Ta 1Àx Nb x 2 O 5 films produced by atomic layer deposition: Temperature dependent dielectric spectroscopy and room-temperature I V characteristics

JOURNAL OF APPLIED PHYSICS VOLUME 90, NUMBER 9 1 NOVEMBER 2001 Ta 1Àx Nb x 2 O 5 films produced by atomic layer deposition: Temperature dependent dielectric spectroscopy and room-temperature I V characteristics M. Strømme a) and G. A. Niklasson Department of Materials Science, The Ångström Laboratory, Uppsala University, P.O. Box 534, S-751 21 Uppsala, Sweden M. Ritala, M. Leskelä, and K. Kukli b) Department of Chemistry, P.O. Box 55, FIN-00014 University of Helsinki, Finland Received 9 April 2001; accepted for publication 27 July 2001 Temperature dependent ac dielectric spectroscopy and room-temperature I V characterization were performed on atomic layer deposited (Ta 1 x Nb x ) 2 O 5 films. The high frequency permittivity, as well as the dc conductivity of the films, were found to increase with increasing Nb content. The conduction mechanism in the mixed Ta Nb oxide films was of the Poole Frenkel type, while the high field conduction in pure Ta 2 O 5 was space-charge limited. The activation energy for dc conduction was higher in mixed Ta Nb oxides compared to pure Ta 2 O 5 and Nb 2 O 5 films. Irreversible changes in the conduction mechanism took place upon heat treatment above a certain irreversibility temperature. This temperature was higher for the mixed oxides than for the binary ones. 2001 American Institute of Physics. DOI: 10.1063/1.1405837 I. INTRODUCTION High permittivity metal oxides are needed for replacing SiO 2 in complementary metal oxide semiconductor CMOS structures and dynamic random access memories DRAM. For this reason, binary oxides such as ZrO 2, 1 Ta 2 O 5, 2,3 Nb 2 O 5, 2,4 7 TiO 2, 3 or Al 2 O 3 Refs. 3 and 8 as well as materials of mixed composition such as Hf Sn Ti O Ref. 9 have been under investigation. TiO 2 and Nb 2 O 5 are characterized by the highest relative dielectric constant 35 120 among binary oxides. However, the performance of these materials in polycrystalline or amorphous form suffers from a markedly increased conductivity due to slight deviations in stoichiometry, in particular, oxygen deficiency. 10,11 In addition, higher permittivity in dielectrics has been shown to be accompanied with higher dielectric losses and conductivity. 12 It is desirable to achieve proper composition and structure of the dielectric materials at low processing temperatures, thus avoiding interface reactions and allowing one to apply low temperature postdeposition annealings, if any. At low processing temperatures, the resulting films are often weakly crystallized or amorphous. Nevertheless, in some cases, e.g., for Ta 2 O 5 layers, amorphicity can be a desired quality, because such films possess the largest homogeneity without grain boundaries acting as channels for the leakage current. Atomic layer deposition ALD a low-temperature method related to chemical vapor deposition is based on the exploitation of highly reactive metal and oxygen precursors, 13 enabling the lowering of reaction barriers and formation of thin solid films in the temperature range of 80 C 500 C. 14 ALD is realized via sequential introduction of either precursor in the gas phase into the reaction zone. Monomolecular layers of the precursors adsorb and react at the substrate surface successively, resulting in the layer-bylayer growth. Some other qualities characteristic of ALD are excellent conformality on trenched substrates 15 and good thickness and composition uniformity. 16 Due to the stepwise, cyclic, introduction of the precursors, intrinsic advantages of the ALD comprise also a simple control of film thickness with nanometer accuracy as well as uniform dopant concentration throughout the film thickness. ALD is an appropriate technique for processing high permittivity materials, for example films of ZrO 2 Ref. 1 and Al 2 O 3, 8 Ta 2 O 5 HfO 2, and Ta 2 O 5 ZrO 2 nanolaminates 17,18 or various homogeneous mixture oxides and solid solutions. 19 ALD growth of Ta 2 O 5 Ref. 20 and Nb 2 O 5 Ref. 21 has been realized as well, resulting in pentoxide films stoichiometric within the accuracy limits of Rutherford backscattering analysis. The composition of Ta 2 O 5 Nb 2 O 5 mixture films as well as nanolaminates has been adjusted simply by changing the number and ratio of the deposition cycles of either binary oxide layers. 16,22 The real time monitoring of the surface reactions between metal precursor and oxygen source has indicated a sharp transfer between the regimes of Ta 2 O 5 and Nb 2 O 5 growth, 23 allowing to rely on the completeness of the surface reactions and removal of the excess precursor from the reaction zone. Niobium additives in the tantalum oxide have resulted in an increase in the permittivity. 16,22,24 A remarkable increase in permittivity from 25 to 165 was earlier obtained in crystalline Nb 2 O 5 :Ta solid solution powders processed at high temperatures. 25 Leakage current characteristics were not rea Electronic mail: maria.stromme@angstrom.uu.se b Also at: Institute of Experimental Physics and Technology, University of Tartu, Tähe 4, EE-51010 Tartu, Estonia. 0021-8979/2001/90(9)/4532/11/$18.00 4532 2001 American Institute of Physics

J. Appl. Phys., Vol. 90, No. 9, 1 November 2001 Strømme et al. 4533 ported. In the ALD (Ta 1 x Nb x ) 2 O 5 films, improved permittivity values were achieved without increasing high field leakage currents 16 when the Nb concentration did not exceed 25 cation %. The ohmic, or low field, conductivity was however not investigated. An addition of small amounts of Ta 2 O 5 to the Nb 2 O 5 matrix stabilizes the dielectric performance of Nb 2 O 5 and reduces the conductivity, though simultaneously reducing the relative dielectric constant. 26 This study has been aimed at a more detailed investigation of the electric properties of Ta 2 O 5, Nb 2 O 5, and (Ta 1 x Nb x ) 2 O 5 films. Temperature dependent ac dielectric spectroscopy was applied to obtain information about the behavior of the frequency dependent permittivity and I V characteristics were examined to learn about the conduction mechanisms governing the high field currents in the films. II. FILM PREPARATION The Nb 2 O 5 and Ta 2 O 5 films and (Ta 1 x Nb x ) 2 O 5 mixture films have been grown at 325 C in a hot-wall flow-type F120 ALD reactor Microchemistry Ltd.. Nitrogen was used as a precursor carrier and purge gas. The pressure in the reactor was about 10 mbar. Ta OC 2 H 5 5 Aldrich and Nb OC 2 H 5 5 ABCR were evaporated from open boats held at 100 C and 90 C, respectively, inside the reactor. Water vapor was generated in an external reservoir at room temperature and led into the reactor through needle and solenoid valves. The carrier gas flow was 400 sccm. The films were grown applying atomic layer deposition recipes well established separately for Ta 2 O 5, 20 Nb 2 O 5, 21 and (Ta 1 x Nb x ) 2 O 5 solid solutions. 22 Pulse time length of 0.2 s was used for metal precursors. The water exposure time was 2.0 s. Purge time periods of 0.5 s were used between metal and oxygen precursor pulses in order to avoid the gas phase reactions and remove the volatile byproducts of surface reactions. The relative cation content in mixture films was modified by varying the ratio between the amounts of successive Nb 2 O 5 and Ta 2 O 5 deposition cycles. The total film thickness was kept between 93 and 165 nm. Nb Ta mixed oxides were capped on both sides by ca 5 10 nm thick amorphous Ta 2 O 5 buffer layers to avoid crystallization of the films. III. FILM CHARACTERIZATION Film thicknesses were calculated, using the method described by Ylilammi and Rantaaho 27 from optical transmission spectra measured by a Hitachi U-2000 spectrophotometer within a wavelength range of 370 1100 nm. Film structure was determined by means of a Philips MPD 1880 powder x-ray diffractometer using Cu K radiation. The relative cation contents were measured by a Link ISIS energy dispersive x-ray spectrometer EDX installed in a Zeiss DSM 962 scanning electron microscope. The composition as well as the film thickness of the samples under study are listed in the Table I. For measuring dielectric properties, 5 cm 5 cm soda lime glasses covered with 200 nm thick Al 2 O 3 barrier layer and patterned indium-tin oxide ITO electrodes were used as film substrates. An array of aluminum electrodes was evapo- TABLE I. Thicknesses of (Ta 1 x Nb x ) 2 O 5 layer in the glass/al 2 O 3 / ITO/ Ta 1 x Nb x ) 2 O 5 /Al capacitors studied and temperatures at which irreversible changes occur in the dielectric response are presented. Composition Film thickness nm Irreversibility temperature C Ta 2 O 5 165 80 Ta 0.6 Nb 0.4 2 O 5 93 160 Ta 0.52 Nb 0.48 2 O 5 135 200 Ta 0.22 Nb 0.78 2 O 5 140 200 Nb 2 O 5 165 120 rated onto the film surface, enabling electrical measurements over a substrate area of 10 cm 2, approximately. The effective single electrode area was 29.4 mm 2. The films were not heat treated after the deposition. The following characterization results, thus, describe the properties of the films in their asdeposited state. The Nb 2 O 5 films, which are amorphous on glass substrates, 21 were crystallized remarkably when deposited directly onto polycrystalline ITO surface showing the reflections of either monoclinic, 21 orthorhombic or hexagonal TT Nb 2 O 5. 16 The crystallization of the Nb 2 O 5 was inhibited by depositing a 10 nm thick Ta 2 O 5 layer between ITO and Nb 2 O 5. Accordingly, all films in this study are amorphous except the Nb 2 O 5 films. The electrode area, 29.4 mm 2, is large. Nevertheless, the films have been quite defect free and possessed stable dielectric characteristics. Some of the capacitors were broken down while the current voltage I V characteristics were examined at high fields. Nevertheless, most often the breakdown was self-healing, so that the capacitance permittivity was usually possible to measure repeatedly, also after the I V measurements. This refers to the high quality of the films in their as-deposited state, allowing one to rely on the pinhole-free structure and a very low lateral density of defects. IV. TEMPERATURE DEPENDENT AC DIELECTRIC SPECTROSCOPY A. Measurement procedures For the ac dielectric measurements, a Novocontrol broadband dielectric converter was used together with a Solartron 1260 frequency response analyzer. The sample cell was a Quiet CHUCK dc hot chuck system with temperature control. Tungsten needle electrodes contacted the ITO electrode underneath the films, exposed in one point by scratching, and the aluminum electrode on top of the films. The amplitude of the sinusoidal voltage used in the measurements was 50 mv. The temperature was varied between 24 C and 200 C and the frequency range was approximately 1 mhz to 10 MHz. The measurements were performed at increasing temperatures in the following manner. First, the roomtemperature dielectric spectrum was recorded. Then, the temperature was increased by approximately 20 C and a spectrum was measured. Between each frequency scan at

4534 J. Appl. Phys., Vol. 90, No. 9, 1 November 2001 Strømme et al. FIG. 1. Real a and imaginary b parts of the dielectric permittivity of (Ta 1 x Nb x ) 2 O 5 films as obtained by dielectric spectroscopy at room temperature are shown. elevated temperatures, a room temperature dielectric spectrum was recorded to check the reversibility of the measurements. B. Results 1. Room temperature data Figures 1 a and 1 b display the real ( ) and imaginary ( ) part of the dielectric permittivity of the investigated films at room temperature. Several features are common for all films. a At high frequencies, there is a peak in and a rapid decrease in due to the ITO electrode. By modeling the presence of the ITO electrode with a series resistance and subtracting this resistance from the total response, these artifacts are suppressed. b In a frequency region about 10 4 10 5 Hz, the high frequency dielectric constant can be extracted from the data following the procedure outlined in Ref. 28. is thus a measure of the dielectric constant of the material in the MHz region. c At intermediate frequencies, there is a relaxation peak in and a step in which is shifted towards increasing frequencies with increasing Nb content. d At low frequencies, 1 and levels out to a constant value. This is the characteristic behavior for dc conduction and a conductivity dc can be extracted from the data. 28 It should be noted that the measurement equipment is able to perform accurate measurements of only when the ratio between and takes a value in a limited interval. This is the reason why are presented for lower frequency values than the data. The results of the room-temperature dielectric measurements are summarized in Fig. 2. Figure 2 shows, as a function of the Nb concentration, a, the position in frequency of the relaxation peak in b, the magnitude of the step in connected with the relaxation peak c, and dc d. For the pure Nb 2 O 5 film, and the step in could only be very roughly estimated because dc conduction is present at very high frequencies. In Fig. 2 d, we notice that different values of dc are presented for the Ta 2 O 5 and the Ta 0.60 Nb 0.40 2 O 5 films. By using different contacts on these samples, we sometimes measured different conductivity values. All other parameters, however, were found to be the same. From Fig. 2, we find that the high frequency dielectric constant increases with increasing Nb concentration, as observed earlier, 16 and that the increase is very rapid when the Nb content changes from 40 to 48 cation %. We also observe that as the relaxation peak in moves towards higher frequencies with increasing Nb concentration, the peak becomes stronger. The general trend for the low field dc conductivity is an increase with increasing Nb concentration. 2. Temperature dependent data Figure 3 shows an example of a series of dielectric spectra taken at different temperatures on the Ta 0.52 Nb 0.48 2 O 5 sample. As can be seen from Fig. 3, the and data are shifted towards higher frequencies with increasing temperature. In Fig. 3 we notice that the curve taken at 200 C is different from the rest of the curves; the low frequency value is lower. The room-temperature spectrum recorded after every temperature scan below 200 C was identical to the room-temperature scan taken on the as-deposited film. After the 200 C scan was made, though, we were not able to reproduce the initial room-temperature curve. An irreversible change in the film had occurred at 200 C. Corresponding irreversibility temperatures are given in Table I for all films in this study. The general feature for all room-temperature scans made after the irreversibility temperature is that the relaxation peak is weaker and the dc conductivity is higher than in the as-deposited film. For temperatures below the irreversibility temperature, the activation energies for the process governing the relaxation peak and the dc conductivity were extracted by using the masterplot technique. 29 The activation energies, that were found to be constant in the reversible temperature range, are shown in Fig. 4. From Fig. 4, it is obvious that the dielectric relaxation process giving rise to a peak in has an activation energy around 0.6 ev

J. Appl. Phys., Vol. 90, No. 9, 1 November 2001 Strømme et al. 4535 FIG. 3. Dielectric spectra taken at different temperatures on a Ta 0.52 Nb 0.48 2 O 5 film are shown. Panel a and b show the real and imaginary part of the permittivity, respectively. FIG. 2. High frequency permittivity a, position in frequency of the relaxation peak in b, magnitude of step in connected with relaxation peak c, and dc conductivity d shown vs Nb content in (Ta 1 x Nb x ) 2 O 5 films. independent of film composition. For the dc conduction process, the activation energy is relatively low for the pure Ta 2 O 5 and Nb 2 O 5 films and high for the (Ta 1 x Nb x ) 2 O 5 films. V. I V CHARACTERIZATION A. Measurement procedures For the I V measurements, an ECO Chemie Autolab/ GPES interface was used together with the same sample cell as for the dielectric spectroscopy. The maximum current that can be measured reliably using this equipment is 35 ma and the maximum potential that can be applied is 10 V. All measurements were performed at room temperature. The potential across the films was stepped to successively higher values between 0.01 and 10 V, and the current flowing through the films 30 s after each potential step, was measured. In a time interval of 30 s, the transient current in the films under study has had time to relax and take a constant value. Before presenting the results on the various films, we will have a brief look at the conduction mechanisms of interest for our samples. B. Conduction mechanisms In this section, we will recapitulate some background to space-charge and Poole Frenkel conduction since these are the mechanisms most likely governing the I V response presented in the next section.

4536 J. Appl. Phys., Vol. 90, No. 9, 1 November 2001 Strømme et al. Here, e is the electron charge, k is the Boltzmann constant, N n is the trap concentration per unit energy at the thermal equilibrium Fermi level F 0,d is the film thickness, and 0 and 0 are the static dielectric constant of the material and the permittivity of free space, respectively. Accordingly, the trap concentration at an arbitrary energy W is given by 32 N t W N n exp W F 0 kt t. 3 The shift in the Fermi level F due to the charge forced into the insulator by the applied voltage is given by 31 F kt t K ln V, where K is a voltage independent constant. 4 2. Poole Frenkel conduction FIG. 4. Activation energies for dc conduction process and the relaxation peak process in (Ta 1 x Nb x ) 2 O 5 films vs Nb content are shown. 1. Space-charge limited conduction A basic requirement for an electrical conduction process to be space-charge limited is that the contribution of thermally generated carriers is smaller than that of injected carriers. 30 For this to be possible, the dielectric sample must be relatively free from traps. Trap densities much higher than 10 24 m 3 are sufficient to reduce the space-charge limited currents to almost unmeasurable values. 31 As the potential across a dielectric sample with ohmic contacts is increased, the current normally starts out being carried by thermally excited carriers. At sufficiently high voltages, however, the injected charges may be the dominant current carriers. The current in the low potential region is proportional to the applied potential; it follows Ohm s law. In the high potential space-charge region, the current has a power-law potential dependence with an exponent equal to 2 in the case when there are no electron traps present or when the traps are shallow, i.e., when they are located at an energy larger than kt the thermal energy above the Fermi level. In amorphous insulators, where traps are distributed in energy rather than localized at a precisely defined energy, the square potential dependence of the current is no longer observed, and the actual potential dependence can be used to extract information about the nature of the traps present in the material. For example, when the current I depends on the potential V as I V T t /T 1, 1 with T t T, the trap distribution is exponential cf. Eq. 3 and information about trap densities can be obtained. 31,32 Here T is the lattice temperature and T t is a parameter characterizing the trap distribution. The crossover voltage from the ohmic conduction region to the space-charge limited region is, in this case, given by 32 V X ekt tn n d 2 0 0. 2 It is very likely that defects are embedded in the matrix of amorphous dielectrics. If the defect centers can be ionized, the electrons made available in this process are free for conduction until they become trapped by other ionized sites. If we assume that this donor density is such that the coulombic fields of the sites are just beginning to overlap, it can be shown that the field dependent conductivity is proportional to exp(ae 1/2 ) where A is a field independent constant. 33 Conduction with such a field dependence is known as Poole Frenkel conduction. This expression for the conductivity neglects current flowing against the applied electric field and it is therefore valid at high fields only. Taking into account the reverse current, it can be shown that the exponential conductivity should be replaced by a sinh term to give the total current as 34 I 2Ae x 1/2 0eE 2m* exp W i e kt sinh 3/2 E 1/2 1/2 n kt 0 i. 5 Here, A is the cross sectional area of the sample, x 0 is half the distance between neighboring impurities, E is the applied electric field, W i is the energy required to ionize a donor under zero field conditions, m* is the effective electron mass, and n i is the density of donors. For low fields, Eq. 5 simplifies to Ohm s law, while for high fields, the sinh term can be replaced by an exponential term to give the Poole Frenkel relationship. In the high-field region, where the sinh term in Eq. 5 is replaced by an exponential term, the pre-exponential factor is proportional to E 1/2. This field dependence of the prefactor means that after emission the carriers travel a distance which is proportional to E 1/2 before they are recaptured in an empty donor. 35 In another model, the prefactor is proportional to E or V, 36 and the distance traveled before recapture, thus, is proportional to the electric field. The high field current in this model has the expression 36 I VA 0 d exp e 3/2 E 1/2 kt 0 1/2, where 0 is the low-field conductivity. With a donor concentration higher than 10 24 ionized donors/m 3, Eq. 5 no longer holds because of significantly 6

J. Appl. Phys., Vol. 90, No. 9, 1 November 2001 Strømme et al. 4537 To get a first clue as to what conduction mechanism is governing a measured I V response curve, it has proven very successful to study a log log plot of vs log E. 38 The parameter is defined according to d ln d 1/E nce1 n. 7 Here, is the field dependent conductivity of the sample and C is a constant at constant temperature. The parameter n indicates what kind of conduction mechanism controls the measured current. For space-charge limited conduction n 0, for Poole Frenkel conduction it takes a value of 0.5, while for Poole s law n 1. 38 Figure 5 shows the measured current response curves as well as the calculated conductivities as a function of the applied electric field. We see that at low fields, E 10 6 V/m, the current response is ohmic with constant conductivities for all materials in this study except for the pure Ta 2 O 5 film. For this film, the current in the low-field region is within the range of the lower detection limit of the instrument, thus, being a bit noisy and producing uncertain conductivity values. However, it should be noted that a shallow minimum in the field dependence of dc has been observed before in Ta 2 O 5 films. 39 This feature is characteristic of disordered materials with hopping conduction. At higher fields, we find that the conductivities increase with increasing field strengths. To investigate the conduction mechanisms controlling the current in the high-field region, we start by applying the method described by Eq. 7 to our data. Accordingly, Fig. 6 presents a plot of vs E. We note that in Fig. 6 a the plot is made with linear axis, while in the other plots the axes are logarithmic. The parameter n extracted by fitting Eq. 7 to the data is presented in each plot. Next, we analyze the current response of each film by applying the theory suggested by the obtained n value. FIG. 5. Current response a and calculated conductivity b as a function of applied electric field for (Ta 1 x Nb x ) 2 O 5 films are shown. Ta 2 O 5, Ta 0.52 Nb 0.48 2 O 5, Nb 2 O 5, (Ta 0.60 Nb 0.40 2 O 5, and Ta 0.22 Nb 0.78 2 O 5. overlapping potential wells. In this case, one would expect the current to follow Poole s law in which the logarithm of the current is proportional to E. 37 C. Results 1. Ta 2 O 5 Figure 6 a shows that n 0, which indicates that the conduction mechanism in the high-field region is spacecharge limited. We, thus, make two fits to the measured current response; in the low-field region we use Ohm s law and in the high-field region proportionality relation 1 is fitted to the current. The result is shown in Fig. 7. From the fit to relation 1, we find that the parameter T t, characterizing the trap distribution in the film, is equal to 2.58 T here T is room temperature. This shows that the trap distribution in the material is an exponential function of energy (T t T), which is to be expected since we are dealing with an amorphous material. 32 Taking the analysis one step further, we find that the crossover voltage from the ohmic conduction region to the space-charge limited region is approximately equal to 0.5 V. From this V X value and with (0) 136, the low frequency value of as obtained from dielectric spectroscopy, we find that N n equals 2.1 10 24 traps per ev and m 3 at the thermal equilibrium Fermi level. We observe space-charge limited conduction in the Ta 2 O 5 film for voltages between 0.5 and 10 V, which was the highest attainable voltage in this study. Using Eq. 4, we find that the Fermi level is moved upwards in the band gap by 0.19 ev as the potential is increased from 0.5 to 10 V. Integrating Eq. 3, we find that there are 2.4 10 24 traps per m 3 in the material in this energy range. 2. Ta 0.60 Nb 0.40 2 O 5 Figure 6 b shows that n 0.47 which indicates that the conduction mechanism most likely is Poole Frenkel in the high-field region. Equation 5, which describes both high- and low-field current response in a material with electron donors not too closely spaced is, thus, fitted to the current measured in the Ta 0.60 Nb 0.40 2 O 5 film. The result is shown in Fig. 8 a. It is immediately seen that the fit is very good over the three orders of magnitude in the electric field covered by our measurement, thus leaving no doubt as to the

4538 J. Appl. Phys., Vol. 90, No. 9, 1 November 2001 Strømme et al. FIG. 6. The parameter, obtained from I V measurements on (Ta 1 x Nb x ) 2 O 5 films, vs applied field is shown. The parameter n is extracted by fitting solid line the data to Eq. 7. Note that the axes in a are linear and in b e logarithmic. conduction mechanism being Poole Frenkel. From the fitting parameters used, we find an value of 18.4, which is somewhat lower than the value of 24.0 obtained from dielectric spectroscopy. Inserting for W i in Eq. 5, the dc conduction activation energy of 0.52 ev obtained in Sec. IV B, we find the density of donors in terms of the effective, m*, and the free electron mass, m e as 3/5 m* n i 1.15 1019 m e per m 3. If we instead use Eq. 6, where the high field preexponential factor, discussed in Sec. V B, is proportional to E or V instead of E 1/2, we obtain the fit shown in Fig. 8 b. Here, the current is divided by the applied potential and plotted vs the square root of the electric field. From this fit we extract a value of 23.7 in very good agreement with the dielectric spectroscopy value. 3. Ta 0.52 Nb 0.48 2 O 5 Figure 6 c shows that n 0.69 which is somewhat higher than what is expected for Poole Frenkel conduction. It is nonetheless closer to the Poole Frenkel n value of 0.5 than to that of Poole s law (n 1). A fit to Eq. 5, Fig. 9 a, shows that this equation describes the current response well over almost three orders of magnitude in applied electric field. From the fitting parameters, used we find an value of 16.4, which is substantially lower than the value of 40.4 obtained from dielectric spectroscopy. Inserting for W i in Eq. 5, the dc conduction activation energy of 0.56 ev obtained in Sec. IV B, we find the density of donors as n i 6.54 10 21 (m*/m e ) 3/5 per m 3. Now, using Eq. 6, we obtain the fit shown in Fig. 9 b. From this fit, we extract a value of 37.9 in much better agreement with the dielectric spectroscopy value than the value found from the fit in Fig. 9 a.

J. Appl. Phys., Vol. 90, No. 9, 1 November 2001 Strømme et al. 4539 FIG. 7. Current response squares in Ta 2 O 5 vs applied potential is shown. The response is fitted to relation 1 in the high potential region solid line and to Ohm s law in the low potential region dashed line. 4. Ta 0.22 Nb 0.78 2 O 5 Figure 6 d shows that n 0.50 which, again, divulges that a Poole Frenkel type conduction mechanism is responsible for the behavior of the current in the high-field region. We thus analyze the current response for all applied fields with Eq. 5, Fig. 10 a. As for samples B and C, we find a very good fit to this equation. The fitting parameters give an value of 17.1, which just like in the case of the Ta 0.52 Nb 0.48 2 O 5 film is much lower than the high frequency permittivity value of 43.4 obtained from dielectric spectroscopy. Inserting for W i in Eq. 5, the dc conduction activation energy of 0.49 ev, we get the density of donors as 3/5 m* n i 5.58 1021 m e per m 3. When using the current model where the high-field preexponential factor is proportional to E Eq. 6, we get the fit shown in Fig. 10 b and can extract an value of 35.3. As for the Ta 0.52 Nb 0.48 2 O 5 film, this value is closer to the dielectric spectroscopy than that found from the fit to Eq. 5. 5. Nb 2 O 5 An n value of 0.32 is extracted from Fig. 6 e, and the conduction mechanism might be Poole Frenkel (n 0.5) or space-charge (n 0) limited in the high-field region. Figure 11 a shows a fit to relation 1 and to Ohm s law while Eq. 5 is fitted to the current response in Fig. 11 b. By assuming that the conduction mechanism is space-charge limited Fig. 11 a, we find that the current in the high-field region is proportional to V 1.48. This potential dependence is, to our knowledge, not in agreement with any model describing space-charge limited currents in real solid materials. Analyzing the current response by instead assuming Poole Frenkel conduction at high fields, we find from the fit in Fig. FIG. 8. a Current response circles in Ta 0.60 Nb 0.40 2 O 5 vs applied field is shown. The response is fitted to Eq. 5 solid line. b Current divided by potential circles vs square root of applied field is presented. The response is fitted to Eq. 6 solid line. The high frequency permittivity, as obtained from the fits, are displayed. 11 b an value of 31.0 which again is much lower than the approximate value of 130 obtained from dielectric spectroscopy. Inserting W i 0.20 ev, as obtained from dielectric spectroscopy, in Eq. 5 we get a density of donors as low as n i 1.44 10 13 (m*/m e ) 3/5 per m 3. From the fit in Fig. 11 c to Eq. 6, an value of 163 is extracted. VI. DISCUSSIONS The results obtained by dielectric spectroscopy show that the high frequency permittivity of the (Ta 1 x Nb x ) 2 O 5 films increases with increasing Nb content. The increase was found to be most rapid at a Nb concentration of 40 48 cation %. For the Ta 2 O 5 and the Ta 0.60 Nb 0.40 2 O 5 films, we found that the dc conductivity had different values in different capacitors measured, while properties such as high frequency permittivity, relaxation peak strength, and activation energies remained the same. No further investigations were made on the different dc conductivity values of these films. It may be possible that exposure to high currents or temperatures under certain circumstances can change the room-temperature con-

4540 J. Appl. Phys., Vol. 90, No. 9, 1 November 2001 Strømme et al. FIG. 9. a Current response diamonds in Ta 0.52 Nb 0.48 2 O 5 vs applied field is shown. The response is fitted to Eq. 5 solid line. b Current divided by potential diamonds vs square root of applied field is presented. The response is fitted to Eq. 6 solid line. The high frequency permittivity, as obtained from the fits, are displayed. FIG. 10. a Current response triangles in Ta 0.22 Nb 0.78 2 O 5 vs applied field is shown. The response is fitted to Eq. 5 solid line. b Current divided by potential triangles vs square root of applied field is shown. The response is fitted to Eq. 6 solid line. The high frequency permittivity, as obtained from the fits, are displayed. ductivity. Earlier studies 40 have shown that switching and negative resistance may occur in Ta 2 O 5, i.e., the dc value can switch to a low conductivity state after exposing the material to high electric fields. The large increase in the high frequency permittivity at a Nb concentration around 40 48 cation % suggests that a more thorough study should be performed on how to lower the conductivity at these Nb contents. If a lower dc conductivity can be achieved, such Ta Nb oxides would serve as excellent materials for DRAM or CMOS structures. The ohmic dc conductivity also increases with Nb doping cf. Fig. 2 d. The reason for this is either an increase in the number of charge carriers, or in the mobility, or both, as Nb is added to the Ta Nb oxide matrix. The activation energy for the dc process does however not follow such a simple Nb content dependence. We found that the pure pentoxide films had lower activation energies than the Ta Nb mixed oxides. This, of course, is closely linked to the conduction processes observed at higher fields. We now discuss the conduction mechanisms present in our films. The ac measurements give information on the dc conductivity and relaxation process, while the I V measurements give evidence on the high-field conduction in the films. It should be noted that conduction mechanisms at low and high fields are not necessarily identical. The relaxation peak was present in all samples with nearly the same activation energy. This strongly suggests a common origin for the relaxation process. The conduction processes in the mixed oxides are the easiest to interpret. It was clearly established that the highfield conduction mechanism is of Poole Frenkel type. In fact, the whole I V curve from low to high fields is in excellent agreement with Eq. 5. Furthermore, the activation energies of dc conductivity and the relaxation peak were very similar. This is readily understood from the Poole Frenkel mechanism. The dc conductivity is due to charge carriers excited from donor levels to the conduction band,

J. Appl. Phys., Vol. 90, No. 9, 1 November 2001 Strømme et al. 4541 FIG. 11. a Current response triangles in Nb 2 O 5 vs applied potential is shown. The response is fitted to relation 1 in the high potential region solid line and to Ohm s law in the low potential region dashed line. b Current response triangles vs applied field is shown. The response is fitted to Eq. 5 solid line. c Current divided by potential triangles vs square root of applied field is shown. The response is fitted to Eq. 6 solid line. The high frequency permittivity, as obtained from the fits, are displayed in panels b and c. while the relaxation peak is due to emitted charge carriers being subsequently retrapped. It is found that the donor levels are situated approximately 0.5 to 0.6 ev below the conduction band. In this picture, the strength of the relaxation peak is a function of the donor density and the capture and emission rates of the charge carriers, while the peak position would depend on the emission rates. 41 It is, however, not clear whether the field dependence of the pre-exponential factor in the expression for the high-field current should be E 1/2 or just E. Assuming the latter gives values that match the values obtained from dielectric spectroscopy better than the values obtained assuming an E 1/2 field dependence. The E 1/2 dependence predicts values lower than those from dielectric spectroscopy. The E 1/2 field dependence can, however, not be excluded by this simple argument. The dielectric spectroscopy is extracted at frequencies around 1 MHz. At higher frequencies there may be other processes taking place that lowers substantially. This is a direct consequence of the fact that the refractive index measured in the visible frequency range 16 is much lower than at 1 MHz. What value of should be used in a Poole Frenkel analysis depends on the response time of the medium surrounding the donor. If the medium is able to polarize within the emission time, an obtained at lower frequencies than that signifying the emission time should be used, and if the medium does not have time to polarize during the emission time, a lower value taken at a higher frequency is appropriate. 42 For the pure Ta 2 O 5 film, we found that the high-field conduction mechanism was space-charge limited with an exponential distribution of traps below the conduction band. We argue that the low-field conduction mechanism is different, though, and most likely due to hopping around the Fermi level. There are two features that support this interpretation: First the minimum in the field dependence of the dc conductivity and secondly the small activation energy are typical of hopping processes. It appears that the contribution of injected carriers to the conductivity increases with the applied field and dominates above the threshold voltage V X. The dc conductivity exhibits a much smaller activation energy than the relaxation peak. This means that they are due to different processes. The relaxation process is most likely due to emission of charge carriers to the Fermi level, or the conduction band, followed by retrapping. For the pure Nb 2 O 5 film, we were not able to clearly establish the conduction mechanism. Analyzing the current response with a Poole Frenkel model, we obtained a suspiciously low density of impurity sites, while assuming that the current response was space-charge limited, we got an unphysically low T t. We may notice that there are only current response data available for applied fields up to 10 7 V/m for the pure Nb 2 O 5 film. This is because the film is so conductive that the maximum current detection limit for the instrument used is reached at a lower field strength than for the rest of the films in this study. The results of the Poole Frenkel analysis should, however, not change very much with more current data at high fields since the analysis already is based on fields spanning over more than two orders of magnitude. For the space-charge model, the situation is different. When extracting T t it might well be that current data in the transi-

4542 J. Appl. Phys., Vol. 90, No. 9, 1 November 2001 Strømme et al. tion region between ohmic and space-charge conduction were used and that a larger value of T t would be obtained with accessibility to currents measured at higher fields. Hence, the current response at high fields might well be space-charge limited also for the pure Nb 2 O 5 film. With this limited information on high-field currents, other conduction mechanisms can, of course, not be ruled out. The activation energy for dc conduction is lower than that for the relaxation process, which is similar to the situation in Ta 2 O 5. However, the dc conductivity level as well as the strength and position of the relaxation peak are so different in Nb 2 O 5 that no conclusions regarding the mechanisms can be drawn from this. Finally, we comment on the origin of the donor states being responsible for the relaxation peaks and, in some cases, the Poole Frenkel conduction. Sawada and Kawakami 43 have found that oxygen vacancies give rise to localized states about 0.8 ev below the band gap in crystalline Ta 2 O 5. This was inferred from bandstructure calculations using the local density approximation, which systematically underestimates energy gaps. However, it was argued that oxygen vacancy states could be more shallow in disordered tantalum oxide. Another possibility is that the donor states in our films are associated with the presence of OHgroups or chemisorbed water. We have found that the strength of the relaxation peak decreases when the samples are subjected to heat treatment above the irreversibility temperature. This shows that the number of donors decreases and that this is most likely due to an oxidation process. VII. SUMMARY The high frequency permittivity of our ALD (Ta 1 x Nb x ) 2 O 5 films was found to increase with increasing Nb concentration. The same was true for the ohmic dc conductivity. In the mixed Ta Nb oxide films the conduction mechanism could be described by Poole Frenkel conduction. In pure Ta 2 O 5, the high-field conduction was spacecharge limited while the conduction mechanism in pure Nb 2 O 5 could not be specified. The activation energy for dc conduction was higher in mixed Ta Nb oxides compared to pure Ta 2 O 5 and Nb 2 O 5 films. Irreversible changes in the conduction mechanism took place upon heat treatment above a certain temperature. This irreversibility temperature was higher for the mixed oxides than for the binary oxides. ACKNOWLEDGMENTS This study was financially supported by the Swedish Natural Science Research Council, the J. Gust. Richert Foundation, the Academy of Finland, the Finnish National Technology Development Agency, and the Estonian Science Foundation Grant No. 4205. The authors appreciate the access to the EDX measurement facilities at the Electron Microscopy Unit at the University of Helsinki. 1 M. Copel, M. Gribelyuk, and E. 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