PROCEDURE FOR OBTAINING SAID CERAMIC MATERIAL DESCRIPTION

The invention relates to a high temperature piezoelectric BiSc0 ₃ -PbTi0 ₃ ceramic material of formula Bi _{1 -x+} 2 _y Pb _x- 3 _y Sc _1-x Ti _x 0 ₃ , wherein x ranges from 0.64 to 0.68 and y ranges from 0.01 to 0,025, which includes a point defect engineering for enhanced voltage response. Furthermore, the invention refers to a procedure for obtaining said ceramic material by conventional sintering of nanocrystalline powders synthesized by mechanochemical activation of a stoichiometric mixture of precursors. Finally, the invention also relates to the use of the chemically engineered BiSc0 ₃ -PbTi0 ₃ ceramic material as part of sensing devices and magnetic sensing devices.

STATE OF ART

BiSc0 ₃ -PbTi0 ₃ is the most promising system among general formula BiM0 ₃ -PbTi0 ₃ , wherein M is a trivalent cation in octahedral coordination, perovskite solid solutions with enhanced electromechanical response at ferroelectric morphotropic phase boundaries (MPB), and high Curie temperature. This material is being extensively investigated as an alternative to state of the art Pb(Zr,Ti)0 ₃ (PZT) for expanding the operation temperature of high sensitivity piezoelectric ceramics beyond 200°C up to 400°C. Specifically, the binary system (1 -x)BiSc0 ₃ -xPbTi0 ₃ presents a MPB between ferroelectric polymorphic phases of rhombohedral R3m and tetragonal PAmm symmetry at x~0.64, composition for which the Curie temperature T _c is « 450°C, while piezoelectric coefficients d ₃₃ of -450 pC N ^"1 are typically achieved after poling. This T _c is 100°C above that of Pb(Zr,Ti)0 ₃ , likewise d ₃₃ that also significantly exceeds the figure of « 245 pC N ^"1 for ceramics of the latter material at its own MPB. Moreover, the charge piezoelectric coefficient is comparable to those of available commercial high sensitivity piezoelectric ceramics of chemically engineered PZT.

However, and in spite of the theoretically expanded operation temperature range enabled by the high Curie temperature, BiSc0 ₃ -PbTi0 ₃ cannot be directly used in most applications. This is the case of sensing technologies like accelerometers, vibration monitoring, hydrophones or magnetic field sensors, for which the voltage piezoelectric coefficient g ₃₃ rather than the charge piezoelectric coefficient d ₃₃ is the key parameter. This coefficient is equal to d ₃₃ times the reciprocal permittivity that is very low (or the permittivity very high) in poled BiSc0 ₃ -PbTi0 _3. Therefore, and for the reasons stated above, it is needed to develop new BiSc0 ₃ -PbTi0 ₃ materials, optimized for specific applications.

DESCRIPTION OF THE INVENTION

The present invention discloses a perovskite piezoelectric BiSc0 ₃ -PbTi0 ₃ material of formula Bi _1-x+2y Pb _x-3y Sc _1-x Ti _x 0 ₃ , wherein x ranges from 0.64 to 0.68 and y ranges from 0.01 to 0.025, which exhibits an enhanced electrochemical response at a perovskite morphotropic phase boundary between polymorphs of rhombohedral R3m and tetragonal PAmm symmetries, high Curie temperature, and a point defect engineering for enhanced voltage response.

Furthermore, the high temperature, high sensitivity and enhanced voltage response piezoelectric ceramic of the present invention is a dense, and highly homogenous fine grained microstructure with an average grain size that can be tailored from 1 .0 μιη up to 2.5 μιη. Specifically, ceramic materials with x=0.64 and y=0.01 have a Curie temperature of 395°C, and a g ₃₃ coefficient of 4.8x10 ^"2 VmN ^"1 .

Moreover, the present invention discloses a procedure for obtaining said ceramic material that refers to its preparation by conventional sintering of nanocrystalline powders synthesized by mechanochemical activation of precursors in a high energy planetary mill. This procedure, based on highly reactive powders, allows the suppression of Bi ₂ 0 ₃ and PbO volatilization during the high temperature sintering, so that stoichiometric mixtures of the precursors (Bi ₂ 0 ₃ , Sc ₂ 0 ₃ , PbO, Ti0 ₂ and Mn ₂ 0 ₃ ) can be used, while avoiding the necessity of controlling the atmosphere during the final thermal treatment by burying the green bodies in powder during the sintering. This is very advantageous for an accurate control of composition and phase coexistence while point defects are engineered, which cannot be reproducibly achieved by conventional ceramic technologies such as solid state synthesis by heating of precursors. A first aspect of the present invention relates to a piezoelectric ceramic material characterized in that it has

• the general formula Bi _1-x+ 2yPb _x- 3 _y Sci- _x Ti _x 0 ₃ , wherein x ranges from 0.64 to 0.68 and y ranges from 0.01 to 0.025;

· perovskite single phase placed at a morphotropic phase boundary between polymorphs of rhombohedral R3m and tetragonal PAmm symmetries; and

• microstructure with average grain size of between 1 .0 μιη and 2.5 μιη.

The piezoelectric ceramic material of the present invention has perovskite single phase placed at a morphotropic phase boundary between polymorphs of rhombohedral R3m and tetragonal PAmm symmetries that is the responsible for the high piezoelectric response and, moreover, it includes an engineered point defect for enhanced voltage response that makes the material suitable to be used in sensing technologies. The term "engineered point defect" refers to a point defect that is introduced at or around a single lattice point of the perovskite, concretely by the controlled substitution of Bi ³⁺ for Pb ²⁺ in the A-site (cuboctahedral coordination) of the AB0 ₃ perovskite, along with the formulation of one Pb vacancy per each two substitutions for charge compensation. This controlled substitution that does not require additional chemical species, but the introduction of an A-site non-stoichiometry, results in significant lattice stiffening and thus, a decrease of the dielectric permittivity and elastic compliance. Moreover, this is achieved while the material is maintained at the morphotropic phase boundary between polymorphs of rhombohedral R3m and tetragonal PAmm symmetries, known to be required for high piezoelectric response.

Additionally, microstructure is also controlled during the point defect engineering, so that optimized dense, homogenous fine-grained microstructure is obtained for optimal mechanical properties. In a preferred embodiment, the piezoelectric ceramic material of the present invention has a dense and homogeneous fine grained microstructure with average grain size of between 1 .0 μιη and 1 .5 μιη. In another preferred embodiment of the present invention, the piezoelectric ceramic material mentioned above has the formula Bi _1-x+ 2yPb _x- 3 _y Sci- _x Ti _x 0 ₃ , wherein x=0.64 and y=0.01. A second aspect of the present invention relates to a procedure for obtaining the piezoelectric ceramic material mentioned above, characterized in that it comprises the following steps:

a) synthesis of a nanocrystalline powder of formula Bi _1-x+ 2 _y Pb _x- 3 _y Sc _1-x Ti _x 0 ₃ , wherein x ranges from 0.64 to 0.68 and y ranges from 0.01 to 0.025 by mechanical activation of a stoichiometric mixture of Bi ₂ 0 ₃ , Sc ₂ 0 ₃ , PbO and Ti0 ₂ ; and b) sintering of the nanocrystalline powder obtained in step (a) at a temperature range of between 1 100 °C and 1 150 °C.

Step (a) is preferably performed in a planetary mill at 300 rpm for 20 h.

In a preferred embodiment, step (a) is performed for the synthesis of a nanocrystalline powder of formula Bi _1-x+ 2yPb _x- 3 _y Sci- _x Ti _x 0 ₃ , wherein x is 0.64 and y is 0.01 .

A third aspect of the invention refers to a piezoelectric ceramic composite comprising · the piezoelectric ceramic material according to any of claims 1 to 3; and

• a magnetostrictive material.

The term "magnetostrictive material" refers to a ferromagnetic material that changes its shape or dimensions during the process of magnetization.

In a preferred embodiment, the magnetostrictive material that forms the piezoelectric ceramic composite is selected from the list consisting of Terfenol-D, Metglass, spinel oxides of formula AFe ₂ 0 ₄ , wherein A is Ni, Co or a combination thereof, and a combination thereof.

In another preferred embodiment of the present invention, the piezoelectric ceramic composite mentioned above is of particulate, fiber or laminate type.

Another aspect of the present invention refers to the use of the piezoelectric ceramic material as described above, as part of a sensing device. Examples of sensing devices are accelerometers, vibration monitoring or hydrophones. The piezoelectric ceramic material of the present invention is the active element in these sensing technologies.

The last aspect of the invention refers to the use of the piezoelectric ceramic composite described above as part of a magnetic sensing device such as a magnetic transducer. The piezoelectric ceramic composite of the present invention comprising the piezoelectric ceramic material of the present invention and a magnetostrictive material, is the active element in magnetic sensing technologies. Unless otherwise defined, all technical and scientific terms used herein have the same meaning as commonly understood by one of ordinary skilled in the art to which this invention belongs. Methods and materials similar or equivalent to those described herein can be used in the practice of the present invention. Throughout the description and claims the word "comprise" and its variations are not intended to exclude other technical features, additives, components, or steps. Additional objects, advantages and features of the invention will become apparent to those skilled in the art upon examination of the description or may be learned by practice of the invention. The following examples and drawings are provided by way of illustration and are not intended to be limiting of the present invention.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG.1 XRD patterns for Bi _{1 -x+} 2 _y Pb _x- 3 _y Sci _-x Ti _x 0 ₃ with x= 0.64 and y=0, 0.01 and 0.025 ceramic samples, showing the absence of second phases other than perovskite for low concentration of point defects.

FIG.2 Patterns for Bi _{1 -x+} 2 _y Pb _x- 3 _y Sci _-x Ti _x 0 ₃ with x= 0.64 and y=0, 0.01 and 0.025 ceramic samples with improved statistics across the perovskite parent cubic phase 200 diffraction peak, showing polymorphic phase coexistence and thus, location of materials at the morphotropic phase boundary for low concentration of point defects.

FIG.3 Scanning electron microscopy (SEM) images for Bi _{1 -x+} 2 _y Pb _x- 3 _y Sci- _x Ti _x 0 ₃ with x= 0.64 and y=0, 0.01 and 0.025 ceramic samples, showing the dense and homogenous fine-grained microstructure. FIG.4 Scanning electron microscopy (SEM) images and XRD patterns for Bi _1-x+2y Pb _{x- 3y} Sci- _x Ti _x 0 ₃ with x=0.66 and y=0.01 and x=0.64 and y=0.01 ceramic samples, showing the ability of tailoring phase coexistence, while introducing the point defects. FIG.5 Temperature dependences of the dielectric permittivity for Bi _1-x+ 2yPb _x- 3 _y Sci- _x Ti _x 0 ₃ with a) x=0.64 and y=0, 0.01 and 0.025, b) x=0.64, 0.65 and 0.66 and y=0.01 c) x=0.64, 0.66 and 0.68 and y=0.025 ceramic samples, showing the position of the ferroelectric transition (that determines the maximum operation temperature). FIG.6 Ferroelectric hysteresis loops for Bi _1-x+ 2yPb _x- 3 _y Sci- _x Ti _x 0 ₃ with a) x=0.64 and y=0, 0.01 and 0.025, b) x=0.64, 0.65 and 0.66 and y=0.01 c) x=0.64, 0.66 and 0.68 and y=0.025 ceramic samples, showing a limited increase of the domain wall mobility after introducing the point defects. FIG.7 Magnetoelectric voltage coefficient for Bi _1-x+ 2 _y Pb _x- 3 _y Sc _1-x Ti _x 0 ₃ with x=0.64 and y=0 (unmodified) or y=0.01 (chemically modified), showing the enhanced voltage response after introducing the point defects.

EXAMPLES

Preparation of piezoelectric ceramic samples

Perovskite single phase nanocrystalline powders of formula Bi _1-x+ 2 _y Pb _x- 3 _y Sci- _x Ti _x 0 ₃ , wherein

were synthesized by mechanochemical activation of stoichiometric mixtures of analytical grade Bi ₂ 0 ₃ (Aldrich, 99.9% pure), Sc ₂ 0 ₃ (Aldrich, 99.9% pure), PbO (Merck, 99% pure), and Ti0 ₂ (anatase, Cerac, 99% pure) with a Pulverisette 6 model Fritsch planetary mill. In all cases, about 10 g of the mixture of the precursor oxides was initially homogenized by hand in an agate mortar, and placed in a tungsten carbide (WC) jar of 250 ml with seven, 2 cm diameter, 63 g mass each WC balls for activation at 300 r.p.m., for 20 h.

These conditions have been shown to provide perovskite single phase fully crystalline powders with nanometer-scale chemical homogeneity.

About 1 g of nanocrystalline powder was uniaxially pressed into 12 mm diameter pellets, which were then sintered in a closed Al ₂ 0 ₃ crucible inside a furnace. A temperature of 1 100°C, 1 125°C and 1 150°C, a soaking time of 1 h and heating/cooling rates of ± 3 °C min ^"1 were selected.

Note that significant PbO or Bi ₂ 0 ₃ losses did not take place under these conditions, which allowed the use of initial precursor excesses or of sacrificial powder; that is, to bury the green bodies in powder during the thermal treatment, to be avoided. This is essential for an accurate composition and phase coexistence control, specially an issue for the problem addressed. Here point defects are engineered while maintaining full control of the structural and microstructural characteristics. Densification values above 95%, and homogenous fine-grained microstructures with average grain sizes between 1.0 and 2.5 μΐη were consistently achieved.

Characterization of the piezoelectric ceramic samples

Samples for phase and microstructural characterizations were prepared by thinning of ceramics to remove one surface (-100 μΐη), followed by polishing to a mirror finish. A final thermal treatment at 600°C for 2h with ± 0.5 °C min ^"1 was carried out to remove the damage introduced, and to restore the equilibrium polymorphic phase coexistence and domain configurations, which are modified by the shear stresses involved in polishing. Perovskite phase stability during sintering was controlled by X-ray diffraction (XRD) with a Siemens D500 powder diffractometer and Cu K _a radiation (λ=1.5418 A). Patterns were recorded between 20 and 50° (2Θ) with 0.05° (2Θ) step and 5 s counting rate. Slow scans; 0.02° (2Θ) step and 10 s counting rate, were carried out between 43 and 47° (2Θ) across the perovskite parent cubic phase 200 diffraction peak, for the analysis of the ferroelectric distortion and the evaluation of the phase percentages within the morphotropic phase boundary region.

Microstructure was studied with a FEI Nova™ NanoSEM 230 field emission gun scanning electron microscope equipped with an Oxford INCA 250 electron dispersive X-ray spectrometer for chemical analysis.

Ceramic capacitor for electrical and electromechanical characterizations were prepared by thinning discs down 0.5 mm, painting of Ag electrodes on the major faces, and their sintering at 700 °C.

Electrical characterization involved the characterization of dielectric permittivity and ferroelectric hysteresis loops. Dependences of the dielectric permittivity and losses on temperature were measured between room temperature (RT) and 550 °C with a HP4284A precision LCR meter. Measurements were dynamically carried out during a heating/cooling cycle with ±1.5 °C min ^"1 rate at several frequencies between 100 Hz and 1 MHz. Room temperature ferroelectric hysteresis loops were recorded under voltage sine waves of increasing amplitude up to 10 kV with a 0.1 Hz frequency, obtained by the combination of a synthesizer/function generator (HP 3325B) and a high voltage amplifier (TREK model 10/40), while charge was measured with a homebuilt charge to voltage converter and software for loop acquisition and analysis. Subsequently, the ceramic discs were poled for electromechanical characterization. A field of 4 kV mm ^"1 was applied at 100 °C for 15 min, and maintained during cooling down to 40 °C. The longitudinal piezoelectric coefficient d ₃₃ was then measured 24 h after the poling step with a Berlincourt type meter. Also, the transverse piezoelectric coefficient d ₃₁ was obtained by complex analysis of piezoelectric radial resonances of the discs by an automatic iterative method described in C. Alemany et al J Phys D: Appl Phys 1995; 28:945. This procedure also provides the Sn ^E and s _i2 ^E compliances and ε ₃₃ ^σ permittivity of the poled material all in complex form and thus, all mechanical, electrical and electromechanical losses.

Preparation and characterization of a magnetoelectric composite with the piezoelectric ceramic samples Finally, magnetoelectric composites were fabricated with selected Bi _{1 -x+} 2yPb _x-3y Sci- _x Ti _x 0 ₃ compositions. Three-layer structures consisting of one piezoelectric ceramic disc, glued between two Terfenol-D metal alloy pieces (ETREMA Products Inc.) by using a silver loaded epoxy adhesive were built, and their magnetoelectric response characterized. A system comprising a combination of two Helmholtz coils, designed to independently provide a static magnetic field up to 1 kOe to magnetize the material, and an alternate magnetic field up to 10 Oe at 10 kHz to play as stimulus (Serviciencia S.L.), was used, while the magnetoelectric voltage response was monitored with a lock-in amplifier. A 31 geometry was chosen to obtain the transverse magnetoelectric coefficient a ₃₁ as a function of the bias magnetic field H, after normalization to the piezoelectric element thickness (0.5 mm).

Results XRD patterns for Bi _1-x+ 2 _y Pb _x- 3 _y Sci _-x Ti _x 0 ₃ with x= 0.64 and y=0, 0.01 and 0.025 are shown in Figure 1. No second phases in addition to the perovskite one are found in the ceramics with y=0 and 0.01 , while an unidentified small diffraction peak appears in the material with y=0.025, suggesting the substitution of Bi for Pb to be incomplete at this nominal level.

Patterns for Bi _{1 -x+} 2yPb _x- 3 _y Sci- _x Ti _x 0 ₃ with x= 0.64 and y=0, 0.01 and 0.025 with improved statistics across the perovskite parent cubic phase 200 diffraction peak, along with their deconvolution by using three pseudovoigt functions are given in Figure 2. Coexistence of rhombohedral and tetragonal phases is assumed for simplicity, even though the former polymorph is known to be monoclinic. Results clearly indicate all materials with increasing amount of Bi substitution to be within the morphotropic phase boundary region, yet the percentage of rhombohedral phase clearly increases with y, while the tetragonal distortion decreases. Scanning electron microscopy (SEM) images for Bi _1-x+ 2yPb _x- 3 _y Sci- _x Ti _x 0 ₃ with x= 0.64 and y=0, 0.01 and 0.025, this is, the three materials with an increasing amount of Bi substituting for Pb are shown in Figure 3. A homogenous microstructure with average grain size of 2.4, 1.2 and 1 .3 μΐη is obtained for y=0, 0.01 and 0.025, respectively. A distinctive decrease in grain size thus results. In order to obtain a series of ceramic materials with comparable phase coexistence and an increasing concentration of point defects, a new series of ceramics with x=0.65 and y=0.01 , x=0.66 and y=0.01 , x=0.66 and 0.025, and x=0.68 and y=0.025 were processed and characterized.

Microstructure coarsening was also targeted, so sintering experiments at 1 125 and 1 150°C were additionally carried out.

Indeed the material with x=0.66 and y = 0.01 showed phase coexistence comparable to that of x=0.64 and y=0 (reference material with no point defect engineering). Its XRD pattern is shown in Figure 4, along with that of the ceramic with x=0.64 and y=0.01 . Note the slight increase of tetragonal distortion with x. In this way, a material having an A-site non-stoichiometry, while still being at the core of the MPB has been chemically engineered.

The temperature dependences of the dielectric permittivity and losses for the different Bii- _x+2 yPb _x- 3 _y Sci- _x Ti _x 0 ₃ materials are shown in Figure 5, while ferroelectric hysteresis loops are given in Figure 6. Curves corresponding to the initial series of materials with x=0.64 and an increasing level of Bi substitution are given in Fig. 5(a) and 6(a). Note firstly in the former figure (Fig. 6(a)) the continuous decrease of the Curie temperature with y. The dielectric anomaly associated with the ferroelectric transition is observed at a temperature of 450, 395 and 375 °C for y=0, 0.01 and 0.025, respectively. A distinctive increase of the room temperature dielectric permittivity and losses of the unpoled ceramics with increasing Bi substitution is also found. Specific values are given in Table I, along with Tcs.

This simultaneous increase of permittivity and losses might be interpreted as an enhancement of the domain wall mobility. Indeed ferroelectric hysteresis loops show a distinctive decrease of the coercive field, from 2.5 down to 2.0 kV mm ^"1 , when a Bi substitution of y=0.01 is introduced, at the same time that the remnant polarization increases from 40 up to 44 μC cm ^"2 . Furthermore, this takes place in spite of the decrease of grain size that is known to increase the coercive field. However, this effect might also be caused by the observed shift of the phase coexistence towards the rhombohedral polymorph, for its coercive field is lower than that of the tetragonal phase.

Results for the series of ceramics with increasing values of x and y=0.01 are presented in Fig. 5(b) and 6(b). Note in the former figure the continuous increase of the Curie temperature with x. The dielectric anomaly associated with the ferroelectric transition is observed at a temperature of 395, 407 and 415 °C for x=0.64, 0.65 and 0.66, respectively. Specific values are also given in Table I, along with T _c s. Note also that the coercive field increases, such as the ceramic with x=0.66 and y=0.01 that has the same phase coexistence of the reference material (x=0.64 and y=0), has a similar coercive field too. This indicates that the introduced A-site non-stoichiometry does not significantly enhance the domain wall mobility, and actually, a higher remnant polarization is not obtained. Results for the series of ceramics with increasing x and y=0.025 are presented in Fig. 5(c) and 6(c). Note in the former figure the continuous increase of the Curie temperature with x. The dielectric anomaly associated with the ferroelectric transition is observed at a temperature of 375, 388 and 408 °C for x=0.64, 0.66 and 0.68, respectively. Generally speaking, materials with y=0.025 show a systematic decrease of the room temperature permittivity and losses with x, along with an increase of the coercive field. They also show a strong depletion of the dielectric anomaly, and reduced remnant polarization. All these strongly indicate the presence of a grain boundary phase formed by accumulation of point defects, segregated from the perovskite phase that cannot accommodate them. Results also indicate this phenomenon to become increasingly relevant as x increases, which suggest it to be a characteristic of the tetragonal polymorph.

Charge longitudinal and transverse piezoelectric coefficients after poling are given in Table I.

A first result worth commenting on is the very large enhancement of permittivity experienced by the materials at the core of the MPB after poling. The permittivity of the unmodified high sensitivity BiSc0 ₃ -PbTi0 ₃ increased from 835 up to 1486, while dielectric losses did not change. Changes in the stress state or/and in the percentage of phases in coexistence at the MPB upon poling could be responsible for the permittivity increase. This is an issue for voltage response in sensing applications and the reason why the point defect engineering was introduced. Indeed materials with initial phase coexistence shifted towards the rhombohedral side and the A-site non- stoichiometry; those with x=0.64 and y=0.01 and 0.025 do not show a comparable increase.

A second result worth stressing out is the distinctive decrease of the elastic compliances of the poled ceramics after the introduction of the A-site non- stoichiometry. This takes place for the series of materials with x=0.64 and increasing y, and also for the material tailored to remain at the core of the MPB, while incorporating the point defects. This latter material does not only show smaller piezoelectric coefficients and elastic compliances than the unmodified BiSc0 ₃ -PbTi0 ₃ , but also a smaller permittivity after poling, and decreased dielectric, mechanical and electromechanical losses.

Most interesting properties are obtained for the material with x=0.64 and y=0.01 that shows an enhanced remnant polarization and decreased coercive field, which result in piezoelectric coefficients close to those of the unmodified high sensitivity composition, but a significantly reduced dielectric permittivity (see Table I). As a result, its voltage piezoelectric coefficient is much higher than that of the latter material. Specifically, the g ₃₃ coefficient increased from 3.3x10 ^"2 to 4.8x10 ^"2 V m N ^"1 after the incorporation of the A-site-non stoichiometry. This value is also higher than that of commercially available chemically engineered PZTs with g ₃₃ values ranging between 2 and-2.8x10 ^"2 V m N ^"1 . Note that this enhancement is basically a consequence of the reduced permittivity of the poled material, and it is not caused by a decrease of the domain wall contribution, but of the single crystal one, either as a consequence of the modification of the phase coexistence or a direct effect of the presence of the point defects in polarizability. Moreover, the material has also a strongly decreased elastic compliance; one can assume then that the incorporation of the point defects causes an overall increase of the lattice rigidity with a direct effect on polarizability and deformability.

A new application of high sensitivity piezoelectrics is magnetoelectric composites, in which they are combined with magnetostrictive materials to provide magnetoelectricity as a product property. Magnetoelectric transducers are being considered for a range of technologies like high sensitivity, room temperature operation magnetic field sensors. It has been shown that the voltage magnetoelectric coefficient of the simplest two-layer piezoelectric-magnetostrictive structure is given by

(1 )

Where d and q are the piezoelectric and piezomagnetic coefficients of the respective phases of volume v, s are the elastic compliances, and ε the permittivity, while superindex p and m refer to the piezoelectric and magnetic component, respectively. Note that the voltage response not only increases with the reciprocal permittivity, but also with the reciprocal compliances. In Figure 7, the magnetoelectric voltage coefficient of three-layer structures fabricated with Terfenol-D and either the BiSc0 ₃ -PbTi0 ₃ material incorporating the A-site non- stoichiometry (x=0.64 and y=0.01 ), or the reference high sensitivity material (x=0.64 and y=0), are compared. Indeed, the magnetoelectric coefficient is increased from 0.2 up to 1 .05 V cm ^"1 Oe ^"1 by the point defect engineering.

Table I.