Ceramic

Abstract

The present invention relates to a ceramic, to a process for preparing the ceramic and to the use of the ceramic as a dielectric in a capacitor.

Claims

1. A ceramic comprising a solid solution having a tetragonal tungsten bronze structure of general formula:
Sr.sub.2−dCa.sub.e[α].sub.f[β].sub.1−gNb.sub.5−h[γ].sub.hO.sub.15−k wherein: [α] denotes one or more of the group consisting of the rare earth elements and actinides; [β] denotes one or more of the group consisting of the alkali metals and the alkaline earth metals; [γ] denotes one or more of the group consisting of zirconium, hafnium, titanium, manganese, tin, silicon and aluminium; −0.1≤d≤0.2; 0<e≤0.1; 0≤f≤0.2; 0≤g≤0.2; 0≤h≤0.1; f=d+g−e; h≤f; and k denotes an oxygen deficit sufficient to ensure charge balance.

2. A ceramic as claimed in claim 1 which is substantially monophasic.

3. A ceramic as claimed in claim 1, wherein 0<d≤0.2.

4. A ceramic as claimed in claim 1 wherein 0.01≤e≤0.075.

5. A ceramic as claimed in claim 1 wherein 0<f≤0.2.

6. A ceramic as claimed in claim 1 wherein 0.01≤f≤0.15.

7. A ceramic as claimed in claim 1 wherein 0<h≤0.1.

8. A ceramic as claimed in claim 1 wherein 0.01≤h≤0.075.

9. A ceramic as claimed in claim 1 wherein [α] is yttrium (Y) or lanthanum (La).

10. A ceramic as claimed in claim 1 wherein [α] is yttrium (Y).

11. A ceramic as claimed in claim 1 wherein [β] is sodium (Na).

12. A ceramic as claimed in claim 1 wherein [γ] is zirconium (Zr).

13. A ceramic as claimed in claim 1 wherein the solid solution has a tetragonal tungsten bronze structure of general formula:
Sr.sub.2−dCa.sub.eY.sub.fNa.sub.1−gNb.sub.5−hZr.sub.h O.sub.15−k.

14. A ceramic as claimed in claim 1 wherein the solid solution has a tetragonal tungsten bronze structure of formula:
Sr.sub.2−x−yCa.sub.x[α].sub.y[β]Nb.sub.5−y[Y].sub.yO.sub.15 wherein: 0<x≤0.1; and 0≤y≤0.1.

15. A ceramic as claimed in claim 14 wherein 0.01≤x≤0.075.

16. A ceramic as claimed in claim 14 wherein 0<y≤0.1.

17. A ceramic as claimed in claim 14 wherein 0.01≤y≤0.075.

18. A ceramic as claimed in claim 14 wherein the solid solution has a tetragonal tungsten bronze structure of formula:
Sr.sub.2−x−yCa.sub.xY.sub.yNaNb.sub.5−yZr.sub.yO.sub.15.

19. A ceramic as claimed in claim 14 wherein x and y are the same.

20. A ceramic as claimed in claim 1 obtainable by sintering a sinterable form of a mixed metal oxide containing Sr, Ca, [α], [β], Nb and [γ].

21. Use of a ceramic as defined in claim 1 as a dielectric in a capacitor.

Description

EXAMPLE 1

Experimental

[0087] Samples of Sr.sub.2NaNb.sub.5O.sub.15, Sr.sub.2−xCa.sub.x NaNb.sub.5O.sub.15 (x=0.025, 0.05 and 0.075) and Sr.sub.2−x−yCa.sub.xY.sub.yNaNb.sub.5−yZr.sub.yO.sub.15 were prepared using a mixed oxide synthesis. The starting reagents in powder form were strontium carbonate (Aldrich, 99.9%), calcium carbonate (Aldrich, >99%), sodium carbonate (Sigma-Aldrich, 99.95%), niobium oxide (Alfa Aesar, 99.9%), yttrium oxide (Alfa Aesar, 99.9%) and zirconium oxide (Aldrich, 99%). The powders were mixed in appropriate ratios before ball-milling for up to 24 hours using stabilised-zirconia grinding media in isopropanol. Dried powders were calcined at 1200° C. for 6 hours (heating rate 5° C./min) in high purity alumina crucibles. The calcined powders with 2 wt % of binder (Optapix AC112, Zschimmer & Schwarz) were ball milled in water for 24 hours, dried and passed through a 300 μm mesh nylon sieve, before being pressed uniaxially at 100 MPa (for 90 s) in a 1 cm diameter steel die. Pellets were placed on a powder bed of the same composition in high purity alumina crucibles and covered to a depth of ˜1 cm with powder of the same composition. For sintering, the compacted pellets were first heated at 1° C./min to 550° C. and held for 4 hours to burn out the binder. The pellets were then heated at 5° C./min to (for example) 1300° C. or 1350° C. and held at this temperature for 4 hours.

[0088] Densities were measured from pellet dimensions and mass. The theoretical density was obtained from the nominal unit cell contents and measured lattice parameters. Phase analysis by X-ray powder diffraction (XRD) was carried out using a Bruker D8 X-ray powder diffractometer. Unit cell lattice parameters were obtained by Rietveld refinement. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were each used with energy dispersive X-ray capability (EDX) for microstructural evaluation and to provide compositional information. Relative permittivity (ε.sub.r) and loss tangent (tan δ) were measured as a function of temperature at fixed frequencies using a HP4284 LCR meter (Hewlett Packard) for the temperature range 20 to 400° C. For temperatures down to −70° C., an environmental chamber was used (Tenney). Silver electrodes were applied to opposite pellet faces (Sun Chemical, Gwent Electronic Materials).

Results and Discussion

Structural Analysis

[0089] XRD patterns of powders of crushed sintered pellets of Sr.sub.2−XCa.sub.xNaNb.sub.5O.sub.15 (see FIG. 1 for x=0.0 and 0.05) were generally similar to those reported in the literature. A shoulder to the peak at 32.24°2θ corresponds to the main peak in NaNbO.sub.3 (see García-González, E., Torres-Pardo, A., Jiménez, R. and González-Calbet, J. (2007). Structural Singularities in Ferroelectric Sr.sub.2NaNb.sub.5O.sub.15. Chemistry of Materials, 19(14), pp. 3575-3580 and Torres-Pardo, A., Jiménez, R., González-Calbet, J. and García-González, E. (2011). Structural Effects Behind the Low Temperature Nonconventional Relaxor Behavior of the Sr.sub.2NaNb.sub.5O.sub.7s Bronze [online] Inorganic Chemistry, 50(23), pp. 12091-12098).

[0090] There is ambiguity in the literature as to whether diffraction patterns of this type of tungsten bronze can be indexed as tetragonal or on the basis of a larger orthorhombic cell which takes account of a supercell arising from NbO.sub.6 octahedral tilting. The diffraction patterns in FIG. 1 are indexed according to an orthorhombic cell as reported for Sr.sub.2NaNb.sub.5O.sub.15 by García-González [supra]. Solid solution formation involving Ca substitution for Sr (and possibly Na) was confirmed by the slight shift in lattice parameters (see FIG. 2).

[0091] Measured densities were in the range 4.7-4.8 g/cm.sup.3 corresponding to ˜88-92% of the theoretical density. Grain sizes were as expected for a ceramic prepared by a conventional mixed oxide synthesis (typically <7 μm—see FIG. 3).

[0092] Contrast variations in backscattered SEM images in ˜5% of the grains implied a compositional variation which was shown to be due to grains with a higher Na content than the matrix by SEM-EDX elemental mapping (see FIG. 3). These grains are the NaNbO.sub.3 phase detected by XRD.

[0093] The effect of incorporating Y and Zr (assuming equal levels of substitution of Y for Sr and Zr for Nb) in a solid solution Sr.sub.2−x-yCa.sub.xY.sub.yNaNb.sub.5-yZr.sub.yO.sub.15 was investigated for x=0.025 and y=0.025 and for x=0.05 and y=0.05. There was little difference in the XRD peak positions compared to when y=0 (see FIG. 1). However a small amount of zirconia was detected.

[0094] SEM-EDX elemental mapping of a YZr modified sample is presented in FIG. 4. This highlights grains that are Na-rich (similar to those observed for y=0 samples). There were also Zr-rich grains evident which was consistent with the ZrO.sub.2 peaks detected by XRD (see FIG. 1). The latter findings cast doubt on the assumed charge balance mechanism of Zr.sup.4+ ion substitution for Nb.sup.5+ ions (Zr.sub.Nb) to complement an equivalent level of Y.sup.3+ ion substitution for Sr.sup.2+ ions (Ysr). However further microstructural analysis using high resolution S/TEM and EDX (see FIG. 5) revealed the presence of Y and Zr in all matrix grains and confirmed lattice substitutions of both elements into the Sr.sub.2−xCa.sub.xNaNb.sub.5O.sub.15 lattice. The amount of ZrO.sub.2 decreased significantly for a sintering temperature of 1350° C. compared with 1300° C. indicating that (in-part) the free ZrO.sub.2 was a result of incomplete solid state reaction.

[0095] The SEM and S/TEM EDX analysis each provided evidence of ˜ 1 μm zirconia grains in samples sintered at 1300° C. Zirconia was also present but in a reduced amount in samples sintered at 1350° C. There was no evidence of any yttrium containing secondary phase. Therefore the SEM and S/TEM EDX analysis results imply that the defect chemistry is more complex than is assumed by the starting compositional formula. The composition of the main phase will have deviated slightly from the nominal solid solution formula. In the sintered ceramic, charge balance could involve mainly Sr (Ca) and Na substitutions but with more limited B site substitution than first anticipated. This would have the effect of generating a surplus of Na, Sr and Zr ions which is consistent with the phase assembly detected. A separate study of phase equilibria and defect chemistry would be required to understand structure-property relationships in detail and to establish the optimum ceramic processing conditions. Thus the combined results of XRD, SEM and S/TEM-EDX suggest the ceramic products with the most useful dielectric properties may best be represented by the general formula Sr.sub.2−dCa.sub.eY.sub.fNa.sub.1−gNb.sub.5-hZr.sub.hO.sub.15−k where f=d+g−e and h<f.

Dielectric Properties

[0096] Sr.sub.2NaNb.sub.5O.sub.15 is characterised by two dielectric peaks in the temperature range of interest (see FIG. 6a). The higher temperature peak at 300° C. of similar tungsten bronzes is reported to correspond to the formation (on cooling) of a supercell involving out of plane (ab) octahedral tilts. These subtle structural changes generate ferroelectric behaviour. Reasons for the lower temperature peak at 0° C. are less well understood. It coincides with a change in thermal expansion coefficient suggesting that the composition is ferroelastic in character although no corresponding structural deviation has been detected (see Toledano, J. and Pateau, L. (1974). Differential thermal analysis of ferroelectric and ferroelastic transitions in barium sodium niobate. Journal of Applied Physics, 45(4), pp. 1611-1614).

[0097] For Sr.sub.2NaNb.sub.5O.sub.15, the higher temperature peak occurred at 305° C. (T2) and the lower temperature peak occurred at −15° C. (T1) (see FIG. 6a). The latter showed frequency dispersion akin to that of a relaxor ferroelectric. The effect of Ca substitution was to increase the relative magnitude of the lower temperature T1. The T2 peak became slightly more diffuse (see FIGS. 6b and 6c).

[0098] The consequences on ε.sub.r−T response of modifications by Y and Zr at x=0.025 and 0.05 are shown in FIG. 7. For both compositional modifications, the T2 peak diminished in magnitude, most notably for x=0.05, y=0.05. The T2 peak was much more diffuse for x=0.05, y=0.05 and its temperature was reduced.

[0099] As a consequence of these changes, at a composition x=0.025, y=0.025 and for a sintering temperature of 1300° C. (density<90% theoretical), the values of permittivity fell in the range ε.sub.r=1076±14% across the temperature range −55 to 300° C. (where 1076 is the median ε.sub.r value and occurred at 85° C.). However for samples of higher density (˜93% theoretical) obtained by sintering at 1350° C. for 4 h, ε.sub.r=1510±16% from −70° C. to 300° C. The corresponding dielectric loss tangent value was <0.035 from −70° C. to 260° C. (see FIG. 7) increasing to 0.09 between 260° C. and 300° C.

[0100] In samples with higher levels of Y and Zr (ie x=0.05, y=0.05) an excellent combination of dielectric properties was observed in the temperature range −70 to 270° C. This upper temperature would meet the demands of most proposed power electronics applications. The permittivity values relative to the 25° C. point were: ε.sub.r(25C)=1370±14% from −70 to 270° C. The temperature stability specification of capacitors is normally described in terms of the % variation relative to a room-temperature value. Hence the ε.sub.r median value of x=0.05, y=0.05 at 25° C. is highly advantageous. Dielectric loss tangent values were 0.025 from −12 to 290° C. increasing to 0.03 at −32° C. and to 0.038 at −70° C. (see FIG. 7).

[0101] The foregoing dielectric properties of CaYZr modified Sr.sub.2NaNb.sub.5O.sub.15 ceramics are summarised in Tables 1a and 1b below.

[0102] Based on the premise that some of the Na.sub.2O in the ceramic may have been lost due to evaporation during sintering, the effect of adding excess Na.sub.2CO.sub.3 to the starting mixture was investigated. Additions of 2 wt % and 4 wt % caused an increase in the peak temperatures T1 and T2 and in the ε.sub.r values (see FIG. 8). This implies that defect structures were moderated by such additions. This hypothesis was confirmed by the lower dielectric losses in the high temperature regime (>250° C.) with tan δ decreasing to below 0.025. This implies a lower contribution from electrical conduction mechanisms. However the variability in ε.sub.r between −55 and 250-300° C. increased to beyond 15%.

Conclusions

[0103] A very promising bismuth- and lead-free ceramic compositional system for new Class II dielectric materials has been demonstrated to have high and stable relative permittivity and low dielectric loss over a very wide temperature range of −70° C. (or lower) to 270-300° C. for the nominal solid solution series: Sr.sub.2−x-yCa.sub.xY.sub.yNaNb.sub.5-yZr.sub.yO.sub.15. For x=0.05, y=0.05 the value of relative permittivity at 25° C. was 1370 with only ±14% variation in values for temperatures between −70 and 270° C. Dielectric loss tangent values were <0.03 except for the range −32 to −70° C. where they rose slightly to 0.038. High resolution scanning transmission electron microscopy with energy dispersive X-ray analysis confirmed Ca, Y and Zr substitution in the crystal lattice of the parent tungsten bronze Sr.sub.2NaNb.sub.5O.sub.15 was achieved but the presence of minor amounts of zirconia and sodium niobate phases imply the composition of the main phase deviates slightly from the nominal solid solution formula (although the amount of secondary phases decreased when sintering temperature was increased from 1300 to 1350° C. due to an increased degree of solid state reaction). The properties indicate that these bismuth- and lead-free tungsten bronze niobates are excellent candidates as high temperature capacitor materials. Thermodynamic calculations predict they are compatible with nickel electrode multilayer ceramic capacitor co-firing techniques.

TABLE-US-00001 TABLE 1 Summary of key dielectric properties of ceramics of Sr.sub.1.95C.sub.0.025Na.sub.1.0Y.sub.0.025Zr.sub.0.025Nb.sub.4.975O.sub.15 (sintered at 1350° C.) and Sr.sub.1.90Ca.sub.0.05Na.sub.1.0Y.sub.0.05Zr.sub.0.05Nb.sub.4.95O.sub.15 (sintered at 1300° C.) a) Temperature Range −70 to 270° C.: summary 1 kHz dielectric data Sample ε.sub.r median ±% ε.sub.r .sub.−70 tanδ.sub.max −70- Code Intended Product Composition density (T) to 270° C. 270° C. tanδ <0.035 tanδ <0.03 tanδ <0.025 x = 0.025 Sr.sub.1.95C.sub.0.025Na.sub.1.0Y.sub.0.025Zr.sub.0.025Nb.sub.4.975O.sub.15 ~92% 1510 16% 0.045 −70-260° C. −70-250° C. −60-245° C. y = 0.025 (85° C.) x = 0.05 Sr.sub.1.90Ca.sub.0.05Na.sub.1.0Y.sub.0.05Zr.sub.0.05Nb.sub.4.95O.sub.15 ~93% 1370 14% 0.038 −50-381° C. −32-370° C. −12-290° C. y = 0.05 (25° C.) b) Temperature Range −70 to 300° C.: summary 1 kHz dielectric data (sintering T = 1350° C. and 1300° C. respectively) Sample ε.sub.r median ±% ε.sub.r −70 tanδ.sub.max −70 Code Intended Product Composition density (T) to 300° C. to 300° C. tanδ <0.035 tanδ <0.03 tanδ <0.25 x = 0.025 Sr.sub.1.95Ca.sub.0.025Na.sub.1.0Y.sub.0.025Zr.sub.0.025Nb.sub.4.975O.sub.15 ~92% 1510 16% 0.09 −70-260° C. −70-250° C. −60-245° C. y = 0.025 (85° C.) x = 0.05 Sr.sub.1.90Ca.sub.0.05Na.sub.1.0Y.sub.0.05Zr.sub.0.05Nb.sub.4.95O.sub.15 ~93% 1271 26% 0.038 −50-381° C. −32-370° C. −12-290° C. y = 0.05 (82° C.)

EXAMPLE 2

Experimental

[0104] In this Example, the chemical formula of the parent niobate phase (Sr.sub.4Na.sub.2Nb.sub.10O.sub.30) is expressed as Sr.sub.2NaNb.sub.5O.sub.15 (SNN) for convenience. The substituted compositions are expressed assuming a solid solution formula Sr.sub.2−2zCa.sub.zY.sub.zNaNb.sub.5-zZr.sub.zO.sub.15. The assumption is that Ca.sup.2+ and Y.sup.3+ substituents will occupy A1/A2 sites and Zr.sup.4+ will occupy Nb.sup.5+ (B) sites. C cites will remain empty. Sample formulations with z=0, 0.025 and 0.05 were prepared using a mixed oxide synthesis. The compositions correspond to a very low level of substitution. Only 1.25 at. % of the Sr.sup.2, (A) sites are substituted by Y.sup.3+ in the composition z=0.025 and 2.5 at. % in the composition z=0.05. For B sites, the levels of substitution of Zr.sup.4+ for Nb.sup.5+ are 0.05 at. % and 1 at. % respectively for compositions z=0.025 and 0.05.

[0105] The starting reagents were strontium carbonate (Aldrich, 99.9%), calcium carbonate (Aldrich, >99%), sodium carbonate (Sigma-Aldrich, 99.95%), niobium oxide (Alfa Aesar, 99.9%), yttrium oxide (Alfa Aesar, 99.9%) and zirconium oxide (Alfa Aesar, 99.7%). Powders were mixed in appropriate ratios before ball-milling for up to 24 hours using stabilised zirconia grinding media in isopropanol. Dried powders were calcined at 1200° C. for 6 hours (heating rate 5° C./min) in high purity alumina crucibles. The calcined powders with the addition of 2 wt. % of binder (Optapix AC112, Zschimmer & Schwarz) were ball milled in water for 24 hours, dried and passed through a 300 μm mesh nylon sieve, before pressing uniaxially at 100 MPa (for 90 s) in a 1 cm diameter steel die. After uniaxial pressing, the green pellets were isopressed (200 MPa for 5 minutes) in an isostatic press (Stanstead fluid power, Essex, UK). Binder burn-out was performed at a heating rate of 1° C./min to a dwell temperature of 550° C. and held for 5 hours. Sintering was carried out after embedding the pellets in a powder of the same composition. Maximum densities were obtained at a sintering temperature of 1300° C. or 1350° C. Dwell times were 4-5 hours. Sintered ceramic densities were measured from measured pellet dimensions and mass. The theoretical density was estimated from the nominal unit cell contents and measured lattice parameters.

[0106] Phase analysis by powder X-ray diffraction (XRD) was carried out using a Bruker D8 X-ray powder diffractometer. Unit cell lattice parameters of an adopted pseudo-tetragonal structure were obtained by full pattern Rietveld refinement using TOPAS 5.0 software (Bruker AXS, Karlsruhe, Germany). In the refinement analysis, the peak shape function was determined by the fundamental parameters of the X-ray diffractometer geometry. The refined parameters are background function coefficient, lattice constant, scale factor and atomic coordination.

[0107] In order to prepare specimens for microstructural characterisation by scanning electron microscopy, ceramic pellets were mounted in epoxy resin (Epothin, Buehler) and ground with P240, P600 and P2500 silicon carbide paper. Subsequent sequential polishing was carried out using Texmet P microcloths with MetaDi 2 diamond suspensions of decreasing particle size: 9 μm, 3 μm and 1 μm. A final polish was carried out with ChemoMet and MasterMet 0.06 am colloidal silica on a Buehler EcoMet 300 grinder/polisher. Chemical etching was carried out with a 2:1 ratio of hydrofluoric acid and concentrated nitric acid for 90 seconds at room temperature.

[0108] Scanning electron microscopy (SEM) was performed using a Hitachi SU8230 high performance cold field emission instrument fitted with an Oxford Instruments Aztec energy dispersive X-ray analysis (EDX) system with 80 mm.sup.2 X-Max SD detector and analysis software. For transmission electron microscopy (TEM), thin sample lamellae were prepared via the in-situ lift-out method using a FEI Helios G4 CX Dual Beam—High resolution monochromated, field emission gun, scanning electron microscope (FEG-SEM) with precise Focused Ion Beam (FIB). In the Dual Beam microscope, 500 nm of platinum (Pt) was electron beam deposited (at 5 kV, 6.4 nA for the electron source) onto the surface of the target area. This was followed by a second Pt layer (1 μm) using the FIB (at 30 kV, 80 pA for the liquid Ga ion source). An initial lamella was cut (by the FIB at 30 kV, 47 nA), before a final cut-out was performed (at 30 kV, 79 nA). Final thinning and polishing of the lamellae to electron transparency was performed with a low energy ion beam (5 kV, 41 pA). The lamellae were attached, using ion beam deposited Pt onto a copper FIB lift-out grid (Omniprobe, USA) mounted within the SEM chamber (in-situ) ready for transfer to the TEM. The lamellae were imaged using a FEI Titan Themis.sup.3 300 kV TEM fitted with a SuperX EDX system and Velox processing software.

[0109] For electrical measurements, silver electrodes were applied to opposite pellet faces (Sun Chemical, Gwent Electronic Materials). Relative permittivity, ε.sub.r, and loss tangent (tan δ) were measured at low-field as a function of temperature at fixed frequencies using a Hewlett Packard, HP4284 LCR analyser. An environmental chamber was used for lower temperatures down to −65° C. (TJR; Tenney Environmental-SPX, White Deer, Calif.). Ferroelectric hysteresis measurements were carried out using a HP33120A function generator in combination with a HVA1B high voltage amplifier (Chevin Research, Otley, UK), using a sinusoidal electric field waveform with a frequency of 2 Hz. The measured electric field-time and current-time waveforms were processed to yield polarisation-electric field (P-E) loops and effective complex permittivity values using the method described by M. Stewart, M. G. Cain, D. A. Hall, Ferroelectric hysteresis measurement & analysis, National Physical Laboratory Report CMMT(A), 152[1] (1999).

Results and Discussion

[0110] Full-pattern refinements of X-ray powder diffraction data for crushed sintered pellets are shown in FIG. 9. The secondary phase NaNbO.sub.3 was present in all three samples. Increasing the calcination and sintering times failed to eliminate the extra phase. Thus even for unmodified SNN the notional formula Sr.sub.2NaNb.sub.5O.sub.15 may be inaccurate. For example, the Na rich secondary phase may be due to Sr.sup.2+ occupancy of a fraction of the perceived Na.sup.+ sites, giving a formula Sr.sub.2+xNa.sub.1−2xNb.sub.5O.sub.10. The monoclinic ZrO.sub.2 secondary phase was identified only in the z=0.05 sample. All phases were included in the Rietveld refinements.

[0111] No convincing evidence was found from XRD of weak extra superlattice reflections at around 20 °2θ or 37 °2θ which others have observed with the aid of electron diffraction and attributed to an orthorhombic unit cell (Space Group Im2a). The lack of any distinct supercell reflections in the XRD patterns prompted indexing on tetragonal axes and refinement of the data on the basis of space group P4bm. Crystallographic data refined on P4bm are summarised in Table 2. The modifications by Ca.sup.2+, Y.sup.3+, Zr.sup.4+ produced a slight contraction in cell volume (see Table 2) consistent with solid solution formation.

TABLE-US-00002 TABLE 2 Summary of (pseudo) tetragonal lattice parameters, goodness of fit, R.sub.wp and phase fractions from Rietveld analysis for Sr.sub.2−2zCa.sub.zY.sub.zNaNb.sub.5—.sub.zZr.sub.zO.sub.15. R.sub.wp % Composition a (Å) c (Å) V (Å.sup.3) % NaNbO.sub.3 z = 0 .sup. 12.365(90 3.8958(2) 595.73(3) 4.58 5.0 z = 0.025 12.362(9) 3.8920(1) 594.84(9) 4.99 7.4 z = 0.05 12.363(0) 3.8861(1) 593.97(3) 4.27 7.3 (+2.5 ZrO.sub.2)

[0112] Scanning electron micrographs of polished and etched sections of z=0 and z=0.05 are shown in FIG. 10. Observed grain sizes were similar (<10 μm) for both compositions. Densities were 92-93% of estimated theoretical values. The possibility of elemental segregation within the grains was investigated by SEM-EDX and TEM-EDX. For X7R BaTiO.sub.3 based capacitor materials, a core-shell grain structure brought about by a variety of additive oxides is responsible for inducing a temperature stable permittivity response from −55° C. to 125° C. Thus it was important to establish if a comparable microstructure-strain mechanism was responsible for flattening the ε.sub.r−T response of SNN. The SEM-EDX analysis for z=0.05 showed no elemental gradation within grains (see FIG. 11). The existence of secondary grains of sodium niobate and zirconia identified in XRD patterns with grain sizes of ˜5 μm and ˜1 μm respectively was confirmed by the SEM-EDX analysis. There was also some evidence from EDX for the presence of Sr in the sodium niobate grains. More detailed analysis using TEM-EDX confirmed an absence of core-shell grain structures, or indeed any form of elemental gradation within individual grains (see FIG. 12).

[0113] The relative permittivity-temperature (ε.sub.r−T) response of the parent tungsten bronze Sr.sub.2NaNb.sub.5O.sub.15 ceramic (SNN) is presented in FIG. 13a. The higher temperature dielectric peak (at 305° C.) is denoted T.sub.2. For other tungsten bronzes, this dielectric anomaly is reported to correspond to the formation (on cooling) of a supercell which induces ferroelectric behaviour: hence T.sub.2 represents the Curie point. Structural correlations are less well understood in the context of the lower temperature dielectric peak T.sub.1, which occurs at −14° C. in SNN (1 kHz) and shows frequency dispersion similar to a relaxor ferroelectric. An accompanying change in thermal expansion coefficient for related tungsten bronzes implies that the T.sub.1 peak corresponds to a ferroelastic transition but no associated structural deviations have been detected. For z=0 (SNN), the ‘standard’ dielectric peaks created a variation in ε.sub.r of ±22% across the important temperature range −55° C. to 300° C. (see FIG. 13a). This is well outside the R-type±15% stability level required of a Class II capacitor material. Consequently partial substitution of Sr.sup.2+ by Ca.sup.2, was investigated for Sr.sub.2−xCa.sub.xNaNb.sub.5O.sub.15 x<0.1. The ε.sub.r−T responses of SNN and Ca-SNN ceramics of comparable densities were generally similar. Further chemical modifications involving co-substitution of Ca.sup.2+, Y.sup.3+ for Sr.sup.2, and Zr.sup.4+ for Nb.sup.5+ were investigated in an attempt to suppress the temperature variability in permittivity and attain R-type performance. The substituent ions were selected on the basis of ionic radii and valence considerations.

[0114] For the Ca.sup.2+, Y.sup.3+, Zr.sup.4+ modified SNN sample composition z=0.025, the T.sub.2 peak temperature increased to 345° C. from a value of 305° C. in unmodified SNN (at 1 kHz). There was also a decline in the ε.sub.r max value due to increased broadening (see FIG. 13b). For the lower temperature peak, there was very little change in peak temperature T.sub.1 with substituent doping (−18° C. compared to −14° C. for SNN z=0) but there was an increase in frequency dispersion. For sample composition z=0.025 the difference in temperature (AT) of ε.sub.rmax temperatures (Tm) between frequencies 1 kHz and 1 MHz was 25° C. compared with 10° C. for unmodified SNN (z=0).

[0115] For a higher level of chemical substitution (z=0.05) the T.sub.2 anomaly was displaced to 255° C. which is 90° C. below the T.sub.2 peak for z=0.025 (see FIG. 13c). This non-monotonic shift of T.sub.2 with z indicates a complex interplay between substitution level and temperature of the dielectric anomalies which may well relate to alterations in defect structures (possibly affecting NbO.sub.6 tilts in the case of T.sub.2). The T.sub.2 anomaly also became significantly more diffuse as the level of substitution increased. As a result, the ε.sub.rmax value at T.sub.2 was approximately 60% of that observed for unmodified SNN (z=0). There was also an increase in broadening of the T1 peak (see FIG. 13c) but less so than for T2.

[0116] The net effect of these chemical modifications on peak temperatures and peak profiles was to achieve the requisite ε.sub.r±15%, R-type consistency in ε.sub.r over very wide ranges of temperature. For z=0.025, the measured variation in the ε.sub.r data was within ±13% of a median value of 1565 for temperatures extending from −65° C. to 325° C. (the median ε.sub.r value occurred at ˜105° C.). A further improvement in temperature-stability was achieved for higher levels of Ca.sup.2+, Y.sup.3+ and Zr.sup.4+ substitution. The z=0.05 sample composition gave a median value of ε.sub.r=1310 with a ±10% variation from temperatures of −65° C. to 300° C. Very relevant to consideration as a capacitor material, the median value of ε.sub.r in z=0.05 ceramics occurred at 25° C. Comparisons of the 1 kHz ε.sub.r−T plots for z=0, z=0.025 and 0.05 are shown in FIG. 14 to highlight the development of temperature stable permittivity.

[0117] The low-field dielectric loss tangent values at 1 kHz were 0.035 from −65° C. to 320° C. (tan δ≤0.025 from −60° C. to 290° C.) for z=0.025. Losses were slightly higher in the z=0.05 sample with tan δ<0.04. Dielectric data for these 92-93% dense samples are summarised in Table 3.

TABLE-US-00003 TABLE 3 Summary of dielectric data of 92-93% dense Sr.sub.2−2zCa.sub.zY.sub.zNaNb.sub.5−zZr.sub.zO.sub.15: ceramics (1 kHz data) Sample ε.sub.r median ±% ε.sub.r T range Code Intended Product Composition (T) T range tanδ ≤0.035 tanδ ≤0.03 tanδ ≤0.025 z = 0 Sr.sub.2.0Na.sub.1.0Nb.sub.5.0O.sub.15 1733 22% −65 to −65 to −65 to −32 to (277° C.) 300° C. 249° C. 238° C. 223° C. z = 0.025 Sr.sub.1.95C.sub.0.025Na.sub.1.0Y.sub.0.025Zr.sub.0.025Nb.sub.4.975O.sub.15 1565 13% −65 to −65 to −65 to −60 to (110° C.) 325° C. 320° C. 310° C. 290° C. z = 0.05 Sr.sub.1.90Ca.sub.0.05Na.sub.1.0Y.sub.0.05Zr.sub.0.05Nb.sub.4.95O.sub.15 1310 10% −65 to −40 to −20 to +20 to (25° C.) 300° C. 370° C.* 320° C. 270° C. *for z = 0.05, tanδ increased to 0.04 between −40° C. and −65° C.

[0118] The P-E hysteresis loops for all of the compositions were generally similar in appearance and showed clear evidence of ferroelectric character (see FIG. 15a). The maximum polarisation (initially around 13 μC.Math.cm.sup.−2) was reduced and the switching range around the coercive field became wider as z increased from 0 to 0.05. Significant dielectric nonlinearity and loss were evident in the sub-coercive field range (see FIG. 15b). For example, the effective tan δ value at an electric field amplitude of 4 kV.Math.cm.sup.−1 was determined as 0.154 for the undoped SNN, reducing to 0.081 and 0.060 for z=0.025 and 0.05 respectively.

[0119] Nonlinearity was also apparent in the real and imaginary parts of the complex dielectric permittivity (see FIG. 16). The observed behaviour generally departs from the classical Rayleigh Law (linear ε.sub.r−E.sub.max relation) tending towards a quadratic response in the field range up to 15 kV.Math.cm.sup.−1. The degree of nonlinearity was strongly suppressed for z=0.05 which indicates reduced contributions to the electric field-induced polarisation from domain switching mechanisms, consistent with increasing disorder.

[0120] In summary, the primary dielectric parameters of Bi-free and Pb-free dielectric ceramics produced by very low levels of chemical substitution of a tungsten bronze Sr.sub.2NaNb.sub.5O.sub.15 ferroelectric with Ca.sup.2+, Y.sup.3+, Zr.sup.4+ are class leading and very significant in the quest to develop base metal electrode Class II capacitor materials capable of operating over very wide temperature ranges. Future fundamental studies of crystal structure and defect chemistry will be required to elucidate the reasons why such low levels of compositional modification by Ca.sup.2+, Y.sup.3+ and Zr.sup.4+ bring about such a dramatic change in the permittivity response. However even at this early stage, it is possible to exclude core-shell microstructural mechanisms of the type which convert perovskite BaTiO.sub.3 into a X7R temperature stable dielectric. Moreover, the concentrations of Ca.sup.2+, Y.sup.3+, Zr.sup.4+ required to flatten the permittivity response of SNN are far below those required to produce significant broadening of Curie peaks in perovskites due to compositional heterogeneity effects.

Conclusions

[0121] A high permittivity (Class II) ceramic dielectric that offers stable permittivity to >300° C. and which does not contain problematic bismuth or lead oxides is demonstrated. Chemical substitution of Sr.sub.2NaNb.sub.5O.sub.15 by Ca.sup.2+, Y.sup.3+ and Zr.sup.4+ ions results in a material which more than satisfies the technologically important −55° C. to 300° C. temperature range of stable capacitance required for next generation power capacitor materials. For the formulation Sr.sub.2−2zCa.sub.zY.sub.zNaNb.sub.5−zZr.sub.zO.sub.15 where z=0.025, values of ε.sub.r lie in the range 1565±13% for temperatures from −65° C. to 325° C. At a higher substitution (z=0.05) the twin dielectric peaks become even more diffuse giving Er values of 1310±10% from temperatures of −65° C. to 300° C. Dielectric loss tangent values are 0.035 (1 kHz) across the full temperature range of stable permittivity for sample composition z=0.025 and tan δ is 0.025 from −60° C. to 290° C. Limiting dielectric losses were slightly higher (tan δ≤0.04) for z=0.05 samples. These primary dielectric properties are of high impact given the growing demands for next generation Class II capacitors that can operate at temperatures well beyond the limit of existing market-leading BaTiO.sub.3 based capacitors (under 200° C.). The absence of any volatile bismuth oxide component is highly advantageous in the search for industrially-relevant dielectrics for future base metal electrode high temperature multilayer ceramic capacitors.