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Proton related defects in α-BaTiO3:H films based MIM capacitors
Solid State Ionics 180 (2009) 853–856
Contents lists available at ScienceDirect
Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s s i
Proton related defects in α-BaTiO3:H films based MIM capacitors F. El Kamel a,⁎, P. Gonon b a b
Grenoble Electrical Engineering Laboratory (CNRS) 25, avenue des Martyrs 38042 Grenoble, France Microelectronics Technology Laboratory (CEA) 17, avenue des Martyrs 38054 Grenoble, France
a r t i c l e
i n f o
Article history: Received 20 September 2008 Received in revised form 21 January 2009 Accepted 1 February 2009 Keywords: α-BaTiO3:H films MIM capacitors Proton related defects
a b s t r a c t Hydrogenated barium titanate film based metal–insulator–metal (MIM) systems show high dielectric constant and have therefore been proposed as solid state supercapacitors. Hydrogen was incorporated in the dielectric layer during the low temperature deposition process. An investigation of the electrical properties has revealed that hydrogen contributes to the conduction process as mobile ionic species as well as donor trap levels. © 2009 Elsevier B.V. All rights reserved.
1. Introduction
2. Experimental detail
Generally, one tends to avoid the presence of electrical defects in insulator materials which could have deleterious consequences on device performances. In contrast, there are some special cases for which the presence of these “defects” is looked for. For instance, point defects, by leading to a substantial ionic conductivity, are beneficial for the fabrication of supercapacitors [1,2]. They rely on the existence of a high density of mobile ions (ionic conductivity). In fact, we know that the capacitance obtained with conventional capacitors finds its origin in the electronic, ionic and dipolar polarization occurring in the bulk whereas for supercapacitors [1] the capacitance is linearly proportional to the ionic conductivity instead of the dielectric relaxation [3]. In these materials, under the electric field mobile charge carriers can accumulate at the metal–dielectric interface resulting in a space charge layer (double layer) over the Debye length. In such capacitors the double-layer capacitance can reach values up to several μF/cm2 depending on the nature of the dielectric. Then, supercapacitors exhibit the property that large amount of energy can be stored in the double-layer. Recently, all solid state supercapacitors such as hydrated lithium fluoride [1] and yttria stabilized zirconia [2] are examples of ionic conductors investigated for such a purpose [2,4–6]. The formation of solid state supercapacitors reported up to now involves high temperature processing. This constitutes a serious limitation to the integration on packages which can not sustain high temperature. From this viewpoint, in the present work we tried to incorporate protons at low temperature in barium titanate films during the deposition process. Attempts were made to study the electrical defects resulting from hydrogen incorporation.
Barium titanate films (1 μm-thick) were grown by rf magnetron sputtering process using a polycrystalline BaTiO3 target (high purity, 10 cm in diameter and powered at 100 W). Standard sputtering is performed with pure argon gas (100% Ar) at 10− 2 mbar. Substrates were gold coated silicon, previously coated with a chromium adhesion layer. The resulting Au/Cr/Si substrates were mounted on a watercooled substrate holder facing the target (5 cm above the target). Here, hydrogen was introduced in the sputtering gas through the addition of pure H2 (H2-ratio varies from 0 to 25%). Films grown at lower temperature were found to be amorphous as expected by X-Ray Diffraction [3]. Chemical properties of the elaborated films were examined at room temperature using a VG MicroTech X-ray Photoelectron Spectrometer (Mg–Kα) and a Bio-Rad FT–IR spectrometer in diffuse reflectance mode. After deposition, Au-dots (1.77 mm2-area) were evaporated through a shadow mask on the front side of the films. In that way we formed Au/a-BaTiO3:H/Au planar capacitors. Dielectric properties of these capacitors were studied using a Novocontrol impedance analyzer (with an ac voltage of 0.1 V) as a function of temperature and sweeping frequency. Current–voltage measurements were performed using a Keithley 6517A electrometer. Temperature control was realized using a Linkam TMS 94 programmable temperature monitor. During the dielectric and electrical measurements samples were enclosed in a dark shielded cell filled with dry nitrogen.
⁎ Corresponding author. Tel.: +33 476 88 1007; fax: +33 476 88 7945. E-mail address: [email protected] (F. El Kamel). 0167-2738/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2009.02.009
3. Results and discussion Standard films (grown under pure argon) were transparent [7] whereas the transparency of hydrogenated films (BaTiO3:H) diminishes widely with the addition of hydrogen. Their color tends gradually to yellow, then to gray and finally they darken for the strongest
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F. El Kamel, P. Gonon / Solid State Ionics 180 (2009) 853–856
Fig. 1. Room temperature XPS analysis associated to the O-1s electronic transition (oxide and hydroxide components), realized on two samples. One deposited under pure argon gas (reference) and another under hydrogen/argon mixture (20% H2 + 80% Ar). The inset shows the FT–IR spectra realized at room temperature on BaTiO3 films grown under different H2-ratios (0, 10 and 20%) in the sputtering gas.
oxide remains constant. Such behavior may results from the incorporation of hydrogen during the deposition process. Further evidence for hydrogen incorporation can be given by FT–IR reflectance spectroscopy (see inset in Fig. 1) as a broad band around 3500 cm− 1. This IR absorption band is ascribed to the O\H stretching vibration [10,11]. Such feature is absent in standard films (grown under pure argon). The large increase of the O\H stretch magnitude with increasing the H2-ratio in the sputtering gas clearly manifests that a large density of protons has been incorporated into the films during the elaboration process. As it is already established [10,12,13] the hydrogen ionizes to give an electron (donor defect) and a proton (H+) which combines with oxygen of the perovskite lattice to form hydroxide group (OH˙O). Fig. 2 shows an overview on the frequency spectra associated to the real part of permittivity (ε'), its imaginary counterpart (εq) and the ac conductivity (σac) performed at − 60 °C on a-BaTiO3:H film (20% H2 + 80% Ar). Films grown under H2-containing atmosphere display hundred times higher capacitance than the one measured for films grown under pure argon [3]. This feature was explained through an electrode polarization mechanism [14] which arises from proton pileup at the cathode over a Debye length [3]. At low frequency domain, measured capacitance was completely determined by the charge stored in the double layer. As a result, electric double-layer capacitors have been widely regarded as energy storage devices. The double-layer is clearly
hydrogen ratio. This dark color indicates that the hydrogenated films are nonstoichiometric with some oxygen deficiency [8,9]. In addition, chemical properties of the hydrogenated films undergo dramatic alteration. Fig. 1 displays a part of the XPS spectrum associated to the O-1s electronic transition (oxide and hydroxide components) and realized on two films. One deposited under pure argon gas (reference) and another under hydrogen/argon mixing gas (20% H2 + 80% Ar). We show that the oxide component of both films is detected at around 530 eV and is assigned to oxygen in the BaTiO3 lattice. This indicates that the elementary lattice does not undergo any change when hydrogen was added to the sputtering gas. On the other hand, based on the XPS analysis and by comparing the peak-area attributed to the oxide and hydroxide components of the analyzed films (reference and hydrogenated) we can determine the evolution of these components when hydrogen was added to the sputtering gas. We note that films grown under hydrogen containing atmosphere display large density of hydroxide that is 12% higher than the one determined for films grown under pure argon while the density of
Fig. 2. Dielectric permittivity (real and imaginary parts) and conductivity measured at –60° C on BaTiO3 layers sputtered under pure argon gas (empty symbols) and under hydrogen/ argon mixture (full symbols).
Fig. 3. Plot (a) shows the temperature dependence of the loss factor, tanδ(f). Films were grown under hydrogen/argon mixture (20% H2 + 80% Ar). Plot (b) illustrates the evolution of the relaxation frequency as a function of 1000 / T (Arrhenius diagram).
F. El Kamel, P. Gonon / Solid State Ionics 180 (2009) 853–856
Fig. 4. Plot (a) presents the temperature dependence of the leakage current (J–E). In figure (b) we replotted the J–E characteristics according to the Poole–Frenkel model.
evidenced, especially at low temperatures by a dispersive behavior of the dielectric constant accompanied by a relaxation peak in the loss measurements. We believe that the dielectric losses are mainly the result of the ion migration. At high frequencies, ε' relaxes toward the value obtained when films are deposited under pure argon (~20 [14]). Furthermore, σac(f) is also affected by the addition of hydrogen. At low frequency domain, it increases by around 7 orders of magnitude and displays two plateaus separated by a step-like behavior which implies a relaxation process. The high-frequency plateau corresponds to the bulk response and the low-frequency one is assigned to the electrode polarization effect. In order to better illustrate the relaxation process, we plotted in Fig. 3a the frequency dependence of the loss factor, tanδ(f). As we can see, the dielectric loss comes from two mechanisms, resistive loss (power law evolution) and relaxation (peak). The resistive mechanism can be explained by mobile charges transfer at the metal– dielectric interface [9,14] whereas the relaxation process is originated by the double-layer modulation under the alternating field [9,14]. As already evidenced in Fig. 3a, the relaxation frequency (peak position, f0) shifts to high-frequencies side as the temperature increases, according to an Arrhenius law, τ0 = τ0,max exp (Ea / kBT). τ0 = 1 / f0 is the relaxation time, f0 is the peak position of tanδ(f) in the frequency domain, Ea is the activation energy of the relaxation process and kB is the Boltzmann's constant. A plot of f0 against 1000 / T (Fig. 3b) gives the activation energy Ea, of the relaxation process. We can find out that this process shows an Arrhenius-type temperature dependence over the whole temperature domain and is thermally activated with 0.16 eV when temperature is lower than − 75 °C and with 0.24 eV when temperature goes beyond this value. This behavior can be
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compared to the one depicted on the bulk conductivity (extracted from the high-frequency plateau in the σac(f) characteristics). Based on the fact that f0 = (σ / 2πεrε0)(2LD / L)1/2, where σ is the bulk conductivity, εr and ε0 are respectively the real part and the free space permittivity, L is the layer thickness and LD is the Debye length [14], it appears that the bulk conductivity and the relaxation peak have the same activation energy. Such a feature was confirmed experimentally for samples grown under different H2-ratios (13, 17, 20 and 25%) in the sputtering gas. We extracted three activation energies, corresponding to three conduction processes. These processes dominate the dielectric response in three different temperature domains. At temperature below −75 °C, the transport phenomenon is thermally activated with an activation energy of about 0.15 eV. When temperature varies from −75 to − 30 °C, activation energy is close to 0.25 eV. Finally, at temperature above −30 °C, the determined activation energy is around 0.37 eV. Activation energies evaluated from the slope of the Arrhenius plots can give an idea on the transport phenomena taking place at these temperature ranges and on the electrical active defects that are at the origin of the relaxation and conduction processes. As we explained in our previous works [14–16], the mobility of oxygen vacancies occurs at sufficiently high temperature and it is thermally activated with an activation energy in the vicinity of 1 eV. Such a value is much higher than those extracted from the dielectric measurements in the present study. So, the effect of oxygen vacancies can be ruled out from the explanation of the observed dielectric response at low temperature. On the other hand, we know that the incorporation of protons in the BaTiO3 films could significantly affect their electrical properties. In order to speculate the conduction mechanism in dc regime, leakage currents were measured on BaTiO3:H films as a function of the applied field at different temperatures (from − 150 to −50 °C). dc bias was applied to the bottom electrode and the current was measured after 60 s for stabilization concern. Experimental data measured on a sample grown under hydrogen containing atmosphere (17% H2) are shown in Fig. 4a. It is clearly seen that J–E characteristics exhibit a Poole–Frenkel type behavior, which implies that the conduction mechanism can be described by a thermally stimulated emission from a discrete traps distribution. Thus, the leakage current density can be expressed as, qffiffiffiffiffiffiffiffi1 0 −q /t − πqE eopt A JPF = qN c μEexp@ ð1Þ kB T where E is the applied field, q is the elementary charge, kB is Boltzmann's constant, Nc is the effective density of states in the
Fig. 5. Evolution of the effective trap depth as a function of E1/2. We found that the leakage current can be explained by carrier detrapping from shallow traps level ϕ0 localized at 0.145 eV below the conduction band.
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F. El Kamel, P. Gonon / Solid State Ionics 180 (2009) 853–856
conduction band assumed to be around 1018 cm− 3 [17], μ is the electronic mobility, qϕt is the trap level and εopt is the optical permittivity. As stated above, log(J /E) vs. E1/2 should be linear according to the Poole– Frenkel model (Fig. 4b). Then, we guess that the measured leakage current can be explained by carriers detrapping from shallow trap levels localized within the bandgap. In order to determine the activation energy of such process, we plotted the J(1 / T) characteristics (at a fixed electric field). J(1 / T) curves (not shown here) depict a linear behavior with a negative slope which suggests that the conduction process is thermally stimulated. Activation energy and mobility vary respectively from 0.12 and 4.31 × 10− 7 to 0.05 eV and 2 × 10− 6 cm2/Vs when the electric field varies from 0.2 to 3 MV/m. Such energies correspond to the effective trap depth [ϕ0 − Δϕ] which is strongly dependent on the electric field. Based on the slowness of the Poole–Frenkel mechanism, these mobility's values are in a good agreement with the electronic conduction. In the literature, electronic mobility shows a large discrepancy and is found to vary from 10− 3 [18–22] to 10− 10 cm2/Vs [17]. To determine the real depth ϕ0 (at E = 0 V/m) of traps, we plotted the effective trap depths as a function of E1/2 (see Fig. 5). We found that the leakage current can be explained by carrier detrapping from shallow traps level ϕ0 localized at 0.145 eV below the conduction band. It is interesting to note that the same amount of energy is required for activating the relaxation process taking place at temperature below −75 °C. So, it is believed that the loss peak and the leakage current arise from the carriers detrapping from shallow traps localized at around 0.15 eV within the bandgap. Positively charged protons can provide these energetic levels. In less symmetric environments, charged defects in the vicinity of the proton tilt markedly the hydroxide group toward one of the neighboring oxygen ions. Since the hydrogenated samples were deposited under a reduced atmosphere, oxygen vacancies content increases and constitute the main charged defect. Thus, surrounding oxygen vacancies effectively act as positively charged defects that repel the proton. This feature allows the formation of a rather weak directional interaction (hydrogen-bond) between the proton and an adjacent oxygen, O(1)\H··O(2) [23]. Such a site might effectively act as a trap for the proton. In this approach, electron issued from the ionization of hydrogen is weakly bonded and it can be easily activated to the conduction band to increase the dielectric loss and the leakage current [24]. Such a structure generates a shallow donor level below the conduction band. This level was determined at around 0.15 eV using the leakage measurements. Moreover, the hydrogen-bond formation can be responsible of a large protonic conductivity in oxide materials. With the aim of examining protonic conduction, previous reports [25] adopted a model for OH oscillators localized on regular oxygen sites and undergoing a stretching mode alongside the “oxygen–oxygen” direction. This is precisely the situation that allows for proton oscillation to take place involving small ionic displacements (less than an atomic spacing) which require small activation energy. Due to their weakness, the hydrogen bonds can be dissociated even at low temperature with an activation energy around 0.22 eV [26]. This value agrees well with the one (0.25 eV) determined using dielectric measurements at temperature ranging from −75 to − 30 °C. Then, in this temperature range, the dielectric response can be ascribed to the dissociation of hydrogenbonds and consequently to a localized migration (or oscillation) of the proton alongside the hydroxide bond. Such process was previously emphasized and discussed by Weber et al. [25] in ionic-conductor oxides. Finally, the relaxation process considered at temperature higher than − 30 °C for hydrogenated samples was thermally activated with
0.37 eV. It is believed that this amount of energy is required to activate the proton diffusion within the oxide material. The reported value of the activation energy should be strongly affected by the amorphous state of the material but it agrees well with predicted experimental [27,28] and theoretical [27–29] values (0.3–0.6 eV) for the best proton conductors such as Y–BaZrO3 or BaCeO3 [27,28,30]. Based on the experimental observation and on the light of literature mentioned above, the conduction mechanism is predominately of Grotthuss-type [31], i.e. involving proton transfer (hopping) from the hydroxide defects to a nearest neighboring oxygen ion. 4. Conclusions Hydrogen incorporation in barium titanate films during the low temperature deposition process can contribute to the transport phenomena by introducing several features. At temperature higher than −30 °C protons contribute to the conduction mechanism as mobile ionic species. Such mechanism is thermally activated with 0.37 eV. In addition, protons have been proposed to induce n-type conductivity by generating shallow donor levels within the bandgap in titanates. Therefore, the trap we detect at 0.15 eV could be associated with such a donor level. Moreover, in disordered systems, protons were shown to engender hydrogen bonds in the structure. These bonds have to be related to a common feature for the oxides protonic conductors. As a result, the dielectric relaxation observed at temperature ranging from − 75 to − 30 °C could be related to the oscillation of the hydroxide bond due to the destruction of the hydrogen bond. Activation energy for this process is around 0.25 eV. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31]
L. Ma, Y. Yang, Appl. Phys. Lett. 87 (2005) 123503. M. Hendriks, M. Heijman, W. Zyl, J. Elshol, H. Verweij, J. Appl. Phys. 90 (2001) 5303. P. Gonon, F. El Kamel, Appl. Phys. Lett 90 (2007) 232902. J.Y. Kim, I.J. Chung, J. Electrochem. Soc. 149 (2002) A1376. P.G. Romero, M. Chojak, K.C. Gallegos, P.J. Asensio, N.C. Pastor, M.L. Cantú, Electrochem. Commun 5 (2003) 149. F. Lufrani, P. Staiti, Electrochim. Acta 49 (2004) 2683. I.H. Pratt, S. Firestone, J. Vac. Sci. Technol. 8 (1970) 256. S. Furusawa, H. Tabuchi, T. Sugiyama, S. Tao, J. Irvine, Solid State Ionics 176 (2005) 553. F. El Kamel, P. Gonon, F. Jomni, B. Yangui, J. European Ceram. Soc. 27 (2007) 3807. R. Waser, J. Am. Ceram. Soc. 71 (1988) 58. S. Aggarwal, S.R. Perusse, C.W. Tipton, R. Ramesh, H.D. Drew, T. Venkatesan, D.B. Romero, V.B. Podobedov, A. Weber, Appl. Phys. Lett. 73 (1998) 1973. J. Ahn, P. McIntyre, L. Mirkarimi, S. Gilbert, J. Amano, M. Schulberg, Appl. Phys. Lett. 77 (2000) 1378. K. Xiong, J. Robertson, Appl. Phys. Lett. 85 (2004) 2577. P. Gonon, F. El Kamel, J. Appl. Phys. 101 (2007) 073901. F. El Kamel, P. Gonon, L. Ortega, F. Jomni, B. Yangui, J. Appl. Phys. 99 (2006) 094107. F. El Kamel, P. Gonon, F. Jomni, B. Yangui, J. Appl. Phys. 100 (2006) 054107. Y.B. Lin, J.Y. Lee, J. Appl. Phys. 87 (2000) 1841. S. Zafar, R. Jones, B. Jiang, B. White, P. Chu, D. Taylor, S. Gillespie, Appl. Phys. Lett. 73 (1998) 175. C.N. Berglund, W.S. Baer, Phys. Rev. 157 (1967) 358. A.J. Moulson, J.M. Herbert, Electroceramics: Materials, Properties and Applications, Chapman and Hall, London, 1996. D. Keroack, Y. Lepine, J.L. Brebner, J. Phys. C. 17 (1984) 833. J.P. Boyeaux, F.M. Calendini, J. Phys. C 12 (1979) 545. L. Pejov, Chem. Phys. Lett. 376 (2003) 11. L. Dobaczewski, K.B. Nielsen, N. Zangenberg, B.B. Nielsen, A.R. Peaker, V.P. Markevich, Phys. Rev. B. 69 (2004) 245207. G. Weber, S. Kappham, Wöhlecke, Phys. Rev. B. 34 (1986) 8406. J.O.M. Bockris, A.K.N. Reddy, Modern Electrochemistry, Plenum, New York, 1970. W. Münch, K.D. Kreuer, G. Seifertlib, J. Majer, Solid State Ionics 125 (1999) 39. E. Matsushita, A. Tanase, Solid State Ionics 97 (1997) 45. E. Matsushita, T. Sasaki, Solid State Ionics 125 (1999) 31. K.D. Kreuer, Solid State Ionics 125 (1999) 285. K.D. Kreuer, Chem. Mater. 8 (1996) 610.
Contents lists available at ScienceDirect
Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s s i
Proton related defects in α-BaTiO3:H films based MIM capacitors F. El Kamel a,⁎, P. Gonon b a b
Grenoble Electrical Engineering Laboratory (CNRS) 25, avenue des Martyrs 38042 Grenoble, France Microelectronics Technology Laboratory (CEA) 17, avenue des Martyrs 38054 Grenoble, France
a r t i c l e
i n f o
Article history: Received 20 September 2008 Received in revised form 21 January 2009 Accepted 1 February 2009 Keywords: α-BaTiO3:H films MIM capacitors Proton related defects
a b s t r a c t Hydrogenated barium titanate film based metal–insulator–metal (MIM) systems show high dielectric constant and have therefore been proposed as solid state supercapacitors. Hydrogen was incorporated in the dielectric layer during the low temperature deposition process. An investigation of the electrical properties has revealed that hydrogen contributes to the conduction process as mobile ionic species as well as donor trap levels. © 2009 Elsevier B.V. All rights reserved.
1. Introduction
2. Experimental detail
Generally, one tends to avoid the presence of electrical defects in insulator materials which could have deleterious consequences on device performances. In contrast, there are some special cases for which the presence of these “defects” is looked for. For instance, point defects, by leading to a substantial ionic conductivity, are beneficial for the fabrication of supercapacitors [1,2]. They rely on the existence of a high density of mobile ions (ionic conductivity). In fact, we know that the capacitance obtained with conventional capacitors finds its origin in the electronic, ionic and dipolar polarization occurring in the bulk whereas for supercapacitors [1] the capacitance is linearly proportional to the ionic conductivity instead of the dielectric relaxation [3]. In these materials, under the electric field mobile charge carriers can accumulate at the metal–dielectric interface resulting in a space charge layer (double layer) over the Debye length. In such capacitors the double-layer capacitance can reach values up to several μF/cm2 depending on the nature of the dielectric. Then, supercapacitors exhibit the property that large amount of energy can be stored in the double-layer. Recently, all solid state supercapacitors such as hydrated lithium fluoride [1] and yttria stabilized zirconia [2] are examples of ionic conductors investigated for such a purpose [2,4–6]. The formation of solid state supercapacitors reported up to now involves high temperature processing. This constitutes a serious limitation to the integration on packages which can not sustain high temperature. From this viewpoint, in the present work we tried to incorporate protons at low temperature in barium titanate films during the deposition process. Attempts were made to study the electrical defects resulting from hydrogen incorporation.
Barium titanate films (1 μm-thick) were grown by rf magnetron sputtering process using a polycrystalline BaTiO3 target (high purity, 10 cm in diameter and powered at 100 W). Standard sputtering is performed with pure argon gas (100% Ar) at 10− 2 mbar. Substrates were gold coated silicon, previously coated with a chromium adhesion layer. The resulting Au/Cr/Si substrates were mounted on a watercooled substrate holder facing the target (5 cm above the target). Here, hydrogen was introduced in the sputtering gas through the addition of pure H2 (H2-ratio varies from 0 to 25%). Films grown at lower temperature were found to be amorphous as expected by X-Ray Diffraction [3]. Chemical properties of the elaborated films were examined at room temperature using a VG MicroTech X-ray Photoelectron Spectrometer (Mg–Kα) and a Bio-Rad FT–IR spectrometer in diffuse reflectance mode. After deposition, Au-dots (1.77 mm2-area) were evaporated through a shadow mask on the front side of the films. In that way we formed Au/a-BaTiO3:H/Au planar capacitors. Dielectric properties of these capacitors were studied using a Novocontrol impedance analyzer (with an ac voltage of 0.1 V) as a function of temperature and sweeping frequency. Current–voltage measurements were performed using a Keithley 6517A electrometer. Temperature control was realized using a Linkam TMS 94 programmable temperature monitor. During the dielectric and electrical measurements samples were enclosed in a dark shielded cell filled with dry nitrogen.
⁎ Corresponding author. Tel.: +33 476 88 1007; fax: +33 476 88 7945. E-mail address: [email protected] (F. El Kamel). 0167-2738/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2009.02.009
3. Results and discussion Standard films (grown under pure argon) were transparent [7] whereas the transparency of hydrogenated films (BaTiO3:H) diminishes widely with the addition of hydrogen. Their color tends gradually to yellow, then to gray and finally they darken for the strongest
854
F. El Kamel, P. Gonon / Solid State Ionics 180 (2009) 853–856
Fig. 1. Room temperature XPS analysis associated to the O-1s electronic transition (oxide and hydroxide components), realized on two samples. One deposited under pure argon gas (reference) and another under hydrogen/argon mixture (20% H2 + 80% Ar). The inset shows the FT–IR spectra realized at room temperature on BaTiO3 films grown under different H2-ratios (0, 10 and 20%) in the sputtering gas.
oxide remains constant. Such behavior may results from the incorporation of hydrogen during the deposition process. Further evidence for hydrogen incorporation can be given by FT–IR reflectance spectroscopy (see inset in Fig. 1) as a broad band around 3500 cm− 1. This IR absorption band is ascribed to the O\H stretching vibration [10,11]. Such feature is absent in standard films (grown under pure argon). The large increase of the O\H stretch magnitude with increasing the H2-ratio in the sputtering gas clearly manifests that a large density of protons has been incorporated into the films during the elaboration process. As it is already established [10,12,13] the hydrogen ionizes to give an electron (donor defect) and a proton (H+) which combines with oxygen of the perovskite lattice to form hydroxide group (OH˙O). Fig. 2 shows an overview on the frequency spectra associated to the real part of permittivity (ε'), its imaginary counterpart (εq) and the ac conductivity (σac) performed at − 60 °C on a-BaTiO3:H film (20% H2 + 80% Ar). Films grown under H2-containing atmosphere display hundred times higher capacitance than the one measured for films grown under pure argon [3]. This feature was explained through an electrode polarization mechanism [14] which arises from proton pileup at the cathode over a Debye length [3]. At low frequency domain, measured capacitance was completely determined by the charge stored in the double layer. As a result, electric double-layer capacitors have been widely regarded as energy storage devices. The double-layer is clearly
hydrogen ratio. This dark color indicates that the hydrogenated films are nonstoichiometric with some oxygen deficiency [8,9]. In addition, chemical properties of the hydrogenated films undergo dramatic alteration. Fig. 1 displays a part of the XPS spectrum associated to the O-1s electronic transition (oxide and hydroxide components) and realized on two films. One deposited under pure argon gas (reference) and another under hydrogen/argon mixing gas (20% H2 + 80% Ar). We show that the oxide component of both films is detected at around 530 eV and is assigned to oxygen in the BaTiO3 lattice. This indicates that the elementary lattice does not undergo any change when hydrogen was added to the sputtering gas. On the other hand, based on the XPS analysis and by comparing the peak-area attributed to the oxide and hydroxide components of the analyzed films (reference and hydrogenated) we can determine the evolution of these components when hydrogen was added to the sputtering gas. We note that films grown under hydrogen containing atmosphere display large density of hydroxide that is 12% higher than the one determined for films grown under pure argon while the density of
Fig. 2. Dielectric permittivity (real and imaginary parts) and conductivity measured at –60° C on BaTiO3 layers sputtered under pure argon gas (empty symbols) and under hydrogen/ argon mixture (full symbols).
Fig. 3. Plot (a) shows the temperature dependence of the loss factor, tanδ(f). Films were grown under hydrogen/argon mixture (20% H2 + 80% Ar). Plot (b) illustrates the evolution of the relaxation frequency as a function of 1000 / T (Arrhenius diagram).
F. El Kamel, P. Gonon / Solid State Ionics 180 (2009) 853–856
Fig. 4. Plot (a) presents the temperature dependence of the leakage current (J–E). In figure (b) we replotted the J–E characteristics according to the Poole–Frenkel model.
evidenced, especially at low temperatures by a dispersive behavior of the dielectric constant accompanied by a relaxation peak in the loss measurements. We believe that the dielectric losses are mainly the result of the ion migration. At high frequencies, ε' relaxes toward the value obtained when films are deposited under pure argon (~20 [14]). Furthermore, σac(f) is also affected by the addition of hydrogen. At low frequency domain, it increases by around 7 orders of magnitude and displays two plateaus separated by a step-like behavior which implies a relaxation process. The high-frequency plateau corresponds to the bulk response and the low-frequency one is assigned to the electrode polarization effect. In order to better illustrate the relaxation process, we plotted in Fig. 3a the frequency dependence of the loss factor, tanδ(f). As we can see, the dielectric loss comes from two mechanisms, resistive loss (power law evolution) and relaxation (peak). The resistive mechanism can be explained by mobile charges transfer at the metal– dielectric interface [9,14] whereas the relaxation process is originated by the double-layer modulation under the alternating field [9,14]. As already evidenced in Fig. 3a, the relaxation frequency (peak position, f0) shifts to high-frequencies side as the temperature increases, according to an Arrhenius law, τ0 = τ0,max exp (Ea / kBT). τ0 = 1 / f0 is the relaxation time, f0 is the peak position of tanδ(f) in the frequency domain, Ea is the activation energy of the relaxation process and kB is the Boltzmann's constant. A plot of f0 against 1000 / T (Fig. 3b) gives the activation energy Ea, of the relaxation process. We can find out that this process shows an Arrhenius-type temperature dependence over the whole temperature domain and is thermally activated with 0.16 eV when temperature is lower than − 75 °C and with 0.24 eV when temperature goes beyond this value. This behavior can be
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compared to the one depicted on the bulk conductivity (extracted from the high-frequency plateau in the σac(f) characteristics). Based on the fact that f0 = (σ / 2πεrε0)(2LD / L)1/2, where σ is the bulk conductivity, εr and ε0 are respectively the real part and the free space permittivity, L is the layer thickness and LD is the Debye length [14], it appears that the bulk conductivity and the relaxation peak have the same activation energy. Such a feature was confirmed experimentally for samples grown under different H2-ratios (13, 17, 20 and 25%) in the sputtering gas. We extracted three activation energies, corresponding to three conduction processes. These processes dominate the dielectric response in three different temperature domains. At temperature below −75 °C, the transport phenomenon is thermally activated with an activation energy of about 0.15 eV. When temperature varies from −75 to − 30 °C, activation energy is close to 0.25 eV. Finally, at temperature above −30 °C, the determined activation energy is around 0.37 eV. Activation energies evaluated from the slope of the Arrhenius plots can give an idea on the transport phenomena taking place at these temperature ranges and on the electrical active defects that are at the origin of the relaxation and conduction processes. As we explained in our previous works [14–16], the mobility of oxygen vacancies occurs at sufficiently high temperature and it is thermally activated with an activation energy in the vicinity of 1 eV. Such a value is much higher than those extracted from the dielectric measurements in the present study. So, the effect of oxygen vacancies can be ruled out from the explanation of the observed dielectric response at low temperature. On the other hand, we know that the incorporation of protons in the BaTiO3 films could significantly affect their electrical properties. In order to speculate the conduction mechanism in dc regime, leakage currents were measured on BaTiO3:H films as a function of the applied field at different temperatures (from − 150 to −50 °C). dc bias was applied to the bottom electrode and the current was measured after 60 s for stabilization concern. Experimental data measured on a sample grown under hydrogen containing atmosphere (17% H2) are shown in Fig. 4a. It is clearly seen that J–E characteristics exhibit a Poole–Frenkel type behavior, which implies that the conduction mechanism can be described by a thermally stimulated emission from a discrete traps distribution. Thus, the leakage current density can be expressed as, qffiffiffiffiffiffiffiffi1 0 −q /t − πqE eopt A JPF = qN c μEexp@ ð1Þ kB T where E is the applied field, q is the elementary charge, kB is Boltzmann's constant, Nc is the effective density of states in the
Fig. 5. Evolution of the effective trap depth as a function of E1/2. We found that the leakage current can be explained by carrier detrapping from shallow traps level ϕ0 localized at 0.145 eV below the conduction band.
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conduction band assumed to be around 1018 cm− 3 [17], μ is the electronic mobility, qϕt is the trap level and εopt is the optical permittivity. As stated above, log(J /E) vs. E1/2 should be linear according to the Poole– Frenkel model (Fig. 4b). Then, we guess that the measured leakage current can be explained by carriers detrapping from shallow trap levels localized within the bandgap. In order to determine the activation energy of such process, we plotted the J(1 / T) characteristics (at a fixed electric field). J(1 / T) curves (not shown here) depict a linear behavior with a negative slope which suggests that the conduction process is thermally stimulated. Activation energy and mobility vary respectively from 0.12 and 4.31 × 10− 7 to 0.05 eV and 2 × 10− 6 cm2/Vs when the electric field varies from 0.2 to 3 MV/m. Such energies correspond to the effective trap depth [ϕ0 − Δϕ] which is strongly dependent on the electric field. Based on the slowness of the Poole–Frenkel mechanism, these mobility's values are in a good agreement with the electronic conduction. In the literature, electronic mobility shows a large discrepancy and is found to vary from 10− 3 [18–22] to 10− 10 cm2/Vs [17]. To determine the real depth ϕ0 (at E = 0 V/m) of traps, we plotted the effective trap depths as a function of E1/2 (see Fig. 5). We found that the leakage current can be explained by carrier detrapping from shallow traps level ϕ0 localized at 0.145 eV below the conduction band. It is interesting to note that the same amount of energy is required for activating the relaxation process taking place at temperature below −75 °C. So, it is believed that the loss peak and the leakage current arise from the carriers detrapping from shallow traps localized at around 0.15 eV within the bandgap. Positively charged protons can provide these energetic levels. In less symmetric environments, charged defects in the vicinity of the proton tilt markedly the hydroxide group toward one of the neighboring oxygen ions. Since the hydrogenated samples were deposited under a reduced atmosphere, oxygen vacancies content increases and constitute the main charged defect. Thus, surrounding oxygen vacancies effectively act as positively charged defects that repel the proton. This feature allows the formation of a rather weak directional interaction (hydrogen-bond) between the proton and an adjacent oxygen, O(1)\H··O(2) [23]. Such a site might effectively act as a trap for the proton. In this approach, electron issued from the ionization of hydrogen is weakly bonded and it can be easily activated to the conduction band to increase the dielectric loss and the leakage current [24]. Such a structure generates a shallow donor level below the conduction band. This level was determined at around 0.15 eV using the leakage measurements. Moreover, the hydrogen-bond formation can be responsible of a large protonic conductivity in oxide materials. With the aim of examining protonic conduction, previous reports [25] adopted a model for OH oscillators localized on regular oxygen sites and undergoing a stretching mode alongside the “oxygen–oxygen” direction. This is precisely the situation that allows for proton oscillation to take place involving small ionic displacements (less than an atomic spacing) which require small activation energy. Due to their weakness, the hydrogen bonds can be dissociated even at low temperature with an activation energy around 0.22 eV [26]. This value agrees well with the one (0.25 eV) determined using dielectric measurements at temperature ranging from −75 to − 30 °C. Then, in this temperature range, the dielectric response can be ascribed to the dissociation of hydrogenbonds and consequently to a localized migration (or oscillation) of the proton alongside the hydroxide bond. Such process was previously emphasized and discussed by Weber et al. [25] in ionic-conductor oxides. Finally, the relaxation process considered at temperature higher than − 30 °C for hydrogenated samples was thermally activated with
0.37 eV. It is believed that this amount of energy is required to activate the proton diffusion within the oxide material. The reported value of the activation energy should be strongly affected by the amorphous state of the material but it agrees well with predicted experimental [27,28] and theoretical [27–29] values (0.3–0.6 eV) for the best proton conductors such as Y–BaZrO3 or BaCeO3 [27,28,30]. Based on the experimental observation and on the light of literature mentioned above, the conduction mechanism is predominately of Grotthuss-type [31], i.e. involving proton transfer (hopping) from the hydroxide defects to a nearest neighboring oxygen ion. 4. Conclusions Hydrogen incorporation in barium titanate films during the low temperature deposition process can contribute to the transport phenomena by introducing several features. At temperature higher than −30 °C protons contribute to the conduction mechanism as mobile ionic species. Such mechanism is thermally activated with 0.37 eV. In addition, protons have been proposed to induce n-type conductivity by generating shallow donor levels within the bandgap in titanates. Therefore, the trap we detect at 0.15 eV could be associated with such a donor level. Moreover, in disordered systems, protons were shown to engender hydrogen bonds in the structure. These bonds have to be related to a common feature for the oxides protonic conductors. As a result, the dielectric relaxation observed at temperature ranging from − 75 to − 30 °C could be related to the oscillation of the hydroxide bond due to the destruction of the hydrogen bond. Activation energy for this process is around 0.25 eV. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31]
L. Ma, Y. Yang, Appl. Phys. Lett. 87 (2005) 123503. M. Hendriks, M. Heijman, W. Zyl, J. Elshol, H. Verweij, J. Appl. Phys. 90 (2001) 5303. P. Gonon, F. El Kamel, Appl. Phys. Lett 90 (2007) 232902. J.Y. Kim, I.J. Chung, J. Electrochem. Soc. 149 (2002) A1376. P.G. Romero, M. Chojak, K.C. Gallegos, P.J. Asensio, N.C. Pastor, M.L. Cantú, Electrochem. Commun 5 (2003) 149. F. Lufrani, P. Staiti, Electrochim. Acta 49 (2004) 2683. I.H. Pratt, S. Firestone, J. Vac. Sci. Technol. 8 (1970) 256. S. Furusawa, H. Tabuchi, T. Sugiyama, S. Tao, J. Irvine, Solid State Ionics 176 (2005) 553. F. El Kamel, P. Gonon, F. Jomni, B. Yangui, J. European Ceram. Soc. 27 (2007) 3807. R. Waser, J. Am. Ceram. Soc. 71 (1988) 58. S. Aggarwal, S.R. Perusse, C.W. Tipton, R. Ramesh, H.D. Drew, T. Venkatesan, D.B. Romero, V.B. Podobedov, A. Weber, Appl. Phys. Lett. 73 (1998) 1973. J. Ahn, P. McIntyre, L. Mirkarimi, S. Gilbert, J. Amano, M. Schulberg, Appl. Phys. Lett. 77 (2000) 1378. K. Xiong, J. Robertson, Appl. Phys. Lett. 85 (2004) 2577. P. Gonon, F. El Kamel, J. Appl. Phys. 101 (2007) 073901. F. El Kamel, P. Gonon, L. Ortega, F. Jomni, B. Yangui, J. Appl. Phys. 99 (2006) 094107. F. El Kamel, P. Gonon, F. Jomni, B. Yangui, J. Appl. Phys. 100 (2006) 054107. Y.B. Lin, J.Y. Lee, J. Appl. Phys. 87 (2000) 1841. S. Zafar, R. Jones, B. Jiang, B. White, P. Chu, D. Taylor, S. Gillespie, Appl. Phys. Lett. 73 (1998) 175. C.N. Berglund, W.S. Baer, Phys. Rev. 157 (1967) 358. A.J. Moulson, J.M. Herbert, Electroceramics: Materials, Properties and Applications, Chapman and Hall, London, 1996. D. Keroack, Y. Lepine, J.L. Brebner, J. Phys. C. 17 (1984) 833. J.P. Boyeaux, F.M. Calendini, J. Phys. C 12 (1979) 545. L. Pejov, Chem. Phys. Lett. 376 (2003) 11. L. Dobaczewski, K.B. Nielsen, N. Zangenberg, B.B. Nielsen, A.R. Peaker, V.P. Markevich, Phys. Rev. B. 69 (2004) 245207. G. Weber, S. Kappham, Wöhlecke, Phys. Rev. B. 34 (1986) 8406. J.O.M. Bockris, A.K.N. Reddy, Modern Electrochemistry, Plenum, New York, 1970. W. Münch, K.D. Kreuer, G. Seifertlib, J. Majer, Solid State Ionics 125 (1999) 39. E. Matsushita, A. Tanase, Solid State Ionics 97 (1997) 45. E. Matsushita, T. Sasaki, Solid State Ionics 125 (1999) 31. K.D. Kreuer, Solid State Ionics 125 (1999) 285. K.D. Kreuer, Chem. Mater. 8 (1996) 610.