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Systematic tuning of the conduction mechanisms in ferroelectric thin films
This content has been downloaded from IOPscience. Please scroll down to see the full text. 2013 J. Phys.: Condens. Matter 25 495901 (http://iopscience.iop.org/0953-8984/25/49/495901) View the table of contents for this issue, or go to the journal homepage for more
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IOP PUBLISHING
JOURNAL OF PHYSICS: CONDENSED MATTER
J. Phys.: Condens. Matter 25 (2013) 495901 (13pp)
doi:10.1088/0953-8984/25/49/495901
Systematic tuning of the conduction mechanisms in ferroelectric thin films D Levasseur1,2 , E Bouyssou2 , R De Paolis3 , A Rousseau1 , F Coccetti3 , G Guegan2 , S Payan1 and M Maglione1 1 2 3
CNRS, Univ. Bordeaux, ICMCB, UPR 9048, F-33600 Pessac, France ST Microelectronics, 16 rue Pierre et Marie Curie, F-37100 Tours, France CNRS, Univ. Toulouse, LAAS, F-31400 Toulouse, France
E-mail: [email protected]
Received 17 July 2013, in final form 10 September 2013 Published 6 November 2013 Online at stacks.iop.org/JPhysCM/25/495901 Abstract We have investigated the macroscopic and microscopic properties of large sets of Ba0.7 Sr0.3 TiO3 thin films including several substitution rates of manganese. Thanks to a high degree of control of the processing parameters at each stage we have been able to find a link between the dc leakage current and the low and high frequency dielectric permittivity and losses. We supplemented these macroscopic observations with in depth investigations of the defect states through x-ray photoelectron spectroscopy. We found that both the leakage current and the extrinsic dielectric parameters arise from a large density of charged point defects related to oxygen vacancies. At the outer surfaces of the films, the density of such charged defects is so high that it can raise the Fermi level to close to the conduction band. Such degradation of the films’ performance can be relieved by appropriate manganese substitution for the titanium host ions. Such doping is able to move back the Fermi level to close to the center of the bandgap thus changing the conduction process from interfacial Schottky to bulk Poole Frenkel and decreasing the extrinsic losses. This beneficial effect was already inferred in ceramics and thin films but we have established a clear link between the macroscopic parameters and the microscopic defect state. This model can be transferred to many high permittivity oxides. (Some figures may appear in colour only in the online journal)
1. Introduction
able to reduce the detrimental influence of residual oxygen vacancies which are donor centers [8–12]. When brought to thin films, the same materials raised the same questions: what is the contribution of oxygen vacancies to the residual currents; what about the contribution of interfaces; is the surface more defective than the bulk [13–19]? At the same time, the dielectric properties of ceramics and thin films of the same materials were also investigated to improve their performance [20]. The reduction of losses mostly in the high frequency range (f > 1 GHz) is a real challenge to enable practical application of ferroelectric thin films in microwave devices [21]. This was achieved thanks again to the control of charged defects and also to mixing with very low loss materials like MgO [22–24]. So as to avoid unnecessary complexity resulting from the contribution of ferroelectric domains and domain walls to these losses, the
Finding a link between the static electrical conductivity and the dynamic losses of dielectric materials is not only a key for their actual applications but also a very basic problem which is largely unsolved. In the case of very high permittivity oxides like ferroelectrics, both phenomena have been extensively investigated. Focusing on BaTiO3 -derived materials, the knowledge and control of residual conductivity in pure and substituted ceramics have been improved over the years [1–7]. The main tool was and is still the compensation of unwanted charged centers through accurate control of the defect chemistry. Including the oxygen stoichiometry, this is a very hard task which has resulted in still active debates after decades of investigation. The main trend is that the substitution of acceptor ions (e.g. Mn3+ on Ti4+ sites) is 0953-8984/13/495901+13$33.00
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c 2013 IOP Publishing Ltd Printed in the UK & the USA
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current density at different temperatures were carried out in order to determine its time dependence (J–t). For each temperature, five different capacitors were tested in order to get more representative results. The electrical field dependence of the leakage current was then studied, denoted J–E for convenience. For each wafer eight capacitors were tested, and a very good reproducibility was w obtained; therefore the following results are always representative of the whole wafer. The data processing will be explained later on. The refractive indices of the films were determined by spectroscopic ellipsometry, using a Nanofilm EP3 Imaging Ellipsometer. The 1/9 data were recorded in the visible range, from 502 to 720 nm. Using the EP4Model software, a Cauchy law was used to fit the data and determine the refractive index dispersion n(λ) of each film. The root mean square errors obtained after fitting were always lower than 3. A VG Scientific 220 i-XL ESCALAB spectrometer was used for the XPS analysis with a monochromatized Al Kα source (hν = 1486.6 eV) at 70 W. The spot size was about 200 µm in diameter. In depth XPS analysis was carried out using argon etching. At each step, full spectra (0–700 eV) were recorded with a constant pass energy of 150 eV and high resolution spectra at a constant pass energy of 20 eV. The high resolution spectra were then fitted and deconvoluted using the AVANTAGE software provided by ThermoFisher Scientific. Finally, the dielectric properties at low frequencies were studied on integrated MIM structures, with 0.063 mm2 Pt top electrodes. The capacitance–voltage (C–V) and dielectric loss–voltage (tan δ–V) characteristics of the MIM capacitors were investigated at 30 ◦ C using a Cascade Microtech probe station and an Agilent 4294A Impedance Analyzer. These C–V curves were measured with a 10 kHz/0.1 V ac signal, and the dielectric permittivity was then calculated using capacitance, and the dimensional parameters: the thickness d obtained from cross section SEM images and top electrode area A. Concerning the RF characterization process, one-port reflection measurements were performed on MIM capacitors at room temperature by means of an Anritsu 37397C VNA with internal bias tee and Picoprobes with 150 µm GSG coplanar pitch. An SOL calibration procedure was applied beforehand. Subsequently, 1 GHz permittivity was computed using an equivalent circuit consisting of a capacitor and a resistor in series. More specifically, a differential method was adopted in order to allow the removal of the most prominent parasitic effects [32, 33]. Finally, the BST permittivity as a function of the electric field was extracted from the capacitance values associated with every different applied voltage.
focus is more on the solid solution Bax Sr1−x TiO3 (BST) which is paraelectric at room temperature for some given composition x. However, the exact origin of these dielectric losses, whether intrinsic or extrinsic, is still a matter of debate and the contribution of charged defects to the extrinsic losses is still to be probed. In this paper, we focus on dc conductivity and high frequency dielectric properties in BST films substituted with manganese. As already mentioned, such films have already been investigated either for their leakage current [25–29] or for their high frequency dielectric properties [30, 31]. The first contribution of this paper is to systematically address both these issues in the same films. Another interest in our strategy was to investigate films of microelectronics industry grade so as to ensure reproducibility over sets of thousands of samples. This enabled us to ensure relevant statistics in our experimental results. By doing so, we were able to show that the dc leakage current gradually changes from interface limited Schottky to bulk limited Poole Frenkel type as the atomic concentration of Mn is raised from 0.1% to 2%. Such a trend was linked to the down shift of the Fermi level from the bottom of the conduction band toward the center of the bandgap by means of x-ray photoelectron spectroscopy (XPS). We then mapped out the dielectric dispersion from 1 kHz up to 1 GHz and we could show that the low frequency permittivity at f < 1 GHz is artificially raised by the residual charged defects for undoped films, and is gradually suppressed under Mn doping. This clearly shows a strong link between the dc current and the dielectric properties. Based on our investigation of the electronic state, we suggest that this common origin results from the Ti3+ states associated to oxygen vacancies. We can further state that these defects are all the more relevant when they are accumulated at the interfaces between the films and the metal electrodes. We thus can draw a consistent picture of the electric and dielectric behavior of Mn doped BST films. These results are of great interest for tunable microwave devices, where a very high level of understanding is needed in both dielectric and leakage current properties, in order to comply with the specifications of such components.
2. Experimental procedure The Ba0.7 Sr0.3 Ti1−δ Mnδ O3 (BST70:Mn) precursor sols used in this study were purchased from Mitsubishi Material Corporation, Japan, with δ = 0, 0.001, 0.005, 0.01, and 0.02. Thin films from these sols were deposited by spin coating on industrial Pt/Si substrates. For each sample, the solution was coated on the substrate to form a layer, this layer was pyrolyzed, and the process was repeated in order to obtain the desired thickness. Finally the whole film was post-annealed under oxygen at 800 ◦ C. In order to investigate the leakage currents of the films, integrated metal–insulator–metal (MIM) structures were processed with 0.004 mm2 Pt top electrodes. These capacitors were characterized through a Cascade Microtech probe station connected to a Keithley 2636A SourceMeter Unit. Constant voltage stress (CVS) measurements of the
3. Results and discussion 3.1. Leakage current study In order to properly investigate the leakage currents of the samples, the current versus time behavior was first studied, providing important information such as current 2
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density levels, and the true leakage current region is shortened to ∼5 s at 125 ◦ C and suppressed at 165 ◦ C. After the flat range ¬, the current starts to increase until a maximum value Jmax is reached; this regime is called the ‘resistance degradation’ (). According to Waser et al [2], this phenomenon is highly related to the presence of oxygen vacancies in the material. After some time under a constant electric field, the charges of such defects can start to migrate toward the cathode. Thus, the spatial modification of the carrier concentrations between the two electrodes will change the conductivity of the film. The accumulation of charged oxygen vacancies close to the cathode induces modifications of the band structure and thereby interfacial energy barriers. As a consequence, the current density increases. Once the degradation current has reached its maximum Jmax , the current density starts to decrease (®). Bouyssou et al attributed this phenomenon to a trapping effect and space charge reorganization and they proposed a mathematical function describing this phenomenon [35]. In figure 1(b), the effect of Mn doping on the J–t characteristics at 200 ◦ C can be seen. This temperature has been chosen in order to better highlight Jmax . It is shown that the Mn doping strongly decreases the maximum level Jmax reached after the resistance degradation, until a Mn concentration of 1%. It seems, however, that when the Mn content is too high, the breakdown process happens more rapidly, as shown for the BST70:Mn2% at around 400 s at 200 ◦ C under 750 kV cm−1 . Zafar et al proposed a quantitative model for the increase of the current densities, based on the assumption that the current increase during the resistance degradation is mainly due to decrease of the effective barrier height [36]. The maximum decrease of the barrier, denoted 18max , is given by
Figure 1. The time dependence of the current density J obtained for 750 kV cm−1 electric field stress (a) for the undoped BST70 film at different temperatures (85, 125, 165 and 200 ◦ C), (b) at 200 ◦ C for each doping rate.
18max = kT ln(Jmax /Jinit )
(1)
where Jmax is the maximum current density and Jinit is the value obtained by extrapolating the current density at t = 0. The computed 18max decreases from 0.17 eV for the undoped sample to 0.08 eV for the BST70:Mn1%. This stabilization of the effective barrier height can be ascribed to a decreased density of free charges under Mn doping. Because of the extensive leakage current evolution as a function of time, special care must be taken subsequently in the J–E spot measurements, in order to avoid any experimental artifacts.
relaxation, resistance degradation and breakdown. Then the leakage current versus electric field was analyzed. These measurements allowed us to study the conduction mechanisms in the BST thin films, and the influence of the Mn doping. 3.1.1. Study of the time dependence. We first investigated the time dependence of the leakage currents of the capacitors. In figure 1(a), the current density of the undoped BST sample as a function of time is presented, under a 750 kV cm−1 external field for different temperatures from 85 to 200 ◦ C In the initial time period, no decrease in current due to the dielectric relaxation is observed, as opposed to many studies [34]; this is likely to be due to the high temperature and high field that we applied. Some typical trends can be drawn from the curves numbered from ¬ to ® in figure 1(a) [34, 35]. For the lowest temperature of 85 ◦ C, the current starts with a flat time dependence corresponding to the so-called true leakage current (¬) for ∼500 s. When the temperature is increased, the plots are globally shifted toward higher current
3.1.2. Study of the electric field dependence. In order to remove time artifacts, the following characterization protocol has been established to determine the influence of the Mn concentration on the leakage current. To avoid relaxation current during the J–E measurement at low temperatures, a short constant voltage stress (CVS) was carried out for 100 s at 85 ◦ C with regard to the flat ‘true’ leakage current dependence observed in figure 1. Because the relaxation is a reversible process, the wafers were then quenched at ambient temperature by contact on a marble plaque so as to keep the material in the stressed state. 3
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Figure 2. Current density as a function of electric field for the undoped BST70 film at different temperatures (from 30 to 200 ◦ C). The voltage was applied to the top electrode. Inset: the schematic staircase profile of the dc voltage applied during J–E measurements; td is the delay time, tm the measuring time and 1V the step height.
Figure 3. Current density as a function of electric field at 30 ◦ C of undoped BST and Mn doped BST from 0.1% to 2%. The voltage was applied to the top electrode.
ones, Poole Frenkel and space charge limited currents (SCLCs) are reported. Concerning the tunneling of free charges through the whole film or only a part of the film, e.g. a grain boundary energy barrier, the mechanism seems unlikely regarding the thickness of the film and the weak band bending at grain boundaries for thin films [16]. In addition, the strong temperature dependence of the leakage current cannot usually be observed with tunneling models. Concerning the SCLC, a low barrier height is required to guarantee sufficient charge injection. According to Dietz et al SCLC is unlikely at metal contacts on high-gap semiconductors or insulators [16]. In view of these concerns about tunneling and SCLC, our study will therefore focus only on Schottky and Poole Frenkel (PF) mechanisms. In order to determine whether the leakage current can be associated to a Schottky regime or a Poole Frenkel injection, the physical parameters involved in the mechanism must be determined. Hence, for each sample the two following protocols, based on the J–E measurements obtained at different temperatures (e.g. figure 2), have been applied.
After the pre-stress, the current versus voltage characteristics were measured for each capacitor at different temperatures from 30 to 200 ◦ C, with a typical staircase profile presented in the inset of figure 2. The voltage was applied for 400 ms (td ) before measuring the current and the step height (1V) was 0.25 V. The behavior of the MIM capacitors was symmetrical, that is to say the same results were obtained whether the voltage was applied to the top or bottom electrode. Typical results obtained are shown in figure 2, where the current density versus electric field of the undoped BST70 film biased on the top electrode is presented. A strong dependence of the current density on the temperature can be noted. For each sample a similar set of plots has been obtained. The influence of increase of the doping rate at 30 ◦ C is presented in figure 3. Increase of the Mn content dramatically decreases the leakage current, as expected, by almost two decades for the BST70:Mn1% and more than three decades in the case of BST70:Mn2%. In addition, it can be seen that the shape of the J–E plots is modified by the Mn insertion; therefore a modification in the conduction mechanism can be expected.
Schottky regime In the case of a Schottky mechanism, the conduction is governed by charges thermally injected at the interface, by lowering the barrier height due to the applied field and image force. The current evolution is described by the equation q qEext 80 − 4πε ε 0 i (2) J = A∗ T 2 exp − kB T
3.1.3. Study of the conduction mechanisms. Many results on the conduction mechanisms in perovskite-type titanate thin films, including SrTiO3 , BaTiO3 , (Ba,Sr)TiO3 and Pb(Zr, Ti)O3 , have previously been reported [15–17, 35, 37–39], and the main physical models describing the leakage currents in the films have been discussed. In the case of an MIM capacitor, the leakage current is either controlled by the interfaces with the electrodes or by the bulk of the film. In most cases, one of them is highly dominant. For the interface-controlled mechanisms, tunneling and Schottky regimes can be found, and concerning the bulk-controlled
where A∗ , q, ε0 and kB are respectively Richardson’s constant, the electron charge, the vacuum permittivity and Boltzmann’s constant, T is the temperature, Eext the external electric field, ϕ0 the barrier height, and εi the relative permittivity at optical 4
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Figure 4. J/T 2 as a function of the square root of the bias, at different measurement temperatures from 30 to 200 ◦ C, for the undoped BST70 film. The voltage was applied to the top electrode. The black lines correspond to the exponential fits y = a ebx , and the equations obtained are written next to the corresponding temperatures.
frequencies [40]. This equation can be turned into " √ # J βSc V = α exp kB T T2
Figure 5. Arrhenius plot, ln(J/T 2 ) as a function of 1000/T, for the undoped BST film. The voltage was applied to the top electrode. The black lines correspond to the linear fits y = ax + b, and the equations obtained are written next to the corresponding biases.
Poole Frenkel effect In the case of the Poole Frenkel emission, the conduction is governed by carriers that are thermally emitted from trapped centers under a strong electric field. The current versus temperature and electric field follows the equation q qEext 8t − πε 0 εi J = qµEext Nc exp − (7) kB T
(3)
where 80 α = A∗ exp − and kB T (4) r q . βSc = 4π ε0 εi d √ Then, by plotting TJ2 = f ( V) for each measured temperature, and using an exponential fit y = a ebx with a and b as fitting parameters, as shown in figure 4, from b and equation (3) the parameter βSc for each temperature can be determined. Then, from equation (4), the optical dielectric permittivity εi of each sample has been computed. Because the effective Richardson’s constant is highly dependent on the material and the electrode [16], the barrier height cannot be determined directly from α (equation (4)). Thus, the barrier height has been calculated using the Arrhenius law, applied to this case, J −Ea = C exp (5) 1 kB T T2
where µ and Nc are respectively the carrier mobility and the carrier density in the conduction band, 8t is the trap barrier height [41]. Turning this equation into " √ # βPF V J 0 = α exp (8) V kB T where α0 =
qµNc 8t exp − d kB T r q βPF = π ε0 εi d
and (9)
√ = f ( V) was plotted for each measured temperature. Then, using an exponential fit y = a ebx and equation (8), the parameter βPF of each measured temperature can be determined, and so from equation (9), computation of εi of each sample has been processed. As µ and Nc are unknown values, ϕt cannot be directly obtained, and Arrhenius curves were plotted. J −Ea = C2 exp (10) V kB T J V
where C1 is a constant and Ea is the activation energy. From equation (3), this activation energy can be written as s qV Ea = 80 − . (6) 4π ε0 εi d
where C2 is a constant and Ea is the activation energy. Therefore the activation energy can be written as s qV Ea = 8t − . (11) π ε0 εi d
The validity of equation (5) is readily seen in figure 5 where ln(J/T 2 ) is plotted versus 1/T. From the fits of these curves, Ea (V) can be extracted and then the barrier height 80 from equation (6). 5
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Figure 6. Current density as a function of bias of the undoped BST70 sample for different temperatures from 30 to 175 ◦ C. The scatter plots correspond to the experimental values, and the black lines correspond to the fits with (a) the Schottky injection law and (b) the Poole Frenkel emission law. The voltage was applied to the top electrode.
From a fit using equation (11), Ea (V) can be extracted and then ϕt the trap barrier height (equation (11)). To sum up, in the cases of both the Schottky and the Poole Frenkel mechanisms, a couple of parameters can be extracted, the dielectric permittivity at optical frequencies εi and the interfacial or trap barrier height, respectively 80 and 8t . From these parameters, the J–E characteristics predicted by the Schottky and Poole Frenkel equations, respectively equations (2) and (7), can be plotted for a given temperature. Knowing the thickness of the thin film and the constant values, these plots can be fitted to the experimental characteristics using the effective Richardson’s constant A∗ as a fitting parameter in the case of Schottky and µNc in the case of Poole Frenkel. An example of these curves is given in figure 6 for the undoped BST sample. As can be seen, both the Schottky and Poole Frenkel laws fit the experimental data fairly well, for voltages higher than 5 V. The low voltage part of the fit was not taken into account, due to the nonlinearity of the capacitance, and to band profile evolution in early applied voltages such as in the Ohmic regime [42]. So as to validate these results, the values of εi should satisfy the equation εi = n2
3.1.4. Influence of the Mn doping in BST70 on conduction mechanisms. The leakage current is decreased on increasing the Mn content in the material, as shown before in figure 3. However, this impurity insertion may also impact the conduction mechanism in the material. Ahn et al found that Mn doping in Ba0.5 Sr0.5 TiO3 films decreases the tunneling current by increasing the depletion layer width near the electrodes [27]. In our case, we observed a strong temperature dependence of the leakage currents, ruling out the possibility of such a mechanism as stated before; however, it seems clear in the literature that acceptor Mn doping influences the interfaces strongly [18, 47]. We focus here only on 0.1 to 1 at.% Mn content. For BST70:Mn2%, the J–E measurements (not shown here) were not reliable, as from 85 ◦ C a deterioration was observed and the ‘true’ leakage current could not be measured any longer because of the early breakdown shown in figure 1(b). Experimental J–V characteristics of BST70:Mn0.1%, BST70:Mn0.5% and BST70:Mn1% and the corresponding simulated plots calculated with the Schottky and Poole Frenkel models are presented in figure 7. For the undoped sample, determination of the dominant mechanism is really not obvious from these charts, since both models seem possible. Besides, a very marked trend can be noticed in the extracted refractive index shown in figure 8(a). For the undoped BST70, the n value calculated from the Schottky model perfectly matches with the experimental index measured by ellipsometry, as shown before. However, when the Mn content increases, this computed refractive index decreases to reach a value of 1.05 for BST70:Mn1%, almost that of vacuum. Additionally, for the index extracted from the Poole Frenkel process, at first for the undoped sample the index value is too large, but it decreases with increase of Mn content, and reaches 2.8 for BST70:Mn1%. The index extracted from PF seems to tend to the measured index values for high doping rates. Therefore, the following hypothesis can be made. For the undoped BST material, the leakage is governed by the Schottky mechanism. For BST70:Mn1%, this mechanism seems impossible considering the very weak extracted value of n, but the Poole Frenkel mechanism could be much more
(12)
where n is the refractive index of the material [43]. The n values extracted with our protocols were 2.09 and 6.85 respectively for Schottky and Poole Frenkel, when the voltage was applied to the top electrode, and 2.18 and 7.36 respectively when applied to the bottom electrode. However, the measured refractive index from ellipsometry was between 2.06 and 2.18, in the wavelength range of 501.1 to 719.8 nm. Therefore, the dominating conduction mechanism in the undoped film can clearly be attributed to the Schottky process. This is in agreement with previous studies performed in BST70 [44] and other compositions of barium strontium titanates [16, 45], showing that Schottky is the leading mechanism. Dietz et al also reported that an Ohmic contribution to the low-field region shifted to a Schottky effect for higher voltage [46], explaining the low-field misfit in figure 6. 6
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Figure 7. Current density as a function of bias of (a) BST70:Mn0.1%, (b) BST70:Mn0.5% and (c) BST70:Mn1% for different temperatures from 30 to 175 ◦ C. The colored scatter plots correspond to the experimental values, the continuous black lines to the Schottky fits and the dashed black lines to the Poole Frenkel fits. V+ means that the voltage was applied to the top electrode.
Figure 8. Physical parameters, extracted from the Schottky and Poole Frenkel models, as a function of Mn content from 0 to 1%. (a) Refractive index (the measured index is plotted in a continuous line as a guide for the eyes but has only been measured for each doping rate), (b) interfacial barrier height from the Schottky mechanism, (c) trap barrier height from the Poole Frenkel mechanism. +te and +be respectively mean that the voltage was applied to the top or the bottom electrode.
similar to the example shown in figure 2 was obtained. Then, the two protocols presented in part 3.1.3 were applied to these Jmix plots, and a couple of refractive indices was determined with the Schottky and PF models for each κ value. It turns out that these refractive indices perfectly match with those determined by experiment as shown in figure 9. Moreover, this means that each mixing coefficient can be related to some Mn content. According to the results, with 0.07 at.% of Mn in the material, 90% of the leakage current would be governed by the Poole Frenkel mechanism, whereas when the Mn doping reaches 0.4%, 99% of the conduction would be led by Poole Frenkel. This result should be taken with caution: BST70:Mn1% with a refractive index of 2.8 may still have conduction due to Schottky injection, even if it is dominated by Poole Frenkel. The idea here is just to show through the extracted optical index that a trend is obtained to confirm a mix between the two mechanisms.
relevant giving a refractive index of 2.8. For intermediate Mn content, a mix of Schottky and Poole Frenkel mechanisms would lead the conduction in the film as proposed by Mihara et al [38]. In figures 8(b) and (c) the evolution of the barrier heights and trap barrier heights versus Mn content, extracted from the Schottky and Poole Frenkel modeling, is presented. In both cases, when the doping rate is increased, the barriers are raised. These results are in good agreement with the leakage current decrease observed in figure 3. So as to confirm this Schottky/Poole Frenkel mixing assumption, we introduced the equation Jmix = κJSc + (1 − κ) JPF
(13)
where Jmix , JSc and JPF are respectively the measured leakage current density, the current density due to Schottky emission and the one due to Poole Frenkel injection; κ is the mixing coefficient between the two mechanisms. Therefore, each value of κ between 0 < κ < 1, Jmix should correspond to the leakage current for a Mn doping rate between 0 and 1%. From equation (13), Jmix was calculated for different coefficients κ, using for JSc the experimental data for the undoped sample, and for JPF data for BST70:Mn1%. For each mixing coefficient, a set of J–E plots at different temperatures
3.2. X-ray photoelectron spectroscopy study In order to find the free charges and electronic levels involved in the leakage current behavior, XPS study was performed 7
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Figure 9. Refractive indices for different κ values, obtained from Jmix plots, with both the Schottky and Poole Frenkel laws. They have been placed on the index versus Mn content plot, obtained from experimental current densities.
Figure 11. XPS valence band spectra of undoped BST, BST70:Mn0.5% and BST70Mn1% after an 8 nm depth etching. The positions of the valence band maxima are indicated in the figure.
pronounced for the undoped sample, is greatly reduced by the Mn doping. Therefore, considering that the Ar+ etching leads to the same damaging effect on the chemical bonds regardless of the doping rate, the Mn doping is decreasing the Ti3+ concentration. However, one can notice that the Ti 2p peak maximum is not shifted by doping. The Fermi level position in the bandgap, lying in the middle in the case of an intrinsic material, can be raised by the presence of Ti3+ acting as donors [18]. X-ray excited valence band spectra of the three samples are presented in figure 11. The valence band maximum (VBM) can be determined by a linear extrapolation of the leading edge of the valence band emission. The VBM is at 3.2 eV for the undoped BST and 2.8 eV and 2.4 eV for the Mn0.5% and Mn1% doped films respectively, after the first etching step of 8 nm. The uncertainty can be estimated at ±0.1 eV. Because the measurements were performed ex situ, no barrier height could be calculated at the interface with the Pt electrode; however, the evolution of the energy of the VBM as a function of the etched depth could be plotted, as shown in figure 12. The energy of the Fermi level has been set to 0 eV; the energy of the valence band energy is then calculated by EVB = EF − VBM and the conduction band minimum energy is obtained with ECB = EVB + Egap using a gap value of 3.3 eV [51]. The main observation that can be made is that the Fermi level is lowered in the bandgap on increasing the Mn content. This may be related with the decrease of Ti3+ and oxygen vacancy content in the material, both having energy levels close to the conduction band. In addition, in the undoped sample, one can see that the Fermi level is closer to the conduction band near the top surface. In our understanding, this can be attributed to a higher concentration of Ti3+ in this region. The Mn doping decreases this effect near the top surface, decreasing the in depth band shift. These XPS data thus confirm that oxygen vacancy related defects are contributing to the leakage current in BST films
Figure 10. XPS Ti 2p3/2 spectra of undoped BST, BST70:Mn0.5% and BST70Mn1% after an 8 nm depth etching.
on three samples, namely undoped BST, BST70:Mn0.5% and BST70:Mn1%. In depth XPS was carried out between the top surface and bottom electrode by means of Ar+ etching. Figure 10 displays the Ti 2p spectra of the three thin films after the first etching step of 8 nm. The typical 459 eV peak corresponding to Ti4+ can be seen. It is well known that ion sputtering may lead to some artifacts in the investigated material [48], however a qualitative comparison can provide relevant information. In the spectra, a noticeable emission at the low binding energy side of the peak maximum occurs. This emission at 457 eV can be attributed to Ti3+ ion formation [49]. Although this Ti3+ ion might be partially due to the argon etching [50], a clear trend of its intensity with the Mn content is observed. The peak shoulder, very 8
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Figure 12. Fermi level position in the bandgap as a function of the relative etched depth for BST70, BST70:Mn0.5% and BST70:Mn1%. The black square dots are the measured valence band energies at different etching steps. The red circular dots correspond to the calculated conduction band energies, determined for each etching level with a 3.3 eV bandgap. The relative etched depth equal to 1 corresponds to the whole film thickness.
which can be improved under Mn doping. We will next see how such doping also improves the dielectric properties of the same films. 3.3. Dielectric properties The dielectric properties of the films were investigated as a function of external dc electric field at 30 ◦ C from 100 Hz up to 1 GHz. The permittivity measured under an electric field at low frequency (LF) is shown in figure 13 at 30 ◦ C, which is in the paraelectric state [52]. A decrease of the zero-field permittivity and of the tunability is observed when the Mn doping rate is raised (figure 13(a)). This nonlinearity, up to 80% under 600 kV cm−1 for the undoped BST70, is almost linearly decreased to 62% for BST70:Mn2%. The radio frequency permittivity measured at 1 GHz under different biases is compared to the low frequency charts in figure 14. One can notice that for the undoped and 0.1% Mn doped BST samples, the low-field RF permittivity value is much lower than the 10 kHz one whereas for the higher doping rate a good continuity is observed between LF and RF. Moreover, the low-field RF permittivity values of the undoped, Mn0.1% and Mn0.5% BST are substantially the same. It is well known that at high frequency, defects, such as electronic and ionic space charges, do not contribute to the permittivity anymore [53]. A relaxation of the permittivity can be observed between low frequencies (102 –103 Hz) and RF (108 Hz). By controlling this phenomenon, BaTiO3 supercapacitors can be obtained [54, 55]. Artemenko et al attributed this effect to oxygen-deficiency-related defects and Ti3+ centers [55]. Consequently, it is sound to assume that in figure 14, the higher permittivity at low frequency of the undoped BST70 sample is related to a similar extrinsic effect. Following the XPS data (figure 10), this extrinsic contribution can also be related to Ti3+ in the undoped film. When the Mn content is increased, the extrinsic contribution is reduced (see figure 14, BST:Mn0.1%), until it disappears for the Mn0.5% and Mn1% where no Ti3+ could be observed.
Figure 13. Dielectric measurements at 10 kHz as a function of electric field for undoped and Mn doped BST70 from 0.1% to 2%. (a) Permittivity versus electric field plot, (b) dielectric loss versus electric field plot.
Thus, the zero-field RF permittivity which does not change under Mn doping corresponds to the intrinsic values of the permittivity. 9
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Figure 14. Comparison of the low and high frequency permittivity under electric field (respectively 10 kHz and 1 GHz) for the undoped BST, and Mn0.1%, Mn0.5%, and Mn1% doped samples. The vertical scale is the same for each chart. Extrinsic contributions to the low frequency permittivity are observed for the undoped and Mn0.1% samples.
Moreover, a dielectric loss increase under high electric field can be observed for undoped and 0.1%Mn doped BST70 in figure 13(b). This effect could be related to free carrier transport in the material under dc field, thus related to conduction [60], and its disappearance from 0.5%Mn doping is coherent with the decrease of the leakage currents observed in figure 3. We thus confirm that the density of oxygen vacancy related defects affects the low frequency extrinsic permittivity and losses, both of them disappearing on doping with Mn. Such a beneficial effect has already been reported in BST thin films and ceramics [30, 61] but a direct link between leakage current, electronic state of titanium and intrinsic RF permittivity could not be found. We next discuss the origin of this beneficial contribution of manganese.
It can be noticed in figure 14 that for the BST70:Mn1%, the zero-field intrinsic permittivity is lower than for the other samples. The dielectric permittivity is directly related to the polarizability, which is decreased by breaking the long range order of dipoles when the Mn ions are inserted in B-sites. This result can also be related to the paraelectric to ferroelectric transition profiles of the films where the transition peak intensity decreases and shifts toward lower temperatures for Mn content higher than 1% [52] (not shown here). This effect has been widely reported for Mn as well as for other dopants [56, 57] and can be explained as follows. The temperature and field dependence of the permittivity ε can be written as ε(T, E) = εL (T) − εNL (T)E2 + o(E4 ),
(14)
with εL and εNL the linear and nonlinear parts of the permittivity respectively. A Landau analysis shows that the temperature dependence of these two parameters is as follows [58]: K1 T − TC K2 εNL (T) = . (T − TC )4 εL (T) =
3.4. Discussion It is well known that the concentration of electronic charge carriers leading to conduction in the material and consequently the Fermi level is influenced by the oxygen partial pressure during the annealing, mostly due to the formation of oxygen vacancies (VO ) and free electrons (e− ). The phenomenon can be explained by the following three equilibriums [62]:
and (15)
Therefore, in the paraelectric phase, the closer to the Curie temperature the measurement is carried out, the higher the nonlinearity is. Considering that when the Mn content is increased the Curie temperature is lowered, this can explain why the Mn doping decreases both the linear and the nonlinear permittivities. The dielectric losses of the films measured as a function of electric field can be seen in figure 13(b). At first, when the Mn content is increased, tan δ is decreased, reaching a minimum of ∼0.9% for the BST70:Mn1%. The decline in the loss is most likely due to the charge compensation of oxygen vacancies again making the link with the density of Ti3+ . For the higher doping rate BST70:Mn2%, the losses are increased to 1.1%. Indeed, it is well known that for high concentration, the Mn ions are not soluble anymore [29, 59], and the properties are degraded [57].
OO(lattice) ⇔ VO + 1/2O2 VO ⇔ V•O + e− − V•O ⇔ V•• O +e . In our case, even if the films are annealed in an oxygen atmosphere, there is no doubt that our material contains such defects. In BST, in order to account for the electro-neutrality of the material, the electrons provided by oxygen vacancies can be equilibrated by the reduction of Ti4+ , Ti4+ + 1e− ⇔ Ti3+ . Hence, the Ti3+ ion becomes a kind of oxygen vacancy marker in the undoped material. In the case of the undoped BST70 film, the presence of such reduced Ti3+ has been highlighted by the XPS 10
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Figure 15. Schematic representation of the electronic structure of defects in the bandgap. (a) For the undoped BST, the configuration is favorable to Schottky injection. (b) For the BST70:MN0.1%, the electronic configuration is compatible with Schottky/PF mixed conduction behavior. (c) For the BST70:Mn1%, the electronic configuration is favorable to Poole Frenkel emission.
measurements presented in figure 10. As a result, the Fermi level is raised to near the conduction band (figure 12). This is corroborated on one hand by the extrinsic contributions to the permittivity at low frequency shown in figure 14 which have also been associated with the presence of Ti3+ [54, 55], and on the other hand by the strong resistance degradation shown in the J–t measurements in figure 1(b) that is related to oxygen vacancy charge migration. Therefore, these results directly explain the high dielectric losses observed in figure 13(b), and the high conductivity of the undoped BST shown in figures 1 and 3. When Mn ions are introduced into the material, a significant decrease of the conduction and the dielectric losses has been observed in figures 1, 3 and 13(b). To explain this result, we based our analysis on the assumption that in BST, the Mn ions are located in a Ti4+ site as definitively assigned by Siegel and Muller [63]. Therefore, the Mn4+ ion can be converted into lower valence states Mn3+ and Mn2+ [12, 64, 65] to compensate the charge brought from the V•• O,
The decrease of the leakage current by increasing the Mn doping rate in the material is associated with a change in the conduction mechanisms (see section 3.1.4). In order to explain this phenomenon the following model is proposed. So as to relate the Ti3+ and Mn3+ concentrations with the trend of the leakage current mechanism shifting from Schottky toward Poole Frenkel as previously shown in section section 3.1.4, we propose to sketch a simplified band diagram. Regarding the undoped sample, a barrier 80 at the interface of about 0.6 eV has been computed from the Schottky model as shown in figure 8(b). For this sample, a large presence of Ti3+ has been directly or indirectly detected through different characterization techniques. According to Moretti et al the Ti–VO centers can induce shallow electron traps, which can modulate the conductivity [66]. Actis et al have shown by calculations that Ti3+ has donor levels localized about 0.1–0.2 eV below the conduction band minimum (CBM) [67]. Recently Mitra et al determined through computing that the V•• O state is the most stable among the oxygen vacancies, with an energy level lying 0.28 eV below the CBM in SrTiO3 [68]. In figure 15(a), we drew a schematic band diagram using these data. It is shown that these energy levels very close to the CBM can easily carry electrons through the film once the electrons have been emitted by the electrode. In other words, only the attenuation of the electrode–film barrier arising from the electrode image–force interaction with the field is controlling the conduction, which is coherent with the interface limited Schottky injection model. Moreover, deeper trap centers can be present in this undoped film, which hardly take part in the conduction, such as Fe3+ impurities, known to be present in many titanates, even in pure crystals [63]. These Fe3+ trap levels in BaTiO3 have been localized 1.6 eV below the CBM [67]. When manganese is introduced into the material, the reduction of Mn4+ into Mn3+ can easily compensate the oxygen vacancy free charges. As a result the concentration of Ti3+ , which is much more unstable, is greatly decreased. In the literature, different calculation results can be found for the energy of Mn trap levels. According to Selme [69], in SrTiO3 when Mn4+ is converted into Mn3+ , the additional electron is localized around 1 eV below the CBM. Moretti [64] found that Mn3+ traps an electron around 0.7 eV below the CBM
Mn4+ + 1e− ⇔ Mn3+ Mn3+ + 1e− ⇔ Mn2+ . Because the energy levels of the Ti3+ are much closer to the conduction band than the Mn3+ , this manganese compensation is more favorable. Therefore, when the Mn concentration is increased, the Ti3+ concentration is decreased, as confirmed by the XPS analysis in figure 10. The Fermi level is lowered into the bandgap, and gets very close to the middle of it, which means that the material shows a more intrinsic behavior. Additional information about the decrease of Ti3+ defects is also pointed out by the dielectric measurements in figure 14, where the extrinsic contribution to the low frequency permittivity is decreased for the BST70:Mn0.1%, and suppressed from BST70:Mn0.5% to higher rates. The time dependence of the leakage current (figure 1(b)) confirms this trend by the decrease of the resistance degradation phenomenon, associated with decrease of the free charge concentration. Finally, when the Mn content is too high, here 2 at.%, a degradation of the properties is observed: an increase of the dielectric losses (figure 13(b)) and an early breakdown process (figure 1(b)). It thus seems that for this doping rate too many defects are created in the system. 11
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References
in BaTiO3 , but according to them, these 3d states are shifted downwards by the presence of a surrounding oxygen vacancy. Therefore, in the case of the BST70:Mn1%, due to the very low concentration of Ti3+ , the barrier energy at the interface 80 becomes lower than the Mn3+ trap height energy 8t , as shown in figure 15(c). This is consistent with a conduction model limited by the bulk part of the film, such as Poole Frenkel, as found in section 3.1.4. For the intermediate doping rate, a mix of Schottky and Poole Frenkel models has been suggested. If we consider, as shown in figure 15(b), that for low Mn concentration, both Ti3+ and deeper trap levels are participating in the bulk conduction, both interfaces and deep trap levels could be limiting the conduction.
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4. Conclusions In this study, we found that it was possible to tune the leakage current mechanism in Ba0.7 Sr0.3 Ti1−δ Mnδ O3−y films by increasing the Mn concentration. As expected, the Mn doping decreased the low frequency dielectric losses at zero-field effectively, as well as the leakage current under an external electric field. In addition, by studying the conduction mechanism in the material and the refractive index, we demonstrated that the conduction is first governed by the Schottky emission at the interface for the undoped BST sample, and an increase of Mn concentration leads to modification of the mechanism, mixing Schottky and Poole Frenkel processes. Then for the BST70:Mn1%, the bulk-like conduction mechanism is mainly due to Poole Frenkel injection. In order to find the origin of this phenomenon, a complementary XPS investigation was performed which evidenced, by a relative comparison, a decrease of Ti3+ with increase of Mn concentration. This led to lowering of the Fermi level into the bandgap, bringing the material closer to intrinsic behavior. This result has been corroborated by additional results on the time dependence of the leakage currents and dielectric measurements at low and high frequencies which evidence indirectly the presence of oxygen vacancies and Ti3+ . Besides the beneficial effect of Mn doping already reported in BST thin films and ceramics [8, 25, 30, 26], we have shown here a direct link between the leakage current, the electronic state of titanium and the intrinsic RF permittivity. Finally, a simplified model has been proposed to explain the evolution of the conduction mechanism, using the theoretical electronic level configuration of the system.
Acknowledgments The authors would like to thank Christine Labrug`ere from ICMCB for the XPS analysis, Dr R Schafranek for the useful discussion about XPS results and Mitsubishi Material Corporation for the sol–gel solutions. Part of this work was supported by the French National Agency for Research under contract ABSYS2 and by the Aquitaine Regional Council. 12
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[39] Mihara T and Watanabe H 1995 Japan. J. Appl. Phys. 34 5674 [40] Sze S and Ng K 2006 Physics of Semiconductor Devices (New York: Wiley) [41] Frenkel J 1938 Phys. Rev. 54 647–8 [42] Stolichnov I and Tagantsev A 1998 J. Appl. Phys. 84 3216–25 [43] Simmons J G 1967 Phys. Rev. 155 657–60 [44] Dietz G W, Schumacher M, Waser R, Streiffer S K, Basceri C and Kingon A I 1997 J. Appl. Phys. 82 2359–64 [45] Zafar S, Jones R E, Jiang B, White B, Kaushik V and Gillespie S 1998 Appl. Phys. Lett. 73 3533–5 [46] Dietz G W and Waser R 1997 Thin Solid Films 299 53–8 [47] Copel M, Baniecki J D, Duncombe P R, Kotecki D, Laibowitz R, Neumayer D A and Shaw T M 1998 Appl. Phys. Lett. 73 1832–4 [48] Hofmann S 1998 Rep. Prog. Phys. 61 827 [49] G¨opel W, Rocker G and Feierabend R 1983 Phys. Rev. B 28 3427–38 [50] Cord B and Courths R 1985 Surf. Sci. 152/153 1141–6 [51] Robertson J and Chen C W 1999 Appl. Phys. Lett. 74 1168–70 [52] Levasseur D, Rousseau A, De Paolis R, Coccetti F, Guegan G, Payan S and Maglione M 2013 Mn doping study in Ba0.7 Sr0.3 TiO3 , from the ceramic to the thin film at press [53] Coelho R 1979 Physics of Dielectrics for the Engineer (New York: Elsevier Scientific) [54] Chung U C, Elissalde C, Mornet S, Maglione M and Estourn`e C 2009 Appl. Phys. Lett. 94 072903 [55] Artemenko A, Elissalde C, Chung U C, Estourn`es C, Mornet S, Bykov I and Maglione M 2010 Appl. Phys. Lett. 97 132901
[56] Fan Y, Yu S, Sun R, Li L, Yin Y, Wong K W and Du R 2010 Appl. Surf. Sci. 256 6531–5 [57] Liang R H, Dong X L, Chen Y, Cao F and Wang Y L 2005 Ceram. Int. 31 1097–101 [58] Lines M E and Glass A M 1977 Principles and Applications of Ferroelectrics and Related Materials ed M E Lines and A M Glass (Oxford: Clarendon) [59] Langhammer H, M¨uller T, Polity A, Felgner K H and Abicht H P 1996 Mater. Lett. 26 205–10 [60] Fang L, Shen M, Yang J and Li Z 2006 Solid State Commun. 137 381–6 [61] Kim R Y, Kim H S, Lim M H, Kim H G and Kim I D 2004 Integr. Ferroelectr. 66 195–204 [62] Chan N H, Sharma R and Smyth D 1981 J. Am. Ceram. Soc. 64 556–62 [63] Siegel E and M¨uller K A 1979 Phys. Rev. B 19 109–20 [64] Moretti P and Michel-Calendini F M 1987 Phys. Rev. B 36 3522–7 [65] Duverger E, Jannot B, Michel-Calendini F M and Maglione M 1992 Ferroelectrics 125 301–6 [66] Moretti P and Michel-Calendini F M 1988 J. Opt. Soc. Am. B 5 1697–701 [67] Actis M and Michel-Calendini F 1997 Int. J. Quantum Chem. 61 657–64 [68] Mitra C, Lin C, Robertson J and Demkov A A 2012 Phys. Rev. B 86 155105 [69] Selme M O and Pecheur P 1985 J. Phys. C: Solid State Phys. 18 551
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Systematic tuning of the conduction mechanisms in ferroelectric thin films
This content has been downloaded from IOPscience. Please scroll down to see the full text. 2013 J. Phys.: Condens. Matter 25 495901 (http://iopscience.iop.org/0953-8984/25/49/495901) View the table of contents for this issue, or go to the journal homepage for more
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IOP PUBLISHING
JOURNAL OF PHYSICS: CONDENSED MATTER
J. Phys.: Condens. Matter 25 (2013) 495901 (13pp)
doi:10.1088/0953-8984/25/49/495901
Systematic tuning of the conduction mechanisms in ferroelectric thin films D Levasseur1,2 , E Bouyssou2 , R De Paolis3 , A Rousseau1 , F Coccetti3 , G Guegan2 , S Payan1 and M Maglione1 1 2 3
CNRS, Univ. Bordeaux, ICMCB, UPR 9048, F-33600 Pessac, France ST Microelectronics, 16 rue Pierre et Marie Curie, F-37100 Tours, France CNRS, Univ. Toulouse, LAAS, F-31400 Toulouse, France
E-mail: [email protected]
Received 17 July 2013, in final form 10 September 2013 Published 6 November 2013 Online at stacks.iop.org/JPhysCM/25/495901 Abstract We have investigated the macroscopic and microscopic properties of large sets of Ba0.7 Sr0.3 TiO3 thin films including several substitution rates of manganese. Thanks to a high degree of control of the processing parameters at each stage we have been able to find a link between the dc leakage current and the low and high frequency dielectric permittivity and losses. We supplemented these macroscopic observations with in depth investigations of the defect states through x-ray photoelectron spectroscopy. We found that both the leakage current and the extrinsic dielectric parameters arise from a large density of charged point defects related to oxygen vacancies. At the outer surfaces of the films, the density of such charged defects is so high that it can raise the Fermi level to close to the conduction band. Such degradation of the films’ performance can be relieved by appropriate manganese substitution for the titanium host ions. Such doping is able to move back the Fermi level to close to the center of the bandgap thus changing the conduction process from interfacial Schottky to bulk Poole Frenkel and decreasing the extrinsic losses. This beneficial effect was already inferred in ceramics and thin films but we have established a clear link between the macroscopic parameters and the microscopic defect state. This model can be transferred to many high permittivity oxides. (Some figures may appear in colour only in the online journal)
1. Introduction
able to reduce the detrimental influence of residual oxygen vacancies which are donor centers [8–12]. When brought to thin films, the same materials raised the same questions: what is the contribution of oxygen vacancies to the residual currents; what about the contribution of interfaces; is the surface more defective than the bulk [13–19]? At the same time, the dielectric properties of ceramics and thin films of the same materials were also investigated to improve their performance [20]. The reduction of losses mostly in the high frequency range (f > 1 GHz) is a real challenge to enable practical application of ferroelectric thin films in microwave devices [21]. This was achieved thanks again to the control of charged defects and also to mixing with very low loss materials like MgO [22–24]. So as to avoid unnecessary complexity resulting from the contribution of ferroelectric domains and domain walls to these losses, the
Finding a link between the static electrical conductivity and the dynamic losses of dielectric materials is not only a key for their actual applications but also a very basic problem which is largely unsolved. In the case of very high permittivity oxides like ferroelectrics, both phenomena have been extensively investigated. Focusing on BaTiO3 -derived materials, the knowledge and control of residual conductivity in pure and substituted ceramics have been improved over the years [1–7]. The main tool was and is still the compensation of unwanted charged centers through accurate control of the defect chemistry. Including the oxygen stoichiometry, this is a very hard task which has resulted in still active debates after decades of investigation. The main trend is that the substitution of acceptor ions (e.g. Mn3+ on Ti4+ sites) is 0953-8984/13/495901+13$33.00
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current density at different temperatures were carried out in order to determine its time dependence (J–t). For each temperature, five different capacitors were tested in order to get more representative results. The electrical field dependence of the leakage current was then studied, denoted J–E for convenience. For each wafer eight capacitors were tested, and a very good reproducibility was w obtained; therefore the following results are always representative of the whole wafer. The data processing will be explained later on. The refractive indices of the films were determined by spectroscopic ellipsometry, using a Nanofilm EP3 Imaging Ellipsometer. The 1/9 data were recorded in the visible range, from 502 to 720 nm. Using the EP4Model software, a Cauchy law was used to fit the data and determine the refractive index dispersion n(λ) of each film. The root mean square errors obtained after fitting were always lower than 3. A VG Scientific 220 i-XL ESCALAB spectrometer was used for the XPS analysis with a monochromatized Al Kα source (hν = 1486.6 eV) at 70 W. The spot size was about 200 µm in diameter. In depth XPS analysis was carried out using argon etching. At each step, full spectra (0–700 eV) were recorded with a constant pass energy of 150 eV and high resolution spectra at a constant pass energy of 20 eV. The high resolution spectra were then fitted and deconvoluted using the AVANTAGE software provided by ThermoFisher Scientific. Finally, the dielectric properties at low frequencies were studied on integrated MIM structures, with 0.063 mm2 Pt top electrodes. The capacitance–voltage (C–V) and dielectric loss–voltage (tan δ–V) characteristics of the MIM capacitors were investigated at 30 ◦ C using a Cascade Microtech probe station and an Agilent 4294A Impedance Analyzer. These C–V curves were measured with a 10 kHz/0.1 V ac signal, and the dielectric permittivity was then calculated using capacitance, and the dimensional parameters: the thickness d obtained from cross section SEM images and top electrode area A. Concerning the RF characterization process, one-port reflection measurements were performed on MIM capacitors at room temperature by means of an Anritsu 37397C VNA with internal bias tee and Picoprobes with 150 µm GSG coplanar pitch. An SOL calibration procedure was applied beforehand. Subsequently, 1 GHz permittivity was computed using an equivalent circuit consisting of a capacitor and a resistor in series. More specifically, a differential method was adopted in order to allow the removal of the most prominent parasitic effects [32, 33]. Finally, the BST permittivity as a function of the electric field was extracted from the capacitance values associated with every different applied voltage.
focus is more on the solid solution Bax Sr1−x TiO3 (BST) which is paraelectric at room temperature for some given composition x. However, the exact origin of these dielectric losses, whether intrinsic or extrinsic, is still a matter of debate and the contribution of charged defects to the extrinsic losses is still to be probed. In this paper, we focus on dc conductivity and high frequency dielectric properties in BST films substituted with manganese. As already mentioned, such films have already been investigated either for their leakage current [25–29] or for their high frequency dielectric properties [30, 31]. The first contribution of this paper is to systematically address both these issues in the same films. Another interest in our strategy was to investigate films of microelectronics industry grade so as to ensure reproducibility over sets of thousands of samples. This enabled us to ensure relevant statistics in our experimental results. By doing so, we were able to show that the dc leakage current gradually changes from interface limited Schottky to bulk limited Poole Frenkel type as the atomic concentration of Mn is raised from 0.1% to 2%. Such a trend was linked to the down shift of the Fermi level from the bottom of the conduction band toward the center of the bandgap by means of x-ray photoelectron spectroscopy (XPS). We then mapped out the dielectric dispersion from 1 kHz up to 1 GHz and we could show that the low frequency permittivity at f < 1 GHz is artificially raised by the residual charged defects for undoped films, and is gradually suppressed under Mn doping. This clearly shows a strong link between the dc current and the dielectric properties. Based on our investigation of the electronic state, we suggest that this common origin results from the Ti3+ states associated to oxygen vacancies. We can further state that these defects are all the more relevant when they are accumulated at the interfaces between the films and the metal electrodes. We thus can draw a consistent picture of the electric and dielectric behavior of Mn doped BST films. These results are of great interest for tunable microwave devices, where a very high level of understanding is needed in both dielectric and leakage current properties, in order to comply with the specifications of such components.
2. Experimental procedure The Ba0.7 Sr0.3 Ti1−δ Mnδ O3 (BST70:Mn) precursor sols used in this study were purchased from Mitsubishi Material Corporation, Japan, with δ = 0, 0.001, 0.005, 0.01, and 0.02. Thin films from these sols were deposited by spin coating on industrial Pt/Si substrates. For each sample, the solution was coated on the substrate to form a layer, this layer was pyrolyzed, and the process was repeated in order to obtain the desired thickness. Finally the whole film was post-annealed under oxygen at 800 ◦ C. In order to investigate the leakage currents of the films, integrated metal–insulator–metal (MIM) structures were processed with 0.004 mm2 Pt top electrodes. These capacitors were characterized through a Cascade Microtech probe station connected to a Keithley 2636A SourceMeter Unit. Constant voltage stress (CVS) measurements of the
3. Results and discussion 3.1. Leakage current study In order to properly investigate the leakage currents of the samples, the current versus time behavior was first studied, providing important information such as current 2
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density levels, and the true leakage current region is shortened to ∼5 s at 125 ◦ C and suppressed at 165 ◦ C. After the flat range ¬, the current starts to increase until a maximum value Jmax is reached; this regime is called the ‘resistance degradation’ (). According to Waser et al [2], this phenomenon is highly related to the presence of oxygen vacancies in the material. After some time under a constant electric field, the charges of such defects can start to migrate toward the cathode. Thus, the spatial modification of the carrier concentrations between the two electrodes will change the conductivity of the film. The accumulation of charged oxygen vacancies close to the cathode induces modifications of the band structure and thereby interfacial energy barriers. As a consequence, the current density increases. Once the degradation current has reached its maximum Jmax , the current density starts to decrease (®). Bouyssou et al attributed this phenomenon to a trapping effect and space charge reorganization and they proposed a mathematical function describing this phenomenon [35]. In figure 1(b), the effect of Mn doping on the J–t characteristics at 200 ◦ C can be seen. This temperature has been chosen in order to better highlight Jmax . It is shown that the Mn doping strongly decreases the maximum level Jmax reached after the resistance degradation, until a Mn concentration of 1%. It seems, however, that when the Mn content is too high, the breakdown process happens more rapidly, as shown for the BST70:Mn2% at around 400 s at 200 ◦ C under 750 kV cm−1 . Zafar et al proposed a quantitative model for the increase of the current densities, based on the assumption that the current increase during the resistance degradation is mainly due to decrease of the effective barrier height [36]. The maximum decrease of the barrier, denoted 18max , is given by
Figure 1. The time dependence of the current density J obtained for 750 kV cm−1 electric field stress (a) for the undoped BST70 film at different temperatures (85, 125, 165 and 200 ◦ C), (b) at 200 ◦ C for each doping rate.
18max = kT ln(Jmax /Jinit )
(1)
where Jmax is the maximum current density and Jinit is the value obtained by extrapolating the current density at t = 0. The computed 18max decreases from 0.17 eV for the undoped sample to 0.08 eV for the BST70:Mn1%. This stabilization of the effective barrier height can be ascribed to a decreased density of free charges under Mn doping. Because of the extensive leakage current evolution as a function of time, special care must be taken subsequently in the J–E spot measurements, in order to avoid any experimental artifacts.
relaxation, resistance degradation and breakdown. Then the leakage current versus electric field was analyzed. These measurements allowed us to study the conduction mechanisms in the BST thin films, and the influence of the Mn doping. 3.1.1. Study of the time dependence. We first investigated the time dependence of the leakage currents of the capacitors. In figure 1(a), the current density of the undoped BST sample as a function of time is presented, under a 750 kV cm−1 external field for different temperatures from 85 to 200 ◦ C In the initial time period, no decrease in current due to the dielectric relaxation is observed, as opposed to many studies [34]; this is likely to be due to the high temperature and high field that we applied. Some typical trends can be drawn from the curves numbered from ¬ to ® in figure 1(a) [34, 35]. For the lowest temperature of 85 ◦ C, the current starts with a flat time dependence corresponding to the so-called true leakage current (¬) for ∼500 s. When the temperature is increased, the plots are globally shifted toward higher current
3.1.2. Study of the electric field dependence. In order to remove time artifacts, the following characterization protocol has been established to determine the influence of the Mn concentration on the leakage current. To avoid relaxation current during the J–E measurement at low temperatures, a short constant voltage stress (CVS) was carried out for 100 s at 85 ◦ C with regard to the flat ‘true’ leakage current dependence observed in figure 1. Because the relaxation is a reversible process, the wafers were then quenched at ambient temperature by contact on a marble plaque so as to keep the material in the stressed state. 3
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Figure 2. Current density as a function of electric field for the undoped BST70 film at different temperatures (from 30 to 200 ◦ C). The voltage was applied to the top electrode. Inset: the schematic staircase profile of the dc voltage applied during J–E measurements; td is the delay time, tm the measuring time and 1V the step height.
Figure 3. Current density as a function of electric field at 30 ◦ C of undoped BST and Mn doped BST from 0.1% to 2%. The voltage was applied to the top electrode.
ones, Poole Frenkel and space charge limited currents (SCLCs) are reported. Concerning the tunneling of free charges through the whole film or only a part of the film, e.g. a grain boundary energy barrier, the mechanism seems unlikely regarding the thickness of the film and the weak band bending at grain boundaries for thin films [16]. In addition, the strong temperature dependence of the leakage current cannot usually be observed with tunneling models. Concerning the SCLC, a low barrier height is required to guarantee sufficient charge injection. According to Dietz et al SCLC is unlikely at metal contacts on high-gap semiconductors or insulators [16]. In view of these concerns about tunneling and SCLC, our study will therefore focus only on Schottky and Poole Frenkel (PF) mechanisms. In order to determine whether the leakage current can be associated to a Schottky regime or a Poole Frenkel injection, the physical parameters involved in the mechanism must be determined. Hence, for each sample the two following protocols, based on the J–E measurements obtained at different temperatures (e.g. figure 2), have been applied.
After the pre-stress, the current versus voltage characteristics were measured for each capacitor at different temperatures from 30 to 200 ◦ C, with a typical staircase profile presented in the inset of figure 2. The voltage was applied for 400 ms (td ) before measuring the current and the step height (1V) was 0.25 V. The behavior of the MIM capacitors was symmetrical, that is to say the same results were obtained whether the voltage was applied to the top or bottom electrode. Typical results obtained are shown in figure 2, where the current density versus electric field of the undoped BST70 film biased on the top electrode is presented. A strong dependence of the current density on the temperature can be noted. For each sample a similar set of plots has been obtained. The influence of increase of the doping rate at 30 ◦ C is presented in figure 3. Increase of the Mn content dramatically decreases the leakage current, as expected, by almost two decades for the BST70:Mn1% and more than three decades in the case of BST70:Mn2%. In addition, it can be seen that the shape of the J–E plots is modified by the Mn insertion; therefore a modification in the conduction mechanism can be expected.
Schottky regime In the case of a Schottky mechanism, the conduction is governed by charges thermally injected at the interface, by lowering the barrier height due to the applied field and image force. The current evolution is described by the equation q qEext 80 − 4πε ε 0 i (2) J = A∗ T 2 exp − kB T
3.1.3. Study of the conduction mechanisms. Many results on the conduction mechanisms in perovskite-type titanate thin films, including SrTiO3 , BaTiO3 , (Ba,Sr)TiO3 and Pb(Zr, Ti)O3 , have previously been reported [15–17, 35, 37–39], and the main physical models describing the leakage currents in the films have been discussed. In the case of an MIM capacitor, the leakage current is either controlled by the interfaces with the electrodes or by the bulk of the film. In most cases, one of them is highly dominant. For the interface-controlled mechanisms, tunneling and Schottky regimes can be found, and concerning the bulk-controlled
where A∗ , q, ε0 and kB are respectively Richardson’s constant, the electron charge, the vacuum permittivity and Boltzmann’s constant, T is the temperature, Eext the external electric field, ϕ0 the barrier height, and εi the relative permittivity at optical 4
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Figure 4. J/T 2 as a function of the square root of the bias, at different measurement temperatures from 30 to 200 ◦ C, for the undoped BST70 film. The voltage was applied to the top electrode. The black lines correspond to the exponential fits y = a ebx , and the equations obtained are written next to the corresponding temperatures.
frequencies [40]. This equation can be turned into " √ # J βSc V = α exp kB T T2
Figure 5. Arrhenius plot, ln(J/T 2 ) as a function of 1000/T, for the undoped BST film. The voltage was applied to the top electrode. The black lines correspond to the linear fits y = ax + b, and the equations obtained are written next to the corresponding biases.
Poole Frenkel effect In the case of the Poole Frenkel emission, the conduction is governed by carriers that are thermally emitted from trapped centers under a strong electric field. The current versus temperature and electric field follows the equation q qEext 8t − πε 0 εi J = qµEext Nc exp − (7) kB T
(3)
where 80 α = A∗ exp − and kB T (4) r q . βSc = 4π ε0 εi d √ Then, by plotting TJ2 = f ( V) for each measured temperature, and using an exponential fit y = a ebx with a and b as fitting parameters, as shown in figure 4, from b and equation (3) the parameter βSc for each temperature can be determined. Then, from equation (4), the optical dielectric permittivity εi of each sample has been computed. Because the effective Richardson’s constant is highly dependent on the material and the electrode [16], the barrier height cannot be determined directly from α (equation (4)). Thus, the barrier height has been calculated using the Arrhenius law, applied to this case, J −Ea = C exp (5) 1 kB T T2
where µ and Nc are respectively the carrier mobility and the carrier density in the conduction band, 8t is the trap barrier height [41]. Turning this equation into " √ # βPF V J 0 = α exp (8) V kB T where α0 =
qµNc 8t exp − d kB T r q βPF = π ε0 εi d
and (9)
√ = f ( V) was plotted for each measured temperature. Then, using an exponential fit y = a ebx and equation (8), the parameter βPF of each measured temperature can be determined, and so from equation (9), computation of εi of each sample has been processed. As µ and Nc are unknown values, ϕt cannot be directly obtained, and Arrhenius curves were plotted. J −Ea = C2 exp (10) V kB T J V
where C1 is a constant and Ea is the activation energy. From equation (3), this activation energy can be written as s qV Ea = 80 − . (6) 4π ε0 εi d
where C2 is a constant and Ea is the activation energy. Therefore the activation energy can be written as s qV Ea = 8t − . (11) π ε0 εi d
The validity of equation (5) is readily seen in figure 5 where ln(J/T 2 ) is plotted versus 1/T. From the fits of these curves, Ea (V) can be extracted and then the barrier height 80 from equation (6). 5
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Figure 6. Current density as a function of bias of the undoped BST70 sample for different temperatures from 30 to 175 ◦ C. The scatter plots correspond to the experimental values, and the black lines correspond to the fits with (a) the Schottky injection law and (b) the Poole Frenkel emission law. The voltage was applied to the top electrode.
From a fit using equation (11), Ea (V) can be extracted and then ϕt the trap barrier height (equation (11)). To sum up, in the cases of both the Schottky and the Poole Frenkel mechanisms, a couple of parameters can be extracted, the dielectric permittivity at optical frequencies εi and the interfacial or trap barrier height, respectively 80 and 8t . From these parameters, the J–E characteristics predicted by the Schottky and Poole Frenkel equations, respectively equations (2) and (7), can be plotted for a given temperature. Knowing the thickness of the thin film and the constant values, these plots can be fitted to the experimental characteristics using the effective Richardson’s constant A∗ as a fitting parameter in the case of Schottky and µNc in the case of Poole Frenkel. An example of these curves is given in figure 6 for the undoped BST sample. As can be seen, both the Schottky and Poole Frenkel laws fit the experimental data fairly well, for voltages higher than 5 V. The low voltage part of the fit was not taken into account, due to the nonlinearity of the capacitance, and to band profile evolution in early applied voltages such as in the Ohmic regime [42]. So as to validate these results, the values of εi should satisfy the equation εi = n2
3.1.4. Influence of the Mn doping in BST70 on conduction mechanisms. The leakage current is decreased on increasing the Mn content in the material, as shown before in figure 3. However, this impurity insertion may also impact the conduction mechanism in the material. Ahn et al found that Mn doping in Ba0.5 Sr0.5 TiO3 films decreases the tunneling current by increasing the depletion layer width near the electrodes [27]. In our case, we observed a strong temperature dependence of the leakage currents, ruling out the possibility of such a mechanism as stated before; however, it seems clear in the literature that acceptor Mn doping influences the interfaces strongly [18, 47]. We focus here only on 0.1 to 1 at.% Mn content. For BST70:Mn2%, the J–E measurements (not shown here) were not reliable, as from 85 ◦ C a deterioration was observed and the ‘true’ leakage current could not be measured any longer because of the early breakdown shown in figure 1(b). Experimental J–V characteristics of BST70:Mn0.1%, BST70:Mn0.5% and BST70:Mn1% and the corresponding simulated plots calculated with the Schottky and Poole Frenkel models are presented in figure 7. For the undoped sample, determination of the dominant mechanism is really not obvious from these charts, since both models seem possible. Besides, a very marked trend can be noticed in the extracted refractive index shown in figure 8(a). For the undoped BST70, the n value calculated from the Schottky model perfectly matches with the experimental index measured by ellipsometry, as shown before. However, when the Mn content increases, this computed refractive index decreases to reach a value of 1.05 for BST70:Mn1%, almost that of vacuum. Additionally, for the index extracted from the Poole Frenkel process, at first for the undoped sample the index value is too large, but it decreases with increase of Mn content, and reaches 2.8 for BST70:Mn1%. The index extracted from PF seems to tend to the measured index values for high doping rates. Therefore, the following hypothesis can be made. For the undoped BST material, the leakage is governed by the Schottky mechanism. For BST70:Mn1%, this mechanism seems impossible considering the very weak extracted value of n, but the Poole Frenkel mechanism could be much more
(12)
where n is the refractive index of the material [43]. The n values extracted with our protocols were 2.09 and 6.85 respectively for Schottky and Poole Frenkel, when the voltage was applied to the top electrode, and 2.18 and 7.36 respectively when applied to the bottom electrode. However, the measured refractive index from ellipsometry was between 2.06 and 2.18, in the wavelength range of 501.1 to 719.8 nm. Therefore, the dominating conduction mechanism in the undoped film can clearly be attributed to the Schottky process. This is in agreement with previous studies performed in BST70 [44] and other compositions of barium strontium titanates [16, 45], showing that Schottky is the leading mechanism. Dietz et al also reported that an Ohmic contribution to the low-field region shifted to a Schottky effect for higher voltage [46], explaining the low-field misfit in figure 6. 6
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Figure 7. Current density as a function of bias of (a) BST70:Mn0.1%, (b) BST70:Mn0.5% and (c) BST70:Mn1% for different temperatures from 30 to 175 ◦ C. The colored scatter plots correspond to the experimental values, the continuous black lines to the Schottky fits and the dashed black lines to the Poole Frenkel fits. V+ means that the voltage was applied to the top electrode.
Figure 8. Physical parameters, extracted from the Schottky and Poole Frenkel models, as a function of Mn content from 0 to 1%. (a) Refractive index (the measured index is plotted in a continuous line as a guide for the eyes but has only been measured for each doping rate), (b) interfacial barrier height from the Schottky mechanism, (c) trap barrier height from the Poole Frenkel mechanism. +te and +be respectively mean that the voltage was applied to the top or the bottom electrode.
similar to the example shown in figure 2 was obtained. Then, the two protocols presented in part 3.1.3 were applied to these Jmix plots, and a couple of refractive indices was determined with the Schottky and PF models for each κ value. It turns out that these refractive indices perfectly match with those determined by experiment as shown in figure 9. Moreover, this means that each mixing coefficient can be related to some Mn content. According to the results, with 0.07 at.% of Mn in the material, 90% of the leakage current would be governed by the Poole Frenkel mechanism, whereas when the Mn doping reaches 0.4%, 99% of the conduction would be led by Poole Frenkel. This result should be taken with caution: BST70:Mn1% with a refractive index of 2.8 may still have conduction due to Schottky injection, even if it is dominated by Poole Frenkel. The idea here is just to show through the extracted optical index that a trend is obtained to confirm a mix between the two mechanisms.
relevant giving a refractive index of 2.8. For intermediate Mn content, a mix of Schottky and Poole Frenkel mechanisms would lead the conduction in the film as proposed by Mihara et al [38]. In figures 8(b) and (c) the evolution of the barrier heights and trap barrier heights versus Mn content, extracted from the Schottky and Poole Frenkel modeling, is presented. In both cases, when the doping rate is increased, the barriers are raised. These results are in good agreement with the leakage current decrease observed in figure 3. So as to confirm this Schottky/Poole Frenkel mixing assumption, we introduced the equation Jmix = κJSc + (1 − κ) JPF
(13)
where Jmix , JSc and JPF are respectively the measured leakage current density, the current density due to Schottky emission and the one due to Poole Frenkel injection; κ is the mixing coefficient between the two mechanisms. Therefore, each value of κ between 0 < κ < 1, Jmix should correspond to the leakage current for a Mn doping rate between 0 and 1%. From equation (13), Jmix was calculated for different coefficients κ, using for JSc the experimental data for the undoped sample, and for JPF data for BST70:Mn1%. For each mixing coefficient, a set of J–E plots at different temperatures
3.2. X-ray photoelectron spectroscopy study In order to find the free charges and electronic levels involved in the leakage current behavior, XPS study was performed 7
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Figure 9. Refractive indices for different κ values, obtained from Jmix plots, with both the Schottky and Poole Frenkel laws. They have been placed on the index versus Mn content plot, obtained from experimental current densities.
Figure 11. XPS valence band spectra of undoped BST, BST70:Mn0.5% and BST70Mn1% after an 8 nm depth etching. The positions of the valence band maxima are indicated in the figure.
pronounced for the undoped sample, is greatly reduced by the Mn doping. Therefore, considering that the Ar+ etching leads to the same damaging effect on the chemical bonds regardless of the doping rate, the Mn doping is decreasing the Ti3+ concentration. However, one can notice that the Ti 2p peak maximum is not shifted by doping. The Fermi level position in the bandgap, lying in the middle in the case of an intrinsic material, can be raised by the presence of Ti3+ acting as donors [18]. X-ray excited valence band spectra of the three samples are presented in figure 11. The valence band maximum (VBM) can be determined by a linear extrapolation of the leading edge of the valence band emission. The VBM is at 3.2 eV for the undoped BST and 2.8 eV and 2.4 eV for the Mn0.5% and Mn1% doped films respectively, after the first etching step of 8 nm. The uncertainty can be estimated at ±0.1 eV. Because the measurements were performed ex situ, no barrier height could be calculated at the interface with the Pt electrode; however, the evolution of the energy of the VBM as a function of the etched depth could be plotted, as shown in figure 12. The energy of the Fermi level has been set to 0 eV; the energy of the valence band energy is then calculated by EVB = EF − VBM and the conduction band minimum energy is obtained with ECB = EVB + Egap using a gap value of 3.3 eV [51]. The main observation that can be made is that the Fermi level is lowered in the bandgap on increasing the Mn content. This may be related with the decrease of Ti3+ and oxygen vacancy content in the material, both having energy levels close to the conduction band. In addition, in the undoped sample, one can see that the Fermi level is closer to the conduction band near the top surface. In our understanding, this can be attributed to a higher concentration of Ti3+ in this region. The Mn doping decreases this effect near the top surface, decreasing the in depth band shift. These XPS data thus confirm that oxygen vacancy related defects are contributing to the leakage current in BST films
Figure 10. XPS Ti 2p3/2 spectra of undoped BST, BST70:Mn0.5% and BST70Mn1% after an 8 nm depth etching.
on three samples, namely undoped BST, BST70:Mn0.5% and BST70:Mn1%. In depth XPS was carried out between the top surface and bottom electrode by means of Ar+ etching. Figure 10 displays the Ti 2p spectra of the three thin films after the first etching step of 8 nm. The typical 459 eV peak corresponding to Ti4+ can be seen. It is well known that ion sputtering may lead to some artifacts in the investigated material [48], however a qualitative comparison can provide relevant information. In the spectra, a noticeable emission at the low binding energy side of the peak maximum occurs. This emission at 457 eV can be attributed to Ti3+ ion formation [49]. Although this Ti3+ ion might be partially due to the argon etching [50], a clear trend of its intensity with the Mn content is observed. The peak shoulder, very 8
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Figure 12. Fermi level position in the bandgap as a function of the relative etched depth for BST70, BST70:Mn0.5% and BST70:Mn1%. The black square dots are the measured valence band energies at different etching steps. The red circular dots correspond to the calculated conduction band energies, determined for each etching level with a 3.3 eV bandgap. The relative etched depth equal to 1 corresponds to the whole film thickness.
which can be improved under Mn doping. We will next see how such doping also improves the dielectric properties of the same films. 3.3. Dielectric properties The dielectric properties of the films were investigated as a function of external dc electric field at 30 ◦ C from 100 Hz up to 1 GHz. The permittivity measured under an electric field at low frequency (LF) is shown in figure 13 at 30 ◦ C, which is in the paraelectric state [52]. A decrease of the zero-field permittivity and of the tunability is observed when the Mn doping rate is raised (figure 13(a)). This nonlinearity, up to 80% under 600 kV cm−1 for the undoped BST70, is almost linearly decreased to 62% for BST70:Mn2%. The radio frequency permittivity measured at 1 GHz under different biases is compared to the low frequency charts in figure 14. One can notice that for the undoped and 0.1% Mn doped BST samples, the low-field RF permittivity value is much lower than the 10 kHz one whereas for the higher doping rate a good continuity is observed between LF and RF. Moreover, the low-field RF permittivity values of the undoped, Mn0.1% and Mn0.5% BST are substantially the same. It is well known that at high frequency, defects, such as electronic and ionic space charges, do not contribute to the permittivity anymore [53]. A relaxation of the permittivity can be observed between low frequencies (102 –103 Hz) and RF (108 Hz). By controlling this phenomenon, BaTiO3 supercapacitors can be obtained [54, 55]. Artemenko et al attributed this effect to oxygen-deficiency-related defects and Ti3+ centers [55]. Consequently, it is sound to assume that in figure 14, the higher permittivity at low frequency of the undoped BST70 sample is related to a similar extrinsic effect. Following the XPS data (figure 10), this extrinsic contribution can also be related to Ti3+ in the undoped film. When the Mn content is increased, the extrinsic contribution is reduced (see figure 14, BST:Mn0.1%), until it disappears for the Mn0.5% and Mn1% where no Ti3+ could be observed.
Figure 13. Dielectric measurements at 10 kHz as a function of electric field for undoped and Mn doped BST70 from 0.1% to 2%. (a) Permittivity versus electric field plot, (b) dielectric loss versus electric field plot.
Thus, the zero-field RF permittivity which does not change under Mn doping corresponds to the intrinsic values of the permittivity. 9
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Figure 14. Comparison of the low and high frequency permittivity under electric field (respectively 10 kHz and 1 GHz) for the undoped BST, and Mn0.1%, Mn0.5%, and Mn1% doped samples. The vertical scale is the same for each chart. Extrinsic contributions to the low frequency permittivity are observed for the undoped and Mn0.1% samples.
Moreover, a dielectric loss increase under high electric field can be observed for undoped and 0.1%Mn doped BST70 in figure 13(b). This effect could be related to free carrier transport in the material under dc field, thus related to conduction [60], and its disappearance from 0.5%Mn doping is coherent with the decrease of the leakage currents observed in figure 3. We thus confirm that the density of oxygen vacancy related defects affects the low frequency extrinsic permittivity and losses, both of them disappearing on doping with Mn. Such a beneficial effect has already been reported in BST thin films and ceramics [30, 61] but a direct link between leakage current, electronic state of titanium and intrinsic RF permittivity could not be found. We next discuss the origin of this beneficial contribution of manganese.
It can be noticed in figure 14 that for the BST70:Mn1%, the zero-field intrinsic permittivity is lower than for the other samples. The dielectric permittivity is directly related to the polarizability, which is decreased by breaking the long range order of dipoles when the Mn ions are inserted in B-sites. This result can also be related to the paraelectric to ferroelectric transition profiles of the films where the transition peak intensity decreases and shifts toward lower temperatures for Mn content higher than 1% [52] (not shown here). This effect has been widely reported for Mn as well as for other dopants [56, 57] and can be explained as follows. The temperature and field dependence of the permittivity ε can be written as ε(T, E) = εL (T) − εNL (T)E2 + o(E4 ),
(14)
with εL and εNL the linear and nonlinear parts of the permittivity respectively. A Landau analysis shows that the temperature dependence of these two parameters is as follows [58]: K1 T − TC K2 εNL (T) = . (T − TC )4 εL (T) =
3.4. Discussion It is well known that the concentration of electronic charge carriers leading to conduction in the material and consequently the Fermi level is influenced by the oxygen partial pressure during the annealing, mostly due to the formation of oxygen vacancies (VO ) and free electrons (e− ). The phenomenon can be explained by the following three equilibriums [62]:
and (15)
Therefore, in the paraelectric phase, the closer to the Curie temperature the measurement is carried out, the higher the nonlinearity is. Considering that when the Mn content is increased the Curie temperature is lowered, this can explain why the Mn doping decreases both the linear and the nonlinear permittivities. The dielectric losses of the films measured as a function of electric field can be seen in figure 13(b). At first, when the Mn content is increased, tan δ is decreased, reaching a minimum of ∼0.9% for the BST70:Mn1%. The decline in the loss is most likely due to the charge compensation of oxygen vacancies again making the link with the density of Ti3+ . For the higher doping rate BST70:Mn2%, the losses are increased to 1.1%. Indeed, it is well known that for high concentration, the Mn ions are not soluble anymore [29, 59], and the properties are degraded [57].
OO(lattice) ⇔ VO + 1/2O2 VO ⇔ V•O + e− − V•O ⇔ V•• O +e . In our case, even if the films are annealed in an oxygen atmosphere, there is no doubt that our material contains such defects. In BST, in order to account for the electro-neutrality of the material, the electrons provided by oxygen vacancies can be equilibrated by the reduction of Ti4+ , Ti4+ + 1e− ⇔ Ti3+ . Hence, the Ti3+ ion becomes a kind of oxygen vacancy marker in the undoped material. In the case of the undoped BST70 film, the presence of such reduced Ti3+ has been highlighted by the XPS 10
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Figure 15. Schematic representation of the electronic structure of defects in the bandgap. (a) For the undoped BST, the configuration is favorable to Schottky injection. (b) For the BST70:MN0.1%, the electronic configuration is compatible with Schottky/PF mixed conduction behavior. (c) For the BST70:Mn1%, the electronic configuration is favorable to Poole Frenkel emission.
measurements presented in figure 10. As a result, the Fermi level is raised to near the conduction band (figure 12). This is corroborated on one hand by the extrinsic contributions to the permittivity at low frequency shown in figure 14 which have also been associated with the presence of Ti3+ [54, 55], and on the other hand by the strong resistance degradation shown in the J–t measurements in figure 1(b) that is related to oxygen vacancy charge migration. Therefore, these results directly explain the high dielectric losses observed in figure 13(b), and the high conductivity of the undoped BST shown in figures 1 and 3. When Mn ions are introduced into the material, a significant decrease of the conduction and the dielectric losses has been observed in figures 1, 3 and 13(b). To explain this result, we based our analysis on the assumption that in BST, the Mn ions are located in a Ti4+ site as definitively assigned by Siegel and Muller [63]. Therefore, the Mn4+ ion can be converted into lower valence states Mn3+ and Mn2+ [12, 64, 65] to compensate the charge brought from the V•• O,
The decrease of the leakage current by increasing the Mn doping rate in the material is associated with a change in the conduction mechanisms (see section 3.1.4). In order to explain this phenomenon the following model is proposed. So as to relate the Ti3+ and Mn3+ concentrations with the trend of the leakage current mechanism shifting from Schottky toward Poole Frenkel as previously shown in section section 3.1.4, we propose to sketch a simplified band diagram. Regarding the undoped sample, a barrier 80 at the interface of about 0.6 eV has been computed from the Schottky model as shown in figure 8(b). For this sample, a large presence of Ti3+ has been directly or indirectly detected through different characterization techniques. According to Moretti et al the Ti–VO centers can induce shallow electron traps, which can modulate the conductivity [66]. Actis et al have shown by calculations that Ti3+ has donor levels localized about 0.1–0.2 eV below the conduction band minimum (CBM) [67]. Recently Mitra et al determined through computing that the V•• O state is the most stable among the oxygen vacancies, with an energy level lying 0.28 eV below the CBM in SrTiO3 [68]. In figure 15(a), we drew a schematic band diagram using these data. It is shown that these energy levels very close to the CBM can easily carry electrons through the film once the electrons have been emitted by the electrode. In other words, only the attenuation of the electrode–film barrier arising from the electrode image–force interaction with the field is controlling the conduction, which is coherent with the interface limited Schottky injection model. Moreover, deeper trap centers can be present in this undoped film, which hardly take part in the conduction, such as Fe3+ impurities, known to be present in many titanates, even in pure crystals [63]. These Fe3+ trap levels in BaTiO3 have been localized 1.6 eV below the CBM [67]. When manganese is introduced into the material, the reduction of Mn4+ into Mn3+ can easily compensate the oxygen vacancy free charges. As a result the concentration of Ti3+ , which is much more unstable, is greatly decreased. In the literature, different calculation results can be found for the energy of Mn trap levels. According to Selme [69], in SrTiO3 when Mn4+ is converted into Mn3+ , the additional electron is localized around 1 eV below the CBM. Moretti [64] found that Mn3+ traps an electron around 0.7 eV below the CBM
Mn4+ + 1e− ⇔ Mn3+ Mn3+ + 1e− ⇔ Mn2+ . Because the energy levels of the Ti3+ are much closer to the conduction band than the Mn3+ , this manganese compensation is more favorable. Therefore, when the Mn concentration is increased, the Ti3+ concentration is decreased, as confirmed by the XPS analysis in figure 10. The Fermi level is lowered into the bandgap, and gets very close to the middle of it, which means that the material shows a more intrinsic behavior. Additional information about the decrease of Ti3+ defects is also pointed out by the dielectric measurements in figure 14, where the extrinsic contribution to the low frequency permittivity is decreased for the BST70:Mn0.1%, and suppressed from BST70:Mn0.5% to higher rates. The time dependence of the leakage current (figure 1(b)) confirms this trend by the decrease of the resistance degradation phenomenon, associated with decrease of the free charge concentration. Finally, when the Mn content is too high, here 2 at.%, a degradation of the properties is observed: an increase of the dielectric losses (figure 13(b)) and an early breakdown process (figure 1(b)). It thus seems that for this doping rate too many defects are created in the system. 11
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References
in BaTiO3 , but according to them, these 3d states are shifted downwards by the presence of a surrounding oxygen vacancy. Therefore, in the case of the BST70:Mn1%, due to the very low concentration of Ti3+ , the barrier energy at the interface 80 becomes lower than the Mn3+ trap height energy 8t , as shown in figure 15(c). This is consistent with a conduction model limited by the bulk part of the film, such as Poole Frenkel, as found in section 3.1.4. For the intermediate doping rate, a mix of Schottky and Poole Frenkel models has been suggested. If we consider, as shown in figure 15(b), that for low Mn concentration, both Ti3+ and deeper trap levels are participating in the bulk conduction, both interfaces and deep trap levels could be limiting the conduction.
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4. Conclusions In this study, we found that it was possible to tune the leakage current mechanism in Ba0.7 Sr0.3 Ti1−δ Mnδ O3−y films by increasing the Mn concentration. As expected, the Mn doping decreased the low frequency dielectric losses at zero-field effectively, as well as the leakage current under an external electric field. In addition, by studying the conduction mechanism in the material and the refractive index, we demonstrated that the conduction is first governed by the Schottky emission at the interface for the undoped BST sample, and an increase of Mn concentration leads to modification of the mechanism, mixing Schottky and Poole Frenkel processes. Then for the BST70:Mn1%, the bulk-like conduction mechanism is mainly due to Poole Frenkel injection. In order to find the origin of this phenomenon, a complementary XPS investigation was performed which evidenced, by a relative comparison, a decrease of Ti3+ with increase of Mn concentration. This led to lowering of the Fermi level into the bandgap, bringing the material closer to intrinsic behavior. This result has been corroborated by additional results on the time dependence of the leakage currents and dielectric measurements at low and high frequencies which evidence indirectly the presence of oxygen vacancies and Ti3+ . Besides the beneficial effect of Mn doping already reported in BST thin films and ceramics [8, 25, 30, 26], we have shown here a direct link between the leakage current, the electronic state of titanium and the intrinsic RF permittivity. Finally, a simplified model has been proposed to explain the evolution of the conduction mechanism, using the theoretical electronic level configuration of the system.
Acknowledgments The authors would like to thank Christine Labrug`ere from ICMCB for the XPS analysis, Dr R Schafranek for the useful discussion about XPS results and Mitsubishi Material Corporation for the sol–gel solutions. Part of this work was supported by the French National Agency for Research under contract ABSYS2 and by the Aquitaine Regional Council. 12
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