Thin Solid Films 33 (1998) 313. [17] A.-L. Barabasi, Comput. Mater. Sci. 6 (1996) 127. [18] E. Matthijs, S. Langerock, E. Michailovna, L. Heerman, J. Electroanal. Chem. 570 (2004) 123. [19] A.I. Oliva, R. Castro-Rodrıguez, O. Ceh, P. Bartolo-Perez, F. Caballero-Briones, V. Sosa, Appl. Surf. Sci. 148 (1999) 42. [20] H.P. Klug, L.E. Alexander, X-ray Diffraction Procedures for Polycrystalline and Amorphous Materials, John Wiley, New York, NY, 1959. [21] V. Kumar, S.Kr. Sharma, T.P. Sharma, V. Singh, Opt. Mater. 12 (1999) 115. ˜ Phys. Rev. B 60 (1999) 4758. [22] A. Iribarren, R. Castro-Rodriguez, V. Sosa, J.L. Pena, [23] B.A. Gizhevski, Yu.P. Sukhorukov, A.S. Moskvin, N.N. Loshkareva, E.V. Mostovshchikova, A.E. Ermakov, E.A. Kozlov, M.A. Uimin, V.S. Gaviko, Zh. Éksp. Teor. Fiz. 129 (2006) 336 [JETP 102 (2006) 297]. [24] N.A. Mohemmed Shanid, M. Abdul Khadar, Thin Solid Films 516 (2008) 6245. [25] B. Krishnamurthy, R.E. White, H.J. Ploehn, Electrochim. Acta 47 (2002) 3775. [26] H. Raebiger, S. Lany, A. Zunger, Phys. Rev. B 76 (2007) 045209. [27] Gary Hodes, Chemical Solution of Deposition Semiconductor Films, Marcel Dekker, New York, 2002.
Contents lists available at ScienceDirect
Electrochimica Acta journal homepage: www.elsevier.com/locate/electacta
Evidence and analysis of parallel growth mechanisms in Cu2 O films prepared by Cu anodization F. Caballero-Briones a,b,∗,1 , A. Palacios-Padrós a , O. Calzadilla c , Fausto Sanz d,a,b,∗∗ a
Department of Physical Chemistry, Universitat de Barcelona, Martí i Franquès 1, 08028 Barcelona, Spain CIBER-BBN, María de Luna 11, 50018 Zaragoza, Spain Facultad de Física, Universidad de La Habana, San Lázaro y L, Colina Universitaria, 10400 Vedado, La Habana, Cuba d Institute for Bioengineering of Catalonia (IBEC), Edifici Hèlix, Baldiri i Reixac 15-21, 08028 Barcelona, Spain b c
a r t i c l e
i n f o
Article history: Received 29 July 2009 Received in revised form 8 September 2009 Accepted 8 October 2009 Available online 20 October 2009 Keywords: Cuprous oxide Anodic films Reflectance Thickness Band gap Urbach tail parameter Dissolution Growth mechanism
a b s t r a c t We have studied the preparation of Cu2 O films by copper anodization in a 0.1 M NaOH electrolyte. We identified the potential range at which Cu+ dissolution takes place then we prepared films with different times of exposure to this potential. The morphology, crystalline structure, band gap, Urbach energy and thickness of the films were studied. Films prepared with the electrode unexposed to the dissolution potential have a pyramidal growth typical of potential driven processes, while samples prepared at increasing exposure times to dissolution potential present continuous nucleation, growth and grain coalescence. We observed a discrepancy in the respective film thicknesses calculated by coulometry, atomic force microscopy and optical reflectance. We propose that anodic Cu2 O film formation involves three parallel mechanisms (i) Cu2 O nucleation at the surface, (ii) Cu+ dissolution followed by heterogeneous nucleation and (iii) Cu+ and OH− diffusion through the forming oxide and subsequent reaction in the solid state. © 2009 Elsevier Ltd. All rights reserved.
1. Introduction Cu2 O is a p-type semiconductor with a band gap around 2 eV that has recently aroused interest because of its potential applications in random access memories [1], low-cost solar cells [2], and gas sensors [3] among other devices. The electrochemical fabrication of Cu2 O has been usually reported by electrodeposition onto transparent conductive oxides [2,4–7]. Copper electrochemistry and Cu2 O growth in alkaline media have been extensively studied [8–12]: the main anodic peak observed around U = −400 mV (vs silver/silver chloride electrode, henceforth SSC) in the potentiodynamic curves is associated to the formation of Cu2 O. Additionally, at potentials negative with respect to Cu2 O formation, the development of Cu+ hydrous species and OH− adsorption has been described. Copper (I) soluble species,
∗ Corresponding author at: Department of Physical Chemistry, Universitat de Barcelona, Martí i Franquès 1, 08028 Barcelona, Spain. Tel.: +34 934039239; fax: +34 934039239. ∗∗ Corresponding author at: Department of Physical Chemistry, Universitat de Barcelona, Martí i Franquès 1, 08028 Barcelona, Spain. E-mail addresses: [email protected] (F. Caballero-Briones), [email protected] (F. Sanz). 1 On leave from: CICATA-IPN Altamira, Km 14.5 Carr. Tampico-Pto Industrial, 89600 Altamira, Mexico. 0013-4686/$ – see front matter © 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.electacta.2009.10.031
namely (Cu2 O2 H)− aq can also be produced. The Cu2 O growth starts from the reconstructed OH-populated surface [11]. Nevertheless only few works focus on the preparation of thick Cu2 O anodic films and although soluble Cu+ formation has been repeatedly suggested, we did not found reports with experimental evidence of copper dissolution and its effects on film properties. In a recent contribution [13] we reported an anodization method of Cu0 in alkaline media that allows the preparation of Cu2 O films as thicker as 100 nm. We studied the electronic properties of the films and proposed an improved electronic band diagram of the Cu|Cu2 O|electrolyte interface. Herewith, a Cu+ dissolution potential range forming Cu+ species in solution before the Cu2 O growth potential range has been identified by current enhancement with the addition of a complexing agent for Cu+ and then, several films with different exposure times to the dissolution potential have been prepared. The contribution of this exposure to the film properties is analyzed in terms of optical, microstructural and thickness changes. The band gap and Urbach parameter were modulated as achieved. Moreover, the consideration of this dissolution effect brings new information on the growth mechanism of Cu2 O films. 2. Experimental details Cu2 O layers were prepared onto polycrystalline Cu disks in a 0.1 M NaOH electrolyte by an electrochemical routine described
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F. Caballero-Briones et al. / Electrochimica Acta 55 (2010) 4353–4358
Fig. 1. Potentiodynamic response of the Cu electrode in the 0.1 M NaOH electrolyte. Inset depicts the enhanced current response in the dissolution peak at around U = −450 mV, when NH4 NO3 was added.
elsewhere [13]. Briefly, the electrochemical procedure involves a series of potential steps that lead to different processes, i.e. the reduction of the native oxide, the formation of an OH-adsorbed submonolayer, a dissolution stage and oxide formation. The growth was performed using a PGSTAT 12 Autolab potentiostat in the three electrode configuration, using a homemade glass cell with
the Cu disk as working electrode placed at the bottom of the cell, a platinum coil as auxiliary electrode and a true reference SSC electrode. Potentiodynamic curves were used to monitor the electrode state and to identify the working potentials. In Fig. 1 is shown the cyclic voltammetry of the Cu electrode in the 0.1 M NaOH electrolyte where the main oxidation peak at U = −400 mV is observed. At around U = −550 mV a shoulder that corresponds to the OH-adsorption processes is distinguished. To characterize the dissolution potential, 500 l of a 0.01 M ammonia nitrate complexing solution was added to the electrolyte solution in order to enhance the current response due to Cu+ dissolution, as observed in the inset of Fig. 1, featuring the Cu+ dissolution peak at around U = −450 mV. The Cu2 O films have been prepared by exposing the electrode to the dissolution potential, Ud = −450 mV, during selected times (td = 0 s, td = 50 s, td = 100 s, td = 300 s, td = 500 s and td = 1000 s), and then to the growth potential, Ug = −400 mV, for 60 min. Current versus time transients were also registered to assess the mechanisms of Cu2 O growth. The amount of dissolved Cu+ and film thickness were calculated by the transferred charge at the potentials Ud and Ug respectively. The morphology of the films was studied by atomic force microscopy (AFM) in the intermittent contact mode with a multimode microscope from Veeco with Si cantilevers of 35 N/m spring constants. The roughness (Rrms ) and the average height were obtained from 2 m × 2 m AFM images. Grazing incidence X-ray diffraction (GIXRD) patterns were obtained using a Siemens D500 X-ray diffractometer with the detector pivoting on the 2 circle, with unfiltered Cu K␣ radiation and an incidence angle of 0.5◦ .
Fig. 2. (a) 2 m × 2 m AFM images of the films prepared at different exposition times at the dissolution potential Ud = −450 mV; (b) 500 × 500 nm AFM images of the td = 0 s and td = 1000 s films in the contour representation to better appreciate the nature of the grain formation. Z scales were equalized for image comparison.
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Fig. 3. Roughness and average height of the films, calculated from 2 m × 2 m AFM images. Inset: standard deviation of the roughness and average height as function of time. Continuous lines are only for guideline.
Phases were identified using the PDF database [14]. The diffuse reflectance measurements, from which the thickness, band gap and Urbach tail parameter were calculated, had been performed in a Shimadzu UV-2101 PC Spectrometer with an integrating sphere, using a polished Cu substrate as a reference. The optical properties of the Cu2 O thin films in the spectral range of interest, 300–800 nm, were fitted by three types of electronic excitations: band gap transitions, interband transitions from the bulk of the valence band into the bulk of the conduction band, and intraband transitions of the electrons in the conduction band. These transitions were modeled with standard formulae available in the SCOUT 98 program [15]. For the band gap transitions, we used the Tauc-Lorentz model [16] that has been proposed to model the band gap. The transitions into the upper half of the conduction band are represented by an harmonic oscillator. The Drude Model was used to calculate the contribution of free electrons to the susceptibility. All the procedures used for the fitting are implemented in the SCOUT 98 program. The experimental data were fitted with a standard deviation better than 10−4 . 3. Results 3.1. Morphology Fig. 2a presents 2 m × 2 m AFM images of the different films prepared at several exposure times. In the absence of a dissolution pathway, grains have a pyramidal shape with a wide size distribution. When the electrode is largely exposed to the dissolution potential, the grains evolve to a more irregular shape and grains become more definite and their size increases with the exposure time. Fig. 2b shows in more detail the morphology of the films td = 0 s and td = 1000 s. In the image corresponding to td = 0 s film the pyramids evolving from the bottom of the layer are evident, constituting a typical morphology of a diffusion limited process where planar diffusion hinders some of the nuclei to grow [17,18]; for the td = 1000 s film, the coalescence and grain size distribution observed in the respective image suggest continuous formation of nucleus and grain growth. Fig. 3 shows the morphology parameters, Rrms and average height as function of exposure time at the dissolution potential. They are typical values from at least five images obtained in different zones of the sample. We report the roughness from the 2 m × 2 m AFM images because at this size Rrms does not depend on the length scale used to probe it and the surface can be described with a single number [17]. In Fig. 3 it is observed that both average height and roughness increase steadily from td = 0 s to td = 1000 s. Similar behavior of the roughness values has been observed in
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Fig. 4. Crystallite size in the (1 1 1) Cu2 O direction vs exposure time at the dissolution potential.
chemical bath deposited films at long deposition times [19] and is attributable to increasing film compactness. The values of Rrms of the films prepared with exposure to the dissolution potential, reach a similar value to that of film td = 0 s only after 300 s of exposure. Dispersions of the parameters, taken as the standard deviation, are presented in the inset. Roughness dispersion tends to reduce from film td = 300 s indicating that for longer exposure times, grain growth becomes the dominant mechanism not formation of new nuclei. Correspondingly, height dispersion does not reduce until growth contribution uniforms grain sizes. 3.2. Structure and optical properties The cuprite phase was identified in all the films by X-ray diffraction, without evidence of CuO. The main cuprite peak at a Bragg angle around 2 = 36.4◦ is always observed and only in the td = 0 s and the td = 1000 s films the (2 0 0) Cu2 O reflection becomes slightly noticeable, thus lattice parameters of the prepared samples were not calculated because the lack of other distinct diffraction peaks and due to the indetermination of the Bragg angle in the GIXRD measurements. Even so, the reported bulk Cu2 O lattice parameter is 0.427 nm and that of Cu is 0.361 nm, leading to a lattice mismatch of about 18%. The generated epitaxial stress between the Cu2 O and Cu lattices is released with a tilt of 5◦ of the Cu2 O lattice with respect to that of Cu and a lattice compression of 1% [11,12]. The crystallite size of the films can be estimated from the Scherrer formula [20]: =
K ˇ cos
(1)
Here, = mean crystallite dimension, K = shape factor ∼0.9, = wavelength (1.54 Å), ˇ = line broadening (FWHM), = Bragg angle, taken for this samples as the position of the powder pattern of the (1 1 1) reflection of Cu2 O (2 = 36.417◦ ) due to the indetermination of the peak position in the GIXRD measurement. Fig. 4 plots the crystallite size versus exposure time at the dissolution potential. The crystallite size of the films increases with the exposure time from 9 to 14 nm, and it is only at 300 s that it reaches a similar size (around 12 nm) than when the electrode was not exposed. The increase in the crystallite size would affect the stress in the films, but in the absence of more diffraction peaks it is not possible to apply the Williamson-Hall treatment to determine it. Nevertheless, as it is mentioned below, the stress effect will be included in the disorder analysis quantified by the Urbach tail parameter.
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Fig. 5. Diffuse reflectance spectra from the prepared samples. Inset shows a transformed spectra used for calculation of the band gap energy (detail in the text).
The UV–vis reflectance spectra of the samples are shown in Fig. 5, where it is observed a reduction in the reflectance in the nonabsorption zone, from 650 to 800 nm that contains the information from which the film thickness was extracted (shown in Fig. 8) [15]. The inflection zone where the band gap is located is around 500 nm and below it the intraband transitions can be observed as two peaks around 300 and 350 nm. The inset in Fig. 5 shows the calculation of the direct band gap from the reflectance spectra as proposed by Kumar et al. [21]. For a direct transition the band gap energy (Eg ) is obtained from the linear zone intercept with the energy axis. In the inset, after the absorption zone, the so-called Urbach region manifests itself as an exponential tail where the information of the disorder is enclosed [22]. From the absorption versus energy plot, the fitting of this exponential region below the gap gives the value of the Urbach energy (E0 ). In Fig. 6, Eg and E0 are plotted versus the exposure time at the dissolution potential. The reported bulk value for Cu2 O is Eg = 2.10 eV [23]. For the studied films Eg values are around 2.07–2.10 eV for the films between td = 0 s and td = 300 s and then increases to about 2.20 eV at td = 500 s. The Urbach energy shows a similar behavior for the samples exposed to the dissolution potential, i.e. E0 increases with time, but for t = 0 a high value (ca. 400 meV) is measured. The increase in Eg would be explained by the formation of nanoparticulated Cu2 O [23], but this seems to be in contradiction with the crystallite size determined by GIXRD. This blue shift has
Fig. 6. Eg and E0 plot of the prepared films vs the exposure time at dissolution potential.
Fig. 7. Semilogarithmic plot for the current transients recorded at the growth potential, Ug = −400 mV.
also been observed in electrodeposited Cu2 O films without quantum confinement possibilities and it has been explained by the formation of potential barriers at the surface due to the preparation of films by repeating deposition and oxidation [24]. This explanation is also valid in our case, because it supports the dual nucleation-growth mechanism suggested in the morphology analysis. The detailed analysis of E0 values in Cu2 O films related to the contribution of the different sources of disorder is subject of a more extended contribution (in preparation), nevertheless is interesting to point out some issues to explain the present observations. As mentioned before, the Urbach energy is a sum of several factors of disorder: crystalline structure defects, grain boundaries (directly related with grain size), stress and electronic interactions [22]. Then, film growth mechanisms that lead to a lack of uniformity, strong dispersion of particle sizes and structural stress would direct to increments in the Urbach energy depending, of course, on the deposition conditions. In our case, we observe strong correlation of the increase in roughness and size dispersion with the time of exposure at the dissolution potential. The suggestion of parallel mechanism in anodic Cu2 O film growth gains more evidences with this analysis. 3.3. Growth mechanism In Fig. 7 the current versus time curves registered at the growth potential, Ug = −450 mV, are shown. For the td = 0 s film curve the typical shape of a nucleation process is observed, with the spike minimum around 60 s. The current then increases due to nucleation and growth of Cu2 O to reach the maximum value after 1800 s of growth and then reduce slowly as the diffusion layer expands. Chronoamperometric curves obtained with increasing dissolution times have the spike minimum at the very beginning of the growth time and after the current maximum is reached it tends to decrease rapidly again. This reduction could be caused by electrode conduction drop due to the thicker Cu2 O film grown. It is interesting to observe that the curves of the films with td = 500 s and td = 1000 s present secondary current peaks near the end of the process confirming that concurrent processes are taking place. The film prepared without exposure to the dissolution potential, td = 0, reaches its maximum current value at 1900 s, compared with the one with a dissolution time of 1000 s that reaches it at 36 s. For the case of td = 0 s the growth depends on OH− ion diffusion to the electrode surface. The film td = 1000 represents a completely different situation, an important amount of Cu+ is now present in the electrode vicinity, thus additionally to a rapid and extensive nucle-
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Fig. 8. Film thickness calculated from reflectance measurements, from the charge transfer calculated during the growth potential application and estimated from the average heights from AFM.
ation process, growth caused by heterogeneous nucleation from the incoming Cu+ ions is expected. Fig. 8 presents the comparison between the thicknesses calculated using three different magnitudes: the transferred Faradaic charge, the average height calculated from the AFM images and the thickness calculated from the optical reflectance measurements. There are two striking features: the first one is the huge discrepancy between the thicknesses calculated optically and by coulometry. This is consequence of parallel mechanisms of potentiostatic growth and chemical deposition; as mentioned above, the first one depends only on OH− diffusion to the electrode, while dissolved Cu+ is readily available in the second. The second striking feature is that the sample prepared without exposition to dissolution potential, td = 0, results much thinner by the charge transfer calculation than optically. The calculated “charge transfer thickness” for td = 0 fairly corresponds to the AFM value, accounting for the “external” growth. This second feature evidences that an important amount of the Cu2 O film grows underneath the electrode surface by another concurrent mechanism. We propose that this mechanism arises from diffusion of oxygen and copper species
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through the growing oxide layer that subsequently react in the solid state. Following the work of Krishnamurthy et al. [25] we propose this mechanism to be potential driven and the copper and oxygen species could be copper vacancies and oxygen interstitials as they are reported to be the main defects in the bulk Cu2 O [26]. It is expected that such a mechanism would cause an important amount of structural stress and electronic defects. This would explain the abnormally high E0 observed in the td = 0 film. On the other side, the reduction in thickness observed at the higher times could be explained because the electric field reduction that would reduce ion migration through the oxide layer. Nevertheless, we believe that more experiments are still needed to clarify the identity of the diffusing species and their role in the charge transfer, to better understand the kinetics of film growth. With the aforementioned evidences of morphology, current transient analysis and thickness, we propose that additionally to the electrochemical process for Cu2 O formation and the diffusionsolid state reaction, there is another mechanism that can be described with the following considerations. • Dissolved Cu+ and OH− reach the supersaturation level at electrode vicinity. • Cu2 O nuclei form on the electrode surface and grow upon ion incorporation. • Homogeneous nucleation does not take place: any Cu2 O precipitation nor porous layer. An scheme of the proposed growth mechanisms is shown in Fig. 9a and b. Fig. 9a shows the scheme of the well-known chemical deposition process [27] that forms Cu2 O nuclei growing from the Cu+ and OH− ions present in solution. This mechanism accounts for the surface inhomogeneity in grain size and roughness increase. Fig. 9b depicts the electrochemical formation of Cu2 O and the solid state reaction mechanism, with Cu+ reacting with the OH− ions in the solution to form an hydroxylated Cu(I) surface [11,13], then after OH− dehydration, oxygen and Cu+ migrate in opposite directions to form Cu2 O that grows beneath the electrode surface [12]. This migration would cause a high density of defects and stress that would be responsible for the lower crystallinity observed. The “lost charge” deduced by the analysis of the td = 0 s film thickness is partially due to this mechanism, but as mentioned before, further studies are needed to verify the role of defects in charge transfer.
Fig. 9. Schemes of (a) chemical deposition of Cu2 O from dissolved Cu+ in alkaline media. (b) Place exchange mechanism of Cu2 O growth under applied potential; electrochemical reaction takes place at electrode surface.
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4. Conclusions We prepared Cu2 O films by Cu anodization using an electrochemical program that includes electrode exposition to Cu+ dissolution potential at selected times. The films were crystalline Cu2 O, with an energy band gap that was tailored from 2.07 to 2.21 eV. From the analysis of the film morphology and thickness and from the current transients during the film growth we propose that the Cu2 O film growth mechanism includes concurrent processes: a pure electrochemical step, an heterogeneous nucleation of Cu2 O from the dissolved Cu+ and OH− , and the solid state reaction between diffusing Cu+ and oxygen that causes Cu2 O formation underneath the electrode surface. Acknowledgements FCB recognizes IPN-México for COTEPABE License; APP acknowledges financial support through an UB-Collaboration Grant at the Physical Chemistry Department. OCA acknowledges a UBVisiting Professor Aid. The support of the Scientific Technical Services of the University of Barcelona is kindly recognized, particularly the Units of X-ray Diffraction, Molecular Spectroscopy and Nanometric Techniques. This work was partially financed by the Project CTQ-2007-68101-C02-01 from the Ministerio de Educación y Ciencia, Spain. References [1] R. Dong, D.S. Lee, W.F. Xiang, S.J. Oh, D.J. Seong, S.H. Heo, H.J. Choi, M.J. Kwon, S.N. Seo, M.B. Pyun, M. Hasan, H. Hwang, Appl. Phys. Lett. 90 (2007) 042107. [2] S.S. Jeong, A. Mittiga, E. Salza, A. Masci, S. Passerini, Electrochim. Acta 53 (2008) 2226.
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