Synthesis and Characterization of a New Ferroelectric with Low Lead Content, a High Curie Temperature, and a High Piezoelectric Response

A new solid solution (1−x)Bi(Fe2/8Ti3/8Mg3/8)O3–(x)PbTiO3 (BFTM‐PT) is synthesized and the electromechanical properties are measured. This system is defined as a low‐lead material with ferroelectric/piezoelectric behavior and a morphotropic phase boundary (MPB) that leads to enhanced properties. The MPB is located between x = 0.30 and 0.35 and coincides with a structural phase transition and a sharp increase in the piezoelectric response. The system demonstrates ferroelectric hysteresis where x = 0.325 displays the best properties with a maximum polarization of 39 µC cm−2 and a remnant polarization of 26 µC cm−2. The range of compositions has high Curie temperature (Tc), ranging from 625–650 °C. Materials with a Tc above 400 °C typically have a low d33 of <50 pC N−1 at room temperature. However, BFTM‐PT has a higher d33 that most other compositions with a Tc in this range, with the highest being 145 pC N−1 for x = 0.375. The d33 drops off above 100 °C, but doping studies can be done in the future to stabilize the piezoelectric response at higher temperatures. These outstanding properties open the possibility of new transducer applications, in particular ones requiring high temperature and high power.


Synthesis and Characterization of a New Ferroelectric with Low Lead Content, a High Curie Temperature, and a High Piezoelectric Response
Thomas Rowe, Brooke N. Richtik, and Michelle Dolgos* DOI: 10.1002/aelm.202200910 machinery, vehicles, or human movement to then power devices such as wireless sensor networks in remote locations, wearable biomedical devices, or mobile electronics. The most highly utilized piezoelectric material is currently the perovskite Pb(Zr 1−x Ti x )O 3 , (PZT), due a morphotropic phase boundary (MPB) that causes an enhancement of d 33 , the piezoelectric response at x ≈ 0.5, and its ability to be easily doped to tune the material for a wide range of applications. [3,4] However, PZT contains 60 wt% lead, which is a toxic element and in addition, PZT has a low operating temperature due to a phase transition at its Curie temperature (T c ) of 350 °C, preventing it's use in high temperature applications such as structural health monitoring in harsh conditions, pressure sensors and other high temperature, high power transducer applications.
This study replaced the end member PbZrO 3 (PZ) with BiFe 2/8 Ti 3/8 Mg 3/8 O 3 (BFTM) where the goal was to increase the operating temperature of piezoelectric devices. BFTM is unique in that it is one of just three perovskites with only bismuth on the A-site that can be synthesized under ambient conditions. [5] Generally, the Bi-cation is too small to stabilize the perovskite phase when it is the sole A-cation and therefore can only be synthesized at high pressures. However, in BFTM, the B-site cations have +3, +4, and +2 oxidation states respectively, creating a flexible O-framework that more easily allows the Bicoordination environment to be satisfied. BFTM was chosen to replace PZ in the PZT solid solution because BFTM has a high T c of 730 °C so with PT having a T c of 495 °C the T c of the resulting solid solution is expected to fall somewhere in between its end members. This contrasts with PZ, which has a lower T c than PT, resulting in a phase transition temperature of ≈350 °C at the MPB of PZT. Additional benefits of using BFTM as an end member is that it decreases the amount of lead in the material. Bismuth, like lead, also has a lone pair of electrons which helps to maintain long range ferroelectric ordering, rather than shifting to a relaxor ferroelectric with short range polar nanodomains, as is often observed when making substitutions with a variety of cations that prefer different coordination environments. [6] In this study, we synthesized solid solutions of (1−x)BiFe 2/8 Ti 3/8 Mg 3/8 O 3 − (x)PbTiO 3 , (BFTM-PT) which has a morphotropic phase boundary, a much

Introduction
Piezoelectric materials convert between mechanical and electrical energy and vice versa so they are used in a plethora of electronic devices for sensor and actuator applications including cell phones, sonar equipment, engine knock sensors, pressure sensors, diesel fuel injectors, medical devices, and many more. [1,2] One of the developing applications is to use them as an energy harvesting device to capture the energy from natural repetitive motions generated by vibrations in www.advelectronicmat.de lower wt% of lead (22% in BFTM-PT compared to 60% in PZT), a higher phase transition temperature of ≈650 °C for use in high temperature applications, and high d 33 value compared to most materials with a similar T c .

Results and Discussion
X-ray diffraction measurements indicate this material to be a solid solution with no visually distinct impurities in any of the compositions synthesized. The diffraction patterns can be seen in Figure 1, and they show that there are multiple phase transitions present. Figure 1b shows the full range of structural phase transitions. The Bi-rich side starts off in rhombohedral (R) R3c, which is the structure of BFTM, and is indicated by splitting of the (104) and (210) peaks as well as the (202) and (006) peaks. As more PT is added, a mixture of the R and monoclinic (M) phase occurs. The space group is either Cm or Cc, but these cannot be differentiated with X-ray data alone due to the distortions in the oxygen positions. Neutron data has been collected and is being analyzed for a future study. A full structural description of this phase will be reported in a follow up paper. Between x = 0.30 and 0.35, a single M phase can be observed. Scanning Electron Microscopy (SEM) images taken within this compositional range also do not indicate a secondary phase as seen in Figure 2. The histogram of the SEM data shows a single size distribution among grains present, along with no visually obvious differences in grain shape. If an impu-rity or secondary phase were present in large amounts, a bimodal distribution of grain sizes would be produced as in lanthanum doped NBT studied by Eaksuwanchai et al. [7] The predominant grain size in BFTM-PT is slightly less than one micron which is on the small side for traditionally sintered ceramics, but can be explained through extensive ball milling of the starting materials and the utilization of the two-step sintering process, which has been used for high density and small grain growth in ceramic materials. [8] On the Pb-rich side of this structure, two phases are again detected, the M phase and tetragonal (T) P4mm, then as the amount of PT increases, the structure changes to single phase P4mm, which is noted by the splitting of the (101) and (110) peaks. The polarization direction varies in each of these structures going from the [111] direction in R3c on the Bi-rich side and the [001] direction in P4mm on the Pb-rich side of the MPB. The M structure has a polarization of [110] which provides a bridge between the R and T phases, and is often referred to a polarization rotation. [9,10] This structural sequence of phase transitions is somewhat similar to PZT where the PZ end starts in the antiferroelectric, orthorhombic Pbam phase, then changes to R3c before reaching the MPB where monoclinic Cm is observed, and finally transforms to P4mm on the Pb-rich side of the MPB. [3] Because the transition sequence in BFTM-PT and PZT is similar, the structural data of BFTM-PT suggests that the monoclinic structure found between 0.30 and 0.35 could be an MPB, and therefore a region of the phase diagram with enhanced piezoelectric and ferroelectric properties. While BFTM-PT and PZT are structurally similar, www.advelectronicmat.de their dielectric, ferroelectric, and piezoelectric properties are different, especially in terms of the Curie temperature, which is rather low in PZT at ≈350 °C and therefore limits its use in high temperature applications.
Electrical properties of BFTM-PT were measured on dense ceramic pellets (>95%). Figure 3 shows the temperature dependent dielectric permittivity and loss data where a sharp phase transition from a ferroelectric to paraelectric state is observed between 630 and 650 °C depending on the composition. These T c values lie in between those of the end members BFTM at 730 °C and PT at 495 °C. Table 1 shows a summary of the permittivity and loss data for the compositions x = 0.25-0.40. There is a trend in the room temperature permittivity at 1 kHz where x = 0.25 on the Bi-rich end has a value of 357, which increases to 758 at x = 0.35 then decreases to 598 at x = 0.40. The loss values at this same temperature and frequency vary from 0.029-0.096 depending on the composition. They show no trend that corresponds to the permittivity data, but these loss values are typical for ferroelectric materials. At higher temperatures near T c , the loss increases significantly and varies from 0.158-0.295. This large increase indicates that the material is likely conductive at high temperatures, which is due to the volatile A-cations and/or the tendency for iron to change oxidation state, leading to A-site or oxygen vacancies, therefore increasing the conductivity at high temperatures. These dielectric properties contrast with PZT, which has a T c of 350 °C. PZT can be easily tailored and shows a range of T c values from 300-400 °C through doping, [11,12] but it still cannot maintain the polar structure up to the T c values found for BFTM-PT. The T c of BFTM-PT can also be compared to other well-known PZT competitors. One well studied lead-free piezoelectric K 0.5 Na 0.5 NbO 3 has a T c that falls between that of PZT and BFTM-PT and ranges from 350 to 450 °C, depending on the dopant incorporated into the structure. [13,14] Other ferroelectrics such as BiFeO 3 -BaTiO 3 (BF-BT) solid solutions [15][16][17][18][19] and the family of bismuth Aurivillius phases (Bi 2 O 2 )(A n−1 B n O 3n+1 ) (where n = the number of perovskite blocks between bismuth oxide layers) have T c values that range from 550-750 °C which is similar to what is observed in BFTM-PT. [20][21][22] The polarization (P) and strain (S) were also measured as a function of electric field (E) on dense ceramic pellets to determine the ferroelectric and piezoelectric properties, respectively.

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A summary of these properties can be found in Table 2. Figure 4a shows the composition in the R region of the phase diagram (Bi rich end) does not have a saturated P−E loop and is similar to what is observed in the parent compound BFTM. [5] This loop is common for ferroelectric materials that have a high T c and/or high coercive field. When moving into the 2-phase region of R and M phases and through to the T phase on the Pb-rich side, all those loops became saturated, indicating an alignment of the dipoles. Table 2 summarizes the ferroelectric and piezoelectric properties. It shows that both the maximum polarization (polarization at saturation), P m , and remnant polarization (polarization that remains after the field is removed), P r increase up through x = 0.325, then slightly decreases as more PT is added. The values for P m and P r at this composition are 39 and 26 µC cm −2 , which are on par with what is found in PZT. [11,23] The coercive field (E c ) of BFTM-PT is fairly high, varying between 38 and 54 kV cm −1 , making these materials hard ferroelectrics. Hard ferroelectrics more readily retain long range ordering and are more difficult to switch, making these materials suitable for applications that require high levels of electrical excitation and/or mechanical stress such as high voltage or high power generators and transducers. [1,24] Undoped PZT shows a lower E c of ≈20 kV cm −1 , which is expected because materials with a lower T c typically have a lower E c value as well. [25] Doping can change the E c of PZT in addition to its T c where soft dopants, consisting of electron donors decrease the E c and hard dopants consisting of electron acceptors increase E c . This easy tunability of PZT has been developed over decades through the use of dopants (hard vs soft) and mechanical synthesis techniques, which implies that these techniques could also be utilized to alter the properties of BFTM-PT and tailor it for specific applications. [26][27][28][29] Both bipolar and unipolar loops, S-E, were measured to determine the electromechanical strain properties as seen in Figure 4b,c. Figure 4b shows bipolar strain versus electric field loops. These loops demonstrate ferroelectric back-switching as indicated by a negative strain component and they also show an increased strain percentage within the MPB region. The strain response was also studied under unipolar electric field cycling to investigate its potential use for actuator applications. (Figure 4c) The highest strain achieved is 0.19% in the x = 0.325 composition, which is consistent with the observed excellent dielectric and ferroelectric properties of this composition. This value is a large increase compared to the strain in other BFTM based materials (0.04-0.08%) [6,30] and even displays a slight increase compared to PT (0.15%) [11] indicating, along with the dielectric and ferroelectric data, the presence of an MPB in this region of the phase diagram. The strain values are comparable to some doped compositions of PZT (0.15-0.3% on average) [12] although due to the vast amount of research on strain engineering and dopants on the properties of PZT it can display a wide range of values. [11,[31][32][33] As mentioned previously, Bi-based Aurivillius phases show a T c similar to BFTM-PT, but their strain values are not typically measured due to the Aurivillius structure [22] only being a 2 dimensional ferroelectric and thus the traditional method of measuring the piezoelectric response along the c-axis usually results in values an order of magnitude smaller.
The indirect piezoelectric response at high field, d 33 * , was determined by calculating the maximum strain divided by the applied field within the unipolar S-E data while the direct piezoelectric response at low field, d 33 , was measured by poling the sample and measuring the response with a Berlincourt meter. As shown in Figure 5, there is an increase in both the high field and low field piezoelectric response around x = 0.325. At this composition, the d 33 value is 100 pC N −1 and the d 33 * value is 190 pm V −1 . In the monoclinic region, the d 33 and d 33 * diverge, with the d 33 * becoming significantly higher than d 33 which could potentially occur due to an electric field induced phase transition as in Bi 0.5 Na 0.5 TiO 3 -BaTiO 3 -K 0.5 Na 0.5 NbO 3 solid solutions. [34] E-field induced phase transitions have large d 33 * values and a suppressed d 33 . If this type of phase transition is indeed present within the monoclinic region of the phase diagram only, and the d 33 values are lowered because of this phenomenon, that could explain why the d 33 is the highest in the x = 0.375 composition. At x = 0.375, the d 33 and d 33 * values converge again which indicates that there would not be an E-field induced phase transition in this region. Another theory for why the d 33 of x = 0.375 could be higher than in the monoclinic region because this composition is mixed phased and therefore has a contribution from each phase.
The piezoelectric response, d 33 , is typically inversely proportional to T c , so materials with T c values around room temperature can have very high d 33 values-on the order of 600-1000 while materials with high T c (>500 °C), typically have a very low piezoelectric response, with d 33 values <50 pC N −1 . [20,21] The thermal depolarization of the piezoelectric constant for the x = 0.325 composition was also measured as seen in Figure 6. This measurement showed an onset in the thermal depolarization above 100 °C until the sample completely depoled at 500 °C. While it is expected that the material loses its piezoelectric response before the phase transition at T c , further work involving dopants can also be done to raise the onset while also lowering the dielectric loss. [35] The d 33 values in BFTM-PT are much higher than what is typically seen for materials with a similar T c as shown in Figure 7. The piezoelectric response in BFTM-PT is on par with recent high T c materials found in the literature like BF-BT, [15][16][17][18][19] but is significantly higher than the well-known piezoelectric Aurivillius phases such as Bi 4 Ti 3 O 12 and PbBi 4 Ti 4 O 15 which have d 33 values in the range of 10-30 pC N −1 . [20][21][22] The reason for the differences in the piezoelectric response between the Aurivillius materials and the perovskites BFTM-PT, BF-BT, BF-PT, [36,37] BScO 3 -PT [38] is that the Aurivillius materials are not a solid solution, so there is no MPB present to enhance the piezoelectric response as in these bismuth-based perovskite materials. Therefore, in order to maximize the piezoelectric response in high T c materials, it  This study is not the first to use BFTM as an end member in a solid solution to try to mimic or even improve upon the properties of PZT. The first solid solution synthesized was BFTM-BaTiO 3 (BFTM-BT) [6] which like BFTM-PT has R3c and P4mm end members, which represent the structures of PZT on either side of its MPB. While the solid solution of BFTM-BT has end members with the same structure as BFTM-PT, the electromechanical properties are different. In BFTM-BT, the structure changes from R3c to a pseudocubic R3m structure that extends from x = 0.95-0.05 so there isn't technically an MPB which typically occurs over a very narrow composition range. The properties of BFTM-BT are in between that of a ferroelectric and a relaxor ferroelectric. Relaxors are a subset of ferroelectrics where the long-range ferroelectric ordering is disrupted by local disorder and forms regions of electric dipole ordering in polar nano domains (PNRs) where those PNRs are not correlated to each other. This phenomenon results in different properties than traditional ferroelectrics and therefore are used in different applications such as energy storage. [39,40] One notable feature of BFTM-BT is the permittivity as a function of temperature. The data show two peaks, one broad peak at lower temperatures and a second sharper peak at a higher temperature. The study suggests that the high temperature peak is from a phase transition while the low temperature peak is from the dielectric relaxation due to the reorientation of the polarization direction within the PNRs caused by thermal fluctuations. In addition, the ferroelectric and piezoelectric properties of BFTM-BT are very poor. The origin of this behavior in BFTM-BT comes from the different preferred cation environments of bismuth and barium. Bismuth has a lone pair of electrons, resulting in a distorted polyhedral environment where barium prefers a symmetric coordination environment, thus causing a structural strain which disrupts the long-range ferroelectric ordering and causes the onset of relaxor behavior.
The next study on BFTM investigated the ternary solid solutions between BFTM-LaFeO 3 -La(Mg0.5Ti0.5)O 3 (BFTM-LF-LMT). [41] LF was used as it has been known to improve the piezoelectric properties of BiFeO 3 , however iron often results in lossy dielectric behavior, so LMT was added to counterbalance the increased loss due to the presence of iron. In this phase diagram, several compositions displayed a new ferroelectric phase with the Pmc2 1 structure, which is rarely found as a perovskite. This space group has a complex set of displacements, but overall shows a ferroelectric distortion along the [110] direction. These solid solutions did display ferroelectric and piezoelectric properties, but the d 33 values were very small at ≈0.25 pC N −1 . Unfortunately, between the ferroelectric phase of the BFTM end member and the new Pmc2 1 structure, there is a range of compositions where phase separation occurs and results in a mixed phase perovskite and Aurivillius structure, so as in BFTM-BT, there is no MPB present with an enhancement of the piezoelectric response.
Solid solutions between BFTM-CaTiO 3 (BFTM-CT) were able to successfully bridge the gap between the ferroelectric phases observed in the BFTM-LF-LMT system. [30] The BFTM-CT solid solution showed a structural transition from the parent R3c structure to a wide mixed phase region (x = 0.05-0.20) of R3c and orthorhombic Pna2 1 in what can be defined as the MPB. On the CT rich side of the MPB, the structure was single phase Pna2 1 . In this solid solution there is no bridging polarization direction as in PZT and BFTM-PT and the polarization changes from the [111] direction in the R phase directly to the [001] direction in the O phase. The electromechanical properties show that the mixed phase region can be considered an MPB. The ferroelectric hysteresis loops were able to be fully saturated and show well developed loops with the best recording a maximum and remnant polarization of 49 and 44 µC cm −2 respectively. This can be compared to BFTM-PT with values of 39 and 26 µC cm −2 respectively. This region of the phase diagram in BFTM-CT displays an enhanced piezoelectric response with a d 33 value of around 50 pC N −1 , which is lower than that found in BFTM-PT, but BFTM-CT also has a higher T c of 840 °C, so

Figure 7.
A demonstration of the inverse relationship between piezoelectric effect and Curie temperature. Adapted with permission. [17] Copyright 2020, American Ceramic Society (ACERS), published by Wiley.
www.advelectronicmat.de the lower d 33 is not unexpected. The piezoelectric response in BFTM-PT is larger due to the inverse relationship between T c and d 33 as described above. The T c values within the BFTM-CT MPB are actually higher than that of BFTM because the other end member CaTiO 3 stabilizes in the Pnma space group which is centrosymmetric and therefore does not display ferroelectricity. By replacing non-ferroelectric CT, with ferroelectric PT, which also has a lower T c than BFTM, the piezoelectric response within the MPB of BFTM-PT is larger.

Conclusion
Overall, it appears that the crystal structure of the non-BFTM end member has little effect on the piezoelectric properties in all these BFTM systems. BFTM-BT and BFTM-PT both have end members with the same space groups as in PZT, but BFTM-BT forms a more of a relaxor-type ferroelectric because of the strain created by the large size difference and coordination environment differences between Ba 2+ and Pb 2+ . There is a phase separation between the ferroelectric structures in BFTM-LFO-LMT likely due to the increased amount of Fe 3+ , which typically distorts along the [111] direction, perhaps restricting the accessible ferroelectric structures in this solid solution. In BFTM-CT, the CT end member is isostructural to LFO, but Ti 4+ has been shown to displace locally along multiple directions that aren't necessarily coordinated to the A-cation, thus increasing the possible ferroelectric structures, which allowed for the formation of an MPB despite CT having a centrosymmetric structure. BFTM-PT and BFTM-CT both display MPBs despite Ca 2+ and Pb 2+ having different coordination environments. However, the d 33 value of BFTM-CT is much lower. This is partly due to the differences in T c but the chemistry could also be a contributing factor. Perhaps the different coordination environments of Bi 3+ which has a lone pair of electrons and Ca 2+ which does not, so while both solid solutions show ferroelectric ordering, on a local scale that might be beginning to break down. BFTM-CT is likely to have weaker long range ferroelectric ordering compared to BFTM-PT. This ordering is important to the enhancement of domain wall motion, which is an extrinsic mechanism of piezoelectric strain. [42,43] With Bi 3+ and Pb 2+ having similar preferred coordination environments, the strong long-range ordering can enable more domain wall motion, therefore increasing the extrinsic effects on the measured piezoelectric strain, resulting in a high d 33 .

Experimental Section
Synthesis: Samples of (1−x)BFTM-(x)PbTiO 3 solid solutions were prepared from x = 0.1-0.5 using standard solid-state methods. The precursor oxides and carbonates were combined in stoichiometric amounts and ground in an agate mortar and pestle. The precursors used were Bi 2 O 3 (Sigma Aldrich 99%), Fe 2 O 3 (Sigma Aldrich 99%), TiO 2 (Sigma Aldrich 99%), Magnesium carbonate hydroxide hydrate (C 4 Mg 4 O 12 · H 2 MgO 2 · xH 2 O) (Sigma Aldrich 99%), and PbCO 3 (Sigma Aldrich 99%). The mixed powders were first ball milled for 30 h at 350 rpm (15 min on, 5 min rest, and 15 min in the reverse direction) in a Planetary Ball Mill (Fritsch Pulverisette) with 25 mL of ethanol and eight, 10 mm diameter, yttria-stabilized zirconia balls. The resulting powder was used to sinter dense ceramic pellets. The powder was dried and mixed with a 3 wt% polyvinyl butyral (PVB, Butvar B-98, Sigma Aldrich) binder/ethanol solution. The powder was then pressed in a Carver uniaxial press using a die to form pellets with a diameter of 10 mm and thickness of ≈1.5 mm. The pellets were placed inside alumina crucibles with sacrificial powder of the same composition and sintered utilizing a two-step sintering process. This includes heating at 300 °C for 1 h with a 10 °C min −1 ramp rate as a binder burnout stage, then continuing to the first holding temperature of 1100 °C for 5 min and then ramping down also at 10 °C min until a second holding temp of 800 °C for 6 h after which the oven ramps down again at 10 °C min until room temperature. All physical properties were measured on pellets with densities greater than 95% of their crystallographic value as determined by a Mettler Toledo Archimedes kit.
Ceramic Preparation: Ceramic pellets for physical property measurements were carefully polished to appropriate thicknesses of ≈1 mm for dielectric measurements and <0.5 mm for strain (S-E) and polarization (P-E) measurements as a function of electric field. The pellets were polished to a mirror finish using a LaboPol-5 (Struers) with #400, #800, #1200, and #4000 SiC foils (Struers) sequentially. For S-E and P-E measurements, electrodes of silver conductive paint (SPI Supplies) were applied to the parallel faces of each pellet before the pellet was heated to 600 °C for 30 min to cure the silver electrodes. Platinum paint (SPI Supplies) was applied to the pellets used in the dielectric measurements and sintered to the polished pellets at 350 °C for 1 h with a ramp rate of 1 °C min −1 , then ramped at 5 °C min −1 up to 1000 °C for 1 h and then cooled to room temperature at 5 °C min −1 .
Physical Properties Measurements: Dielectric permittivity was measured using an HP 4192A LF Impedance Analyzer, a NorECsAS Probostat, and a Carbolite tube furnace. The measurements were performed on unpoled samples during cooling as a function of temperature and frequency from 1 kHz to 1 MHz and from 800 °C down to room temperature. Strain versus electric field (S-E) and polarization versus electric field (P-E) measurements were performed on a Radiant Precision Premier II, at a frequency of 1 Hz, with the sample pellet submerged in an insulating silicone oil. The S-E measurements were run using 2 loops, with the average being used as the final data set. The strain measurements utilized an optical displacement sensor (MTI-2100). Poling was performed on the Radiant Precision Premier II by applying a DC bias to the sample at 100 kV cm −1 for 30 min. The direct piezoelectric effect was measured on a Berlincourt-type d 33 meter (APC International, Ltd. YE2730A). The depolarization temperature measurements were performed on a poled pellet which was subsequently heated to and then held at each corresponding temperature for 30 min and then left to cool to room temperature before measuring the direct piezoelectric effect.
Scanning Electron Microscopy: The imaging was done on a Zeiss Σigma VP scanning electron microscope utilizing a working distance (WD) of 4.4 mm as well as utilizing the Extra High Tension (EHT) of 10 kV for higher resolution. The Variable Pressure Scanning Electron (VPSE) detector was also utilized in order to lower the amount of charging on the surface of the material. All scans were done on a sintered and dense ceramic pellet.
X-Ray Diffraction: Phase purity of each sample was verified using X-ray diffraction (XRD) on a Rigaku Miniflex 600 using Cu Kα radiation (λ = 1.541862 Å) and scanning a 2θ range from 10° to 60° at a rate of 5° min −1 . Synchrotron powder X-ray diffraction data was also collected at Argonne National Laboratory on the 11-BM beamline using their mail-in program. The 11-BM beamline utilized a 0.458165 Å wavelength scanning from −6° to 28° (2θ) with a step size of 0.001° and 0.01″ per step. Pawley refinements to determine lattice parameters were performed using Topas Academic software.

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Science by Argonne National Laboratory under Contract No. DE-AC02-06CH11357. The authors would like to thank Saul Lapidus for assisting with the synchrotron powder diffraction data collection through the mail-in program at the Advanced Photon Source and Hannah Dokken, the visiting Inorganic Chemistry Exchange student from University of Alberta for helping to prepare some of the samples for properties measurements.