Smart Materials ResearchVolume 2012 (2012), Article ID 426048, 7 pagesdoi:10.1155/2012/426048

Research Article

Epitaxial Piezoelectric Pb() Thin Films on Silicon for Energy Harvesting Devices

A. Sambri,1,2 D. Isarakorn,³ A. Torres-Pardo,⁴ S. Gariglio,¹ Pattanaphong Janphuang,³ D. Briand,³ O. Stéphan,⁴ J. W. Reiner,⁵ J.-M. Triscone,¹ Nico F. de Rooij,³ and C. H. Ahn⁵

¹Department of Condensed Matter Physics, (DPMC), University of Geneva, 24 Quai Ernest-Ansermet, 1211 Geneva 4, Switzerland²Dipartimento di Scienze Fisiche & CNR-SPIN, Università degli Studi di Napoli Federico II, Complesso Universitario di Monte S. Angelo, Via Cintia, 80126 Napoli, Italy³The Sensors, Actuators and Microsystems Laboratory, Institute of Microengineering (IMT), Ecole Polytechnique Fédérale de Lausanne (EPFL), Rue Jaquet-Droz 1, P.O. Box 526, 2002 Neuchâtel, Switzerland⁴Laboratoire de Physique des Solides, Université Paris-Sud, CNRS-UMR 8502, 91405 Orsay, France⁵Department of Applied Physics, Yale University, P.O. Box 208284, New Haven, CT 06520-8284, USA

Received 20 December 2011; Accepted 27 January 2012

Academic Editor: Mohammed Es-Souni

Copyright © 2012 A. Sambri et al. This is an open access article distributed under the Creative Commons Attribution License, which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

Abstract.

We report on the properties of ferroelectric Pb(Zr0.2Ti0.8)O₃ (PZT) thin films grown epitaxially on (001) silicon and on the performance of such heterostructures for microfabricated piezoelectric energy harvesters. In the first part of the paper, we investigate the epitaxial stacks through transmission electron microscopy and piezoelectric force microscopy studies to characterize in detail their crystalline structure. In the second part of the paper, we present the electrical characteristics of piezoelectric cantilevers based on these epitaxial PZT films. The performance of such cantilevers as vibration energy transducers is compared with other piezoelectric harvesters and indicates the potential of the epitaxial approach in the field of energy harvesting devices.

1. Introduction

In the last decade, studies on epitaxial ferroelectric thin films have led to many interesting results and exciting discoveries [1]. The possibility of tailoring or even enhancing some physical parameters via epitaxial strain engineering [2–4] has suggested the idea of the exploitation of such thin films for several technological applications [5]. Nevertheless, the benefits of the epitaxial approach on the performances of ferroelectric thin film-based devices have to compensate the hurdles related to the epitaxial growth on industrial substrates such as silicon, the modern technological platform. It is in fact well known that the basic requirements for the epitaxy, that is, a good lattice match between substrate and film and a reciprocal chemical stability, are not easily fulfilled in the case of oxide growth on silicon. Beside the difference in lattice parameters and thermal expansion coefficients, the main problem is the surface reactivity of silicon to oxygen, with the formation of an amorphous layer of silicon dioxide that hints any further epitaxy. Moreover, the cations of most ferroelectric compounds interdiffuse into the silicon substrate, forming spurious extra phases at the interface [6]. In order to overcome such difficulties, a suitable buffer layer is needed that acts as a barrier for cations migration and as a structural template for the growth of the ferroelectric epitaxial film [7–9].

PZT is one of the most investigated ferroelectric materials, due to its high values of remnant polarization and piezoelectric coefficients. In the bulk form, it displays a complex phase diagram versus the Ti/Zr content: for the stoichiometry of our choice, that is, Pb()O₃ (PZT 20/80), it is ferroelectric with a tetragonal structure up to a transition temperature of 460°C [10]. It is worth stressing that for ferroelectric materials the mechanical boundary conditions may affect substantially the ferroelectric properties because of the strong strain-polarization coupling present in these compounds [11–13]. Thus, the epitaxial strain often makes the properties of ferroelectric thin films substantially different from the bulk counterpart.

2. Heterostructure Growth and Characterization

The investigated epitaxial heterostructure is composed of a ferroelectric PZT layer grown on top of a metallic SrRuO₃ (SRO) film, used as bottom electrode. This bilayer is deposited onto a thin SrTiO₃ (STO) film epitaxially grown on a 2-inch (001) silicon wafer. The choice of STO as buffer layer is determined by the fact that it has the same crystalline structure (perovskite) and a good lattice match with SRO and PZT. STO also acts as a barrier for Pb diffusion into the silicon wafer [14–16]. The STO layer can reach a very high degree of crystalline perfection on silicon when the growth is performed by molecular beam epitaxy through a complex multistep process, where fixed amounts of the elemental materials have to be carefully deposited at specific conditions of substrate temperature and pressure. The deposition process we used is detailed in [17, 18]. In situ Reflection High Energy Electron Diffraction (RHEED) is the key tool in order to monitor and control the whole deposition procedure, since every step of the deposition corresponds to a specific diffraction pattern linked to a particular surface reconstruction. Figure 1 shows the evolution of the RHEED pattern during the growth of a STO film on silicon.

(a)

(b)

(c)

(d)

(e)

(f)

Figure 1: Evolution of RHEED pattern during the growth of a STO layer. The patterns shown are taken with the electron beam along the direction. (a) As-received Si wafer (001), the diffuse background indicates the presence of a native amorphous layer of SiO₂; (b) crystalline Si surface after a thermal treatment in vacuum: the appearance of fractional spots (0,1/2) due to Si dimers indicates the complete removal of SiO₂, (c) 2× surface reconstruction corresponding to the deposition of half monolayer (ML) of Sr in vacuum at high temperature, (d) 3× reconstruction corresponding to the deposition of half ML of Sr in vacuum at low temperature, (e) 1 ML of SrO, (f) 3 ML of crystalline STO.

The SRO and PZT layers are subsequently grown by rf magnetron sputtering as described in [19]. In the following, a detailed investigation of the crystalline structure of a PZT(100 nm)/SRO(30 nm)/STO(6 nm)/Si stack is given. X-ray structural characterization of the heterostructure has been performed with a high-resolution PANalytical X’Pert diffractometer, equipped with a four-bounce asymmetric Ge(220) monochromator for Cu radiation. Through -2 and -scan diffractograms, the following out-of-plane orientation and cube-on-cube arrangement of the oxides stack on silicon have been determined: PZT[001]//SRO[001]//STO[001]//Si[001] and PZT[100]//SRO[100]//STO[100]//Si[110] (see also [13]).

Transmission electron microscopy (TEM) characterization, performed on a JEOL JEM2010 microscope, shows at the silicon-STO interface the presence of an amorphous layer of SiO₂, which clearly did not affect the epitaxy of the oxides multilayer (Figure 2).

Figure 2: Cross-sectional TEM image, revealing the presence of 4.5 nm of amorphous SiO₂ located between the silicon substrate and the epitaxial oxide stack.

Most likely then, the formation of this SiO₂ layer occurs after the growth of the STO/Si interface template through oxygen diffusion, thus not hindering the epitaxy [20–22].

Interestingly, the TEM investigation also revealed the presence in the PZT layer of crystallographic domains with the -axis aligned in the plane of the film, so-called -domains. Figure 3 shows that their width is about 20 nm. We note that the overlap of the more intense SRO reflections with the diffraction peaks of the PZT a-domains most probably prevents their detection in the X-ray -2 scans. The origin of the formation of these domains can be related to the strain arising from the lattice and thermal mismatch of the film with the substrate during the cool-down across the ferroelectric transition [23, 24]. Figure 3(a) shows a pattern of domains in the PZT layer: they repeat periodically every ~105 nm. The domain wall between and domains runs along the diagonal of the tetragonal unit cell and the difference between the and axis ( = 1.046) [13] gives rise to a distortion of the neighboring lattices on each side of a 90° domain wall. The misalignment of the growth planes between a and c domains is clearly visible in Figure 3(c) and can be estimated to be ~3°, in good agreement with the calculated value of ~2.6°.

(a)

(b)

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Figure 3: (a) Cross-sectional TEM image, revealing the presence of periodic -domains, (b) HRTEM of a single -domain surrounded by a -axis oriented regions. (c) Details of the 90° domain wall.

Due to the different polarization orientations of the and domains, respectively, perpendicular and parallel to the film-substrate interface, we can detect their presence by piezoelectric force microscopy (PFM). In PFM, a conducting tip is scanned in contact with the surface while an ac voltage, lower than the coercive field of the material, is applied between the tip and the sample. From the detection of the in-phase and out-of-phase deformation signals in the vertical (V) and lateral (L) PFM modes [25], it is possible to map the presence of and domains.

Figure 4(a) shows a topography scan of a PZT surface, revealing a nRMS roughness of 5 nm. The V-PFM scan shown in the middle panel is sensitive to the polarization direction of the c-domains of the film, here shown as reddish area, while the yellowish lines correspond to a-domains, with zero contrast. Conversely, through L-PFM it is possible to detect the piezoelectric response of -domains, with the restriction that domains parallel to the scan direction cannot be detected, as clearly shown in Figure 4(c). Combining the V-PFM and L-PFM measurements, the PZT film appears to be a very dense network of small a-domains embedded in a matrix of c-axis domains. As can be seen, the phase response detected in the V-PFM and L-PFM is unrelated to the topographic features detected simultaneously by the AFM scan. Figure 5 shows a detail of the PZT -domains network, on a 1 × 1 μm² area.

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(b)

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Figure 4: (a) Topography scan on a 20 × 20 μm² area, (b) V-PFM and (c) L-PFM: phase response of the vertical and lateral PFM measurements on the same area and schematic of the corresponding detected domains.

Figure 5: V-PFM scan on a 1 × 1 μm² area: zoom of the -domains network.

According to some reports, applying an electric field induces a lateral motion of thin a-domains [26–33]. Even though some of these results are still under debate, such a motion would give rise to an extrinsic contribution to the piezoelectric response of the material and an enhancement of the conversion coefficients. -domains switching would indirectly enhance the strain state in the film and its piezoelectric response. Indeed, while the estimation of the PZT coefficient by PFM measurement yields a value of about 50 pm/V (in agreement with the data reported for bulk PZT 20/80), the indirect estimation of from the deflection of a microfabricated piezoelectric cantilever gives a higher value (details in the next section).

We investigated the ferroelectric behavior of these epitaxial films grown on silicon through a series of polarization hysteresis measurements, performed by a TF analyzer 2000. To perform the measurements, Cr/Au top electrodes were prepared with a size of 100 × 100 μm². The polarization versus voltage loop, shown in Figure 6(a), reveals a remnant polarization of about 70 μC cm⁻² and no significant leakage current up to 16 V. Figure 6(b) displays the measurement of the capacitance versus applied bias, revealing the dielectric tunability of our epitaxial PZT. In the approximation of a plane capacitor, we estimate a dielectric constant at zero field to be about 150.

(a)

(b)

Figure 6: (a) Polarization versus applied voltage (PE loop) (full red circle) and corresponding switching current (empty blue circle). (b) Plot of the capacitance versus applied bias. The difference in the coercive field in the PE loop measurement and in the capacitance cycle is due to a different electrical configuration, top-top contacts for PE loos while top-bottom for the capacitance.

3. Energy Harvesting Devices

The development of transducers allowing energy from mechanical vibrations to be generated has advanced rapidly during the past few years. Several review articles [34, 35] discuss the principles and advantages of each conversion method. Among these, piezoelectricity based devices have received large attention due to their high power density and ease of integration compared to other transduction techniques [36–38]. Several piezoelectric silicon micromachined energy scavengers have been proposed so far, most of these devices being based on polycrystalline PZT (poly-PZT) or aluminum nitride (AlN) films. Energy harvesting devices involving epitaxial PZT thin films are however less common [39–41]. In this study we report on the characteristics of MEMS based vibrations energy scavengers made of an epitaxial PZT film on silicon cantilevers.

A typical device is shown in Figure 7: it is based on a cantilever fabricated using micropatterning techniques optimized for the oxide layers deposited on silicon. The details of the microfabrication process are presented in [19].

Figure 7: Optical image of an epi-PZT cantilever (1000 × 2500 × 7 μm³) with a Si proof mass (1000 × 500 × 230 μm³). The inset shows the Si proof mass on the back side of the cantilever. The effective volume of the final device is 0.1325 mm³.

The piezoelectric coefficient has been estimated from the displacement at the free end of a cantilever. The deflection of this cantilever, coated with 100 nm thick epi-PZT films, has been measured as a function of dc voltage using an interferometric profiler (Wyko NT110). The piezoelectric coefficient value is estimated to be 135.4 ± 7.1 pm/V [19]. For a PZT thickness of about 500 nm, we measure a residual stress after the film growth of about 60 MPa, according to stress measurement performed using a Tencor FLX-2320A system. In order to evaluate the performance of our device, the cantilever’s resonant frequency has been determined using a Polytec MVS-400 laser Doppler vibrometer, obtaining a value of 2.3 ± 0.1 kHz. The power generation performance of the device has been investigated with a shaker (Bruel and Kjaer type 4811) driven by a vibration exciter control (type 1050) and a power amplifier (type 2712) by applying an acceleration as a mechanical input. The device is connected with various resistive loads () and the current generated under different acceleration levels is recorded with a multimeter (Agilent 34411A). The corresponding average power is calculated by the relation: . The output current, average power, and voltage as a function of resistive loads are shown in Figure 8.

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Figure 8: (a) Average power, (b) output current, and (c) corresponding output voltage of an epi-PZT harvester versus resistive load.

In order to compare this performance with other MEMS energy harvesters, these output data are normalized as following: the output power per square acceleration (g²), the current and voltage per acceleration (g). From a single device with an optimal resistive load of 4.7 kΩ, a maximum output power of 14 μW/g² with 60 μA/g output current is obtained, thus resulting in an output voltage of 0.28 V/g. The maximum power density is as high as 105 μW/(g² mm³). As shown in details in [41], the comparison of the electrical output characteristics between this epitaxial-PZT energy harvester and other piezo-harvesting devices based on poly-PZT and AlN films [42–46] shows that the former provides higher current at smaller resistive load. Usually, energy harvesting devices with a low optimal load resistance are highly desirable because they can generate high output current, and also because their impedance can be easily matched to standard electronic devices. Energy harvesting devices based on poly-PZT and AlN films demonstrated useful power generation, but their high impedance limits the output current. Moreover, even though the power generated by the epi-PZT is similar to that of AlN, the epi-PZT harvester exhibits the highest power density, which is of high interest when realizing miniaturized devices. Unlike the polycrystalline film based devices, the device reported here can generate high power and current with usable voltage, while maintaining low optimal resistive load. For piezoelectric energy harvesters, the expressions for the figure of merit for the power, voltage and current are, respectively: , , , where is the effective piezoelectric coefficient. From the measured piezoelectric coefficient , it is possible to estimate a value of 18.2 ± 0.9 C m⁻² [47], a value which is significantly higher than what has been reported in the past [48]. Such large piezoelectric coefficient together with a low dielectric constant is among the key parameters to realize high performance piezoelectric energy harvesters.

4. Conclusions

This paper reports on the structural and physical properties of epitaxial PZT thin films deposited on silicon and on the electrical performance of cantilevers microfabricated from such heterostructures. We have shown that the PZT thin films exhibit excellent ferroelectric characteristics with a remnant polarization of 70 μC cm⁻². The structural microscopic investigation by TEM and PFM has revealed the presence of -domains embedded in a mostly oriented -axis film. This coexistence of domains with different crystallographic orientation could be at the origin of the large piezoelectric response estimated from the electromechanical behavior of the cantilevers. The performances of our epitaxial piezo-transducers fabricated for energy harvesting indicate that these generate higher power and current with usable voltage requiring lower optimal resistive load as compared to piezoelectric harvesters realized with polycrystalline PZT or AlN films.

Acknowledgments

The authors would like to thank M. Lopes for his technical assistance and P. Zubko for many fruitful discussions. This work was supported by the Swiss National Science Foundation through the National Center of Competence in Research, Materials with Novel Electronic Properties, MaNEP and Division II, and by European Union through the project “Nanoxide.” D. Isarakorn and P. Janphuang are grateful to the Ministry of Science and Technology, Thailand, for the financial support for their Ph.D. studies at EPFL IMT-NE SAMLAB. A. Torres-Pardo is grateful to the Spanish Ministry of Education and Science for the support through MEC/Fulbright postdoctoral grant. Work at Yale was supported by NSF DMR 1006256 and MRSEC DMR 1119826.

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