Colossal Dielectric Permittivity in (Nb+Al) Codoped Rutile TiO2

Jun 25, 2015 - The search for novel, high capacitance dielectric materials has long been driven by the need for device miniaturization and energy stor...
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Colossal Dielectric Permittivity in (Nb+Al) Codoped Rutile TiO2 Ceramics: Compositional Gradient and Local Structure Wanbiao Hu,† Kenny Lau,† Yun Liu,*,† Ray L. Withers,† Hua Chen,‡ Lan Fu,§ Bill Gong,∥ and Wayne Hutchison⊥ †

Research School of Chemistry, ‡Centre for Advanced Microscopy, §Research School of Physics and Engineering, The Australian National University, Canberra, ACT 2601, Australia ∥ Solid State & Elemental Analysis Unit, The University of New South Wales, Kensington, Sydney, NSW 2052, Australia ⊥ School of Physical, Environmental and Mathematical Sciences, Canberra Campus of The University of New South Wales, Canberra, ACT 2610, Australia S Supporting Information *

ABSTRACT: (Nb+Al) codoped rutile TiO2 ceramics with nominal composition Ti4+0.995Nb5+0.005yAl3+0.005zO2, z = (4−5y)/3 and y = 0.4, 0.5, 0.6, 0.7, and Ti4+0.90Nb5+0.05Al3+0.05O2 have been synthesized. The resultant samples in ceramic pellet form exhibit a colossal dielectric permittivity (>∼104) with an acceptably low dielectric loss (∼10−1) after optimization of the processing conditions. It is found that a conventional surface barrier layer capacitor (SBLC) effect, while it contributes significantly to the observed colossal permittivity, is not the dominant effect. Rather, there exists a subtle chemical compositional gradient inward from the pellet surface, involving the concentration of Ti3+ cations gradually increasing from zero at the surface without the introduction of any charge compensating oxygen vacancies. Instead, well-defined Gr ± 1/3[011]* satellite reflections with the modulation wave-vector q = 1/3[011]r* and sharp diffuse streaking running along the Gr ± ε[011]* direction from electron diffraction suggest that the induced additional metal ions appear to be digested by a locally intergrown, intermediate, metal ion rich structure. This gradient in local chemical composition exists on a scale up to ∼ submillimeters, significantly affecting the overall dielectric properties. This work suggests that such a controllable surface compositional gradient is an alternative method to tailor the desired dielectric performance.

1. INTRODUCTION

codoped rutile materials of this class in order to further optimize the performance of such materials. Al3+ is in the same main group (13) as In3+ and seems an ideal alternative acceptor dopant ion, given its high natural abundance and low cost. Indeed, preliminary studies have shown that high dielectric permittivity can be obtained when Al3+ replaces In3+ in (M3+, Nb5+) a codoped rutile.4 However, it is important to bear in mind that the ionic radii of Al3+ and In3+ ions (as well as any Ti3+ ions) in the six-coordinate, pseudo-octahedral Ti4+ cation environment of the rutile type average structure, differ quite significantly as follows: rAl3+(= 67.5 pm) < rTi4+ (= 74.5 pm) < rTi3+ (= 91 pm) < rIn3+ (= 94 pm). Bond valence sum (BVS) calculations5,6 indicate that replacing a Ti4+ ion by an Al3+ ion in the Ti4+O2 rutile type average structure gives a 10% underbonded, apparent valence (AV) for the Al ion of +2.6929, certainly not necessitating any neighboring oxygen vacancies. By contrast, replacing a Ti4+ ion by an In3+ ion in the Ti4+O2 rutile structure leads to a very significantly (∼75%) overbonded In3+

The search for novel, high capacitance dielectric materials has long been driven by the need for device miniaturization and energy storage.1−3 The development of new, high-performance, colossal dielectric permittivity (εr > 103) materials represents a significantly better approach to device miniaturization in comparison to approaches based on the geometric optimization of pre-existing capacitor materials. We recently discovered one such new, high-performance colossal permittivity materials system in the form of acceptor (In3+) and donor (Nb5+) codoped rutile TiO2.4 The material exhibited a high dielectric permittivity (on the order of 104) and low dielectric loss (2 × 104) and low tan δ’s ( ∼1 kHz) but does not fully recover in the low frequency region (< ∼1 kHz) once the bias is back to 0 V. We also investigated the impact of DC voltages (from −30 V to +30 V) on the dielectric properties of (Nb+Al) codoped TiO2 at 1 kHz and 1 MHz frequencies respectively (Figure 3c). Again, there is no noticeable voltage dependency of the permittivity and loss tangent at both measured frequencies. The tan δ measured at 1 kHz increases slightly when the DC voltage is applied, but the relative amplitude is still less than 10% higher. This increase can be associated with limited space charges. It differs from grain-boundary, interfacial polarization dominated systems where numerous space charges are accumulated at grain boundaries, resulting in a significant increase in both permittivity and loss when a small DC bias is applied.11,12 In the current (Nb+Al) codoped rutile TiO2 case, interfacial polarization arising from the grain boundary effect (or the internal barrier layer capacitor effect - IBLC) can only be observed at high temperatures (see Supporting Information, Figure S1) due to the hopping or transport of thermally activated electrons or space charges.13−15 We therefore conclude that the primary origin of the observed colossal dielectric permittivities apparent in our (Nb+Al) codoped rutile TiO2 ceramics is not fundamentally related to, or at least not dominated by, grainboundary-related interfacial polarization effects.

To gain further insight into the nature of the observed colossal dielectric response, the ceramic pellet prepared under the y = 0.5//Tcal = 1500 °C conditions was heavily polished. It is found that when the thickness of the pellet is decreased from 1.6 mm to 0.8 mm, both the measured dielectric permittivity as well as the loss tangent dramatically increase, as shown in Figure 1b. (A similar effect was also seen for the y = 0.6//Tcal = 1400 °C sample). This cannot be attributed to a surface contact effect because the pellet is always well polished using 1200-grid sand paper before coating on the electrodes. EPR and XPS results (Figure S2, Supporting Information) show that while the Nb5+ concentration remains unchanged on sample thinning, the concentration of the Ti3+ (Ti4+ + 3d1 electron) species gradually increases from an initial concentration of zero on sample thinning. Surprisingly, it was found that the Ti3+ concentration gradually increases on moving in from the initial sample surface, i.e., there is a chemical compositional gradient for Ti3+. Furthermore, the 3d1 electron associated with the presence of Ti3+ ions seems not to be tightly pinned by local defect states as the measured dielectric loss tangent is not as low as in the case of the (Nb+In) codoped TiO2 rutile, especially for the inner parts of the original sample. This compositional gradient (or coring) effect becomes even more noticeable when the (Nb+Al) dopant level is further increased. X-ray powder diffraction (XRPD) from a Ti0.90(0.5Nb +0.5Al)0.10O2 (sintered at 1400 °C) sample indicates a single phase rutile-type average structure. Scanning electron microscope (SEM) characterization of the same sample using backscattered electron (BSE) imaging coupled with energy dispersive X-ray spectroscopy (EDXS) indicates a homogeneously distributed Nb/Al/Ti molar ratio of very close to the nominal 5:5:90 even at such a high dopant level. Furthermore, the Nb/Al/Ti molar ratio does not change between the outer and inner sections of the pellets. XPS and/or EPR characterization indicates that the Al3+ and Nb5+ concentrations remain unchanged as a function of pellet thickness. The appearance and systematic increase in concentration of Ti3+ relative to Ti4+ cations (Figure S3, Supporting Information) on moving in from the initial sample surface is again apparent in the Ti3+ signals from the XPS spectra.16,17 This systematic increase in the concentration of Ti3+ relative to Ti4+ cations on moving in from the initial sample surface is accompanied by a gradual darkening of the pellet appearance (see Figure 4) when the sample thickness is thinned from 1.6 mm to 0.8 mm, consistent with the formation of a Ti3+ concentration gradient layer. Note that while the dielectric permittivity of the resultant sample gradually increases with decreasing pellet thickness, the dielectric loss tangent also significantly increases, i.e., deteriorates (see Figure 4). This time, no significant difference in the measured dielectric permittivities and tan δ’s were observed when either Ag or Au was used as electrode, respectively. Similarly to the 0.5% (Al+Nb) codoped case, grain-boundary interfacial polarization effects could only be observed at high temperatures (see Supporting Information, Figure S4). The origin of the greatly different dielectric properties upon thinning is thus most likely associated with the color change caused by the increasing reduction of Ti4+ to Ti3+ and the associated compositional gradient and defect states in (Nb+Al) codoped TiO2. It is noted that a similar coring effect, i.e., a compositional Ti4+/ Ti3+ gradient, appears in traditional core structured material, induced by the reduction of Ti4+ and attributed to associated oxygen vacancies,18 which is, however, different from the present cases. Given that the Al3+ and Nb5+ concentrations remain 4937

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Figure 5. Raman spectra of (Nb+Al) codoped rutile ceramics with Ti0.90(0.5Nb+0.5Al)0.10O2 composition: (a) full-wavenumber spectra and (b) enlarged spectra of the Eg and A1g bands.

Figure 4. Dielectric properties of (Nb+Al) codoped TiO2 rutile ceramics with Ti0.90(0.5Nb+0.5Al)0.10O2 composition. Au electrodes were used for these dielectric property measurements. Images of the pellets are also shown. The top pellet is the unpolished, as-prepared pellet. The following pellets were gradually polished to decreasing thicknesses, from 1.6 mm to 0.8 mm.

particular, it appears that there is almost certainly no oxygen vacancy in the (Nb+Al) codoped TiO2 samples. If there is no oxygen vacancy to compensate for the increasing Ti 3+ concentration upon thinning, charge balance can only be maintained by increasing the total cation to oxygen ratio from the ideal 1:2 to 1 + x/(4 − x):2, where the ratio of Ti3+/Ti4+ is assumed to be x:0.90 − x. The overall stoichiometry for charge balance then becomes {(1 + x/(4 − x))[(Ti4+0.90‑xTi3+x)(Nb5+0.05Al3+0.05)]}O2. The maximum value for x of 0.03127 detected by XPS (Figure S3, Supporting Information), for example, then corresponds to an excess metal ion concentration of 0.788% right in the middle of the pellet. The question now becomes how might this very small amount of excess cations be locally incorporated into the rutile type average structure? The way this is done must necessarily involve at least face-sharing M2O6 (M for a metal or cation) octahedral units given that the rutile structure type represents the maximum filling of the pseudo-octahedral cation sites in a close to hcp anion array without any such face-sharing octahedral units (see e.g. Figures 6d and e). Indeed, the insertion of a single isolated interstitial Al3+ ion into the normal rutile structure (see Figure 6g) automatically leads to three consecutive face-sharing M3O9 structural units and to neighboring O2 and O3 ions that are now 4 coordinate rather than the usual 3 coordinate (see Figure 6g) and hence lead to very overbonded O2 and O3 ions, particularly in the latter case as a result of the distorted AlO6 octahedral shape in the unrotated ideal rutile structure. (The apparent valence, AV, of the O3 ions was calculated as 2.93 valence units using softBV,6 severely overbonded). Furthermore, the addition of a single interstitial Al3+ ion simultaneously requires 3 of the 4 neighboring Ti3 and Ti4 ions to reduce from Ti4+ to Ti3+ in order to maintain charge balance, also leading to significant overbonding (the calculated AV of the neighboring Ti3 and Ti2 ions calculated as AV(Ti3) = 3.626 and AV(Ti2) = 3.628, again severely overbonded). Clearly, the initial local cluster shown in Figure 6g, corresponding to an isolated interstitial M3+ ion, is highly energetically unfavorable. The situation can be significantly improved via local octahedral rotation of the two Ti3O6 and 1 Al1O6 octahedra that are face-sharing (Figure 6g) in the

unchanged as the Ti3+ to Ti4+ concentration increases upon thinning, as determined by EPR and XPS (Figures S2 and S3, Supporting Information), it was at first expected that a proportionate amount of compensating oxygen vacancies would also need to be introduced for the purposes of charge balance as follows: [Ti4+0.90‑xTi3+x][Nb5+0.05Al3+0.05]O2‑x/2. The maximum value for the Ti3+ to Ti4+ concentration, in the middle of the pellet, was determined by XPS (S3) to be x/(0.90 − x) = 0.036, corresponding to x = 0.03127. As Raman spectroscopy is known to be sensitive to local chemical bonding and, in particular, to variations in the local environment, e.g., of oxygen ions around the Ti ions in TiO6 octahedra, it is employed here to probe for the presence or otherwise of oxygen vacancies. Figure 5 shows Raman spectra collected on pellets of different thicknesses. Given that both the Eg and A1g modes in oxygen-deficient, TiO2‑x rutile are known to be very sensitive to the oxygen ion concentration,19,20 it was initially surprising that (1) no changes in peak positions or full width at half maxima (fwhm) indicative of oxygen vacancies that could be observed (Figure 5a), especially for the two strongest Eg and A1g modes21 at ∼445 cm−1 and ∼618 cm−1, respectively (Figure 5b); and (2) the Raman peak positions of the Eg and A1g modes of (Nb+Al) codoped TiO2 are nearly the same as those of undoped rutile nanopolycrystals, single crystals,21,22 and polycrystalline rutile microspheres (see Figure S4, Supporting Information).23 For instance, when the ratio of [O]/[Ti] in a slightly oxygendeficient rutile decreases from 2.0 to 1.99, 1.98, and 1.97, the position of the Eg mode systematically shifts toward lower wavenumbers, from 447 cm−1 to 443, 437, and 433 cm−1, respectively.24 Such shifts, however, are clearly not observed for the current (Nb+Al) sample, in contrast with the behavior observed for the (Nb+In) codoped samples (Figure S5, Supporting Information). This strongly suggests that (Nb+Al) codoped TiO2 behaves quite differently from the (Nb+In) codoped case, at least as far as defect chemistry is concerned. In 4938

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Figure 6. (a) Narrowest structural element of the corundum (Al2O3) structure type in projection along [100] rutile. (b and c) Single layer (perpendicular to the projection axis direction) of the Al2O3 structure type along the (b) [110] direction, (c) the equivalent [-1,-3,2] direction for Ti4O7. The red plane in b is the (1,-1,2) plane of the corundum structure. Panel d shows the single layer (perpendicular to the projection axis direction) of the TiO2 rutile structure along the [100] (top) and [001] (bottom) directions, respectively. (e) Single layer (again perpendicular to the projection axis direction) of the TiO2 rotated rutile structure. The relationship of the rotated rutile structure to the ideal rutile structure is obtained by rotating the latter TiO6 octahedra ∼ ±12.3° around the ideal rutile c direction in a co-operative fashion (cf. the bottom part of e with that of d). Note that such a rotation converts the originally buckled oxygen atom (010 lattice planes in d into an hcp oxygen array in e). The red planes in a, b, and c correspond to the (011) plane of the rotated rutile structure. (f) Proposed (011) twin boundary. (g) Isolated Al interstitial inserted into unrotated rutile in projection along [100]. Note that the neighboring O2 and O3 oxygen ions are now 4 coordinate rather than the usual 3 coordinate.

manner shown at the bottom of Figure 6d to e. Figure 6c shows an example of how this can happen locally where face-sharing in pairs takes place. This leads to a strong driving force for the interstitial ions to cluster along the {011} planes of the ideal rutile structure type as shown in Figures 6a,b,c and f. While we cannot rule out isolated interstitials as shown in Figure 6g, it is considered more likely that defect clustering will eventually lead to {011} twin boundaries as shown in Figure 6f. If this is the case, then the M2O6 face sharing octahedral units are most likely either M3+2O6 units, i.e., Ti3+Al3+O6 units, or M3+3O9 units, i.e., Ti3+2Al3+O9 units; see Figure 7f. Under the former assumption, the stoichiometry of the 10% (Nb+Al) codoped TiO2 sample at x = 0.03127 can be written in the following form:

[Ti4 +0.8756(Ti, Al)3 +0.0740 Nb5 +0.0504 ]1 [(Ti, Al)3 + ]0.00788 O2 or 0.9764[Ti4 +0.8968(Ti, Al)3 +0.0516 Nb5 +0.0516]1 O2 + 0.0158{(Ti, Al)3 +2 O3}

That is, the overall stoichiometry suggests that the average rutile structure type might well be very occasionally intergrown with structural elements of the corundum structure type, (Ti,Al)2O3 (see Figure 6b) or other intermediate, Al-rich phases such as the monoclinic AlTi7O12 structure proposed in Figure 7f below. Note that this proposed structure consists of alternating (011)r slabs of rutile (in blue) and Al-rich (011)r slabs and hence might well intergrow with the rutile structure type itself. 4939

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Figure 7. (a−c) EDP’s of 10% (Nb+Al) doped rutile TiO2 taken from the (Al,Ti)3+-rich middle of the sample pellet along (a) [−103]r, (b) [0−14]r, and (c) [1−11]r indexed with respect to the rutile structure type (subscript r). Note the clear presence of weak G ± 1/3[011]r* type satellite reflections (arrows in a and b but not present in c), where G represents the Bragg reflections of the underlying rutile average structure type. EDP’s of this type are likely associated with (Al,Ti)3+-rich regions and were only found infrequently. (d and e) Simulated EDP’s of a proposed AlTi7O12 pseudomonoclinic, (Al,Ti)3+-rich, intermediate phase (see f), showing clear evidence of the same G ± 1/3[011]r* satellite reflections, where G represents the Bragg reflections of the underlying rutile average structure type. The green octahedra and blue octahedra in f denote AlO6 and TiO6 octahedra, respectively.

The narrowest such structural element is shown in Figure 6a and consists of corner-connected M3+2O6 units running along an (011) plane of the rutile structure type. (The two octahedra per M3+2O6 unit in Figure 6a are drawn in distinct colors to show that the rutile blocks on either side of such an (011) boundary (see also Figure 6f) are not in the same orientation as one another. They are in fact related by a b glide symmetry operation, {mx|1/2b} in the current setting). Note that this narrowest basic structural element occurs in both the M2O3 corundum structure type (cf. the red lines in Figures 6a and b) as well as reduced rutile structures such as Ti4O7 (shown by the red lines in Figure 6c). If regularly spaced, such planar (011) twin boundaries (and the associated displacive relaxation) give rise to cation excess, rutilerelated superstructures, or more properly NiAs related superstructure phases (see e.g. Figures 6b,c and also Figures 7a,b). If isolated (Figure 6f), they give rise to diffuse streaking along the [011]* directions of reciprocal space (see e.g. Figure 7c), as shown by fast Fourier transforms of high resolution lattice images taken around (011) twin boundary regions (Figure 8). It is intriguing that the boundary shown in Figure 8a appears significantly thicker than that shown in Figure 8b and may indicate that the initial (011) twin boundary acts as a nucleation site for a rutile, or more correctly, the NiAs type related superstructure phase. While no evidence could be found for the presence of Al2O3 or Ti2O3 grains from either X-ray powder diffraction or from electron diffraction, clear evidence was found for a closely related, monoclinic Pb11 (am = ar, bm = br − cr, c = 3cr; am* = ar*, bm* = br*, and cm* = 1/3(br* + cr*); m for monoclinic, r for rutile) intermediate structure type via electron diffraction (see e.g. Figures 7a and b). Note that Gr ± 1/3[311]* ≡ Gr′ ± 1/3[011]* in Figure 7a and that Gr ± 1/3[341]* ≡ Gr′ ± 1/3[011]* in Figure 7b. It is intriguing that the measured reciprocal space dimensions of the strong Bragg reflections, |Gr|, in Figures 7a and b, were the same as those determined using the ideal rutile type lattice parameters to within measurement accuracy, ∼1%. This suggests that the intermediate structure type may be able to intergrow coherently with the rutile itself, presumably along (011) lattice

Figure 8. (a and b) Typical high resolution lattice images of typical {011} twin boundaries. (c and d) Corresponding fast Fourier transforms showing diffuse streaking along the ⟨011⟩* direction of reciprocal space. The d-spacings of the 011 and −101 lattice fringes in a and b are 2.492 Å, while that of the 110 lattice fringes is 3.251 Å.

planes (see Figure 7f). The same intermediate structure type has indeed previously been reported as occurring at the boundary region between a rutile matrix and an alumina precipitate in a lightly alumina doped rutile and to grow coherently on the rutile matrix.25,26 Likewise, Putnis27 has reported the existence of the same intermediate transition phase in the case of a lightly doped hematite (Fe2O3) phase. While no direct evidence of (011) intergrowth between the intermediate and rutile phases was found, (011) planar boundaries of variable apparent width giving 4940

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Chemistry of Materials diffuse streaking along the 011* directions of reciprocal space were found (see e.g. Figures 8a and b below). The above (011) planar faults were found in the same grain. However, grains were found without any such planar faulting or associated diffuse streaking along [011]* suggesting that the distribution of the excess metal ions is not homogeneous from grain to grain and may require significant additional annealing at lower temperatures in order to homogenize. The latter sharp diffuse streaking has commonly been attributed to 1/2⟨0-11⟩ {011} “anti-phase boundaries” in slightly reduced, i.e., Ti3+-rich, TiO2‑x rutile rather than the twin domains suggested in Figure 6g. We acknowledge that this is a possibility but believe that twin boundaries as a mechanism to accommodate excess metal cations are more likely in the current case. (In either event, it is clear that the boundaries/defects are associated with reduced rutiles, i.e., with the presence of excess Ti3+ ions.) If so, the excess Al3+ and Ti3+ ions may indeed be incorporated into the dominant rutile type average structure in the form of relatively narrow platelets or Guinier Preston (GP) zones (coherently intergrown microstructures within the parent rutile matrix) of the above intermediate type structure, as suggested by Blanchin et al. in the Al2O3-doped rutile.25,26 The combination of a bond valence approach with electron diffraction revealing tripling along the ⟨011⟩* axes supports the presence of GP zones with Al3+ residing within coherently intergrown corundum plates in the rutile matrix. It therefore negates a model involving point defects because the substitutional Al3+ ions require moving significantly off-center in the octahedral coordination site, away from the oxygen vacancy and toward other oxygen ions to increase their bond valence. This would have particularly dramatic effects on the apical oxygen in the remaining square pyramid and would lead to an unstable structure. A nonpoint defect model suggesting the assumption that the origin of oxygen vacancies in the (Nb+Al)-TiO2 system is simply to compensate for charge imbalances resulting from the reduction of Ti4+ is also implausible, consistent with Raman results. Consequently, all of the local and macroscopical evidence direct to the fact that no oxygen vacancies exist in the (Nb+Al)TiO2 system. Returning to the dielectric properties and variable composition with thickness, Figure 9 schematically demonstrates the observed compositional gradient (Figure 9a) as well as the concentration variation of the Ti3+ species (Figure 9b) along with the homogeneous distribution of the O2− species (Figure 9c). It is thus interesting to approximate the composition of the gradient defect states. Assuming the thickness of the gradient layer is d0 and that there is a homogeneous inner layer with the thickness d1, one can estimate the composition from the outer 4+ 3+ surface to the bulk (dt) to be [Nb5+]0.05Ti4+ 0.85 [Ti1 − x Tix ]0.05 3+ [Al ]0.05 O2−Δ(x = (dt/d0)). It should be especially noted that Δ is a variable associated with the local intermediate state rather than a normal oxygen vacancy as there are no oxygen vacancies in the present case. In general, d1 could be relatively large or very small (e.g., marked by the rectangle in Figure 9b) and could even be zero (e.g., marked by the triangle in Figure 9b), depending on the preparation conditions (e.g., temperature, duration, original pellet thickness, and also the oxygen partial pressure28 etc.). Here, for the 10% (Nb+Al) codoped case, d1 is roughly estimated to be 0.8 mm (or smaller) while d0 is estimated to be 0.4 mm (or larger). The dielectric properties are simultaneously affected by dt, d0, and d1. The existence of dt (equivalently, the compositional gradient layer) would gradually increase the potential energy of charge transport, effectively maintaining the large dielectric

Figure 9. Compositional gradient model: (a) Scheme for a crosssectional view of a pellet, and schemes for (b) the gradient in the concentration of the Ti3+ species and (c) the constant oxygen concentration. d0 denotes the thickness of the gradient layer. d1 denotes the thickness of the homogeneous inner layer. dt, perpendicular to the surface, denotes the distance from the outer surface to the location under consideration. The arrows in b and c denote the concentration change (CTi3+ and CO2−) with thickness. Panel b demonstrates the variation of the Ti3+ concentration (CTi3+) with thickness where the trapezoid-shape (dashed blue) means that CTi3+ could reach a constant value if there is a homogeneous inner layer with thickness equal to d1 (patterned rectangle), while if there is not (i.e., d1 = 0), the triangle (solid red) would denote the gradient in the Ti3+ species.

permittivity and lower energy dissipation. A colossal dielectric response induced by a compositional gradient layer near the materials surface should not be simply classified as a conventional SBLC effect, well-known to form a Schottky barrier within several to tens of nanometers near the pellet surface. The lower free carrier concentration (free electrons induced by the reduction of Ti4+) in the chemical compositional gradient region would significantly broaden the width of the Schottky barrier between the electrode and the dielectric material, thereby lowering the SBLC as well as the dielectric loss. This, therefore, opens a route to balance optimal dielectric permittivity and sufficiently low dielectric loss.

4. SUMMARY AND CONCLUSIONS In summary, the synthetic optimization of (Nb+Al) codoped TiO2 rutile ceramics with the compositions Ti0.995(yNb +zAl)0.005O2 where [z = (4−5y)/3] with variable y = 0.4, 0.5, 0.6, 0.7 and Ti0.9(0.5Nb+0.5Al)0.1O2 sintered in the temperature range from 1300 to 1500 °C has been carried out. All of the samples can exhibit colossal dielectric permittivity. Their dielectric permittivities do not change significantly with different electrodes or DC biases. This suggests that even though the SBLC effect dominates the polarization mechanism that contributes to the observed colossal permittivity, the appearance of the Ti3+ concentration gradient layer with the absence of oxygen vacancies results in a free-electron-poor surface region (with good insulating properties) and an inner free-electron-rich region. Such a surface region blocks further diffusion of the inner free electrons and thus naturally balances both the dielectric permittivity and dielectric loss parameters. The absence of oxygen vacancies in this system is attributed to a locally intergrown, intermediate, metal rich structure growing coherently on the rutile matrix. Such a compositional gradient model opens a new compositional gradient design approach for dielectric ceramics and films. 4941

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Chemistry of Materials

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ASSOCIATED CONTENT

S Supporting Information *

Example complex impedance spectra, Raman, XPS, and EPR. The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.5b01351.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS W.H., K.L., Y.L., and R.L.W. acknowledge the support of the Australian Research Council (ARC) in the form of ARC Discovery projects, and of the Australian National University Connect Ventures in the form of the Discovery Translation Fund. Y.L. also appreciates support from the ARC Future Fellowships program. We thank Dr. Paul Smith for the EPR analysis.

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REFERENCES

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DOI: 10.1021/acs.chemmater.5b01351 Chem. Mater. 2015, 27, 4934−4942