Local Structure Adaptations and Oxide Ionic Conductivity in the Type

May 2, 2018 - However, doping with small concentrations of d0 transition metal cations such as ... (8,16,21,22) Type III has a commensurately modulate...
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Article Cite This: Chem. Mater. 2018, 30, 3387−3394

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Local Structure Adaptations and Oxide Ionic Conductivity in the Type III Stability Region of (1 − x)Bi2O3·xNb2O5 Julia Wind,† Neeraj Sharma,‡ Aleksey A. Yaremchenko,§ Vladislav V. Kharton,§,∥ Douglas A. Blom,⊥ Thomas Vogt,# and Chris D. Ling*,† †

School of Chemistry, The University of Sydney, Sydney 2006, Australia School of Chemistry, The University of New South Wales, Sydney 2052, Australia § Department of Materials and Ceramic Engineering, CICECO − Aveiro Institute of Materials, University of Aveiro, 3810-193 Aveiro, Portugal ∥ Institute of Solid State Physics, Russian Academy of Sciences, Chernogolovka 142432, Moscow Distr., Russia ⊥ NanoCenter & Department of Chemical Engineering and #NanoCenter & Department of Chemistry and Biochemistry, University of South Carolina, Columbia, South Carolina 29208, United States

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‡

ABSTRACT: Starting from a previously published stoichiometric model for the commensurate Type III phase in the (1 − x)Bi2O3· xNb2O5 system, Bi94Nb32O221 (x = 0.254), we have developed a crystal-chemical model of this phase across its solid−solution range 0.20 ≤ x ≤ 0.26. After using annular dark-field scanning transmission electron microscopy to identify the metal sites that support nonstoichiometry, we show that the maximum possible range of that nonstoichiometry is 0.198 ≤ x ≤ 0.262, perfectly consistent with the experimental result. Intersite cation defects on these sites provide some local coordinative flexibility with respect to the surrounding oxygen sublattice, but not enough to create continuous fluorite-like channels like those found in the hightemperature incommensurate Type II phase. This explains the reduced oxide-ionic conductivity of Type III compared to Type II at all temperatures and compositions, regardless of which phase is thermodynamically stable under those conditions. The solid− solution model shows that oxygen disorder and vacancies are both reduced as x increases, which also explains why Type III becomes relatively more stable, and why oxide ionic conductivity decreases, as x increases.

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INTRODUCTION The cubic δ phase of Bi2O3 shows the highest reported oxide ion conductivities of any known solid-state material (∼1 S cm−1 at 750 °C), largely due to the presence of 25% disordered vacancies in its fluorite-type oxygen sublattice.1−4 δ-Bi2O3 is only stable at temperatures between 750 and 825 °C, which rules out its practical use as an electrolyte in solid-oxide fuel cells (SOFCs). However, doping with small concentrations of d0 transition metal cations such as V5+/Nb5+/Ta5+, Cr6+/Mo6+/ W6+, or Tc7+/Re7+ has been shown to stabilize the average fluorite-type structure of δ-Bi2O3 to room temperature.5−14 While this reduces ionic conductivity by reducing the total number of oxygen vacancies and increasing the ordering of the remaining oxygen sublattice, it serves to increase temperature stability and eliminates major structural transformations. Particularly when codoped with rare-earth elements, such stabilized δ-Bi2O3 phases are of genuine interest for practical SOFC applications.15 The present study concerns itself with the Nb5+-doped system, i.e., (1 − x)Bi2O3·xNb2O5. This system contains a variety of fluorite-type δ-Bi2O3-related superstructures depending on x, the relative conductivities of © 2018 American Chemical Society

which can be rationalized in terms of the type of superstructure formed. The most intensively studied phase in the (1 − x)Bi2O3· xNb2O5 system is the solid−solution denoted Type II by Zhou et al.16 (0.062 ≤ x ≤ 0.25). Type II was first recognized by Miida and Tanaka17 as having a (3 + 3)-D incommensurately modulated structure which was further characterized by us8,18 and, more recently, solved and refined against single crystal neutron diffraction data19 in the cubic superspace group Fm 3̅ m (α,α,α)q00(α,−α,−α)q00(−α,α,−α)000 (number 225.3.215.8 using the notation of Stokes, Campbell, and van Smaalen20). Its complex but remarkably high-symmetry structure features displacive and compositional modulation of both anions and cations. The compositional modulation arises due to the ordering of Nb5+/Bi3+ and O2−/vacancies while the displacive modulation can be considered as a Coulombic consequence of the compositional modulation: O2− reacts to Received: February 26, 2018 Revised: May 2, 2018 Published: May 2, 2018 3387

DOI: 10.1021/acs.chemmater.8b00846 Chem. Mater. 2018, 30, 3387−3394

Article

Chemistry of Materials cationic charge by moving from its average fluorite-type position toward Nb5+ and away from Bi3+. At the Nb-rich end of the Type II solid solution, an intermediate phase referred to as Type III is found at lower temperatures.8,16,21,22 Type III has a commensurately modulated 3 × 3 × 7 superstructure of the δ-Bi2O3 subcell with tetragonal I4m ̅ 2 space-group symmetry and lower oxide ionic conductivity.23 The published Type III model provides a reasonable fit to highresolution synchrotron X-ray and neutron powder diffraction data; however, its fixed composition of Bi94Nb32O221 (x = 0.254) contradicts clear experimental evidence for a solid− solution range of 0.20 ≤ x ≤ 0.26.21 The exact range of this flexibility, and the crystal-chemical mechanisms that permit it, remain unclear and are addressed in this work. From a crystal-chemical perspective, both Type II and Type III can be understood in terms of the ordering of Nb dopant cations along ⟨110⟩F directions of the fluorite-type Bi sublattice to form corner-connected chains of NbO6 octahedra. Where these chains meet, they form tetrahedral clusters of NbO6 octahedra reminiscent of the basic structural unit of pyrochlore type, such that the Type II solid−solution can be thought of as representing a smooth transition from fluorite-type δ-Bi2O3 to a (hypothetical) pyrochlore-type Bi2Nb2O7 phase. In Type II, these ordered NbO6 chains delineate continuous, relatively disordered and oxygen-deficient δ-Bi2O3-like “channels” whose width depends on composition (Figure 1a), explaining its high thermal stability and oxide ionic conductivity (∼0.1 S cm−1 at 1000 °C).19 In contrast, Type III features noncontinuous NbO6 chains organized into layers or isolated groupings (Figure

1b).18,23 In Type III, oxide vacancies are all ordered in proximity to Nb cations, explaining its relatively lower oxide ionic conductivity at the same temperatures.19 Over the compositional range in which they coexist, the fact that the high-temperature Type II and low-temperature Type III phases have different Bi/Nb distributions as well as different O/vacancy distributions means that Type II can be easily quenched to low temperatures. Quenched metastable Type II samples will be referred to in the present work as Type II*. While oxide vacancy concentrations remain the same for Types II*, II, and III of a given Bi/Nb ratio,24 the reconstructive Type II* → III → II phase transitions on reheating these quenched samples have obviously detrimental implications for their use in SOFCs. A number of kinetics-based studies22,25 using conductivity and X-ray powder diffraction (XRD) of the Type II → Type III transition have set out to probe its structural subtleties. Type III phase was found to have a smaller unit cell volume compared to the Type II, and the Type III unit cell volume decreases as a function of x, implying that shrinkage occurs during conductivity measurements. Despite this shrinkage, it has been shown that Type III has higher electronic (σe−) and ionic (σion) conductivity than Type II at temperatures below 600 °C for x = 0.25. Interestingly, some measurements suggest that the ionic and total conductivity converge near the temperature limits; this would indicate a completely ionically conducting material.22 The activation energies are 1.14 and 0.82 eV for Type II and Type III, respectively, implying a lower energy barrier for conductivity in Type III at these temperatures. In both cases, the p-type σe− is shown to be the minor contributor to the total conductivity. However, when x = 0.25 at 820 °C, σe− of the Type II phase increases by a factor of 3−4, consistent with a transformation from Type II to Type III. Related work on the (1 − x)Bi2O3·xNb2O5 series where x = 0.05, 0.08, 0.10, and 0.16 26 reports a gradual decrease in oxide ionic conductivity as a function of temperature as the Nb content is increased, but limited details concerning the structure were presented. In the study reported here, we set out to definitively establish the range of the Type III solid−solution by consolidating previous experimental reports and new data and to develop a coherent crystal-chemical model explaining that range. We then used this crystal-chemical model to correlate changing structural features with physical properties, specifically, hightemperature oxide-ionic conductivity.

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EXPERIMENTAL SECTION

Polycrystalline samples of composition x = 0.20, 0.22, 0.24, and 0.26 in the series (1 − x)Bi2O3·xNb2O5 were synthesized by ball-milling stoichiometric quantities of Bi2O3 (Sigma-Aldrich, 5N) and Nb2O5 (Sigma-Aldrich, 4N5) for 24 h. The samples were then placed in closed Pt crucibles and heated in cycles ranging from 700−880 °C for 5−24 h. Phase purity at room temperature (Type II* or Type III) was confirmed by X-ray powder diffraction (XRD) using a Panalytical X’Pert Pro diffractometer with a Cu Kα source and by neutron powder diffraction (NPD) using the instrument D20 at the Institut LaueLangevin (ILL), Grenoble, France,27 with a takeoff angle of 120° and neutron wavelength λ = 1.88 Å. Figure 2 shows a Rietveld fit to the NPD data for the Type III phase at x = 0.26, using the structural model of Ling and Johnson.23 Aberration-corrected scanning transmission electron microscopy (AC-STEM) data in the Z-contrast high-angle annular dark-field (HAADF) imaging modes were collected from powdered samples. Dry powders were lightly ground in an agate mortar and pestle before

Figure 1. (a) Commensurate approximate (8 × 8 × 8 fluorite-type subunit cells) of the incommensurate Type II structure19 and (b) the Type III structure23 in the (1 − x)Bi2O3·xNb2O5 system, x ≈ 0.25. Bi atoms are green, NbOy polyhedra are blue, and O atoms are red. The Bi8 and Nb4 sites in the Type III structure (discussed in detail in the main text) are highlighted in slightly darker and lighter shades, respectively. 3388

DOI: 10.1021/acs.chemmater.8b00846 Chem. Mater. 2018, 30, 3387−3394

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minimize 60 Hz noise in the images, and a pixel dwell time of 15.8 μs was used. The total conductivity (σ) in air was studied by ac impedance spectroscopy (HP4284A precision LCR meter, 20 Hz−1 MHz, Agilent Technologies) using dense bar-shaped samples with porous Pt electrodes. Dilatometric studies were performed in air using an alumina Linseis L70/2001 dilatometer at a constant heating rate of 3 K min−1.

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RESULTS AND DISCUSSION Relating the Type III solid−solution range to its crystal structure requires first of all knowledge of Bi/Nb mixed site occupancies. Unfortunately, the complexity and pseudosymmetry of the Type III structure (50 crystallographically independent atoms in the unit cell of the highest-symmetry ordered model23) make it impossible to unambiguously determine them by Rietveld-refinement against XRD or even neutron powder diffraction data. We therefore used ACHAADF-STEM as the primary structural characterization technique in this study. Figure 3a shows a ⟨010⟩ = ⟨110⟩F zone axis HAADF STEM image of Type III (1 − x)Bi2O3·xNb2O5, x = 0.26. Figure 3b is a portion of the micrograph in Figure 3a along with a structural model ⟨010⟩ projection of the cations from the published structure.23 The labels refer to the crystallographic atomic sites in the published model. In the ⟨010⟩ zone axis projection of the Type III model structure, there are four classes of atomic cation columns: fully Bi, fully Nb, 1:2 Nb:Bi, and 2:1 Bi:Nb. We therefore expect to see differences in the HAADF signal due to the composition differences of these columns. It is evident from the micrograph that there are intensity variations for different columns of atoms in the unit cell. To quantify the HAADF signals for the different types of atomic columns in the ⟨010⟩ zone axis image, we used the integrated signal over the projected columns which has been shown to be relatively insensitive to small focus variations.28,29 To account for the experimentally observed thickness gradient in the micrograph, we subtracted a background signal selected from the nearby electron probe positions not associated with the atomic columns.30 For each type of atomic column in the ⟨010⟩ projected unit cell, an area with a radius of 7.5 pixels (131.26 pm) was selected and the signal integrated.

Figure 2. Rietveld fit to λ = 1.88 Å NPD data for the Type III phase of (1 − x)Bi2O3·xNb2O5 where x = 0.26, using the structural model of Ling and Johnson.23 Red crosses are experimental data, black line is calculated fit, and purple line below is the difference. The cell is tetragonal I-4m2, a = 11.5732(3), c = 38.6903(15) Å. Refinement statistics: Rp = 3.77%, wRp = 4.84%, χ2 = 1.37 for 30 parameters. loading onto holey carbon-coated Cu grids. A JEOL JEM2100F with a CEOS aberration corrector for the electron probe was used at 200 kV to image the sample with a 24 mrad convergence angle. Images were acquired using a Fischione Model 3000 detector with a camera length such that the detector spanned 75−178 mrad. The scanning acquisition was synchronized to the 60 Hz ac electrical power to

Figure 3. (a) HAADF-STEM images along the (a) ⟨110⟩ = ⟨010⟩F zone axis of Type III (1 − x)Bi2O3·xNb2O5, x = 0.26. (b) A portion of this micrograph along with a projection of the cations in the same direction, from the published structure,23 with labels referring to the crystallographic sites in that model. 3389

DOI: 10.1021/acs.chemmater.8b00846 Chem. Mater. 2018, 30, 3387−3394

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the mean integrated intensity observed for the most intense atomic column, Bi3/Bi7) to the expected values for the fully ordered model for Type III Bi94Nb32O221 (x = 0.254).23 A number of discrepancies are apparent. First, we note that in the crystallographic model derived from powder diffraction four columns are significantly offset in projection: Nb2/Nb5 (by 32 pm) Bi2/Bi6 (34 pm), Nb4/Bi6 (25 pm), and Nb5/Bi5 (29 pm). This should make them appear dimer than expected from a purely Z-contrast perspective due to electron channeling (a noncentrosymmetric intensity distribution along the channel direction).31 The Nb2/Nb5 column has an intensity consistent with the other fully Nb column, Nb1/Nb3, which itself has an offset in projection of 10.5 pm. The Bi2/Bi6 columns are in fact dimer than most of the other fully Bi atomic columns. Mixed Nb−Bi column Nb4/Bi6 is also dimer than the majority of the other mixed Nb−Bi columns, although the model structure implies that its composition is more Nb-rich than the other mixed columns and the atoms are offset in this projection, both factors which would tend to suppress the HAADF signal. Finally, the Nb5/Bi5 column exhibits a HAADF signal which is dimer than other columns with the same composition according to the model structure. Clearly, electron channeling in this orientation plays a role in the observed HAADF intensity. The data at this point are not sufficient to definitively decouple the effects of atomic column composition and channeling for atomic columns which have significant offset. We then note that the Bi7/Bi8 column (expected to be 100% Bi), which is not offset, is also too dim. Given that the Bi3/Bi7 and Bi5/Bi9 columns which have similar amounts of offset to the Bi7/Bi8 column are both equally brighter than the Bi7/Bi8 column, this suggests that Bi8 site has some mixed occupancy by Nb. If Bi8 was 100% Bi (fully ordered model), the cation stoichiometry would be Bi94Nb32 (x = 0.254); whereas if Bi8 were 100% Nb, it would be Bi92Nb34 (x = 0.270). This is inconsistent with the experimental Type III solid−solution range 0.20 ≤ x ≤ 0.26, indicating that some Nb sites from the ordered model must simultaneously support Bi occupancy. Noting that Nb2/Nb5 and Nb1/Nb3 are by far the dimmest columns in Table 1, Nb4 is the site most likely to support additional Bi occupancy. If Nb4 was 100% Bi, the Birich end of the solid solution becomes Bi102Nb24 (x = 0.190). AC-HAADF-STEM results concerning cation occupancy are therefore consistent with a solid−solution range 0.190 ≤ x ≤ 0.270. The fact that the experimental range is more restricted than this must then be due to the oxygen sublattice. In considering this, we note that every Nb5+ which replaces a Bi3+ increases the total O stoichiometry. The fully ordered model for Bi94Nb32O221 includes one O vacancy per formula unit to meet this requirement. However, every Nb5+ also increases the number of O vacancies required to achieve local NbO6/NbO4 coordination rather than the average 8-fold cubic coordination of Bi in fluorite-type δ-Bi2O3. Bi8 has cubic coordination, and its symmetry is consistent with tetrahedral coordination, so converting a Bi atom on this site into Nb increases the required number of O vacancies by 8 per formula unit, while also increasing the O stoichiometry by 2. This could only be accommodated up to the point of filling the one oxygen vacancy per formula unit in the fully ordered model of Bi94Nb32O221, i.e., Bi93Nb33O222 or x = 0.262. Higher Nb content on the Bi8 site would require losing Nb from another site (i.e., Nb4).

Figure 4 shows the mean HAADF background-subtracted integrated signal for each of the 14 different atomic columns in

Figure 4. Mean HAADF background-subtracted integrated signal for each of the 14 unique atomic columns in the ⟨110⟩ = ⟨010⟩F projection of the unit cell of Type III (1 − x)Bi2O3·xNb2O5, x = 0.26. Error bars are 95% confidence limits on the mean value.

this projection of the unit cell. Error bars are 95% confidence limits on the mean value. There are broadly four different signal levels observed in Figure 4. The brightest atomic columns are not surprisingly columns consisting of only Bi-containing sites: Bi3/Bi7, Bi5/Bi9, Bi1/Bi2, and Bi4/Bi10. The least intense atomic columns correspond to the two examples of atomic columns in the ⟨010⟩ projection that are fully Nb: Nb2/Nb5 and Nb1/Nb3. Because the HAADF-STEM image was not collected on an absolute scale, for convenience Table 1 compares Z-contrast ADF intensities in the ⟨010⟩ = ⟨110⟩F image (normalized by Table 1. Atom Columns in ⟨010⟩ Zone Axis Projection to the Fully Ordered Model for Type III Bi94Nb32O221 (x = 0.254),23 Fractional Nb Composition, Normalized ADF Intensities from Figure 2a, Offsets of the Atomic Sites in the Model Perpendicular to ⟨010⟩, and Agreement between Measured HAADF Intensities and the Model Composition atomic column

model frac (Nb)

Bi7/Bi8 Bi3/Bi7 Bi5/Bi9 Nb2/Nb5 Nb3/Bi4 Nb1/Bi10 Nb4/Bi1 Bi2/Bi6 Nb4/Bi6 Bi1/Bi2 Bi4/Bi10 Nb1/Nb3 Nb5/Bi5 Nb2/Bi9

0 0 0 1 1/3 1/3 1/3 0 2/3 0 0 1 1/3 1/3

normalized integrated intensity

a offset (pm)

c offset (pm)

agreement with model

± ± ± ± ± ± ± ± ± ± ± ± ± ±

0 −2.65 1.27 0 0 −2.53 5.07 0 0 9.10 −8.30 0 0 −1.73

2.66 2.66 4.13 32.35 −4.97 8.56 0.42 33.55 25.06 8.06 6.90 10.49 28.53 0.31

no yes yes yes yes yes yes no yes yes yes yes no yes

0.79 1.00 0.91 0.45 0.78 0.82 0.83 0.78 0.62 0.92 0.96 0.43 0.63 0.80

0.03 0.02 0.03 0.02 0.02 0.02 0.02 0.03 0.02 0.02 0.02 0.02 0.02 0.02

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Chemistry of Materials Converting a Nb atom on the Nb4 site into a Bi atom decreases the required number of O vacancies by 2.5 per formula unit, while also decreasing the O stoichiometry by 2. This gives a lower end to the solid−solution range at Bi102Nb24O213 (x = 0.190). However, the Bi atoms replacing 4-coordinate Nb would be significantly underbonded, especially with the concomitant reduction in O stoichiometry. Placing Nb on the Bi8 site would increase the O stoichiometry and relieve this underbonding, up to Bi101Nb25O214 or x = 0.198. This gives us a crystal-chemical model of Bi substitution on the Nb4 site, partially compensated for by a smaller degree of Nb substitution on the Bi8. This model predicts a solid solution range 0.198 ≤ x ≤ 0.262, perfectly consistent with the experimental range of 0.20 ≤ x ≤ 0.26.21 It also predicts the increased stability of Type III relative to Type II at higher x, observed experimentally:21 increasing x leads to replacement of the last remaining coordinatively flexible (partially disordered) Bi(III) site, Bi8, with coordinatively rigid Nb(V), while also filling in the last remaining oxygen vacancies. Type III is thus most ordered, and therefore most enthalpically favored, at higher x. Another significant feature of this model is that the compositional flexibility of the Type III phase is limited to intersite cation defects between only two sites: Bi8 and Nb4. These sites are highlighted in Figure 1b. From inspection of the structure, it is clear that while these cation defects will provide some local coordinative flexibility in the surrounding oxygen sublattice, they will not create continuous fluorite-like channels comparable to those found in the Type II phase.19 The model therefore predicts reduced oxide-ionic conductivity of Type III compared to Type II at all temperatures and compositions, regardless of which phase is more thermodynamically stable under those conditions. This is consistent with reduced oxide-ionic conductivity of Type III compared to Type II at all temperatures and compositions, regardless of which phase is more thermodynamically stable under those conditions. Conductivity measurements taken on cooling Type II samples at a range of compositions are shown in Figure 5.

Table 2 presents densities and activation energies for conduction. As the Nb content is increased, the oxide ion Table 2. Activation Energy for Total Conductivity of Type II (1 − x)Bi2O3·xNb2O5 Samples composition

density (g cm−3)

T (K)

x = 0.20 x = 0.22 x = 0.24

7.25 7.12 7.31

x = 0.26

8.31

910−1220 910−1220 1110−1250 910−1110 1090−1220 910−1090

EA (kJ mol−1) 118 127 155 129 157 130

± ± ± ± ± ±

1 1 1 1 2 4

conductivity is reduced and the activation energy exhibits a slight increase. Thus, increasing Nb content requires more energy for conduction. Two activation energy regions were found for the more Nb-rich samples. Higher temperature regions were characterized by higher activation energies. Conductivity measurements follow the trends described by Yaremchenko et al.26 for x = 0.05−0.16 (wholly Type II region in the phase diagram, see Figure 9 of Pirnat et al.21); i.e., as x increases (Nb content increases), conductivity decreases. The activation energies for samples 0.20 ≤ x ≤ 0.26 range from 118 to 157 kJ mol−1 (Table 2), higher than observed for 0.05 ≤ x ≤ 0.16 for which the range was 58−121 kJ mol−1.26 This can be ascribed to the change in Nb and O vacancy content, whereby increasing Nb content reduces the O vacancies by charge compensation. More energy is required to drive conduction and overcome energetic barriers imposed by increasing Nb content and decreasing O vacancies. Castro et al.32 report a lower activation energy of 85 kJ mol−1 for x = 0.25, but Takahasi and Iwahara3 report activation energies of 115 and 135 kJ mol−1 for x = 0.22 and x = 0.30, respectively. Our observations follow the same trend as the latter measurements; namely, increasing Nb content slightly increases activation energies. The result of Castro et al.32 appears to be an outlier, possibly related to differences in sample preparation and sample history, noting that the compositions in their study (like ours) involve two phases with significantly different conductivities (see below). At lower temperatures, a difference in conductivity between the Type II and Type III phases of composition x = 0.25 was also observed by Wang et al.22 To determine the differences in conductivity between the Type II/II* and Type III phases upon heating, ac impedance measurements were taken on heating an x = 0.24 Type II* sample to 860 °C. The sample was held at this temperature for 12 h to allow the Type II* → III transition to occur and then taken to 980 °C, followed by cooling. Figure 6 shows the results, highlighting a reduction in conductivity in the Type III region between 860 and 940 °C. These show that on cooling sufficient time was not allowed for the transition back to Type III. These observations were verified by an ex situ series of experiments with pellets sintered at appropriate temperatures (Figure 7). The Type II* → Type III transition is demonstrated by the splitting of the ⟨311⟩ cubic Type II* peak and its re-formation, showing the presence of the Type III phase between 860 and 940 °C. The Type III phase shows lower conductivity relative to the Type II phase, although their activation energies are similar, ∼155 ± 5 kJ mol−1 (see Table 2).

Figure 5. Temperature dependence of the total conductivity of (1 − x)Bi2O3·xNb2O5 samples in air. 3391

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that at temperatures below 830 °C this component is higher in Type III than Type II. The activation energy we obtain for the Type II phases of composition x = 0.24 and 0.26 below 830 °C is around 130 kJ mol−1, close to the previously reported value of 110 kJ mol−1 for x = 0.25 in this temperature region by Wang et al.22 Interestingly, their activation energy for the Type III phase is smaller, ∼80 kJ mol−1, similar to Castro et al.’s32 findings, suggesting that Castro et al. measured a Type III rather than a Type II phase. Thus, we conclude that at lower temperatures (