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Piezoelectric properties of sputtered PZT films: influence of structure, micro structure, film thickness (Zr,Ti) ratio and Nb substitution
Materials Science in Semiconductor Processing 5 (2003) 123–127
Piezoelectric properties of sputtered PZT films: influence of structure, micro structure, film thickness (Zr,Ti) ratio and Nb substitution D. Remiens*, E. Cattan, C. Soyer, T. Haccart IEMN–DOAE–MIMM-UMR, CNRS 8520, Universit!e de Valenciennes, ZI petite savatte, 59600 Maubeuge, France
Abstract Lead titanate zirconate Pb(Zr,Ti)O3 (PZT) with different (Zr,Ti) ratio and Niobium (Nb) doped PZT thin films were deposited on silicon substrates by R.F. magnetron sputtering. The Zr/Ti ratio varied between the quadratic to the rhombohedric phases. For the thickness influence study, we have chosen PZT films with the 54 46 composition. The Nb concentration incorporated in PZT films varied between 1 at% and 7 at%. The electrical properties (dielectric, ferroelectric and piezoelectric) were evaluated functions of the film thickness, the orientation, the (Zr,Ti) ratio and the Nb content. r 2002 Published by Elsevier Science Ltd. Keywords: Piezoelctric thin films; Sputtering; PZT materials; Electrical properties
1. Introduction Lead zirconate titanate (PZT) thin films have been extensively investigated for many applications [1–3]. An important parameter to develop MEMS applications, is the piezoelectric performance of the films. The piezoelectric properties of PZT ceramics can be modified by the substitution of dopants [4]; lanthanum and niobium ions have been currently used. In thin film form, only very few papers describe the Nb influence on the PZT structural and electrical properties. Usually papers focus on the dielectric and ferroelectric properties [5,6]. For MEMS device applications, it is also imperative to perfectly know the influence of the film characteristics such as: the microstructure, the structure and the interfaces, the film thickness, etc. on the electrical properties. Many papers have been pub were in many cases in contradiction [7,8]. Concerning the evaluation of the e31, d 31 and d 33 coefficients dependence with the *Corresponding author. E-mail address: [email protected] (D. Remiens).
PZT film thickness there is a limited number of papers [9]. The growth and characterization of PZT with different (Zr,Ti) ratio, orientation and thickness and PNZT were presented in this paper.
2. PZT and PNZT films preparation and characterization An R.F. magnetron sputtering system, described previously [10], was used to prepare the films. We have used single oxide targets since they are easy to fabricate and facilitate compositional change. Nb5+ was considered as a donor dopant for PZT materials; it substitutes Ti4+/Zr4+ ions. So, since the niobium incorporation occurs on the B-site; the compensation of the additional positive charge introduced, to achieve electroneutrality, is assumed through A-site, i.e. Pb-site. The films were deposited onto oxidized Si (100) n-type substrates ( ( electrodes. The coated with Ti (100 A)/Pt (1500 A) 15 PZT (with (Zr/Ti) ranging from 85 to 70 30) and PNZT sputtering conditions are summarized in Table 1. An annealing treatment is necessary to crystallize the films
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D. Remiens et al. / Materials Science in Semiconductor Processing 5 (2003) 123–127
Table 1 Typical sputtering and annealing conditions of PZT and PNZT films Target diameter (mm) R.F. power density (W/cm2) Gas pressure (mT) Inter electrodes distance (mm) Substrate temperature (1C) Annealing temperature (1C) Annealing time (mm) Atmosphere
76 3 30 (Ar) 60 150 (ion bombardment) 625 30 Air
in the perovskite phase. We present here only the results relative to the conventional treatment. The annealing conditions were given in Table 1. The films are crystallized in the pure polycristalline perovskite phase but two types of preferred orientation have been obtained: the (1 1 0) or the (1 1 1) orientation. At present, it is difficult to explain why the orientation is different; the sputtering and annealing conditions are identical. We have observed the evolution of the microstructure as a function of the film thickness. The grain size and the grain joints surface of the (1 1 0)-oriented films are large. In comparison, the grain size of the (111)-oriented film is very fine and their distribution is more homogeneous. This increase could be explained by the existence of stresses and the possible grain disorientation with the film thickness [11,12]. XRD Patterns for thin films of PZT with different Zr/ Ti ratio are also analyzed. All the films exhibit a perovskite polycrystalline structure. When xp0.30, the diffractograms display split peaks for the (1 0 1/1 1 0) doublet, related to the tetragonal phase. When x ¼ 35 45 65 and 65; a single peak is observed. This behaviour is in contradiction with result for ceramics (the peaks splitting is observed for xp54 46). The PNZT films are also crystallized in the alone perovskite phase; no pyrochlore (or fluorite for high doping level) phase is detectable. As in PZT films, it is also possible to have (1 1 0)-oriented PNZT. The films are dense and crack-free. We have observed an increase of the grain size for PNZT films [13,14]. This result is in contradiction with the observed reduced grain size with increasing Nb content on bulk PNZT ceramics. These variations of the film properties must have some issues on the electrical properties of the films. We have systematically evaluated these contributions. The experimental set-up and conditions are described elsewhere [10]. The coefficients e31 and d 33 were determined by the embedded beam method [15] and the double-beam laser interferometry [16] respectively.
3. Results and discussion 3.1. Influence of film thickness, orientation and (Zr/Ti) ratio The er evolution (Fig. 1) can be decomposed in two parts: for films thickness thinner than 0.6 mm, er increases linearly. For thicker films er attains a saturation value, for a threshold thickness (eth ), of 920 independently of the film orientation. In our case eth is equal to 0.5–0.6 mm whatever the film orientation. The er variations with the film thickness can be attributed to different factors: the existence of a material with a low er (non-ferroelectric) at the interface between the film and the Pt bottom electrode, the correlation between the grain size and er ; an increase of the grain size, resulting from an increase of the thickness, induces an increase of er [17], the increase of er with the thickness is related to the decrease of the stresses with the thickness [18]. Fig. 2 shows the evolution of er as a function of the Zr/Ti ratio. The maximum value was achieved near the morphotropic phase; er values varied betwenn 120 and 900. The remnant polarization increases with the film thickness and attains a saturation value for films thicker than 0.6 mm whatever the film orientation. The variations of the Ec with the film thickness are presented Fig. 3. Ec decreases with the film thickness and a saturation value of 25 V/cm is measured for films thicker than 0.7 mm. These results are in agreement with those published [19,20]. For very thin film, the presence of an interfacial layer degrades the ferroelectric properties; when the film thickness increases, this layer have a less effect. Different types of hysteresis loops were obtained (Fig. 4) with various Zr/Ti ratios; these films were 0.75 mm thick. For x=0.35, the hysteresis loop appears to be more squared, and exhibits a higher coercive field than for the MPB. Ec decreased with increasing Zr concentration to a minimum of 30 kV/cm at the 60 40 composition.
Fig. 1. Evolution of er function of the film thickness for a (1 1 0) and a (1 1 1)-oriented PZT films.
D. Remiens et al. / Materials Science in Semiconductor Processing 5 (2003) 123–127
The e31eff.rem. evolution (after poling treatment [21,22]) as a function of the film thickness is given in Fig. 5; it is remnant coefficient since the DC poling voltage is removed when we want to acquire the
125
piezoelectric signal. The variations can be decomposed into two parts: an increase for films thinner than 0.6mm and after e31eff.rem. remains constant in the order of 4.5 C/cm2 The increase is directly related to the domain contribution, which increases with the film thickness. The e31eff.rem. saturation value detected for films thicker than 0.6mm was not observed in the d33eff. variation. An example is given Fig. 6; d 33eff.rem. is in the order of 45 and 80 pm/V for PZT films of 1 and 1.7 mm thick, respectively. Fig. 7 shows the evolution of e31eff.rem. function of x: The maximum value of e31 is obtained at the morphotropic phase boundary. 3.2. Influence of Nb doping
Fig. 2. Evolution of er function of the Zr/Ti ratio. Comparison between bulk and film.
Fig. 3. Evolution of Ec function of the film thickness for a (1 1 0) and a (1 1 1) PZT films.
For low doping concentration (o2 at%), er increases with the percentage of doping (Fig. 8); it reaches a maximum value of 1100 for 2 at% doping. The increase of er could be explained by the grain growth of the PNZT films and the decrease by the variation of the lattice parameter. For small grains, the density of
Fig. 5. Evolution of the remnant effective e31 coefficient (e31eff.rem.) function of the film thickness.
Fig. 4. Typical example of PðEÞ hysteresis loops of some PZT films of different Zr/Ti ratio.
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Fig. 6. Typical exemple of d 33 hysteresis loops of PZT films with different thickness.
Fig. 9. Fatigue characteristics of PZT and PNZT (2 at%) films.
Fig. 7. Evolution of e31eff.rem. function of the Zr/Ti ratio.
Fig. 10. Evolution of e31eff.rem. function as a of the Nb concentration.
Fig. 8. Evolution of er and the dielectric loss, tan d; with the niobium concentration.
domains increases and so their mobility is reduced. The change of the domain contribution with the introduction of donor doping in PZT is induced by a reduction of the oxygen vacancies density. An electrical characteristic, which is very sensitive to the oxygen vacancies is the ferroelectric fatigue properties. We have tested the fatigue resistance of PZT and PNZT films (Fig. 9). Whatever the doping level, the fatigue properties were
more improved for PNZT rather than for PZT; the better performance is obtained for a 2 at% doped PZT films; the loss of polarizations appears only after 108 cycles (instead of 106 cycles for PZT films). This improvement can be attributed to the decrease of the oxygen vacancies. The maximum polarizations (PM ) are larger for PNZT than for PZT; PM reaches its maximum value (40 mC/cm2) for a PNZT doped at 2 at%. For large doping level, polarizations slightly decrease, but remain superior to those of PZT films. Ec is found to be independent of the doping level and it is lightly weaker than PZT. Fig. 10 shows the evolution of e31eff.rem. (measured after poling treatment) as a function of the Nb concentration. The introduction of Nb enhances the PZT film properties, probably by an increase of the domain wall mobility, as in the bulk soft PZT ceramics. The maximum value of e31eff.rem. (4.6 C/cm2) has been obtained on PNZT films with Nb level of 1 at%. Very little variation was observed between 1 at% and 2 at%. Fig. 11 shows the evolution of d 33eff. remnant as a function of Nb. d 33eff.rem. is maximal for a 2 at% doped PZT film; it is in the order of 75 pm/V. It is clear that the
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to a modification of the charges: oxygen and lead vacancies.
References
Fig. 11. Evolution of d 33eff.rem. as a function of Nb concentration.
introduction of a few atomic% of Nb in PZT induces an improvement of the d 33 coefficient.
4. Conclusion The main electrical properties variation can be summarized as follow: er increase with the film thickness whatever the film orientation and a saturation value is attained for film thicker than 0.6 mm, the ferroelectric properties is independent of the orientation and a dependence with the film thickness is observed—e31 and d 33 increase with the film thickness and a saturation is observed only for e31. The dielectric constant exhibited a clear dependence on composition with values ranging from 120 to 900. Ec decreased with increasing Zr content to a minimum of 38 kV/cm at 60 : The electrical 40 properties dependence with the thickness can be explained by the existence of an interfacilal layer (a ‘‘dead’’ layer), the domain structure, the pinning centers and the microstructure. The effects of Nb doping in PZT 54 46 films are: er increases for small Nb concentration, PM ; and Pr are improved, d 33eff. and e31eff. increase with low Nb content. The enhancement of the piezoelectric PNZT films properties, as in the bulk material, can be attributed to a variation of the domain wall mobility due
[1] Paz de Araujo CA, McMillan LD, Melnick BD, Cuchiaro JD, Scott JF. Ferroelectrics 1990;104:241. [2] Watanabe S, Suzuki M, Fujiu T. J Vac Sci Technol 1995;B13:1119. [3] Muralt P, Kohli M, Maeder T, Kholkin A, Brooks K, Luthier R, Setter N. Sensors Actuators 1995;A48:157. [4] Jaffe B, Cook WR, Jaffe H. Piezoelectric ceramics. London: Academic, 1971. [5] Barlingay CK, Dey SK. Thin Solid Films 1996;272:112. [6] Klissurka RD, Tagantsev AK, Brooks KG, Setter N. J Am Ceram Soc 1997;80:336. [7] Lian L, Sottos NR. J Appl Phys 2000;87(8):3941–9. [8] Aoki K, Fukuda Y, Numata K, Nishima A, Nishimura A. Jpn J Appl Phys Part 1 1994;33(9B):5155. [9] Nam H-J, Kim H-H, Lee W-J. Jpn J Appl Phys 1998;37(6A):3462–70. [10] Velu G, Remiens D. J Europ Ceram Soc 1999;19:2005. [11] Jaber B, Remiens D,, Thierry B. J Appl Phys 1996;79:1182. [12] Kim JC, Yoon DS, Lee JS, Choi CG, No K. J Mater Res 1997;12(4):1043–7. [13] Taylor DV, Damjanovic D. Appl Phys Lett 2000;76(12): 1615–7. [14] Tuttle BA, Doughty DH, Schwartz RW, Garino TJ, Martinez SL, Goodnow D, Fernandez CL, Tissot RG, Hammeter WF. Ceram Trans 1992;15:179. [15] Griswold EM, Sayer M, Amm T. Can J Phys 1991;69:260. [16] Deshanvre JL, Rey P, Delabouglise G, Labeau M. Sensors Actuators 1992;A33:43–5. [17] Kholkin AL, Wutchrich C, Taylor DV, Setter N. Rev Sci Instrum 1996;65(5):1935. [18] Fujisawa F, Nakashima S, Kaibara K, Shimuzu M, Niu H. Jpn J Appl Phys 1999;38(9B):5392. [19] Okino H, Nishikawa T, Shimuzu M, Horiuchi T, Matsushige K. Jpn J Appl Phys 1999;38:5388. [20] Xu F, Trolier McKinstry S, Ren W, Xu B, Xie ZL. J Appl Phys 2000;89(2):1336. [21] Chen SY, Chen IW. J Am Ceram Soc 1997;36(7A):4451–8. [22] Cattan EE, Haccart T, Remiens D. J Appl Phys 1999; 86(12):7017.
Piezoelectric properties of sputtered PZT films: influence of structure, micro structure, film thickness (Zr,Ti) ratio and Nb substitution D. Remiens*, E. Cattan, C. Soyer, T. Haccart IEMN–DOAE–MIMM-UMR, CNRS 8520, Universit!e de Valenciennes, ZI petite savatte, 59600 Maubeuge, France
Abstract Lead titanate zirconate Pb(Zr,Ti)O3 (PZT) with different (Zr,Ti) ratio and Niobium (Nb) doped PZT thin films were deposited on silicon substrates by R.F. magnetron sputtering. The Zr/Ti ratio varied between the quadratic to the rhombohedric phases. For the thickness influence study, we have chosen PZT films with the 54 46 composition. The Nb concentration incorporated in PZT films varied between 1 at% and 7 at%. The electrical properties (dielectric, ferroelectric and piezoelectric) were evaluated functions of the film thickness, the orientation, the (Zr,Ti) ratio and the Nb content. r 2002 Published by Elsevier Science Ltd. Keywords: Piezoelctric thin films; Sputtering; PZT materials; Electrical properties
1. Introduction Lead zirconate titanate (PZT) thin films have been extensively investigated for many applications [1–3]. An important parameter to develop MEMS applications, is the piezoelectric performance of the films. The piezoelectric properties of PZT ceramics can be modified by the substitution of dopants [4]; lanthanum and niobium ions have been currently used. In thin film form, only very few papers describe the Nb influence on the PZT structural and electrical properties. Usually papers focus on the dielectric and ferroelectric properties [5,6]. For MEMS device applications, it is also imperative to perfectly know the influence of the film characteristics such as: the microstructure, the structure and the interfaces, the film thickness, etc. on the electrical properties. Many papers have been pub were in many cases in contradiction [7,8]. Concerning the evaluation of the e31, d 31 and d 33 coefficients dependence with the *Corresponding author. E-mail address: [email protected] (D. Remiens).
PZT film thickness there is a limited number of papers [9]. The growth and characterization of PZT with different (Zr,Ti) ratio, orientation and thickness and PNZT were presented in this paper.
2. PZT and PNZT films preparation and characterization An R.F. magnetron sputtering system, described previously [10], was used to prepare the films. We have used single oxide targets since they are easy to fabricate and facilitate compositional change. Nb5+ was considered as a donor dopant for PZT materials; it substitutes Ti4+/Zr4+ ions. So, since the niobium incorporation occurs on the B-site; the compensation of the additional positive charge introduced, to achieve electroneutrality, is assumed through A-site, i.e. Pb-site. The films were deposited onto oxidized Si (100) n-type substrates ( ( electrodes. The coated with Ti (100 A)/Pt (1500 A) 15 PZT (with (Zr/Ti) ranging from 85 to 70 30) and PNZT sputtering conditions are summarized in Table 1. An annealing treatment is necessary to crystallize the films
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D. Remiens et al. / Materials Science in Semiconductor Processing 5 (2003) 123–127
Table 1 Typical sputtering and annealing conditions of PZT and PNZT films Target diameter (mm) R.F. power density (W/cm2) Gas pressure (mT) Inter electrodes distance (mm) Substrate temperature (1C) Annealing temperature (1C) Annealing time (mm) Atmosphere
76 3 30 (Ar) 60 150 (ion bombardment) 625 30 Air
in the perovskite phase. We present here only the results relative to the conventional treatment. The annealing conditions were given in Table 1. The films are crystallized in the pure polycristalline perovskite phase but two types of preferred orientation have been obtained: the (1 1 0) or the (1 1 1) orientation. At present, it is difficult to explain why the orientation is different; the sputtering and annealing conditions are identical. We have observed the evolution of the microstructure as a function of the film thickness. The grain size and the grain joints surface of the (1 1 0)-oriented films are large. In comparison, the grain size of the (111)-oriented film is very fine and their distribution is more homogeneous. This increase could be explained by the existence of stresses and the possible grain disorientation with the film thickness [11,12]. XRD Patterns for thin films of PZT with different Zr/ Ti ratio are also analyzed. All the films exhibit a perovskite polycrystalline structure. When xp0.30, the diffractograms display split peaks for the (1 0 1/1 1 0) doublet, related to the tetragonal phase. When x ¼ 35 45 65 and 65; a single peak is observed. This behaviour is in contradiction with result for ceramics (the peaks splitting is observed for xp54 46). The PNZT films are also crystallized in the alone perovskite phase; no pyrochlore (or fluorite for high doping level) phase is detectable. As in PZT films, it is also possible to have (1 1 0)-oriented PNZT. The films are dense and crack-free. We have observed an increase of the grain size for PNZT films [13,14]. This result is in contradiction with the observed reduced grain size with increasing Nb content on bulk PNZT ceramics. These variations of the film properties must have some issues on the electrical properties of the films. We have systematically evaluated these contributions. The experimental set-up and conditions are described elsewhere [10]. The coefficients e31 and d 33 were determined by the embedded beam method [15] and the double-beam laser interferometry [16] respectively.
3. Results and discussion 3.1. Influence of film thickness, orientation and (Zr/Ti) ratio The er evolution (Fig. 1) can be decomposed in two parts: for films thickness thinner than 0.6 mm, er increases linearly. For thicker films er attains a saturation value, for a threshold thickness (eth ), of 920 independently of the film orientation. In our case eth is equal to 0.5–0.6 mm whatever the film orientation. The er variations with the film thickness can be attributed to different factors: the existence of a material with a low er (non-ferroelectric) at the interface between the film and the Pt bottom electrode, the correlation between the grain size and er ; an increase of the grain size, resulting from an increase of the thickness, induces an increase of er [17], the increase of er with the thickness is related to the decrease of the stresses with the thickness [18]. Fig. 2 shows the evolution of er as a function of the Zr/Ti ratio. The maximum value was achieved near the morphotropic phase; er values varied betwenn 120 and 900. The remnant polarization increases with the film thickness and attains a saturation value for films thicker than 0.6 mm whatever the film orientation. The variations of the Ec with the film thickness are presented Fig. 3. Ec decreases with the film thickness and a saturation value of 25 V/cm is measured for films thicker than 0.7 mm. These results are in agreement with those published [19,20]. For very thin film, the presence of an interfacial layer degrades the ferroelectric properties; when the film thickness increases, this layer have a less effect. Different types of hysteresis loops were obtained (Fig. 4) with various Zr/Ti ratios; these films were 0.75 mm thick. For x=0.35, the hysteresis loop appears to be more squared, and exhibits a higher coercive field than for the MPB. Ec decreased with increasing Zr concentration to a minimum of 30 kV/cm at the 60 40 composition.
Fig. 1. Evolution of er function of the film thickness for a (1 1 0) and a (1 1 1)-oriented PZT films.
D. Remiens et al. / Materials Science in Semiconductor Processing 5 (2003) 123–127
The e31eff.rem. evolution (after poling treatment [21,22]) as a function of the film thickness is given in Fig. 5; it is remnant coefficient since the DC poling voltage is removed when we want to acquire the
125
piezoelectric signal. The variations can be decomposed into two parts: an increase for films thinner than 0.6mm and after e31eff.rem. remains constant in the order of 4.5 C/cm2 The increase is directly related to the domain contribution, which increases with the film thickness. The e31eff.rem. saturation value detected for films thicker than 0.6mm was not observed in the d33eff. variation. An example is given Fig. 6; d 33eff.rem. is in the order of 45 and 80 pm/V for PZT films of 1 and 1.7 mm thick, respectively. Fig. 7 shows the evolution of e31eff.rem. function of x: The maximum value of e31 is obtained at the morphotropic phase boundary. 3.2. Influence of Nb doping
Fig. 2. Evolution of er function of the Zr/Ti ratio. Comparison between bulk and film.
Fig. 3. Evolution of Ec function of the film thickness for a (1 1 0) and a (1 1 1) PZT films.
For low doping concentration (o2 at%), er increases with the percentage of doping (Fig. 8); it reaches a maximum value of 1100 for 2 at% doping. The increase of er could be explained by the grain growth of the PNZT films and the decrease by the variation of the lattice parameter. For small grains, the density of
Fig. 5. Evolution of the remnant effective e31 coefficient (e31eff.rem.) function of the film thickness.
Fig. 4. Typical example of PðEÞ hysteresis loops of some PZT films of different Zr/Ti ratio.
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Fig. 6. Typical exemple of d 33 hysteresis loops of PZT films with different thickness.
Fig. 9. Fatigue characteristics of PZT and PNZT (2 at%) films.
Fig. 7. Evolution of e31eff.rem. function of the Zr/Ti ratio.
Fig. 10. Evolution of e31eff.rem. function as a of the Nb concentration.
Fig. 8. Evolution of er and the dielectric loss, tan d; with the niobium concentration.
domains increases and so their mobility is reduced. The change of the domain contribution with the introduction of donor doping in PZT is induced by a reduction of the oxygen vacancies density. An electrical characteristic, which is very sensitive to the oxygen vacancies is the ferroelectric fatigue properties. We have tested the fatigue resistance of PZT and PNZT films (Fig. 9). Whatever the doping level, the fatigue properties were
more improved for PNZT rather than for PZT; the better performance is obtained for a 2 at% doped PZT films; the loss of polarizations appears only after 108 cycles (instead of 106 cycles for PZT films). This improvement can be attributed to the decrease of the oxygen vacancies. The maximum polarizations (PM ) are larger for PNZT than for PZT; PM reaches its maximum value (40 mC/cm2) for a PNZT doped at 2 at%. For large doping level, polarizations slightly decrease, but remain superior to those of PZT films. Ec is found to be independent of the doping level and it is lightly weaker than PZT. Fig. 10 shows the evolution of e31eff.rem. (measured after poling treatment) as a function of the Nb concentration. The introduction of Nb enhances the PZT film properties, probably by an increase of the domain wall mobility, as in the bulk soft PZT ceramics. The maximum value of e31eff.rem. (4.6 C/cm2) has been obtained on PNZT films with Nb level of 1 at%. Very little variation was observed between 1 at% and 2 at%. Fig. 11 shows the evolution of d 33eff. remnant as a function of Nb. d 33eff.rem. is maximal for a 2 at% doped PZT film; it is in the order of 75 pm/V. It is clear that the
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to a modification of the charges: oxygen and lead vacancies.
References
Fig. 11. Evolution of d 33eff.rem. as a function of Nb concentration.
introduction of a few atomic% of Nb in PZT induces an improvement of the d 33 coefficient.
4. Conclusion The main electrical properties variation can be summarized as follow: er increase with the film thickness whatever the film orientation and a saturation value is attained for film thicker than 0.6 mm, the ferroelectric properties is independent of the orientation and a dependence with the film thickness is observed—e31 and d 33 increase with the film thickness and a saturation is observed only for e31. The dielectric constant exhibited a clear dependence on composition with values ranging from 120 to 900. Ec decreased with increasing Zr content to a minimum of 38 kV/cm at 60 : The electrical 40 properties dependence with the thickness can be explained by the existence of an interfacilal layer (a ‘‘dead’’ layer), the domain structure, the pinning centers and the microstructure. The effects of Nb doping in PZT 54 46 films are: er increases for small Nb concentration, PM ; and Pr are improved, d 33eff. and e31eff. increase with low Nb content. The enhancement of the piezoelectric PNZT films properties, as in the bulk material, can be attributed to a variation of the domain wall mobility due
[1] Paz de Araujo CA, McMillan LD, Melnick BD, Cuchiaro JD, Scott JF. Ferroelectrics 1990;104:241. [2] Watanabe S, Suzuki M, Fujiu T. J Vac Sci Technol 1995;B13:1119. [3] Muralt P, Kohli M, Maeder T, Kholkin A, Brooks K, Luthier R, Setter N. Sensors Actuators 1995;A48:157. [4] Jaffe B, Cook WR, Jaffe H. Piezoelectric ceramics. London: Academic, 1971. [5] Barlingay CK, Dey SK. Thin Solid Films 1996;272:112. [6] Klissurka RD, Tagantsev AK, Brooks KG, Setter N. J Am Ceram Soc 1997;80:336. [7] Lian L, Sottos NR. J Appl Phys 2000;87(8):3941–9. [8] Aoki K, Fukuda Y, Numata K, Nishima A, Nishimura A. Jpn J Appl Phys Part 1 1994;33(9B):5155. [9] Nam H-J, Kim H-H, Lee W-J. Jpn J Appl Phys 1998;37(6A):3462–70. [10] Velu G, Remiens D. J Europ Ceram Soc 1999;19:2005. [11] Jaber B, Remiens D,, Thierry B. J Appl Phys 1996;79:1182. [12] Kim JC, Yoon DS, Lee JS, Choi CG, No K. J Mater Res 1997;12(4):1043–7. [13] Taylor DV, Damjanovic D. Appl Phys Lett 2000;76(12): 1615–7. [14] Tuttle BA, Doughty DH, Schwartz RW, Garino TJ, Martinez SL, Goodnow D, Fernandez CL, Tissot RG, Hammeter WF. Ceram Trans 1992;15:179. [15] Griswold EM, Sayer M, Amm T. Can J Phys 1991;69:260. [16] Deshanvre JL, Rey P, Delabouglise G, Labeau M. Sensors Actuators 1992;A33:43–5. [17] Kholkin AL, Wutchrich C, Taylor DV, Setter N. Rev Sci Instrum 1996;65(5):1935. [18] Fujisawa F, Nakashima S, Kaibara K, Shimuzu M, Niu H. Jpn J Appl Phys 1999;38(9B):5392. [19] Okino H, Nishikawa T, Shimuzu M, Horiuchi T, Matsushige K. Jpn J Appl Phys 1999;38:5388. [20] Xu F, Trolier McKinstry S, Ren W, Xu B, Xie ZL. J Appl Phys 2000;89(2):1336. [21] Chen SY, Chen IW. J Am Ceram Soc 1997;36(7A):4451–8. [22] Cattan EE, Haccart T, Remiens D. J Appl Phys 1999; 86(12):7017.