Grain boundary space charge engineering and ceramic processing for radiation detection in ion-conducting ceramics

Abstract

Disclosed herein are methods for setting the radiation sensitivity and operational temperature range of ion-conducting polycrystalline radiation detectors. Disclosed herein are material fabrication methods for setting the ion-conducting polycrystalline microstructure and the grain boundary ion blocking characteristics in ion-conducting polycrystalline materials. The methods may include introducing a selected precursor element into the grain boundaries of the ion-conducting polycrystalline material. These methods may allow for the device detection threshold, dynamic range, and / or operational temperature range to be adjusted for ion-conducting polycrystalline radiation detectors. The methods disclosed herein allow for the design of high-sensitivity and high-performance radiation detectors that are low cost, non-toxic, environmentally robust, scalable, and able to operate in harsh environments (e.g., high temperature, high humidity, pressure, and / or corrosive).

Description

GRAIN BOUNDARY SPACE CHARGE ENGINEERING AND CERAMIC PROCESSING FOR RADIATION DETECTION IN ION-CONDUCTING CERAMICS CROSS-REFERENCE TO RELATED APPLICATION(S)

[0001] The present application claims the priority benefit, under 35 U.S.C.119(e), of U.S. Application No.63 / 652,426, filed on May 28, 2024, which is incorporated herein by reference in its entirety for all purposes. GOVERNMENT SUPPORT

[0002] This invention was made with government support under 22CWDARI00046-01-00 awarded by Department of Homeland Security. The government has certain rights in this invention. BACKGROUND

[0003] High-performance, low-cost nuclear radiation detectors based on semiconductors have broad applicability in homeland security (nonproliferation of nuclear materials), industrial and medical imaging, nuclear energy, radioactive waste monitoring, oil / gas / geothermal exploration, and fundamental scientific research. For example, nuclear nonproliferation officers, first responders, and security personnel need to rapidly detect, identify, and quantify radioactive materials that may be security threats. While state-of-the- art semiconducting nuclear radiation sensing technologies can achieve high performance, they are either too costly (e.g., CdZnTe) or inconvenient (e.g., cryogenic cooling required for high purity germanium, HPGe) for widespread use as portable or large-area device monitors. Furthermore, the detrimental impact of lattice defects on semiconducting performance requires complex crystal growth processing to achieve high-quality materials and, thereby, high sensing efficiency that relies on large electronic carrier mobility-lifetime (^^^^) products.

[0004] Advances in new and emerging robust detector materials are warranted for the development of sensitive and convenient portable detectors and spectrometers at reduced cost. Design considerations may include simplicity of device design and materials processing, high sensitivity, robust operation and ability to distinguish between different radiation sources. The search for alternative defect-insensitive, high effective atomic number (e.g., high-Z) and wide-bandgap material candidates that can be produced with large areas and that are chemically, thermally, and mechanically robust has remained commerciallyelusive. Materials such as CsPbBr3, TlBr, FAPbBr3 / MAPbBr3, MAPbI3, and TlInSe2 have shown promising radiation monitoring properties, but proofs of concepts to-date still often require single crystals, limiting scalability. Moreover, they suffer from a variety of environmental and operational degradation mechanisms given low temperature stability limits and / or those related to ionic defect migration under applied DC potential, resulting in undesirable and irreversible electrochemical reactions at the electrodes that alter the charge transfer and collection efficiency of semiconductor-based sensing devices, ultimately complicating the overall device design and operation.

[0005] State-of-the-art radiation spectrometers generally rely on semiconducting materials constructed with simple photoconductor architectures or more complex p-i-n structures. The electron-hole pairs generated by ionizing radiation are separated by built-in and / or externally applied electric fields and drift towards the electrodes where they are collected. Key material figures of merit include overall dark resistance, which establishes the lower threshold limit of detection, above which sufficiently high concentration of photogenerated carriers can bedetected, and the mobility-carrier lifetime (^^ ∙ ^^) product, which defines how efficientlyphotogenerated charges can be extracted from the device and therefore the signal magnitude. In turn, the mobility and carrier lifetimes in semiconductors depend on the quality of the crystal, as defects can act as scattering or trapping centers that prevent charge collection.

[0006] Single crystal materials with reduced defect densities are thus preferred; unfortunately, this leads to increased fabrication costs and size limitations. For example, high purity germanium remains the gold standard of semiconductor gamma-ray spectrometers in terms of energy resolution, but, owing to its rather small band gap, requires cryogenic cooling to achieve reasonable dark resistances, thus limiting applications where small size and low power are important. Wider band gap materials are more promising for room temperature operation, as they allow more ready distinction of low radiation intensity signals from background thermal generation.

[0007] While CdZnTe (CZT) detectors operate at room temperature and provide high energy resolution, they remain costly. The complex ternary nature of the CZT compounds requires precise composition and temperature control during synthesis to obtain superior electrical transport properties. Structural (e.g., twinning, stacking faults, and grain boundaries) and chemical inhomogeneities (e.g.,Te precipitates) can easily arise during synthesis, thereby limiting overall device performance by quenching or scattering charges generated by gamma irradiation. Significant efforts are needed to control crystal growth to reduce defectgeneration with meticulous harvesting from pristine sections of the boule needed to obtain the best quality crystals. This limits overall large-scale production capabilities and increases cost.

[0008] These examples of engineering challenges for state-of-the-art, room-temperature CZT detectors also translate to next generation candidate materials such as CsPbBr3, TlBr, MAPbBr3, MAPbI3, Tl6SeI4, and TlInSe2 that are currently in research and development phases and exhibit similar synthesis dilemmas. Of the next generation candidates that are most promising, the halides, such as TlBr, CsPbBr3, or MAPbI3, exhibit some of the mostpromising radiation monitoring properties with high ^^ ∙ ^^ products and unique defecttolerance properties that make them less sensitive to point defects. However, they suffer from a broad range of engineering challenges, related to various environmental degradation mechanisms (atmosphere, humidity, and temperature). Degradation mechanisms related to ionic defect migration and electrochemical reactions at the electrodes further affect performance under above band gap irradiation. It is difficult to find semiconducting materials that exhibit the superior transport properties required for accurate sensing while meeting desired cost, production, and operating specifications.

[0009] High-performance, low-cost nuclear radiation detectors based on semiconductors have broad applicability in homeland security (nonproliferation of nuclear materials), industrial and medical imaging, nuclear energy and radioactive waste safety, oil / gas / geothermal exploration, and fundamental scientific research. Thus, the development of radiation sensing technology characterized by low manufacturing costs, while offering high performance gamma ray detection, even under harsh operating conditions, is highly desirable. SUMMARY

[0010] Space charge potential barriers exist at grain boundaries in ion-conducting ceramics, including the ion conducting polycrystalline materials disclosed herein, because the interfaces are net-charged. This occurs due to the unintentional segregation of charge point defects (e.g., intrinsic defects and / or extrinsic foreign elements) from the bulk lattice towards the grain boundary interface. The interfaces may be positively charged or negatively charged. Controlling the sign and magnitude of the net charge at the interface may allow the space charge potential height to be controlled. Disclosed herein are methods for selecting and / or introducing impurities (e.g., a selected precursor element) directly at the grain boundaryinterfaces to intentionally control the net charge of the grain boundary interfaces and, therefore, control the space charge potential barrier heights These methods may allow for the control of the detection threshold, dynamic range, and / or operating temperature of ion- conducting polycrystalline radiation detection devices.

[0011] Disclosed herein are radiation detection devices that include a polycrystalline ion- conducting solid electrolyte instead of a single-crystalline semiconducting material. These devices use illumination-induced reductions in space charge barriers at grain boundaries that may initially block ion conduction, thereby enhancing the ionic current and enabling a purely ionic radiation sensing response that capitalizes on structural defects existing at interfaces, harnessing changes in ionic conductance due to irradiation as the sensing response. These devices may be operated in spectroscopic and dosimetric application modes. Disclosed herein are material fabrication protocols to control the radiation sensitivity, detector dynamic range, and / or operating temperature range of these radiation detection devices. Further disclosed herein are descriptions including the material selection and synthesis protocols to control and optimize these parameters for radiation detection devices.

[0012] In some aspects, the techniques described herein relate to a detector for radiation, the detector including an ion-conducting polycrystalline material to absorb the radiation, wherein the ion-conducting polycrystalline material includes grain boundaries and a precursor element diffused within the grain boundaries, a pair of electrodes, electrically coupled to the ion- conducting polycrystalline material, to apply a voltage across the ion-conducting polycrystalline material, and a sensor, electrically coupled to the pair of electrodes, to measure a change in conductance of the ion-conducting polycrystalline material caused by absorption of the radiation.

[0013] In some aspects, the techniques described herein relate to a detector wherein the grain boundaries are positively charged grain boundaries and wherein the precursor element is a positively charged precursor element.

[0014] In some aspects, the techniques described herein relate to a detector wherein the positively charged precursor element increases the positive charge of the grain boundaries.

[0015] In some aspects, the techniques described herein relate to a detector wherein the grain boundaries are negatively charged grain boundaries and wherein the precursor element is a negatively charged precursor element.

[0016] In some aspects, the techniques described herein relate to a detector wherein negatively charged precursor element increases the negative charge of the grain boundaries.

[0017] In some aspects, the techniques described herein relate to a detector wherein the precursor element increases a space-charge potential barrier of the grain boundaries.

[0018] In some aspects, the techniques described herein relate to a detector wherein the precursor element is substitutionally inserted within the grain boundaries.

[0019] In some aspects, the techniques described herein relate to a detector wherein the precursor element is interstitially inserted within the grain boundaries.

[0020] In some aspects, the techniques described herein relate to a detector wherein the precursor element has an ionic radius at least 20% smaller than an ionic radius of a host element of the ion-conducting polycrystalline material.

[0021] In some aspects, the techniques described herein relate to a detector having a room temperature dark resistance >1010 ohm.

[0022] In some aspects, the techniques described herein relate to a detector having a sensitivity ∆R / R of about 103.

[0023] In some aspects, the techniques described herein relate to a method of diffusing an ion-conducting polycrystalline material with a precursor element, the method including selecting a precursor element based on a charge of an ionic carrier of the ion-conducting polycrystalline material, contacting a film the ion-conducting polycrystalline material with a source of the precursor element, and annealing the source of the precursor element and the film of the ion-conducting polycrystalline material at a temperature between 600 ℃ and 900 ℃ for 2 hours to 48 hours.

[0024] In some aspects, the techniques described herein relate to a method, wherein selecting the precursor element includes comparing an ionic radius of the precursor element to an ionic radius of the ionic carrier of the ion-conducting polycrystalline material and selecting the precursor element with an ionic radius at least 20% smaller than the ionic radius of the ionic carrier of the ion-conducting polycrystalline material.

[0025] In some aspects, the techniques described herein relate to a method wherein the source of the precursor element is a film of the precursor element.

[0026] In some aspects, the techniques described herein relate to a method wherein the film of the precursor element is deposited below the film of the ion-conducting polycrystalline material, the method further including diffusing, by the precursor element, from the film of the precursor element into grain boundaries of the ion-conducting polycrystalline material.

[0027] In some aspects, the techniques described herein relate to a method wherein the film of the precursor element is deposited above the film of the ion-conducting polycrystallinematerial, the method further including diffusing, by the precursor element, from the film of the precursor element into grain boundaries of the ion-conducting polycrystalline material and after annealing the film of the precursor element and the film of the ion-conducting polycrystalline material, etching the film of the precursor element to remove the film of the precursor element.

[0028] In some aspects, the techniques described herein relate to a method wherein the source of the precursor element is a substrate comprising the precursor element and wherein contacting the film the ion-conducting polycrystalline material with the source of the precursor element further comprises depositing the film of the ion-conducting polycrystalline material on a surface of the substrate.

[0029] In some aspects, the techniques described herein relate to a detector for radiation including the ion-conducting polycrystalline material.

[0030] In some aspects, the techniques described herein relate to a method of diffusing an ion-conducting polycrystalline material with a precursor element, the method including mixing particles of the ion-conducting polycrystalline material with particles of the precursor element to form mixed particles, drying the mixed particles at 80 ℃ to 120 ℃ to form coated particles including the ion-conducting polycrystalline material coated with the precursor element, calcinating the coated particles at about 400 ℃ to about 600 ℃ for about 2 hours to partially decompose the precursor element coating of the coated particles, pressing the coated particles to form a consolidated body, and sintering the consolidated body at 600 ºC to 1000 ℃ for 2 hours.

[0031] In some aspects, the techniques described herein relate to a detector for radiation including the ion-conducting polycrystalline material.

[0032] All combinations of the foregoing concepts and additional concepts discussed in greater detail below (provided such concepts are not mutually inconsistent) are part of the inventive subject matter disclosed herein. In particular, all combinations of claimed subject matter appearing at the end of this disclosure are part of the inventive subject matter disclosed herein. The terminology used herein that also may appear in any disclosure incorporated by reference should be accorded a meaning most consistent with the particular concepts disclosed herein. BRIEF DESCRIPTION OF THE DRAWINGS

[0033] The skilled artisan will understand that the drawings primarily are for illustrative purposes and are not intended to limit the scope of the inventive subject matter described herein. The drawings are not necessarily to scale; in some instances, various aspects of the inventive subject matter disclosed herein may be shown exaggerated or enlarged in the drawings to facilitate an understanding of different features. In the drawings, like reference characters generally refer to like features (e.g., functionally similar and / or structurally similar elements).

[0034] FIG.1A shows a schematic diagram of an oxygen ion conducting polycrystalline solid electrolyte thin film supported on an insulating substrate with porous platinum (Pt) surface electrodes under illumination that may be a radiation detector.

[0035] FIG.1B shows schematic spatial potential energy (^^∙∙^^^^^^) / defect concentration (^^^^ , ^^)diagrams representative of the space charge zone at a grain boundary in the dark (on left) and under above band gap illumination (on right). As illustrated, a positively charged grain boundary core may lead to a potential that may induce a depletion of positively charged ionic defects ^^^∙∙^ in the adjacent grains near the interface, resulting in high total film resistance in the dark. Illumination induced reductions in space charge barriers occur as photogenerated negative charge carriers recombine in the grain boundary core, neutralizing the original positive charge, alleviating the depletion of ionic carriers, and thus increasing the film’s conductance.

[0036] FIG.1C shows a schematic diagram of a radiation detector device including an oxygen ion conducting polycrystalline solid electrolyte.

[0037] FIG.1D is a schematic representation of light modulated ionic grain boundary radiation detector device in the dark.

[0038] FIG.1E is a schematic representation of light modulated ionic grain boundary radiation detector device under illumination / irradiation).

[0039] FIG.2A is a graph of the simulated grain boundary space charge potential (eV) as a function of photogeneration rate (in suns) for Gd doped CeO2 at 300 °C as a function of the photogeneration rate for different dark space charge potentials.: 1.00 eV (top curve), 0.75 eV (second curve from the top), 0.50 eV (third curve from the top), and 0.25 eV (bottom curve). The space charge potential is proportional to log10 (RgbR ) where R is the grain bo bulkgbundary resistance and Rbulkis the bulk resistance.

[0040] FIG.2B is a graph showing that the equivalent optoionic ionic response is equal to GLight / GDark where G is conductance. For reference, 1 sun is equivalent to 1017photon carriers cm-2s-1.1.00 eV (top curve), 0.75 eV (second curve from the top), 0.50 eV (third curve from the top), and 0.25 eV (bottom curve). As indicated, the magnitude of the conductance change is directly related to the detector's dynamic range, while the lower conductance change limits define its detection threshold.

[0041] FIG.2C is a schematic of interstitial insertion of precursor elements into a grain boundary.

[0042] FIG.2D is a schematic of substitutional insertion of precursor elements into a grain boundary.

[0043] FIG.3 is a graph showing bulk and grain boundary (GB) conductance as a function of temperature for uninfiltrated, 1.7 at% Al (M7), 2 at% B-infiltrated (M20), and 2 at% Ga. B and Al-infiltrated samples were sintered at 1100ºC, while the uninfiltrated sample was sintered at 1300°C. Results are plotted as 1 / R versus reciprocal thermal energy 1 / kT, where R is resistance.

[0044] FIG.4 is a graph of GB conductance G in the dark (closed symbols) and under constant UV light intensity (2×1018photons / s) (open symbols) as a function of temperature for three samples: B infiltrated and sintered at 1100°C (dark: closed circle, diamond; light: open circle), Al infiltrated and sintered at 1100°C (dark: closed square; light: open square), Al infiltrated sintered at 1300°C (dark: closed triangle; light: open triangle). For context, the maximum light intensity of UV LED utilized as the light source is 1018photons / (cm2s). Results are plotted as conductance G (≡ 1 / R) versus reciprocal thermal energy 1 / kT. The smaller ionic radius elements (B vs. Al < Ce) may result in higher activation energies, while lower sintering temperatures may result in higher overall activation energies.

[0045] FIG.5A is a graph showing the conductance ratio under light and in the dark as a function of photon flux for the Al infiltrated sample (bottom curve) and the B infiltrated sample (top curve). At high light intensity, the response provides the maximum resistance modulation ratio, indicative of dynamic range, whereas at low light intensity, it allows the estimate of the material’s detection threshold. The slope of the curve is associated with detection linearity.

[0046] FIG.5B is a graph of grain boundary and bulk resistance for different levels of LED current (mA) illumination intensity: dark (bottom curve), 0.7% (second curve from the bottom), 7.1% (third curve from the bottom), 45% (fourth curve from the bottom), and 85.7%(fifth curve from the bottom). The grain boundary resistance results in difference GB conductance values due to the optoionic effect.

[0047] FIG.5C is a graph showing the conductance ratio under light and in the dark near room temperature as a function of photon flux for the Al infiltrated sample (bottom curve), the B infiltrated sample (middle curve), and for the B infiltrated sample after correcting for the theoretical dark conductance (top curve). The dark conductance is obtained from the high-temperature extrapolation of measured conductance down to room temperature.

[0048] FIG.6 is a schematic representation of an annealing process for up diffusion where a polycrystalline thin film on an insulating (source) substrate is exposed to subsequent thermal annealing from 600 ℃ to 1200 ℃ for varying periods from 2 hours to 48 hours to enable the diffusion of the substrate elements upwards along the grain boundary while reducing the bulk segregation of unwanted defects. Above 900 ℃, undesirable bulk dopant segregation towards the grain boundaries may occur, which may be detrimental to the grain boundary engineering strategy disclosed herein.

[0049] FIG.7A is a schematic representation of annealing process for up diffusion where a polycrystalline thin film on an insulating (source) substrate is exposed to subsequent thermal annealing from 600 ℃ to 900 ℃ for varying periods from 2 hours to 48 hours to enable the diffusion of the of precursor elements from the substrate upwards along the grain boundary while reducing the bulk segregation of unwanted defects.

[0050] FIG.7B is an overview of a method for the up diffusion of precursor elements from a substrate for a polycrystalline thin film.

[0051] FIG.8A is a schematic representation of annealing and process steps for down diffusion from a top thin source layer of precursor elements. The top source layer of precursor elements may be grown above the polycrystalline films (left schematic) and etched before metallization steps (right schematic).

[0052] FIG.8B is an overview of a method for the down diffusion of precursor elements from thin top source layer for a polycrystalline thin film.

[0053] FIG.9A is a schematic representation of annealing and process steps for up diffusion from a bottom thin source layer of precursor elements. The bottom source layer of precursor elements may be grown on the substrate (first schematic) and pre-annealed to ensure high conformality before growing the polycrystalline films (second schematic) and performing the in-diffusion step (third schematic).

[0054] FIG.9B is an overview of a method for the up diffusion of precursor elements from thin bottom source layer for a polycrystalline thin film.

[0055] FIG.10A is a graph showing dark conductance results for 3 at% Gd doped CeO2 polycrystalline thin films grown on an MgO substrate and annealed at various temperatures (850 ℃ (bottom line), 973 ℃ (second line from the bottom), 1073 ℃, 1173 ℃ (middle line), and 1275 ℃ (second line from the top)) versus an epitaxial sample (top line). Results are plotted as the natural logarithm of inverse resistance (conductance) × temperature vs reciprocal temperature 1 / kT.

[0056] FIG.10B is a graph showing dark conductance results for 3 at% Gd doped CeO2 polycrystalline thin films grown on an Al2O3 substrate and annealed at various temperatures (750 ℃ (third line from the bottom), 850 ℃ (bottom line), 973 ℃ (second line from the bottom), 1073 ℃ (second line from the top), and 1173 ℃ (third line from the top)) versus an epitaxial sample (top line). Results are plotted as the natural logarithm of inverse resistance (conductance) × temperature vs reciprocal temperature 1 / kT.

[0057] FIG.10C is a plot of the extracted space charge potential as a function of annealing history for the 3 at% Gd doped CeO2 polycrystalline thin film grown on an Al2O3 substrate (top curve) and the 3 at% Gd doped CeO2 polycrystalline thin film grown on an MgO substrate (bottom curve).

[0058] FIG.11A is a graph of UV conductance results for 3 at% Gd doped CeO2 polycrystalline thin films grown on a MgO substrate and annealed at various temperatures (850 ℃ (closed triangle), 973 ℃ (square), 1073 ℃ (× symbol), 1173 ℃ (open triangle), and 1275 ℃ (plus sign)) versus an epitaxial sample (diamond). Results are plotted as a natural logarithm of inverse resistance (conductance) × temperature vs reciprocal temperature 1 / kT.

[0059] FIG.11B is a graph of UV conductance results for 3 at% Gd doped CeO2 polycrystalline thin films grown on an Al2O3 substrate and annealed at various temperatures (750 ℃ (triangle), 850 ℃ (x with line through it), 973 ℃ (× symbol), 1073 ℃ (diamond), and 1173 ℃ (plus sign)). Results are plotted as a natural logarithm of inverse resistance (conductance) × temperature vs reciprocal temperature 1 / kT.

[0060] FIG.11C is a graph showing the equivalent optoionic ionic response at a fixed light intensity (approximately 1018photons / (cm2s)), equal to GLight / GDark where G is conductance, as a function of reciprocal temperature 1 / kT for 3 at% Gd doped CeO2 polycrystalline thin films grown on a MgO substrate and annealed at various temperatures (850 ℃ (solidtriangle), 973 ℃ (solid square), 1073 ℃ (× symbol), 1173 ℃ (open triangle), and 1275 ℃ (plus sign)) versus an epitaxial sample (solid diamond).

[0061] FIG.11D is a graph showing the equivalent optoionic ionic response at a fixed light intensity (approximately 1018photons / (cm2s)), equal to GLight / GDark where G is conductance, as a function of reciprocal temperature 1 / kT for 3 at% Gd doped CeO2 polycrystalline thin films grown on an Al2O3 substrate and annealed at various temperatures (750 ℃ (solid triangle), 850 ℃ (x with a line through it), 973 ℃ (× symbol), 1073 ℃ (plus sign), and 1173 ℃ (diamond)).

[0062] FIG.12A is a schematic representation of the diffusion of precursor elements for a bulk ceramic.

[0063] FIG.12B is an overview of a method for the diffusion of precursor elements for a bulk ceramic.

[0064] FIG.13A shows an illustration of a polycrystalline film on a substrate with a columnar GB, initially positively charged. GB enables selective up-diffusion of substrate cations into film. The inset shows a schematic of space charge potential ϕSC(x) profile generated by the positively charged GB core and corresponding majority oxygen vacancy (^^o..) depletion and minority electron n accumulation, following the Mott Schottky model where dopant concentration [^^′Ce ] is assumed to be flat throughout space charge region.

[0065] FIG.13B shows an illustration of impact of selective grain boundary in-diffusion on GB space charge potentials. Magnesium ions (Mg²⁺), and aluminum ions (Al³⁺), respectively decrease and increase the density of positive core charge, modifying grain boundary space charge potentials, space charge width λ, and degree of ionic carrier ^^o..depletion.

[0066] FIG.14A shows a natural logarithm plot of the Arrhenius dependence of conductance G as (1 / R) × T vs.1 / kBT derived from the impedance plots for the film grown on a MgO substrate. Activation energies (Ea, in eV) obtained from slops of curves are displayed in legend.

[0067] FIG.14B shows a natural logarithm plot of the Arrhenius dependence of conductance G as (1 / R) × T vs.1 / kBT derived from the impedance plots for the film grown on a Al2O3 substrate. Activation energies (Ea, in eV) obtained from slops of curves are displayed in legend.

[0068] FIG.15 shows a graph of calculated space charge potential on left axis and extrapolated resistance ratios at 30 °C on right axis. Results are obtained by comparing impedance of polycrystalline films with epitaxial films and using Equation 2 derived fromthe Mott Schottky approximation. All data in figure plotted versus annealing temperature for films grown on MgO (bottom curve) vs. Al2O3 (top curve).

[0069] FIG.16A shows Time-of-Flight Secondary Ion Mass Spectrometry (ToF-SIMS) depth profiles of Al+in GDC films deposited on Al2O3 (left plot) and Mg+in GDC film deposited on MgO (right plot) obtained following film deposition and subsequent annealing treatments at 750 °C (middle curve) and 850 °C (top curve) for 6 hours for Al2O3 and 650 °C (second curve from the bottom), 750 °C (third curve from the bottom), and 850 °C (top curve) for MgO. The vertical dotted line signifies the interface between the GDC film and substrate, corresponding to the point where the signals of Ce+and Al+ / Mg+overlap.

[0070] FIG.16B shows STEM- EDS spectra images collected from a representative grain boundary in GDC grown on Al2O3 (left image) and on MgO (right image) at 1050°C. The top image displays HAADF intensity, while the bottom images show the semiquantitative concentrations of Gd, Mg, Al and Si. The Mg boundary shows strong enrichment of Mg, while the Al grain boundary shows evidence of Al enrichment with both showing evidence of Gd enrichment. In all cases, the enrichment is localized to within approximately 1 nm of the grain boundary core. Scale bars 10 nm.

[0071] FIG.17A shows a schematic model explaining the observed trends associated with redistribution of cations at grain boundaries. From left to right: increasing anneal temperature leads initially to deeper GB penetration of dopants, ultimately leading at intermediate temperatures to saturation of in-diffused cations at the GBs and some grain growth. Further increases in annealing temperatures may lead to substantial grain growth and Gd segregation from within the grains into the GBs.

[0072] FIG.17B shows a schematic of the spatial distributions of space charge potential, bulk dopant concentration, majority oxygen vacancy, and minority electron defect concentrations as ^^^^Ce′becomes sufficiently mobile at progressively higher temperatures (e.g., ≥ 900°C) to segregate towards the grain boundaries, accumulating in the depletion region and contributing negative charge, thereby reducing the barrier height ^^^^^^ and the space charge width, λ , as displayed to the right.

[0073] FIG.18A is a SEM cross-sectional micrograph for Al2O3.

[0074] FIG.18B is a SEM cross-sectional micrograph for MgO.

[0075] FIG.18C is a graph of the XRD patterns of GDC films as grown (about 250 °C) on Al2O3 (bottom curve) and MgO (top curve) respectively.

[0076] FIG.19 is a characterization of GDC film grown on Al2O3 by SEM (upper left image) and TEM (upper right image and lower image) showing the GDC film to be polycrystalline with columnar grains.

[0077] FIG.20A is an overlaid view of XRD spectra for GDC films grown on Al2O3 (left schematic) and annealed from 650 ºC to 1250 ºC.

[0078] FIG.20B is a normalized cascaded view of the XRD patterns for each film from FIG. 20A.

[0079] FIG.20C are rocking curves of the (111) and (200) orientations for each film from FIG.20A as a function of annealing temperature.

[0080] FIG.20D is micrographs of GDC thin films grown on Al2O3 and annealed at different temperatures. Images were collected at 3 keV electron energy, working distance 4 mm, using a secondary electron detector. Scale bars 1 µm.

[0081] FIG.21A is an overlaid view of XRD spectra for 3GDC films grown on MgO (left schematic) and annealed from 650 ºC to 1250 ºC.

[0082] FIG.21B is a normalized cascaded view of the XRD patterns for each film from FIG. 21A.

[0083] FIG.21C are rocking curves of the (111) and (200) orientations as for each film from FIG.21A a function of anneal temperature.

[0084] FIG.21D is micrographs of GDC thin films grown on MgO and annealed at different temperatures. All images were collected at 15 keV electron energy, working distance 12 mm. The 1075°C micrograph was taken with a backscatter detector; the others were taken with a secondary electron detector. Scale bars 1 µm.

[0085] FIG.22A shows a graph of measured grain sizes as a function of annealing temperature for GDC thin films grown on MgO (circles) and Al2O3 (triangles).

[0086] FIG.22B shows an Arrhenius plot of the thermal activation of grain growth. An ideal grain growth exponent of 0.5 and an initial grain size of 30 nm were assumed. An activation energy of 1.32 eV is found for MgO and 1.55 eV for Al2O3.

[0087] FIG.23A shows an example Nyquist plot of the complex impedance response obtained at 350°C under open-circuit conditions for a GDC3 film on an Al2O3 substrate. The smaller plot at the left represents a zoom-in of data near the origin to showcase the low resistance impedances at the more elevated temperatures near the origin at the highest frequencies.

[0088] FIG.23B shows an example Nyquist plot of the complex impedance response obtained at 350°C under open-circuit conditions for a GDC3 film on an MgO substrate. The smaller plot at the left represents a zoom-in of data near the origin to showcase the low resistance impedances at the more elevated temperatures near the origin at the highest frequencies.

[0089] FIG.24 is a semi-log plot of Arrhenius dependence of conductance G as (1 / R) × T vs. 1 / kT. Data points derived from impedance plots obtained for epitaxial films grown on MgO (top line) and Al2O3 (bottom line).

[0090] FIG.25A is a graph of total activation energy vs. annealing temperature for films grown on MgO (bottom curve) versus Al2O3 (top curve).

[0091] FIG.25B is a plot of log10 (RgbRbulk) ^^^^ 1 / ^^^^ used for the fitting space charge potentials with the aid of Equation 2 for films grown on MgO. ^^^^^^^^^^ ~ ^^^^^^^^^^^^^^^^^^^^ assumed for the fitting procedure to calculate the space charge potentials ^^^^(0) listed as a function of temperature in the tables below each plot.

[0092] FIG.25C is a plot of log10 (RgbRbulk) ^^^^ 1 / ^^^^ used for fitting space charge potentials with the aid of Equation 2 for films grown on Al2O3. ^^^^^^^^^^ ~ ^^^^^^^^^^^^^^^^^^^^ assumed for the fitting procedure to calculate the space charge potentials ^^^^(0) listed as a function of temperature in the tables below each plot.

[0093] FIG.26A is a ToF-SIMS depth profile of Al+ (bottom curve, left graph), Ce+ (middle curve, left graph), Gd+ (top curve) in GDC deposited on Al2O3 (left graph) and Mg+ (bottom curve, right graph), Ce+ (middle curve, right graph), Gd+ (top curve, right graph) deposited on MgO (right graph) after annealing at 750 °C for 6 hours. The dotted line signifies the interface between the GDC film and the substrate.

[0094] FIG.26B is a graph of depth profiles of Al+ and Mg+ as in FIG.26B for listed annealing temperatures.

[0095] FIG.27 shows ToF-SIMS spectra profiles of cations for GDC deposited on Al2O3. Depth profiles were obtained for as-deposited samples and samples annealed at varied temperatures for 6 hours.

[0096] FIG.28 shows ToF-SIMS spectra profiles of cations for GDC deposited on MgO. Depth profiles were obtained for as- deposited samples and samples annealed at varied temperatures for 6 hours.

[0097] FIG.29 shows normalized depth profiles of Al+ in GDC film layer for samples annealed at 700 °C (left graph) for 9 hours, 750 °C (middle graph) for 6 hours, and 800 °C (right graph) for 6 hours. The best-fit lines are illustrated as the curved line in each spectrum. Annealing time was adjusted to induce the desired diffusion amount and improve the profile fitting quality.

[0098] FIG.30 shows normalized depth profiles of Mg+in GDC film layers for samples annealed at 650 °C (left graph), 750 °C (middle graph), and 850 °C (right graph) for 6 hours. The best-fit lines are illustrated as the curved line in each spectrum.

[0099] FIG.31 shows normalized depth profiles of Al+in GDC film layers for samples annealed at 1050 °C (left graph), 1150 °C (middle graph), and 1250 °C (right graph) for 30 hours. The best-fit curves are illustrated as the curved lines in each spectrum.

[0100] FIG.32 shows normalized depth profiles of Mg+in GDC film layers for samples annealed at 1050 °C (left graph), 1150 °C (middle graph), 1250 °C (right graph) for 30 hours. The best-fit curves are illustrated as the curved lines in each spectrum.

[0101] FIG.33 shows Arrhenius plots of measured Mg (star) and Al (diamond) diffusivities with fitted activation energies indicated. Grain boundary diffusivities are shown in the left graph and bulk diffusivities are shown in the right graph.

[0102] FIG.34 shows ToF-SIMS spectra profiles of cations for GDC deposited on Al2O3. Depth profiles were obtained for as-deposited samples and samples annealed at varied temperatures for 6 hours.

[0103] FIG.35 shows ToF-SIMS spectra profiles of cations for GDC deposited on MgO. Depth profiles were obtained for as- deposited samples and samples annealed at varied temperatures for 6 hours.

[0104] FIG.36A shows a high-angle annular dark field (HAADF-STEM) overview of a STEM-EDS dataset collected from a representative grain boundary of GDC grown on Al2O3 and annealed at 850°C. The grain boundary travels vertically through the center of the image. The box and arrow indicate the linescan displayed in FIGS.36B and 36C. Scalebar 2 nm.

[0105] FIG.36B shows a summed STEM-EDS spectrum. Al peak is visible. Mo peaks are artifacts from the sample grid, and Pt peaks result from sample preparation.

[0106] FIG.36C shows a graph of semi-quantitative concentrations measured across the grain boundary in FIG.36A and HAADF intensity

[0107] FIG.36D is a zoomed in graph of FIG.36C to display trace elements. Al and Gd are enriched at the grain boundary.

[0108] FIG.37A shows a HAADF-STEM overview of a STEM-EDS dataset collected from a grain boundary of GDC grown on Al2O3 and annealed at 850°C. The grain boundary travels vertically through the center of the image. The box and arrow indicate the linescan displayed in FIGS.37C and 37D. Scalebar 2 nm.

[0109] FIG.37B shows a summed STEM-EDS spectrum. An Al peak is visible. Mo peaks are artifacts from the sample grid, and Pt peaks result from sample preparation.

[0110] FIG.37C shows a graph of semi-quantitative concentrations measured across the grain boundary in FIG.37A and HAADF intensity.

[0111] FIG.37D is a zoomed in graph of FIG.37C to display trace elements. No evidence of chemical changes are found at the grain boundary.

[0112] FIG.38A shows a HAADF-STEM overview of a STEM-EDS dataset collected from a representative grain boundary of GDC grown on Al2O3 and annealed at 1050°C. Grain boundary travels vertically through the center of the image. Arrows indicate the linescan displayed in FIGS.38C and 38D. Scalebar 2 nm.

[0113] FIG.38B shows a summed STEM-EDS spectrum. Mg, Al, and Si peaks are visible. Mo peaks are artifacts from the sample grid, and Pt peaks are a result of sample preparation.

[0114] FIG.38C shows a graph of semi-quantitative concentrations measured across the grain boundary in FIG.38A, as well as HAADF intensity.

[0115] FIG.38D is a zoomed in graph of FIG.38C to display trace elements. Al and Si are enriched at the grain boundary, and there is weak enrichment of Gd.

[0116] FIG.39A is a HAADF-STEM overview of a STEM-EDS dataset collected from a representative grain boundary of GDC grown on MgO and annealed at 1050°C. Grain boundary travels vertically through the center of the image. The box and arrow indicate the linescan displayed in FIGS.38C and 38D. Scalebar 5 nm.

[0117] FIG.39B is a summed STEM-EDS spectrum. Mg, Al, and Si peaks are visible. Mo peaks are artifacts from the sample grid, and Pt peaks are a result of sample preparation.

[0118] FIG.39C is a graph of semi-quantitative concentrations measured across the grain boundary in FIG.39A, as well as HAADF intensity.

[0119] FIG.39D is a zoomed-in graph of FIG.39C to display trace elements. Mg and Si are strongly enriched at the grain boundary. Weak enrichment of Al and Gd is also evidenced.

[0120] FIG.40A is a HAADF-STEM overview of a STEM-EDS dataset collected from a representative grain boundary of GDC grown on MgO and annealed at 1250°C. The grain boundary travels horizontally through the image. MgO substrate is visible on the right of image. The box and arrow indicate the linescan displayed in FIGS.40C and 40D. Scalebar 10 nm.

[0121] FIG.40B is a summed STEM-EDS spectrum. Mg and trace Al are visible. Mo peaks are artifacts from the sample grid.

[0122] FIG.40C is a graph of semi-quantitative concentrations measured across the grain boundary in FIG.40A, as well as HAADF intensity. There is some evidence of oxygen enrichment at the grain boundary.

[0123] FIG.40D is a zoomed-in graph of FIG.40C to display trace elements. Mg may be enriched at the grain boundary. No evidence of Gd enrichment is found.

[0124] FIG.41 shows a physical model explaining spatial distributions of space charge potential, bulk dopant concentration, majority oxygen vacancy, and minority electron defect concentrations in response to combined substrate cation precursor elements up- diffusion and bulk dopant Gd segregation at the GB of polycrystalline films. As samples are annealed at higher temperatures, the grain boundaries may become saturated with precursor elements Al (right) and Mg (left), respectively, while ultimately, bulk dopant Gd segregation at the grain boundary may further reduce the space charge potential for samples with MgO (right), while counteracting the increase observed with GB decoration by Al in samples with Al2O3 (left). The concentration profile and space charge width may follow a restricted equilibrium scenario instead of simple mott Schottky approximation. DETAILED DESCRIPTION

[0125] Disclosed herein are methods for setting the barrier height at grain boundaries of ion-conducting polycrystalline radiation detectors. The methods disclosed herein may improve the radiation sensitivity, operating temperature range, and / or dynamic range of ion- conducting polycrystalline radiation detectors. Disclosed herein are methods for tuning the space charge potential in polycrystalline bulk and thin-film radiation detectors. The methodsdisclosed herein may populate grain boundaries with an appropriate elemental charge to increase the potential barrier at the grain boundaries in the dark. This may be achieved by carefully selecting atomic elements with a low lattice solubility, the appropriate ionic radii and / or oxidation state, and utilizing a series of thermal processing steps to selectively maximize the population of the selected atomic elements within the grain boundary. The methods disclosed herein may also simultaneously reducing the diffusion and / or segregation of unwanted impurities (e.g., oppositely charged elements and / or dopants) towards the grain boundary interface.

[0126] Ion-Conducting Polycrystalline Radiation Detectors

[0127] The irradiation detection devices disclosed herein may include a polycrystalline material. The polycrystalline material may be an oxygen ion (O2-) conducting polycrystalline material. Instead of or in addition to an oxygen ion, the polycrystalline material may include one or more of the following ions: Li+, OH-, Na+, Mg2+, I-, K+, F-, Br-, Ag+, Cu+, Cl-, Al3+, H+. The polycrystalline material may be a cerium (Ce), zirconium (Zr), magnesium (Mg), lithium (Li), sodium (Na), fluorine (F), iodine (I), bromine (Br), silver (Ag), copper (Cu), aluminum (Al), lead (Pb), boron (B), helium (He), cadmium (Cd), gadolinium (Gd), promethium (Pm), samarium (Sm), scandium (Sc), europium (Eu), dysprosium (Dy), indium (In), hafnium (Hf), erbium (Er), bismuth (Bi), barium (Ba), gallium (Ga), or proton containing solid-state polycrystalline compound material. For example, the polycrystalline material may include CeO2, ZrO2, LiLaZrO4, LiLaTiO3, GaO3, Bi2O3, BaZrO3, BaCeO3, a lithium-garnet material (e.g., Li7La3Zr2O12), a sodium super ionic conductor (NASICON) (e.g., Na₁₊^Zr₂Si^P3−xO₁₂, where 0 < x < 3), a lithium phosphate material (e.g., Li1+xTi2−xMx(PO4)3 (where M = aluminum (Al), gadolinium (Ga), indium (In), or scandium (Sc)), or a lithium sulfide material (e.g., lithium germanium phosphorous sulfide (LGPS)). The proton solid-state polycrystalline material may include a perovskite (CaTiO₃) material, including but not limited, to BaCeO3 or BaZrO3.

[0128] The polycrystalline material may also be doped with one or more compounds, including but not limited to, boron (B), calcium (Ca), strontium (Sr), samarium (Sm), ytterbium (Yb), yttrium (Y), lanthanum (La), gadolinium (Gd), and / or lutetium (Lu). The doping may range from about 0.5 atm% to about 40 atm%, for example, about 1 atm% to about 20 atm%. For example, the polycrystalline material may be a Gd doped polycrystalline material, including but not limited to 3 atm% Gd doped CeO2 (3GDC). Alternatively, the polycrystalline material may be Y doped ZrO2 or Ca doped ZrO2.

[0129] The polycrystalline material may have a material density of > 80% density and more preferably > 90% density. For example, the polycrystalline material may have a material density of about 90% to about 95%, and preferably about 92%. The polycrystalline material may have one or more grain boundaries. The grain boundaries may be positively charged. The polycrystalline material may have positively charged grain boundaries for positively charged ions or defects (e.g., oxygen vacancies, Li interstitials, protons (e.g., OH-, H+), Mg interstitials, sodium interstitials, fluorine vacancies, bromine and iodine vacancies, silver interstitials, copper interstitials, aluminum interstitials., chlorine vacancies). Alternatively, the grain boundaries may be negatively charged. The polycrystalline material may have negatively charged grain boundaries for negatively charged ions or defects (e.g., oxygen interstitials, Li vacancies, Mg vacancies, sodium vacancies, fluorine interstitials, bromine and iodine interstitials, silver vacancies, copper vacancies, aluminum vacancies, chlorine interstitials). The spacing between adjacent grain boundaries in the polycrystalline material may range from about 10 nm to about several µm (e.g., about 1 µm to about 10 µm) including all values in between. For example, the grain boundaries in the polycrystalline material may be separated from each other by about 200 nm to about 400 nm. The grain size of the polycrystalline material may be modified via manufacturing through the use of nano powder fabrication, isostatic pressing, and / or rapid sintering, for example.

[0130] The polycrystalline material may be used in a variety of irradiation detection devices, including but not limited to, a large panel for use in security, a small, encapsulated cylinder for use in drilling and / or mining, or a thin film substrate for use on chip miniaturized radiation detection schemes. The polycrystalline material may be in the form of a film, a pellet, and / or a plate for use in a variety of irradiation detection devices.

[0131] Example 1: A Solid Electrolyte Radiation Detector

[0132] FIG.1A shows a schematic diagram of an irradiation detection device 100. The irradiation detection device 100 may be used in chip miniaturized radiation detection devices. The chip miniaturized radiation detection device may range in size from about 1 mm × 1 mm to about 5 mm × 5 mm. The irradiation detection device 100 may include a solid electrolyte film 110 made of any of the polycrystalline materials disclosed herein. In one embodiment, the solid electrolyte film 110 is an oxygen ion solid electrolyte film. For example, the solid electrolyte film 110 may be a CeO2 solid electrolyte film. The solid electrolyte film 110 may also be doped with any of the materials disclosed herein. In one embodiment, the solid electrolyte film 110 may be a gadolinium (Gd) doped solid electrolytefilm. The Gd doping may range from about 0.5 atm% to about 40 atm%, for example, about 1 atm% to about 20 atm%. Preferably, the solid electrolyte film 110 may be a 3 atm% Gd doped CeO2 (3GDC) solid electrolyte film.

[0133] The solid electrolyte film 110 may be prepared as described below in the sample preparation section. Other suitable methods for preparing the solid electrolyte film 110 include, but are not limited to, a traditional ceramic consolidation process (e.g., pressing (uniaxial, isostatic), slip casting, or extrusion), 3D printing, robocasting, pelletization, injection molding, physical vapor deposition (PVD), sputtering, thermal evaporation, pulsed laser deposition, chemical vapor deposition (CVD), atomic layer deposition (ALD), a sol-gel process (e.g., spin coat, dip coat, and / or spray coat), tape casting, slip casting, screen printing, phase inversion, and / or extrusion). The preparation of the solid electrolyte film 110 may also include a heat treatment (e.g., in a furnace and / or UV irradiaton).

[0134] The solid electrolyte film 110 may be < 1 µm thick (e.g., about 100 nm to about 900 nm thick) to about tens of µm thick (e.g., about 10 µm to about 100 µm thick). For example, the solid electrolyte film 110 may be about 100 nm, about 200 nm, about 300 nm, about 400 nm, about 500 nm, about 600 nm, about 700 nm, about 800 nm, about 900 nm, about 1 µm, about 10 µm, about 20 µm, about 30 µm, about 40 µm, about 50 µm, about 60 µm, about 70 µm, about 80 µm, about 90 µm or about 100 µm thick, including any values in between. The solid electrolyte film 110 may be formed in a sheet from about 1 mm × 1 mm to about 5 mm × 5 mm. For example, the solid electrolyte film 110 may be about 1 mm × 1 mm, about 1 mm × 2 mm, about 1 mm × 3 mm, about 1 mm × 4 mm, about 1 mm × 5 mm, about 2 mm × 2 mm, about 3 mm × 3 mm, about 4 mm × 4 mm, or about 5 mm × 5 mm, including all values in between. The solid electrolyte film 110 may also be formed in a disk with a diameter from about 1 mm to about 10 mm. For example, the solid electrolyte film 110 may have a diameter of about 1 mm, about 2 mm, about 3 mm, about 4 mm, about 5 mm, about 6 mm, about 7 mm, about 8 mm, about 9 mm, or about 10 mm, including all values in between. The solid electrolyte film 110 may be circular, semi-circular, rectangular, or any other suitable shape.

[0135] The solid electrolyte film 110 may have a material density of > 80% density and more preferably > 90% density. For example, the solid electrolyte film 110 may have a material density of about 90% to about 95%, and preferably about 92%. The solid electrolyte film 110 may have one or more grain boundaries 112 as shown in FIG.1A. The grain boundaries 112 may be positively charged as shown in FIG.1B. Alternatively, the grainboundaries may be negatively charged. The spacing between adjacent grain boundaries 112 may range from about 10 nm to about several µm (e.g., about 1 µm to about 10 µm) including all values in between. For example, the grain boundaries 112 may be separated from each other by about 200 nm to about 400 nm.

[0136] The irradiation detection device 100 may also include a substrate 120. The solid electrolyte film 110 may be disposed on the substrate 120 via chemical or physical vapor deposition, including but not limited to chemical vapor deposition, sputtering, pulsed laser deposition, or molecular-beam epitaxy deposition. The substrate 120 may support the growth of the solid electrolyte film 110. The substrate 120 may be attached to the underside of the solid electrolyte film 110. The substrate 120 may be made of any electrically insulating material so as not to partially short out current flowing through the solid electrolyte film 110. For example, the substrate 120 may include a MgO, Al2O3, SiO2, SrTiO3, LaAlO3, or SiC substrate.

[0137] The irradiation detection device 100 may also include one or more electrodes 130. The electrodes 130 may be made of any metallic material that may conduct electrons at a conductivity of about > 10-3S / cm and / or any metallic material that may conduct ions at a conductivity of about 10-5S / cm. For example, the electrodes 130 may be platinum (Pt), a platinum alloy, gold, or stainless-steel electrodes. The platinum alloy may include but is not limited to a PtNi alloy (e.g., Pt40Ni60, Pt50Ni50, and / or Pt75Ni25), a PtFe alloy (e.g., Pt3Fe), a PtCo alloys (e.g., Pt3Co and / or PtCo), and / or a PtCu alloys (e.g., Pt0.5Cu0.5). Alternatively, the electrodes 130 may be made of a highly electronically conducting metallic oxide (e.g., indium tin oxide (ITO)) and / or mixed ionic-electronic conducting metallic oxide (e.g., lanthanum strontium cobalt iron oxide (LSCF)). Alternatively, the electrodes may be made of graphite, lithium, lithium cobalt oxide (LiCoO2), lithium manganese oxide (LiMn2O4), lithium iron phosphate (LiFePO4), nickel manganese cobalt oxide (NMC), lithium nickel manganese cobalt oxide (LiNixMnyCozO2), lithium nickel cobalt aluminum oxide (LiNiCoAlO2 or NCA), lithium titanate (Li4Ti5O12), La0.6Sr0.4FeO3−δ (LSF), La0.5Sr0.5Cr0.2Mn0.8O3−δ(LSCrMn), La0.9Sr0.1CoO3−δ(LSC), Ba0.5Sr0.5Co0.8Fe0.2O3-δ(BSCF) , La0.8Sr0.2MnO3 (LSM) , Sm0.5Sr0.5CoO3-δ (SSC) , La0.5Sr0.5MnO3-δ (LSM), LSM with Sc doping, La0.8Sr0.2Sc0.5Fe0.5O3-δ (LSSF), BaCe0.26Ni0.1Fe0.64O3-δ (BCNF10), BaCo0.4Fe0.4Zr0.1Y0.1O3-δ (BCFZY0.1), BaCe0.26Ni0.1Fe0.64O3-δ (BCNF10), NdBa0.5Sr0.5Co1.5Fe0.5O5+δ (NBSCF), BaPr0.8In0.2O3-δ (BPI), palladium (Pd), a palladium alloy (e.g., PdAg and / or PdCu), NaxCoO2, Na3V2(PO4)3, NaxMnO2, NaxTiO2, Na2Ti3O7),K0.3MnO2, K0.55CoO2, KxFe2(CN)6, K2Fe4(CN)6, K2Mn[Fe(CN)6], KxMnO2, CuBr, AgBr, PbBr2, CuBr2, CsPbBr3, CuI, AgI, PbI2, BiI3, and / or SnI4. Preferably, the electrodes 130 are Pt electrodes. The electrodes 130 may be attached to the top surface 111 of the solid electrolyte film 110. Preferably the irradiation detection device 100 includes two electrodes 130. In this embodiment, a first electrode 130a may be a working electrode and the second electrode 130b may be a counter electrode. In another embodiment, the first electrode 130a may be a cathode and the second electrode 130b may be an anode. For example, at least two electrodes 130 may be attached to either side of the solid electrolyte film 110.

[0138] The electrodes 130 may be interdigitated electrodes (e.g., with interwoven fingers to increase the electrode 130 surface area) or planar electrodes. Each electrode 130 may be approximately 100 nm thick. The electrodes 130a and 130b may be the same size (e.g., length, width, or diameter) as the solid electrolyte film 110. The electrodes 130a and 130b may be spaced about 1 µm apart to about > 100 µm apart (e.g., about 100 µm to about 900 µm). If the electrodes 130a and 130b are interdigitated electrodes, the spacing between the digits may range from about 1 µm to about > 100 µm (e.g., about 100 µm to about 900 µm). For example, the spacing between the digits of the interdigitated electrodes may be about 1 µm, about 10 µm, about 20 µm, about 30 µm, about 50 µm, about 60 µm, about 70 µm, about 80 µm, about 90 µm, about 100 µm, about 200 µm, about 300 µm, about 400 µm, about 500 µm, about 600 µm, about 700 µm, about 800 µm, or about 900 µm, including all values in between. The electrodes 130 may be oval, circular, square, or rectangular in shape. For example, each electrode 130 may be a rectangle covering an area of about 5 mm ^ 1 mm on the surface 111 of the solid electrolyte film 110. The electrodes 130 may be deposited on the solid electrolyte film 110 via sputtering. The electrodes 130 may be attached to one or more wires 131 using a silver paste. The wires 131 may be made of any high melting point metal that exhibits corrosion resistance. For example, the wires 131 may be platinum, gold, and / or stainless-steel wires. The electrodes 130 may be used to apply a voltage to the solid electrolyte film 110. The irradiation detection device 100 may also include a thermocouple (not shown) to measure the temperature of the solid electrolyte film 110.

[0139] The irradiation detection device 100 may be used to detect optical illumination 105 or other radiation, including gamma irradiation, UV radiation, X-rays and / or neutrons.

[0140] Example 2: A Solid Electrolyte Pellet / Plate Detector

[0141] FIG.1C shows a schematic diagram of another embodiment of an irradiation detection device 150. The irradiation detection device 150 may be used to detect radiation 105’. The irradiation detection device 150 may be used in a variety of formats. For example, the irradiation detection device 150 may be in the form of a large panel or plate for use in security. The large panel device may range in size from about 2 mm × 2 mm to about 1 m × 1 m. Alternatively, the irradiation detection device 150 may be in the form of a small, encapsulated cylinder for use in drilling and / or mining. The small, encapsulated cylinder device may range in size from about 2 mm in diameter to about 10 mm in diameter and from about 1 cm long to about 5 cm long.

[0142] The irradiation detection device 150 may include a solid electrolyte plate or pellet 110’ made of the polycrystalline material described above. The solid electrolyte pellet 110’ may be made of any of the polycrystalline materials disclosed herein. In one embodiment, solid electrolyte pellet 110’ is an oxygen ion solid electrolyte pellet. For example, the solid electrolyte pellet 110’ may be a CeO2-based solid electrolyte pellet. The solid electrolyte pellet 110’ may also be doped with any of the materials disclosed herein. In one embodiment, the solid electrolyte pellet 110’ may be a gadolinium (Gd) doped solid electrolyte pellet. The Gd doping may range from about 0.5 atm% to about 40 atm%, for example, about 1 atm% to about 20 atm%. Preferably, the solid electrolyte pellet 110’ may be a 3 atm% Gd doped CeO2 (3GDC) solid electrolyte pellet.

[0143] The solid electrolyte pellet 110’ may be prepared as described below in the sample preparation section (e.g., by consolidating a powder through applied pressure and high temperature sintering). Other suitable methods for preparing the solid electrolyte pellet 110’ include, but are not limited to, a traditional ceramic consolidation process (e.g., ressing (uniaxial, isostatic), slip casting, or extrusion), 3D printing, robocasting, pelletization, injection molding, physical vapor deposition (PVD), sputtering, thermal evaporation, pulsed laser deposition, chemical vapor deposition (CVD), atomic layer deposition (ALD), a sol-gel process (e.g., spin coat, dip coat, and / or spray coat), tape casting, slip casting, screen printing, phase inversion, and / or extrusion). The preparation of the solid electrolyte pellet 110’ may also include a heat treatment (e.g., a furnace and / or UV irradiaton).

[0144] The solid electrolyte pellet 110’ may be approximately 100 µm thick to about 100 mm thick, preferably about 500 µm thick to about 10 mm thick. The thickness of the solid electrolyte pellet 110’ may vary depending on the type of polycrystalline material and / or the type of irradiation to be detected. For example, the solid electrolyte pellet 110’may be about 100 µm, about 200 µm, about 300 µm, about 400 µm, about 500 µm, about 600 µm , about 700 µm, about 800 µm, about 900 µm, about 1 mm, about 10 mm, about 20 mm, about 30 mm, about 40 mm, about 50 mm, about 60 mm, about 70 mm, about 80 mm, about 90 mm, or about 100 mm thick, including all values in between. The solid electrolyte pellet 110’ may range in size depending on the type of detector. The solid electrolyte pellet 110’ may be circular, semi-circular, rectangular, square, or any other suitable shape. The solid electrolyte pellet 110’ may range from about 2 mm in diameter to about 10 mm in diameter. For example, the solid electrolyte pellet 110’ may have a diameter of about 2 mm, about 3 mm, about 4 mm, about 5 mm, about 6 mm, about 7 mm, about 8 mm, about 9 mm, or about 10 mm, including all values in between. Alternatively, the solid electrolyte pellet 110’ may range from about 1 cm long to about 5 cm long. For example, the solid electrolyte pellet 110’ may be about 1 cm long, about 2 cm long, about 3 cm long, about 4 cm long, or about 5 cm long, including all values in between. In another embodiment, the solid electrolyte pellet 110’ may range in size from about 2 mm × 2 mm to about 1 m × 1 m. For example, the solid electrolyte pellet 110’ may be about 2 mm × 2 mm, about 2 mm × 5 mm, about 5 mm × 9 mm, about 2 mm × 10 mm, about 5 mm × 5 mm, about 5 mm × 10 mm, about 5 mm × 20 mm, about 5 mm × 30 mm, about 5 mm × 50 mm, about 5 mm × 60 mm, about 5 mm × 70 mm, about 5 mm × 80 mm, about 5 mm × 90 mm, about 5 mm × 100 mm, about 10 mm × 10 mm, about 20 mm × 20 mm, about 30 mm × 30 mm, about 40 mm × 40 mm, about 50 mm × 50 mm, about 50 mm × 100 mm, about 60 mm × 60 mm, about 70 mm × 70 mm, about 80 mm × 80 mm, about 90 mm × 90 mm, about 100 mm × 100 mm, 100 mm × 1 m, about 200 mm × 200 mm, 200 mm × 1 m, about 300 mm × 300 mm, 300 mm × 1 m, about 400 mm × 400 mm, 400 mm × 1 m, about 500 mm × 500 mm, 500 mm × 1 m, about 600 mm × 600 mm, 600 mm × 1 m, about 700 mm × 700 mm, 700 mm × 1 m, about 800 mm × 800 mm, 800 mm × 1 m, about 900 mm × 900 mm, 900 mm × 1 m, or about 1 m × 1 m, including all values in between.

[0145] The solid electrolyte pellet 110’ may have a material density of > 80% density and more preferably > 90% density. For example, the solid electrolyte pellet 110’ may have a material density of about 90% to about 95%, and preferably about 92%. The solid electrolyte pellet 110’ may have one or more grain boundaries 112’ as shown in FIG.1C. The grain boundaries 112’ may be positively charged. Alternatively, the grain boundaries may be negatively charged. The spacing between adjacent grain boundaries 112’ may range from about 10 nm to about several µm (e.g., about 1 µm to about 10 µm) including all values inbetween. For example, the grain boundaries 112’ may be separated from each other by about 200 nm to about 400 nm.

[0146] The irradiation detection device 150 may also include one or more electrodes 130’ as described above. The electrodes 130’ may be made of any suitable metallic material that can be used to collect current and apply a voltage to the solid electrolyte pellet 110’. For example, the electrodes 130’ may be made of any metallic material that may conduct electrons at a conductivity of about > 10-3S / cm and / or any metallic material that may conduct ions at a conductivity of about 10-5S / cm. For example, the electrode 130’ may be platinum (Pt), a platinum alloy, gold, or stainless-steel electrodes. The platinum alloy may include but is not limited to a PtNi alloy (e.g., Pt40Ni60, Pt50Ni50, and / or Pt75Ni25), a PtFe alloy (e.g., Pt3Fe), a PtCo alloys (e.g., Pt3Co and / or PtCo), and / or a PtCu alloys (e.g., Pt0.5Cu0.5). Alternatively, the electrodes 130’ may be made of a highly electronically conducting metallic oxide (e.g., indium tin oxide (ITO)) and / or mixed ionic-electronic conducting metallic oxide (e.g., lanthanum strontium cobalt iron oxide (LSCF). Alternatively, the electrodes may be made of graphite, lithium, lithium cobalt oxide (LiCoO2), lithium manganese oxide (LiMn2O4), lithium iron phosphate (LiFePO4), nickel manganese cobalt oxide (NMC), lithium nickel manganese cobalt oxide (LiNixMnyCozO2), lithium nickel cobalt aluminum oxide (LiNiCoAlO2 or NCA), lithium titanate (Li4Ti5O12), La0.6Sr0.4FeO3−δ (LSF), La0.5Sr0.5Cr0.2Mn0.8O3−δ (LSCrMn), La0.9Sr0.1CoO3−δ (LSC), Ba0.5Sr0.5Co0.8Fe0.2O3-δ (BSCF) , La0.8Sr0.2MnO3 (LSM) , Sm0.5Sr0.5CoO3-δ (SSC) , La0.5Sr0.5MnO3-δ (LSM), LSM with Sc doping, La0.8Sr0.2Sc0.5Fe0.5O3-δ (LSSF), BaCe0.26Ni0.1Fe0.64O3-δ (BCNF10), BaCo0.4Fe0.4Zr0.1Y0.1O3-δ (BCFZY0.1), BaCe0.26Ni0.1Fe0.64O3-δ (BCNF10), NdBa0.5Sr0.5Co1.5Fe0.5O5+δ (NBSCF), BaPr0.8In0.2O3-δ (BPI), palladium (Pd), a palladium alloy (e.g., PdAg and / or PdCu), NaxCoO2, Na3V2(PO4)3, NaxMnO2, NaxTiO2, Na2Ti3O7), K0.3MnO2, K0.55CoO2, KxFe2(CN)6, K2Fe4(CN)6, K2Mn[Fe(CN)6], KxMnO2, CuBr, AgBr, PbBr2, CuBr2, CsPbBr3, CuI, AgI, PbI2, BiI3, and / or SnI4. Preferably, the electrodes 130’ are Pt electrodes.

[0147] In one embodiment, the irradiation detection device 150 may have two electrodes 130a’ and 130b’. In this embodiment, the first electrode 130a’ may be a working electrode and the second electrode 130b’ may be a counter electrode. In another embodiment, the first electrode 130a’ may be a cathode and the second electrode 130b’ may be an anode. Electrodes 130a’ and may be placed on either side of the solid electrolyte pellet 110’ (e.g., sandwiching the solid electrolyte pellet 110’) as shown in FIG.1C. In this embodiment, theelectrodes 130a’ and 130b’ may be spaced approximately 100 µm apart to about 1 cm apart depending on the thickness of the solid electrolyte pellet 110’. For example, the electrodes 130a’ and 130b’ may be spaced about 100 µm, about 200 µm, about 300 µm, about 400 µm, about 500 µm, about 600 µm , about 700 µm, about 800 µm, about 900 µm, about 1 mm, about 10 mm, about 20 mm, about 30 mm, about 40 mm, about 50 mm, about 60 mm, about 70 mm, about 80 mm, about 90 mm, or about 1 cm apart, including all values in between.

[0148] The electrodes 130 may be interdigitated electrodes (e.g., with interwoven fingers to increase the electrode 130 surface area) or planar electrodes. The spacing between the digits of the interdigitated electrodes may be about 1 µm, about 10 µm, about 20 µm, about 30 µm, about 50 µm, about 60 µm, about 70 µm, about 80 µm, about 90 µm, about 100 µm, about 200 µm, about 300 µm, about 400 µm, about 500 µm, about 600 µm, about 700 µm, about 800 µm, or about 900 µm, including all values in between. The electrodes 130a’ and 130b’ may be oval, circular, square, or rectangular in shape. The electrodes 130a’ and 130b’ may be deposited on the solid electrolyte pellet 110’ via sputtering or by painting on an appropriate electrode ink. The electrodes 130a’ and 130b’may be used to apply a voltage to the solid electrolyte pellet 110’. The electrodes 130a’ and 130b’ may be approximately 100 nm thick. The electrodes 130a’ and 130b’ may be the same size (e.g., length, width, or diameter) as the solid electrolyte pellet 110’.

[0149] The detection device 150 may also include a circuit 160. The circuit 160 may include AC source 161. The AC source 161 may be used to operate the detection device 150 and measure the impedance of the detection device 150. The impedance of the detection device 150 may be measured by applying an alternating current (AC) from an AC source 161 with varying frequency (ω) and locking in on the output current to measure amplitude and phase. The current applied by the AC source 161 may range from about 10 mV to about 100 V, including all values in between. For example, the current applied by the AC source 161 may have an amplitude that may range from about 10 mV to about 10 V, or about 50 mV to about 500 mV. The circuit 160 may also include a current meter 162. The current meter 162 may measure the current running through the detection device 150 under the applied AC source 161. The AC source 161 and the current meter 162 may be connected to the electrodes 130a’ and 130b’ through one or more wires 131a’ and 131b’ forming the circuit 160 as shown in FIG.1C. The wires 131a’ and 131b’ may be platinum wires. The wires 131a’ and 131b’ may be attached to the electrodes 130a’ and 130b’ using a silver paste as described above. The circuit 160 may also include a DC source 163. The DC source 163 may be usedto apply a DC overpotential to the circuit 160. The overpotential may range from about 0 V to about 10 V. When the circuit 160 includes the DC source 163, the current meter 162 may also measure any DC overpotential induced current by the DC source 163. The irradiation detection device 150 may also include a thermocouple to measure the temperature of the solid electrolyte pellet 110’.

[0150] Sample preparation:

[0151] Gd0.03Ce0.97O2 (3GDC) powder was synthesized by a solution combustion route, starting from Ce(NO3)3·6H2O (99.99%, Alfa Aesar), Gd(NO3)3:6H2O (Strem Chemicals) (99.99%) precursors and citric acid. The solution was heated on a hot plate, and following gel combustion, the reaction resulted in the formation of a whitish powder. The resulting GDC powders were calcinated at 900°C for 1 hour to obtain the GDC powder crystallized in the fluorite structure. The powder was then pressed into a 30 mm diameter disk with a uniaxial press (1 Tcm-2) and sintered at 1300°C for 8 hours, followed by cooling at 1°C min-1to limit crack formation, resulting in a 3GDC solid electrolyte film pellet with 92% density.100 nm thick platinum electrodes were subsequently DC sputtered on either side of the pellet using Kapton tape to serve as a shadow mask covering a rectangular area of about 5 mm ^ 9 mm. The electrode was subsequently attached to Pt wires in both the furnace setup and custom-built microheater device using a commercial Ag paste purchased from Fuel Cell Materials (#321201) and dried using a hand-held heat gun for 10 minutes and then subsequently dried in a furnace to 450 °C while holding for 1 hour.

[0152] Another suitable sample preparation method may include a powder fabrication method (co-precipitation) followed by surface infiltration and subsequent powder consolidation and sintering. For example, a GDC powder may be synthesized using a coprecipitation method. Cerium nitrate (Ce(NO3)3·6H2O) and gadolinium nitrate (Gd(NO3)3·6H2O) may be dissolved in distilled water to form a solution of gadolinium and cerium in a 3:97 stoichiometric ratio. Ammonium carbonate may then be dissolved in distilled water to reach a concentration of 0.5M. This solution may be used as the precipitation agent and may then poured drop-wise into the stirred nitrate solution at room temperature to elicit precipitation. The resulting precipitate may be filtered from the solution using a filter (e.g., a Buchner filter) connected to a vacuum pump. While filtering, the precipitate may be washed in ethanol to remove surface hydroxyl groups that can lead to the formation of micron-order agglomerations. The filtered powder may be dried for 24 hours at about 100 °C, and calcined at about 750°C for 2 hours to obtain a GDC nanocrystallinepowder. To synthesize an infiltrated powder, the GDC powder may be sonicated in anhydrous ethanol for approximately 15 minutes to break up particle agglomerations. Simultaneously, either aluminum nitrate (Al(NO3)3*9H2O) or gallium nitrate (Ga(NO3)3*8H2O) may be dissolved in ethanol. The GDC suspension and impurity solution may then be combined in a crystallization dish on a hot plate set to about 120 °C until all liquid evaporated while being stirred with a magnetic stirrer. The resulting powder may be calcined at about 400°C for 2 hours. Batches of powder may be uniaxially pressed under about 2000 psi, then isostatically cold pressed under about 30,000 psi in round dies (d = 0.5 in). Samples can then be sintered in a tube furnace at about 1100ºC for 2 hours, about 1200ºC for 2 hours, or about 1300ºC for 2 hours in O2.

[0153] Further details and examples of irradiation detection devices including ion conducting polycrystalline materials can be found in International Application No. PCT / US2023 / 086168, filed December 28, 2023, and in U.S. Application No.19 / 042,616, filed January 31, 2025, which are incorporated herein by reference in their entirety for all purposes.

[0154] Origin of grain boundary space charge potentials

[0155] Space charge potential barriers exist at grain boundaries in ion-conducting ceramics, including the ion conducting polycrystalline materials disclosed herein, because the interfaces are net-charged. This occurs due to the unintentional segregation of charge point defects (e.g., intrinsic defects and / or extrinsic foreign elements) from the bulk lattice towards the grain boundary interface. The interfaces may be positively charged or negatively charged. For example, intrinsic defects may include, but are not limited to, point defects already present in the lattice of the host due to intrinsic reasons and / or intentional ones (e.g., oxygen vacancies in metal oxides, Li interstitials or lithium vacancies in Li conducting oxides, and / or protons in proton conductors) and / or bulk dopants (e.g., acceptor and / or donor dopants) that were intentionally added to the ion-conducting polycrystalline material. Extrinsic foreign elements may include, but are not limited to, impurity elements that may be present unintentionally (e.g., trace amount of cations such as Silicon, aluminum, sodium, calcium, etc.). These extrinsic foreign elements may be present in the ion-conducting polycrystalline bulk material at about < 100 ppm (e.g., 100 ppm, 90 ppm, 80 ppm, 70 ppm, 60 ppm, 50 ppm, 40 ppm, 30 ppm, 20 ppm, and / or 10 ppm, including all values in between). These extrinsic foreign elements may be present in the ion-conducting polycrystalline bulk material in the lattice as either interstitials or substitutional elements, for example.

[0156] Controlling the sign and magnitude of the net charge at the interface may allow the space charge potential height to be controlled. Disclosed herein are methods for selecting and / or introducing impurities (e.g., a selected element) directly at the grain boundary interfaces to intentionally control the net charge of the grain boundary interfaces and, therefore, control the space charge potential barrier heights These methods may allow for the control of the detection threshold, dynamic range, and / or operating temperature of ion-conducting polycrystalline radiation detectors (e.g., detection device 100 / 150).

[0157] FIGS.1D–1E show schematic representations of light modulated ionic grain boundary 112 / 112’ resistance of the radiation detector device 100 / 150. As shown in FIGS. 1D–1E grain boundaries 112 / 112’may act as barriers to the transport of ions in the dark (FIG. 1D, top) preventing the flow of current. Under above band gap illumination 105 (e.g., radiation) that may generate electron-hole pairs (FIG.1E, top) the potential barriers to the grain boundaries 112 / 112’ may decrease, and ultimately collapse under high illumination 105 intensities, allowing for larger flows of ions, commensurate with the radiation dose, to occur. The potential barriers at the interface may lead to a modulation of the mobile defect concentration in the proximity of the interface. Under dark conditions (FIG.1D, bottom) the defect concentration (e.g., oxygen vacancies in the case of an oxygen ion conducting polycrystalline material) may be depleted by orders of magnitude in concentration in the proximity of the interface. Upon above band gap illumination 105 (FIG.1E, bottom), the potential may be reduced and therefore the depletion of defects may be reduced leading to a decreased resistance of the interface to ionic motion, enabling higher currents to run through the radiation detector device 100 / 150 commensurate with the radiation dose (e.g., the illumination 105).

[0158] Grain Boundary Space Charge Potential as a Means to Control the Detection Threshold and / or Dynamic Range of ion-conducting Radiation Detectors

[0159] The physics controlling the optoionic response of an ionic conducting ceramic, including the ion conducting polycrystalline materials disclosed herein, may be similar, in principle, to physics originally derived for more traditional polycrystalline semiconductor devices. Without being bound to a particular theory, the mechanism may rely on the competition of the photo-generation rate G of electron-hole pairs within the grains, the transport rate of photo-generated electrons into the space charge region in terms of the current density (Je), the trapping rate of photo-generated species at the grain boundary core energy traps (c), the hole current density of valence band holes towards the grain boundarycore (Jh), and / or the recombination of trapped electrons with the depleted holes within the grain boundary core (Jr). Accounting for the different rates of defect transport and recombination kinetics, an equation describing the dependence of the space charge potential at the grain boundary core on the photo-generation rate G can be obtained. This equation may also be related to the optoionic response as illustrated in FIGS.2A and FIG.2B.

[0160] As shown in FIGS.2A, it was observed that for all space charge potentials present at grain boundaries, an increasing photogeneration rate may depress the boundary barrier heights, which may lead, in turn, to exponentially increasing conductance G. The higher the initial space charge potential, the lower the onset value of G, for which a noticeable barrier depression may be observed. Moreover, the higher the dark potential, the higher the magnitude of the conductance change under equivalent irradiation. This may also result in a higher dynamic range for higher potential barriers.

[0161] FIG.2A allows for the evaluation of multiple detector-relevant figures of merit, including detection threshold (e.g., the lowest photon flux needed to generate a measurable signal) and dynamic range (e.g., the difference between the maximum and the minimum signals).

[0162] From FIG.2B, it can be seen that higher grain boundary space charge barriers may lead to higher optical sensitivity. Additionally, the onset of generation rate may alter the space charge potential barrier. A smaller onset of generation rate may be observed for higher barriers compared to smaller barriers (e.g., Gonset (0.25 eV) is approximatly10-3Suns versus Gonset (1 eV) is approximately 10-9Suns – a factor of about a million lower), while the magnitude of the response may be more significant for equivalent irradiation intensities. Moreover, higher space charge barriers may be expected to have exponentially larger dark resistivities (e.g., 1 Mohm to about > 100 Tohm), thus resulting in a much broader dynamic range for the radiation detection device 100 / 150.

[0163] Selection Criteria for Precursors Impurities

[0164] As discussed above, carefully selecting appropriate elements or impurities, referred to herein as precursor elements, to be introduced within the grain boundary interfaces of an ion conducting polycrystalline material (e.g., the solid electrolyte film 110 and / or the solid electrolyte pellet 110) may allow the sign and magnitude of the space charge potentials of the grain boundaries (e.g., grain boundaries 112 / 112’) to be controlled. For example, the higher the space charge potentials of the grain boundaries 112 / 112’, the higher the optoionic sensitivity of the grain boundaries 112 / 112’ and thus the radiation detectordevice 100 / 150. Thus, selecting an appropriate precursor element to be introduced within the grain boundary 112 / 112’ interfaces may increase the sensitivity of radiation detector device 100 / 150.

[0165] Therefore, when selecting precursor elements (e.g., impurities), choosing precursor elements expected to create a large potential barrier (e.g., the solid electrolyte film 110 and / or the solid electrolyte pellet 110) may be beneficial. For example, for an ion conducting polycrystalline materials whose majority ionic carriers are net positively charged (e.g., ^^2+^^ , ^^+^^ , ^^^^+, ^^^^+^^ ,^^^^^2^+) the precursor elements introduced at the grain boundaries (e.g., grain boundary 112 / 112’) should preferably have a net positive charge. In other words, for grain boundaries 112 / 112’ that are positively charged, introducing positively charged precursor elements may increase the positive charge of the grain boundaries 112 / 112’ and thus the sensitivity of radiation detector device 100 / 150. Preferably the precursor elements may create a large potential barrier. The sign (e.g., charge) of the potential barrier should preferably be the same as the sign of the majority ionic carriers in order for the barrier to oppose ionic transport outside of the grain boundary. For example, for an oxygen ion conductor whose majority ionic carriers are net positively charged, preferably the precursor elements introduced at the grain boundaries should have a net positive charge. For example, for an ion conductor whose majority ionic carriers are net negatively charged, preferably the precursor elements introduced at the grain boundaries should have a net negative charge.

[0166] Alternatively, where the majority ionic carriers of the ion conducting polycrystalline materials (e.g., the solid electrolyte film 110 and / or the solid electrolyte pellet 110) are negatively charged (e.g.,the precursor elements introduced at the grain boundaries (e.g., grain boundary 112 / 112’) should preferably have a net negative charge. In other words, for grain boundaries 112 / 112’ that are negatively charged, introducing negatively charged precursor elements may increase the negative charge of the grain boundaries 112 / 112’ and thus the sensitivity of radiation detector device 100 / 150.

[0167] The selection of the precursor element may be based on several factors, including, but not limited to, the method of incorporation (e.g., substitutional incorporation or interstitial incorporation), the oxidation state of the precursor element, the oxidation state of the majority ionic carrier of the ion conducting polycrystalline material, the charge polarity of the precursor element, the charge polarity of the majority ionic carrier of the ionconducting polycrystalline material, the solubility of the precursor element in the ion conducting polycrystalline material, the size of the precursor element, the size of the majority ionic carrier of the ion conducting polycrystalline material, and / or the number of sites of the ion conducting polycrystalline material. Each of these factors will be described in more detail below. These factors may govern the choice of precursor element to create controlled space- charge potentials at grain boundaries in the ion conducting polycrystalline materials disclosed herein. These factors may ensure that the precursor element that occupies the intended sites (substitutional or interstitial), carry the desired net charge, and / or preferentially segregate to interfaces rather than dissolve in the bulk of the ion conducting polycrystalline materials disclosed herein. By applying these factors, one can systematically choose precursor elements that occupy the intended sites (substitutional or interstitial) of appropriate size, valence, and / or solubility to engineer large, tunable space-charge potentials at grain boundaries 112 / 112’ across a wide range of ionic conductors (e.g., the ion conducting polycrystalline materials disclosed herein).

[0168] Lattice solubility

[0169] Since as the precursor elements are expected to be selectively introduced and remain at the grain boundaries 112 / 112’, selecting precursor elements with low and / or limited thermodynamic solubility within the host (e.g., the ion conducting polycrystalline material) lattice may favor the segregation of the precursor elements towards the grain boundary 112 / 112’.

[0170] Alternatively, if the precursor elements can be introduced into the grain boundary 112 / 112’, while kinetically being prevented from dissolving in the bulk of the ion conducting polycrystalline material, the solubility of the precursor elements may be less relevant. For example, using rapid quenching and / or controlling the atmosphere may kinetically prevent precursor elements from dissolving in the bulk of the ion conducting polycrystalline material regardless of the solubility of the precursor elements. Rapid quenching may include heating the sample up to a high temperature for a short period of time to achieve sintering and then very rapidly cooling the sample down (e.g., by removing the sample from the high temperature furnace and letting it air cool or plunging it into oil and / or water) to prevent elements from segregating out of the grain boundary interface. The atmosphere may be controlled by sintering under a gas (e.g., argon, Hydrogen) and / or under a vacuum. Atmospheric control may prevent the precursor elements from dissolving into the host lattice and thus may ensure that the precursor elements remain at the grain boundaries.

[0171] Interstitial insertion:

[0172] FIG.2C shows a schematic of interstitial insertion of precursor elements 170 into a grain boundary 112 / 112’. As shown in FIG.2C, interstitial insertion involves the insertion of the precursor element into a space within the grain boundary. The space may not have been previously occupied by another element (e.g., by an element of the host ion- conducting polycrystalline material). While FIG.2C shows the grain boundary 112 / 112’ and the precursor elements 170 as being positively charged, the grain boundary 112 / 112’ and / or the precursor elements 170 may be negatively charged instead.

[0173] Ionic Radius

[0174] If the ionic radius of the precursor element, for example, a cation, is much smaller than the constituent elements of the host lattice, then the precursor element may favorably enter interstitial sites. A constituent element may be any element of the host matrix that represents >10% of the lattice sites in the host matrix. For example, if the host matrix is CeO2, then Ce and O are constituent elements of the host matrix. Preferably, the ionic radius of the precursor element may be about >20% smaller than that of the host cations. For example, the ionic radius of the precursor element may be about 20%, about 25%, about 30%, about 35%, about 40%, about 45%, about 50%, about 55%, about 60%, about 65%, about 70%, about 75%, about 80%, about 85%, about 90%, or about 95%, including all values in between, smaller than the constituent elements of the host lattice.

[0175] In a single-cation host lattice, preferably Rprecursor ≤ 0.70 × Rhost, where R refers to ionic radius.

[0176] In a multi-cation host lattice, preferably Rprecursor ≤ 0.70 × Rmin, where R refers to ionic radius and Rmin refers to radius of the smallest host-cation. Preferably, the ionic radius of the precursor element (Rprecursor) is at least 30% smaller than every host cation. For example, the ionic radius of the precursor element may be about 30%, about 35%, about 40%, about 45%, about 50%, about 55%, about 60%, about 65%, about 70%, about 75%, about 80%, about 85%, about 90%, or about 95%, including all values in between, smaller than every host cation in the multi-cation host lattice.

[0177] Oxidation State

[0178] Ionized elements from periodic table columns 1, 2, 3, and / or 4 (e.g., Li⁺, Mg²⁺, Al³⁺, etc.) naturally possess a positive oxidation state (e.g., carry a positive core charge), and thus result in a positive charge as an interstitial.

[0179] Ionized elements from periodic table columns 5, 6, and / or 7 (e.g., F⁻, Cl⁻, Br⁻, etc.) possess a negative oxidation state (e.g., carry a negative core charge), and thus result in a negative charge as an interstitial.

[0180] Substitutional insertion:

[0181] FIG.2D shows a schematic of substitutional insertion of precursor elements 170 into a grain boundary 112 / 112’. As shown in FIG.2D, substitutional insertion involves replacing an existing element in the host (e.g., ion-conducting polycrystalline material) with the precursor element within the grain boundary. While FIG.2D shows the grain boundary 112 / 112’ and the precursor elements 170 as being positively charged, the grain boundary 112 / 112’ and / or the precursor elements 170 may be negatively charged.

[0182] Ionic Radius

[0183] If the ionic radius of the precursor element, for example a cation, is about the same size as the constituent elements of the host lattice, then the precursor element may favorably enter substitutional sites in the host lattice (e.g., in the ion-conducting polycrystalline material). Preferably the ionic radius of the precursor element, is within about 20% of the ionic radius of the host element(s). For example, the ionic radius of the precursor element may be within about 1%, about 2%, about 3%, about 4%, about 5%, about 6%, about 7%, about 8%, about 9%, about 10%, about 11%, about 12%, about 13%, about 14%, about 15%, about 16%, about 17^%, about 18%, about 19%, or about 20%, including all values in between, of the ionic radius of the host element(s). Preferably (Rprecursor – Rhost) / Rhost ≤ 0.20, where R refers to ionic radius.

[0184] Oxidation State

[0185] If the oxidation state of the precursor element is higher than that of the constituent elements of the host lattice that it substitutes, then a positive charge will be generated. When the precursor element’s oxidation state (n) exceeds that of the host (m), each substitution by a precursor element may contribute +(n–m) to the oxidation state of the grain-boundary core.

[0186] If the oxidation state of the precursor element is smaller than that of the constituent elements of the host lattice that it substitutes, then a negative charge will be generated. When the precursor element’s oxidation state is lower than that of the host, each substitution precursor element may contribute –(m–n) to the oxidation state of the grain- boundary core.

[0187] Oxidation-State and / or Charge Polarity:

[0188] The core of the grain boundary may also be referred to herein as the space- charge core. The space-charge core may have a net charge. The net charge may be due to either intrinsic and / or extrinsic defects segregating to the grain boundary interface. The purposeful insertions of precursor elements into the interface with different charges may allow for the modification of the number and / or amount of charges present at the grain boundary interface. Mobile defects may be present in the host lattice (e.g., in the ion- conducting polycrystalline material) but the concentration of mobile defects may be different in the close vicinity of the grain boundary, because of the space charge effect. Thus, the space charge effect may cause the depletion of mobile defects at the grain boundaries and therefore increase grain boundary resistance.

[0189] For example, positive space-charge cores may oppose positively charged mobile defects (e.g., VO²⁺, H⁺, Li⁺, Ag+). Negative space-charge cores may oppose negatively charged mobile defects (e.g., VLi-, Oi²⁻, VAg-).

[0190] Multi-cations and Site-Competition:

[0191] For interstitials, preferably the precursor element meets the ≥30 % size- reduction criterion (e.g., an ionic radius size of about ≥30 %) with respect to every cation type in the lattice of the ion-conducting polycrystalline material. In other words, preferably the ionic radius of the precursor element is about ≥30 % smaller than the ionic radius of the host lattice cations. This may allow the precursor elements to fit into interstitial sites in the grain boundaries. For example, preferably the precursor element is about 30%, about 35%, about 40%, about 45%, about 50%, about 55%, about 60%, about 65%, about 70%, about 75%, about 80%, about 85%, about 90%, or about 95%, including all values in between, smaller than every cation in the lattice of the ion-conducting polycrystalline material, This may help prevent unintended substitutional uptake of the precursor element.

[0192] For substitutional strategies in perovskites (e.g., BaZrO₃, BaCeO₃, LiLaTiO₃) or garnets (e.g., LLZO (Li₇La₃Zr₂O₁₂)) that may have multiple cation sites (e.g., A- and B- sites), the ±20 % radius window may be evaluated independently for each host site. For example, the ionic radius of the precursor element is within about 1%, about 2%, about 3%, about 4%, about 5%, about 6%, about 7%, about 8%, about 9%, about 10%, about 11%, about 12%, about 13%, about 14%, about 15%, about 16%, about 17%, about 18%, about 19%, or about 20%, including all values in between, of the ionic radius of the host site (e.g., the A- and / or B- host site).

[0193] For substitutional strategies, in systems that may have multiple cations (e.g., ternary, quaternary, and above), preferably the ionic radius differences between the precursor elements and all host elements are evaluated to determine what lattice site the precursor elements may sit on and therefore determine the charge the precursor elements should possess. In other words, the precursor elements may substitute only one of the lattice sites in a multiple cation system. Alternatively, the precursor elements may substitute all of the lattice sites in a multiple cation system.

[0194] Examples

[0195] 1. Oxygen ion conducting Acceptor-Doped Ceria (CeO₂):

[0196] An ion-conducting polycrystalline material of oxygen ion conducting Acceptor-Doped Ceria (CeO₂) may include oxygen vacancies and VO²⁺ may be the primary mobile ionic defect. To raise the space-charge potential barrier and suppress vacancy percolation across the grain boundaries, preferably a positive charge is introduced at the grain boundary core. In other words, preferably the precursor element has a positive charge (e.g., an element from periodic columns 1, 2, 3, and / or 4, including but not limited to, Li⁺, Mg²⁺, Al³⁺, etc.

[0197] For an interstitial precursor element, Rprecursor ≤ 97  picometers (pm) × 0.70 = 68 pm. In other words, preferably the Rprecursor is less than 68 pm. For example, Rprecursor may be 67 pm, 60 pm, 55 pm, 50 pm, 45 pm, 40 pm, 35 pm, 30 pm, 25 pm, 20 pm, 15 pm, 10 pm, or about 5 pm, including all values in between. The interstitial precursor element may include, but is not limited to, B³⁺ (about 27 pm) and / or Al³⁺ (about 53 pm), both of which are well below the 68 pm cutoff and carry +3 charge. Alternatively, the interstitial precursor element may include Ga³⁺ (about 62 pm). While Ga³⁺ is near the 68 pm threshold, it may stably occupy interstitial voids when introduced under O₂-lean, high-temperature atmospheres (e.g., about > 1000 ℃). Additionally, B³⁺, Al³⁺, and / or Ga³⁺ may exhibit limited bulk solubility in fluorite CeO₂, which may encourage segregation of these precursor elements to the grain boundaries under typical sintering and / or hot-pressing conditions.

[0198] For a substitutional precursor element, preferably the ionic radius of the substitutional precursor element is about 70 ppm to about 120 ppm, for example, about 72 ppm to about 118 ppm, or about 77.6 to about 116 pm, including all values in between. The substitutional precursor element may include, but is not limited to, Nb⁵⁺ (about 78 pm) and Ta⁵⁺ (about 78 pm). Nb⁵⁺ and / or Ta⁵⁺ may be substituted on Ce⁴⁺ sites and each may generate a +1 net charge. The ionic radii of Nb⁵⁺ and / or Ta⁵⁺ may be within ±20 % the ionic radius ofCe⁴⁺ (e.g., about 97 pm). Alternatively, the substitutional precursor element may include U⁵⁺ (about 90 pm). U⁵⁺ may be substituted on Ce⁴⁺ sites and may generate a +1 charge per substitution. Additionally, U⁵⁺ may have low solubility in the bulk, which may encourage segregation of U⁵⁺ to the grain-boundary plane.

[0199] 2. Proton-Conducting Perovskites (BaZrO₃, BaCeO₃):

[0200] Hydrated BaZrO₃ and / or BaCeO₃ may both conduct H⁺ as the primary mobile species. A positive boundary core may repel protons and may enhances the overall electrolyte performance of BaZrO₃ and / or BaCeO₃ by reducing intergranular resistivity.

[0201] Interstitial precursor elements for BaZrO₃ (Rprecursor ≤ 72  pm× 0.70 = 50 pm) may include, but are not limited to, Ba²⁺ (about 135 pm), Zr⁴⁺ (about 72 pm), B³⁺ (about 27 pm), and / or Be²⁺ (about 45 pm). B³⁺ and / or Be²⁺ may introduce +2 or +3, respectively, per ion. Ba²⁺ and / or Zr⁴⁺ may introduce +2 or +4, respectively, per ion.

[0202] Substitutional precursor elements for the A-site (e.g., Ba²⁺ = 135 pm ± 20 %), may include, but are not limited to, La³⁺ (about 103 pm) and / or Bi³⁺ (about 117 pm), which may substitute for Ba²⁺. La³⁺ and / or Bi³⁺ may each deliver a +1 net charge.

[0203] Substitutional precursor elements for the B-site (e.g., Zr⁴⁺ = 72 pm ± 20 % and / or Ce⁴⁺ = 97 pm ± 20 %) may include, but are not limited to, Nb⁵⁺ (about 65 pm) and / or Ta⁵⁺ (about 64 pm). Nb⁵⁺ and / or Ta⁵⁺ may each deliver a +1 net charge to the grain boundary core. Additionally, dual-site co-doping (e.g., Ba→La and / or Zr→Ta) may offer independent tuning of A- and / or B-site contributions to the space-charge potential. For example, if the appropriate precursor element is selected, one may be able to selectively introduce elements only on the B site (e.g., by matching the ionic radius of the precursor elements to the B-site ions as described above) or only on the A site (e.g., by matching the ionic radius of the precursor elements to the A-site ions as described above). Additionally, different precursor elements may be substituted on both the A and B site simultaneously (e.g., a first precursor element substitute on the A site and a second precursor element substituted on the B site), provided the ionic radius of each of the A and B precursor elements matches the A and B site, respectively.

[0204] 3. LiLaTiO₃ Perovskite:

[0205] LiLaTiO₃ may support Li⁺ conduction through interstitial channels. A positive core charge may be introduced to impede Li⁺ leakage across grain boundaries, which may enhance net ionic selectivity.

[0206] Interstitial precursor elements for LiLaTiO₃ (Rprecursor ≤ 61 pm × 0.70 = 43 pm) may include, but are not limited to, B³⁺ (about27 pm). B³⁺ is uniquely small enough to occupy octahedral and / or tetrahedral voids in LiLaTiO₃. Additionally, the +3 charge of B³⁺ may sharply raises the local grain barrier charge.

[0207] Substitutional precursor elements for the La³⁺ site (103 pm ± 20 % = about 82pm to about 124 pm), include, but are not limited to, Ce⁴⁺ (about 97 pm). Ce⁴⁺ may introduce +1 per substitution. Substitutional precursor elements for the Ti⁴⁺ site (61 pm ± 20 % = about 49 pm to about 73 pm) include, but are not limited to, V⁵⁺ (about 60 pm, +1) and / or Mo⁶⁺ (about 73 pm, +2). Each of V⁵⁺ and / or Mo⁶⁺ may boost the grain boundary charge.

[0208] 4. Garnet LLZO :

[0209] LLZO’s may on Li-vacancy (VLi-) migration, which carries a net negative charge. To suppress intergranular vacancy percolation, negative grain- boundary cores may be introduced.

[0210] Large anions (e.g., F⁻, Cl⁻) may fail the ≥ 30 % rule for both La³⁺ and / or Zr⁴⁺. Thus, large anions may not preferentially occupy interstitials for LLZO. In other words, precursor elements that may enter the LLZO lattice interstitially and have a negative may be less favored.

[0211] Substitutional precursor elements for the Li⁺ site (76 pm ± 20 % = about 61 pm to about 91 pm) include, but are not limited to, Mg²⁺ (about 72 pm, –1) and / or Zn²⁺ (about 74 pm, –1). Mg²⁺ and / or Zn²⁺ may substitute into Li sites and may introduce a net negative charge. Substitutional precursor elements for the La³⁺ site (103 pm ± 20 % = about 82 pm to about 124 pm) include, but are not limited to, Ca²⁺ (about 100 pm, –1) and / or Sr²⁺ (about 118 pm, –1). Ca²⁺ and / or Sr²⁺ may substitute into La sites and may each impart a –1 charge. Substitutional precursor elements for the Zr⁴⁺ site (72 pm ± 20 % = about 58 pm to about 87 pm) include, but are not limited to, In³⁺ (about 80 pm, –1) and / or Sc³⁺ (about 75 pm, –1). In³⁺ and / or Sc³⁺ may substitute into Zr⁴⁺ sites and may each impart a –1 charge.

[0212] 5. Superionic RbAg₄I₅:

[0213] Ag⁺ mobility may dominate RbAg₄I₅ conduction. By creating positive grain- boundary cores, Ag⁺ migration may be impeded and selectivity under an applied field may be improved.

[0214] Interstitial precursor elements for RbAg₄I₅ (Rprecursor ≤ 115 pm × 0.70 = 80 pm) may include, but are not limited to, Li⁺ (about 76 pm). Li⁺ may occupy interstitial cavities in the rigid I⁻ sublattice and may yield a charge of +1.

[0215] Substitutional precursor elements for the Rb⁺ site (152 pm ± 20 % = about 122 to about 182 pm) include, but are not limited to, Ba²⁺ (about 135 pm, +1). Ba²⁺ may fit well (e.g., substitute into Rb⁺ sites) and introduce a net +1 charge. Substitutional precursor elements for the Ag⁺ site (115 pm ± 20 % = about 92 to about 138 pm) include, but are not limited to, Cd2+(about 95 pm, +1) and / or Hg2+(about 102 pm, +1). Cd2+and / or Hg2+may both substitute into Ag⁺ sites and introduce a net +1 charge.

[0216] In the examples disclosed herein careful control of the processing temperature, atmosphere, and / or precursor element chemistry may assist in kinetically trapping the chosen precursor element at the grain boundaries while also preventing dissolution of the precursor element into the bulk of the ion-conducting polycrystalline material. Also disclosed herein below are experimental protocols, microstructural analyses, and measurements of space-charge potentials and optoionic sensitivity for each precursor– host combination.

[0217] Fabrication Methods

[0218] Successful fabrication may rely on two factors: 1) inserting the desired precursor element 170 directly at the grain boundary and 2) preventing the unintentional segregation of other charged impurities, for example, intentional bulk dopants used to increase the bulk conductivity of the grains, towards the interface that may compete with the charge of the precursor element 170 addition. The objective is to allow for selective insertion of the precursor element 170. This may be achieved by direct insertion of the precursor element 170 and by minimizing temperature and time to prevent alternative impurity or defect segregation at the grain boundaries 112 / 112’. Below, we describe three methods, one for bulk ceramics (e.g., solid electrolyte plate or pellet 110’) and two for thin films (e.g., solid electrolyte film 110), that may allow for direct elemental insertion of precursor element 170 into the grain boundaries (e.g., grain boundaries 112 / 112’), while also reducing bulk segregation of other elements (e.g., charged impurities and / or bulk dopants).

[0219] A. Bulk Ceramics:

[0220] FIG.12A is a schematic representation of the process diffusion of precursor elements 1270 in a bulk ceramic 1210. The bulk ceramic 1210 may be made of any of any of the polycrystalline materials disclosed herein. For the bulk ceramic 1210 the surface of rawpowder particles 1251 of the polycrystalline materials may be coated with the desired precursor elements 1270 and high pressure may be used to consolidate the powder into one or more green bodies 1252. The green bodies 1252 may then be sintered at moderate temperatures (e.g., at about 900 ºC to about 1100 ºC for 2 hours) to densify the material into the bulk ceramic 1210. Sintering at higher temperatures (e.g., > 1100 ºC) may cause the segregation of undesired elements (e.g., bulk dopants) to the grain boundary. This approach may ensure that the desired impurities may be already present at the grain boundary interfaces 1212, since the precursor elements 1270 were already coating the particle 1251 surfaces. This approach may also reduce the segregation of oppositely charged dopant cations from the centers of the powder particles, given their low diffusivities at these moderate temperatures.

[0221] Protocol

[0222] FIG.12B provides an overview of the method 1290 for the diffusion of precursor elements 1270 in a bulk ceramic 1210. First, at step 1291, the raw bulk powder particles 1251 may be sonicated (e.g., to about < 1-micron particle size) in alcohol for about 15 minutes to break up particle agglomerations. Second, at step 1292, the precursor elements 1270 may be dissolved in alcohol. The precursor elements 1270 may be present in a salt form of the precursor elements 1270, for example, in a nitrate form (e.g., aluminum nitrate (Al(NO₃)₃), boron nitride, lithium nitrate (LiNO3), gallium nitrate, etc.). Any nitrogen present in the salt form of the precursor elements 1270 may be burned off during processing of the precursor elements 1270 (e.g., by heating the precursor elements 1270 as described below). Preferably the precursor elements 1270 have a concentration no higher than about 0.5 M. For example, the precursor elements 1270 may have a concentration of about 0.1M, about 0.2M, about 0.3M, about 0.4M, or about 0.5M, including all values in between. Third, at step 1293, the suspended powder and precursor element solutions may be combined (e.g., particles of the ion-conducting polycrystalline material may be mixed into the precursor element solutions) in a crystallization dish and stirred on a hot plate at about 80 °C to about 120 °C until all the liquid has evaporated. At this step, particles of the ion-conducting polycrystalline material that are coated with the precursors may be formed. Fourth, at step 1294, the resulting dried powder may be collected and then hand ground in a mortar and pestle to form a fine powder. Fifth, at step 1295, the ground powder may be ground in a ball mill for about 24 hours and then dried in an oven at about 120 °C. The fourth step 1294 and the fifth step 1295 may ensure proper homogenization is achieved. Sixth, at step 1296, thedried powder may be hand-ground in a mortar and pestle. Seventh, at step 1297, the powder may then be calcined at about 400 ℃ to about 600 ℃ for about 2 hours at step 1298. This may partially decompose the precursor elements on the surface of the powder. Eighth, at step 1299, the powder may be recovered consolidated into a green body 1252 via high-pressure pressing (e.g., uniaxial or isostatic). Ninth, at step 1300, the consolidated body may be sintered at a reduced temperatures for a short duration to produce the bulk ceramic 1210. Preferably, the green bodies 1252 may be sintered at a temperature below 1100 ºC for less than 2 hours. For example, the green bodies 1252 may be sintered at a temperature of about 600 ºC, about 650 ºC, about 700 ºC, about 750 ºC, about 800 ºC, about 850 ºC, about 900 ºC, about 950 ºC, about 1000 ºC, about 1050 ºC, or about 1100 ºC, including all values in between. For example, the green bodies 1252 may be sintered for about 30 minutes, about 1 hour, about 1.5 hours, or about 2 hours, including all values in between. Tenth, at step 1301, surface metallization (e.g., electrodes 1230 and a circuit 1260) may be applied to the bulk ceramic 1210 to measure conductance and operate the radiation detection device 1250. The electrodes 1230 and circuit 1260 may include any of the electrodes and / or circuits described herein.

[0223] Results

[0224] An example where precursor elements Ga (2 at%), Al (1.7 at%, and B (2 at%) were infiltrated on 3 at% Gd doped CeO2 powders and then consolidated following the protocol described above. FIG.3 is a plot of the logarithm of conductance for the bulk and grain boundary for each sample. It can be seen that while the bulk conductance (1 / R) may be weakly affected by the infiltration protocol, the grain boundary resistance may vary by orders of magnitude and may be associated with increased activation energies. These activation energy and resistance increases are consistent with the selective modulation of the grain boundary potential barrier heights. While all three precursor elements are expected to have a 3+ oxidation state and their ionic radius, in general, is about < 20% of that of the host Ce cation, all three precursor elements are therefore expected to enter lattice sites interstitially. The general trend for effectiveness in space charge engineering follows B>Al>Ga infiltrated powders versus uninfiltrated powders, which follows the opposite trend in ionic radius B<Al<Ga<<Ce.

[0225] Next similar temperature dependence conductance curves between the boron and previous aluminum-infiltrated samples were compared under similar UV light intensity conditions (e.g., 1018Photons / (s)). FIG.4 shows that the conductance values under UVintensity, irrespective of temperature, were the same for all the samples (top three curves). However, as can be seen, the dark resistance is dramatically different for each sample, which results in significantly different resistance modulation ratios with UV (×104for B versus ×103for Al at 1100 °C versus 101for Al at 1300 °C) (bottom three curves) and higher operational ranges where a perceptible optoionic ionic response can be obtained. For example, optoionic response may be limited to below 100 ºC for a barrier height of 0.83 eV, while it may extend to 200 ºC for a barrier height of 0.93 eV and to 300 ºC for a barrier height of 1.21 eV. This observation may be correlated with the increasing dark resistivity due to the higher activation energies (1.21eV, 0.93eV, and 0.83eV) associated with higher grain boundary space charge potentials.

[0226] FIG.5A is a plot of the logarithm of the ratio of the conductance under UV and in the dark against the logarithm of the photon flux. This plot is the same as that in FIG. 2B. As shown, the optoionic response for all samples may generally exhibit a linear trend with photon flux without observable saturation, as expected from the theory described in FIG.2B. While this may prevent a saturation point (upper limit response) from being established, the maximum measured conductance modulation magnitude may vary ×26 (Al) versus ×1130 (B) samples for the highest photon flux (e.g., 1018Photons / s), indicating a much higher optoionic response, and dynamic range in the boron samples. The same can be seen for detection linearity when looking at the slopes of the curves varying from 1×10-17(For Al) versus 5×10-15(for B), indicating the much higher sensitivity of the boron sample to photon flux. Extrapolated to lower light intensities, the minimum photon flux required to generate a measurable signal, which measures the detection threshold, can be estimated. As seen in FIG.5A, the B sample (top curve) has a detection threshold ×103smaller than that of the Al sample. The boron-doped sample may be expected to show resistance modulation down to 1014Photons / s versus 1017Photons / (s) for the Al sample. This is a direct example of the increasing radiation sensitivity of these ceramics with higher grain boundary space charge potentials. For example, the ceramics with higher grain boundary space charge potentials disclosed herein may have a sensitivity ∆R / R of about 103. FIG.5B is a graph of 1 / R versus reciprocal thermal energy 1 / kT for the B infiltrated sample at different GB conductance values. FIG.5C is a graph showing the conductance ratio under light and in the dark as a function of photon flux for the Al infiltrated sample (bottom curve), the B infiltrated sample (top curve), and an uninfiltrated sample (middle curve).

[0227] B. Thin film devices on an insulating substrate: In-diffusion of impurity element from a thin layer or from the substrate

[0228] For thin films (e.g., solid electrolyte film 110), the diffusion of elemental impurities along grain boundaries in metal oxides may be many orders of magnitude faster than their corresponding diffusivities in the grains, especially at intermediate temperatures (e.g., about 600 ºC to 900 ºC). This allows one to selectively in-diffuse elements from a solid source along the grain boundaries (e.g., grain boundary 112) of the film (e.g., solid electrolyte film 110) by exposing the grown thin films to moderate temperature and extended time anneals. This may be achieved, for example, by growing thin (e.g., about < 500 nm thick) polycrystalline ion conductive films (e.g., solid electrolyte film 110) on different selected types of insulating binary or ternary metal oxide substrates (e.g., substrate 120 / 920) and subsequently exposing the films to thermal anneals (e.g., about 600 ºC to about 900ºC) for varying periods of time, e.g., from 2 hours to 48 hours to enable the selective in-diffusion of the substrate elements (e.g., the precursor elements) upwards along the grain boundaries while reducing any segregation of unwanted dopants from within the grains towards the grain boundaries. The thin polycrystalline ion conductive films (e.g., solid electrolyte film 110) may be about 5 nm, about 50 nm, about 100 nm, about 200 nm, about 300 nm, about 400 nm, or about 500 nm thick, including all values in between. The substrate 120 / 920 may be a binary or ternary metal oxide substrate, including, but not limited to, a MgO, an Al2O3, a SiO2, a SrTiO3, a LaAlO3, or a SiC substrate as described above. The choice of substrate material may depend on the choice of precursor elements. For example, preferably the substrate material serves as the source of precursor elements that may diffuse up the grain boundaries. Additionally, the substrate (e.g., substrate 120 / 920) may support the thin polycrystalline ion conductive film (e.g., solid electrolyte film 110) as described above. Preferably, the substrate 120 / 920 is electrically insulating.

[0229] FIG.6 is a schematic representation of an annealing process for up diffusion where a polycrystalline thin film 610 is grown on an insulating substrate 620 that includes the precursor elements 670 that is exposed to subsequent thermal annealing from 600 ℃ to 1200 ℃ for varying periods from 2 hours to 48 hours to enable the diffusion of the precursor elements 670 upwards along the grain boundary 612. The polycrystalline thin film 610 may be made of any of the polycrystalline materials disclosed herein. As can be seen in FIG.6, at temperatures above 900 ℃ (e.g., 1200 ℃), there may be undesirable bulk dopant 680segregation towards the grain boundaries 612, which is detrimental to the grain boundary engineering strategy.

[0230] As in the bulk ceramic approach described above, the same selection criteria may be used to select which substrate to act as the source for the up-diffusion process along the grain boundaries (e.g., grain boundary 112). The selection criteria may be designed to modulate the grain boundary space charge potentials as described above. By utilizing an insulating substrate as the source, the ability to readily measure the conductivity of the thin films without electrically shorting through the substrates and with minimal additional processing steps (e.g., metallization) may be ensured. This may also enable rapid implementation into device operation (e.g., into detection device 100).

[0231] Alternatively, the source of precursor elements to be in-diffused into the film (e.g., solid electrolyte film 110), may be from a thin layer of precursor elements deposited on top of the ion conductive film (e.g., solid electrolyte film 110) following its growth or below it prior to initiating the ionic conductive film (e.g., solid electrolyte film 110) growth. The thin layer of precursor elements may be composed of only precursor elements (e.g., a pure metal). Alternatively, the thin layer of precursor elements may include precursor elements and one or more additional elements (e.g., oxygen). For example, the film of precursor elements may be a metal oxide film. In the case of a metal oxide film of precursor elements, since oxygen may already be present in the ion conductive film (e.g., solid electrolyte film 110), this may reduce the diffusion of additional elements from the film of precursor elements. For example, deposition of an Al metal film onto the solid electrolyte film 110 by sputtering or electron beam deposition, followed by an in-diffusion step. The in-diffusion may be done in a vacuum to maintain the Al largely in metallic form or following its oxidation to Al2O3. After the in-diffusion step, any remaining material in the thin layer of precursor elements may be removed by either chemical (e.g., 3:1 HNO3:H2O) and / or physical etching (e.g., ion milling). Additionally, the substrate material (e.g., substrate 120) and / or annealing conditions should be selected to reduce corresponding up-diffusion from the substrate that may interfere with the desired effect of the in-diffusion of the precursor elements from the top surface of the film of precursor elements.

[0232] Alternatively, when depositing the thin layer of precursor elements before growing the polycrystalline films (e.g., solid electrolyte film 110), concerns regarding substrate element diffusion and subsequent removal of the thin layer of precursor elements by chemical / physical etching may be unwarranted as long as the film thickness is selected forthe thin layer of precursor elements to ensure that the source layer is entirely consumed during the diffusion process. Preferably the film is about < 2 nm. For example, the film may be about 0.5 nm, about 0.75 nm, about 1 nm, about 1.1 nm, about 1.2, nm, about 1.3 nm, about 1.4 nm, about 1.5 nm, about 1.6 nm, about 1.7 nm, about 1.8 nm, about 1.9 nm, or about 2.0 nm, including all values in between. This may ensure that no layer of precursor elements is left below the polycrystalline films (e.g., solid electrolyte film 110) after growth, and that electrical measurements can be performed on the polycrystalline film (e.g., solid electrolyte film 110) by avoiding electrical shorting. Moreover, an initial thermal anneal (e.g., at about 600 ºC for about 1 hour) after growing the source layer prior to growing the polycrystalline film layer (e.g., solid electrolyte film 110) may ensure high conformality of the source layer with the substrate and may prevent it from dewetting during the growth of the polycrystalline film (e.g., solid electrolyte film 110).

[0233] There have been studies, where thin metallic or oxide source layers are deposited onto the surface of polycrystalline thin films, which are subsequently thermally annealed to induce in-diffusion of the elements into the grain boundaries. While some of those studies have demonstrated some evidence of resistance modulation, these studies were done on mixed electronic and ionic conductors, not on solid electrolytes as disclosed herein. Furthermore, the impact was limited in scope due to the simultaneous up-diffusion of substrate elements compensating for the original in diffusion.

[0234] Protocol I– up diffusion from a substrate source

[0235] FIG.7A shows a schematic of up diffusion from a substrate 720 source for a polycrystalline thin film 710. FIG.7B shows an overview of a method 790 for up diffusion from a substrate source for a polycrystalline thin film 710. The polycrystalline thin film 710 may be made of any of the polycrystalline materials disclosed herein. First, at step 791, a layer of a polycrystalline thin film 710 is grown on a substrate 720 including the precursor elements 770 at moderate temperatures (e.g., < 600 ºC). The polycrystalline thin film 710 may be < 500 nm thick. For example, the polycrystalline thin film 710 may be about 100 nm, about 200 nm, about 300 nm, about 400 nm, or about 500 nm thick, including all values in between. Alternatively, the polycrystalline thin film 710 may about 500 nm to about 1 µm thick. Second, at step 792, the polycrystalline thin film 710 is annealed at intermediate temperatures (e.g., about 600 ºC to 900 ºC) for an intermediate time (e.g., about 2 hours to 48 hours) to enable diffusion of the precursor elements 770 into the grain boundaries 712 of the polycrystalline thin film 710. In the case of a thicker polycrystalline thin film 710 (e.g., about500 nm to about 1 µm), the polycrystalline thin film 710 may be annealed at a higher temperature (e.g., about 1000 ºC for about 10 hours to about 30 hours). Third, at step 793, surface metallization (e.g., electrodes 730 and a circuit 760) may be applied to the polycrystalline thin film 710 to measure conductance and operate the radiation detection device 700. The electrodes 730 and circuit 760 may include any of the electrodes and / or circuits described herein.

[0236] Protocol II– down diffusion from thin top source layer

[0237] FIG.8A shows a schematic of down diffusion from thin top source layer 871 for a polycrystalline thin film 810. FIG.8B shows an overview of a method 890 for up diffusion from a thin top source layer 871 for a polycrystalline thin film 810. First, at step 891, a layer of a polycrystalline thin film 810 is grown on a substrate 820 at moderate temperatures (e.g., about < 600 ºC). The polycrystalline thin film 810 may be made of any of the polycrystalline materials disclosed herein. The polycrystalline thin film 810 may be < 500 nm thick. For example, the polycrystalline thin film 810 may be about 100 nm, about 200 nm, about 300 nm, about 400 nm, or about 500 nm thick, including all values in between. Second, at step 892, a thin source layer 871 of the precursor elements 870 is grown on top of the polycrystalline thin film 810 by electron beam and / or sputtering at a reduced temperature of about < 600ºC. The thin source layer 871 may be about 10 nm to about 100 nm. For example, the thin source layer 871 may be about 10 nm, about 15 nm, about 20 nm, about 25 nm, about 30 nm, about 35 nm, about 40 nm, about 45 nm, about 50 nm, about 55 nm, about 60 nm, about 65 nm, about 70 nm, about 75 nm, about 80 nm, about 85 nm, about 90 nm, about 95 nm, or about 100 nm, including all values in between. Third, at step 893, the thin source layer 871 may be annealed at intermediate temperatures (e.g., about 600 ºC to about 900 ºC) for intermediate times (e.g., about 2 hours to 48 hours) to enable diffusion of the precursor elements 870 from the thin source layer 871 into the grain boundaries 812. Fourth, at step 894, the surface layer of the thin source layer 871 may be etched away by either chemical and / or physical etching. Fifth, at step 895, surface metallization (e.g., electrodes 830 and a circuit 860) may be applied to measure conductance and operate the radiation detection device 800. The electrodes 830 and circuit 860 may include any of the electrodes and / or circuits described herein.

[0238] Protocol III– up diffusion from thin bottom source layer

[0239] FIG.9A shows a schematic of up diffusion from thin bottom source layer 971 for a polycrystalline thin film 910. FIG.9B shows an overview of a method 990 up diffusionfrom thin bottom source layer 971 for a polycrystalline thin film 910. First, at step 991, a thin source layer 971 of the precursor elements 970 is grown on a substrate 920 by electron beam and / or sputtering at reduced temperatures (e.g., < 600ºC). The thin source layer 971 may be about 10 nm to about 100 nm. For example, the thin source layer 971 may be about 10 nm, about 15 nm, about 20 nm, about 25 nm, about 30 nm, about 35 nm, about 40 nm, about 45 nm, about 50 nm, about 55 nm, about 60 nm, about 65 nm, about 70 nm, about 75 nm, about 80 nm, about 85 nm, about 90 nm, about 95 nm, or about 100 nm, including all values in between. Second, at step 992, the thin source layer 971 of the precursor elements 970 is annealed on the substrate at intermediate temperatures (e.g., about 400 ºC to 600 ºC) for about 1-2 hours. Third, at step 993, a polycrystalline thin film 910 is grown on top of thin source layer 971 at moderate temperatures (e.g., about < 600ºC). The polycrystalline thin film 710 may be made of any of the polycrystalline materials disclosed herein. The polycrystalline thin film 910 may be < 500 nm thick. For example, the polycrystalline thin 910 may be about 100 nm, about 200 nm, about 300 nm, about 400 nm, or about 500 nm thick, including all values in between. Fourth, at step 994, the polycrystalline thin film 910 may be annealed film at intermediate temperatures (e.g., about 600 ºC to about 900ºC) for intermediate times (2 hours to 48 hours) to enable in-diffusion of the precursor elements 970 from the thin source layer 971 into the grain boundaries 912. Fifth, at step 995, surface metallization (e.g., electrodes 930 and a circuit 960) may be applied to measure conductance and operate the radiation device 900. The electrodes 930 and circuit 960 may include any of the electrodes and / or circuits described herein.

[0240] The selection criteria described herein may be used to determine which precursor elements 770 / 870 / 970 will impart a positive or negative charge to the grain boundary 712 / 812 / 912 core. Subsequently, binary or ternary metal oxides with the precursor elements may be selected for the substrates 720 / 820 / 920. Preferably the substrate 720 / 820 / 920 is electrically insulating (e.g., >1010Ohm cm) and non-reactive both with the environment and the polycrystalline films 710 / 810 / 910. Binary metal oxides (e.g., MO2, M2O3, MO3) may be preferable as only the metal cation can be expected to diffuse up the grain boundaries 712 / 812 / 912, while the oxygen anion is already present in abundance in the polycrystalline metal oxide film (e.g., polycrystalline films 710 / 810 / 910) and therefore may not be expected to have an impact.

[0241] Alternatively, ternary elemental substrates (e.g., ABO3, SrTiO3) may be used as the substrates 720 / 820 / 920. However, it is preferable that the various elements that willdiffuse do not counteract each other. For example, MgAlO3 can be expected to result in the diffusion of both Mg and Al up the grain boundaries 712 / 812 / 912, which may charge balance each other, thus limiting the space charge modulation. On the other hand, if SrTiO3 is selected, while both Sr and Ti can be expected to in-diffuse, the diffusivity of the smaller Ti4+cation versus that of the much larger Sr2+ion may ensure that only Ti in-diffusion up the grain boundaries 712 / 812 / 912 may be significant.

[0242] Alternative insulating substrates 720 / 820 / 920 such as nitrides (e.g., BN, AlN, GaN), fluorides (e.g., BF3, AlF3, GaF3), and chlorides may also be used with the methods described herein. However, the up diffusion of the anion may need to be considered.

[0243] Results

[0244] An example where 200 nm polycrystalline thick films of 3at% Gd doped CeO2 were grown on both MgO and Al2O3 substrates and subsequently annealed for 6 hours each at different annealing temperatures varying from about 600 ºC to about 1200 ºC. Mg (ionic radius of about 0.89 Å) is within 20% of the value of Ce (ionic radius of about 0.97 Å), whereas Al (ionic radius of about 0.535 Å) is much smaller. Therefore, Al would be expected to enter interstitially, creating a net 3+ positive charge, whereas Mg could be expected to enter substitutionally, making a net negative charge (Kroger Vink notation.. To isolate the grain boundary effect, an “epitaxial” (grain boundary-free) sample was added to the analysis.

[0245] FIGS.10A and 10B are plots of the inverse resistance – temperature product (T / R) vs inverse temperature (1 / kT) for MgO (FIG.10A) and Al2O3 (FIG.10B). It can be seen that the films grown on MgO systematically show an increase in conductance for higher annealing temperatures, up to a point where the conductance of the sample becomes equal to that of an “epitaxial” sample. As shown in FIGS.10A and 10B, the overall polycrystalline film conductance on MgO (FIG.10A) may gradually approach the epitaxial sample’s values (both in magnitude and activation energy) for increasing annealing temperatures whereas for Al2O3 (FIG.10B) it does not. On the other hand, the annealing behavior for the samples grown on Al2O3 showed a different trend, with the conductance (R-1) initially decreasing up to 900ºC, above which it increases again; however, it may never reach the value of the epitaxial film. In FIG.10C, the grain boundary space charge potential variation is plotted as a function of annealing temperature for the samples grown on both substrates. As shown, in the case of MgO, the space charge potential systematically decreases (from about 0.25 V to about 0 V) with increasing annealing temperature, whereas the space charge potentialinitially increases for Al2O3 (from about 0.25 V to about 0.45 V), reaching a peak around 850 ºC and subsequently decreasing to a saturated value of about 0.22 V.

[0246] FIGS.11A and 11B plot the conductance response for all the samples for MgO (FIG.11A) and Al2O3 (FIG.11B) under constant UV illumination (photon flux of approximately 1018photons / (cm2s)). The overall optoionic response on MgO (FIG.11A) gradually decreases and approaches the epitaxial sample’s values for increasing annealing temperature, whereas for Al2O3 (FIG.11B) it does not. The magnitude of the optoionic response and its temperature dependence are observed to be the same for both sets of films. A general increase in conductance is observed for higher annealing temperatures. However, the variation is smaller (e.g., less than a factor of 10). In FIGS.11C and 11D, the optoionic response, equivalent to the conductance under UV divided by the conductance in the dark, is plotted for MgO (FIG.11C) and Al2O3 (FIG.11D) as it represents the magnitude increase in conductance of the films as a function of illumination intensity. The optoionic response in the set of Al2O3 films is observed to be much larger than the magnitude of the optoionic response in the MgO films at equal temperatures. Moreover, while in the case of the MgO substrate, the optoionic response systematically decreases with increasing annealing temperature, in the case of the films grown on Al2O3, a maximum is observed around 850ºC. This is consistent with the maximum dark resistance and space charge potential observed in FIGS.10A–10C. Moreover, the higher the optoionic response (correlated with higher space charge potential and lower dark conductance), the higher the temperature operational range over which an optoionic response can be obtained, as can be seen in FIGS.11C and 11D as the x-intercept on the x axis line (lower annealing temperature have a smaller 1 / (kT)) x axis intercept, consistent with higher temperatures.

[0247] Grain Boundary Space Charge Engineering of Solid Oxide Electrolytes: Model Thin Film Example

[0248] Grain boundaries (GB) profoundly influence the electrical properties of polycrystalline ionic solids. Yet, precise control of their transport characteristics has remained elusive, thereby limiting the performance of solid-state electrochemical devices. Disclosed herein are methods for the manipulation of space charge controlled ionic grain boundary resistance (e.g., up to 12 orders of magnitude) in metal oxide thin films. We exploit the orders of magnitude higher grain boundary diffusivities of substrate cation elements (e.g., Al from Al2O3 and Mg from MgO) relative to the bulk to modify the grain boundary chemistry, and thereby GB core charge, in a model oxygen ion conducting polycrystallinethin film solid electrolyte, Gd-doped CeO₂. This approach, confirmed jointly by TEM imaging and by extracting the respective GB and bulk diffusivities from measured SIMS profiles, enabled the selective control of the chemistry of the GBs, while minimally modifying grain (bulk) chemistry or film microstructure, thereby ruling out potential effects of microstructure, strain or secondary phases. Broad tuning of GB space charge potentials may be achieved by manipulating GB core charge density by over an order of magnitude, thereby providing a powerful tool for systematic studies of grain boundary phenomena across various functional materials. The implications of such control are far-reaching in achieving new functionality, improving efficiency and longevity of solid-state electrochemical devices (e.g., the radiation detection device 100 / 150 / 700 / 800 / 900 / 1250 disclosed herein).

[0249] Grain boundaries (GB) are features of polycrystalline materials and may influence the electrical behavior of polycrystalline materials. For example, they contribute to reduced electronic mobility in semiconductors due to scattering, cause substantially higher GB resistances compared to the bulk due to the presence of carrier depletion zones adjacent to the interfaces, and serve as charge recombination sites in optoelectronic devices, limiting device sensitivity and efficiency. In ion-conducting materials, such as the ion-conducting polycrystalline materials disclosed herein, GBs may impede ionic transport and induce ohmic losses while also leading to chemical and mechanical instability under extreme loading conditions, degrading device performance over time. In semiconducting ZnO-based varistors and BaTiO3-based positive temperature thermistors, abrupt collapse or creation of space charge barriers at grain boundaries due to over-voltages or over-heating may lead to desired switching behavior. Given the ubiquity and the highly impactful effects of GB barriers to the flow of electronic and ionic charges, it is surprising how little success there has been in controlling their properties to achieve improved device performance and stability. Herein, we demonstrate the ability to systematically control the GB resistance of model oxygen ion conducting solid electrolyte thin films by up to 12 orders of magnitude. This may be achieved by selectively engineering GB chemistry, without impacting the chemistry of the surrounding grains, thereby isolating the role played by the magnitude and sign of the GB core charge.

[0250] The transport characteristics of GBs may be highly sensitive to synthesis and processing conditions and depend on the nature and concentration of impurities, defect types, stoichiometric variations, and / or secondary phases localized at the interfaces. Even in high- purity systems, resistive GBs continue to be primarily induced by space charge effects thatarise from charges trapped at GBs, creating electrostatic barriers that alter mobile charge carrier distributions in adjacent grains. The exact origin of the boundary core charge remains under debate, especially when dealing with ionic conductors. Early models, inspired by semiconductors, point to carrier trapping at mid-gap defect states to explain GB core charge origins. In ionic systems, these charges have been assigned to either intrinsic sources connected with free energy differences in the generation of, e.g., oxygen vacancies in oxygen ion conductors or extrinsic effects due to impurity segregation to the GBs. While common impurities like Al and Si have been correlated with increases in barrier height in ionic systems, dopants used to enhance ionic conductivity in solid electrolytes, for instance, in CeO₂, ZrO₂, (Ba, Sr)TiO₃, BaZrO₃ and BaCeO₃, are known to reduce space charge barriers, thereby reducing GB blocking. Alkaline earth additives and transition metals used as sintering aids have also, at times, been reported to reduce GB space charge barrier heights in such metal oxide systems.

[0251] While our understanding of grain boundary structure and chemistry has advanced over the past decades, a direct link to space charge effects remains elusive. Unintentional impurities complicate the analysis, as they push the limits of detectability and are often limited to only a subset of interfaces. Nevertheless, studies indicate that residual impurities may critically impact space charge properties even in highly purified materials. For example, even in high purity (e.g., < 40 ppm) undoped CeO2 fibers and 0.2 at% Sm- doped CeO₂ bulks samples, grain boundaries may exhibit unexpected levels of extrinsic impurities due to scavenging and accumulation of various cations (up to 2 at% Al, for example). The impurity distribution within the same specimen correlated with a range of space charge potentials, some near zero, while some as high as 1 V, with similar observations being made for SrTiO₃. These results indicate that grain boundary core charges in ionic solids may arise from trace impurities segregating to grain boundary regions.

[0252] Whether the origin of the GB core charge is intrinsic or extrinsic, methods for controlling the potential barriers may help achieve electrochemical devices with superior performance. Traditional grain boundary modification strategies rely on doping the host lattice and promoting dopant segregation during high-temperature treatments or sintering (typically at temperatures > 1200 ºC), but this makes it difficult to separate impacts of changing microstructure and grain chemistry from that of GB chemistry. This may be circumvented by selective chemical modification of GBs following completion of microstructural development.

[0253] Previous studies have shown that GBs in polycrystalline CeO2 films may exhibit higher solubility and diffusion rates at more moderate temperatures than bulk grains (e.g., about 600 ºC to about 800ºC), allowing for selective incorporation of elements into the GBs from a surface layer source via an in-diffusion thermal anneal step. This may allow for controlled chemical modification of the GB interfaces without altering the bulk composition and minimally impacting the microstructure. Systematic studies of how this may influence space charge properties of grain boundary interfaces have nevertheless been limited both in extent and scope. For example, it was previously observed that while Ni in-diffusion in CeO2 thin films reduced space charge potentials, its impact was limited due to unexpected impurity up-diffusion from the MgO substrate. This observation is consistent with a report that suggested that the decreased grain boundary resistance in thin films of Yttria Stabilized Zirconia (YSZ) grown on MgO substrates at 700°C may be due to a reduction in space charge potential, resulting from Mg diffusion into the YSZ film. Similar phenomena utilizing substrates with other compositions have not been investigated, leaving this approach's universality unexamined.

[0254] Disclosed herein the ability to control grain boundary properties in a model oxygen ion conducting system, 3 at% Gd-doped CeO₂ (Gd0.03Ce0.97O1.985 also referred to herein as GDC3) is disclosed. This method relies on growing thin films of GDC3 on MgO and on Al₂O₃ substrates and subsequently exposing them to intermediate temperature anneals to induce selective grain boundary in-diffusion of the respective substrate elements. FIG. 13A illustrates the impact of excess positive charge, and the corresponding impact on the spatial dependence of the space charge potential ^^^^^^(^^), the mobile concentration of doubly charged oxygen vacancies [^^0∙∙], the minority electron density ^^, and the bulk acceptor dopant density [^^ce′], may be assumed to be immobile at the lower annealing temperatures. A thin film 1310 (e.g., a GDC3 thin film) may be grown on a substate 1320 (e.g., a MgO and / or on Al₂O₃ substrate). The thin film 1310 may have one or more grain boundaries 1312. FIG.13B illustrates a thin film 1310a (e.g., a GDC3 thin film) grown on a MgO substrate 1320a (left schematic) and a thin film 1310b (e.g., a GDC3 thin film) grown on a Al₂O₃ substrate 1320b (left schematic). FIG.13B (upper section) illustrates negatively charged precursor elements 1370a on the left and positively charges precursor elements 1370b on the right selectively diffusing up the grain boundary 1312a / 1312b, leading to respectively a net decrease (increase) in the net positive charge at the grain boundary 1312a / 1312b. This leads in turn (see lower section FIG.13B) to a decrease (increase) in ^^^^^^(^^), a decrease (increase)in depletion of [^^0∙∙], and a decrease (increase) in the accumulation of minority electrons, ^^. With this approach we may achieve up to approximately 12 orders of magnitude change in the GDC3’s room temperature extrapolated resistance when compared against nominally epitaxial standards (e.g., GB resistance free). STEM and Secondary Ions Mass spectroscopy were used to verify the selective nature of our grain boundary chemical modification, while XRD and SEM studies confirmed that this is achieved via minimal difference in bulk chemistry, strain, or microstructure. This approach offers evidence for chemical control of grain boundary space charge properties and establishes guidelines that may be applied to solid electrolytes (e.g., the ion-conducting polycrystalline materials disclosed herein), and thus may offer a pathway to enhanced performance for polycrystalline ionic conductors such as the radiation detector device 100 / 150 / 700 / 800 / 900 / 1250 disclosed herein.

[0255] Results

[0256] Polycrystalline GDC3 thin films 1310a and 1310b (approximately 200 nm thick) were prepared by pulsed laser deposition on (0001) Al2O3substrate 1320b and on (001) MgO single-crystal substrate 1320a, then exposed to 6 hour thermal annealing conditions at temperatures ranging from 650 °C to 1250 °C, and subsequently studied electrically by electrochemical impedance spectroscopy. The selection of the substrates was based on the ionic sizes and charges of the cations that make up the substrates 1320a / 1320b relative to the size and charge of the GDC host cation Ce4+with the aim of modulating the net GB 1312a / 1312b core charge (as described herein). The details regarding sample fabrication and electrode geometry are provided below and in section titled Electrode Geometry, while the microstructural characterization can be found in the sections titled Structural Characterization and Structural Evolution of GDC3 films 1310a on MgO substrate 1320a as a function of temperature (XRD / SEM). As confirmed by SEM and TEM (FIGS. 18A–18C and 19), the as-prepared samples are polycrystalline with columnar grains perpendicular to the substrate, with an initial (111) / (100) mixed grain orientation distribution. As shown in FIG.19, the GDC3 films may include a carbon layer (C) that was used for the sample preparation for imaging. This may be expected to facilitate the selective diffusion of substrate elements (e.g., precursor elements 1370a / 1370b) up the grain boundary 1312a / 1312b channels. As shown in FIGS.20A–20D and 21A–21D, XRD and surface SEM analysis may show no differences in structural evolution with annealing temperatures for films 1310a / 1310b grown on the different substrates (e.g., substrates 1320a / 1320b) and no evidence of changes in the presence of secondary phases or strain state which supports theidea that the electrical trends result solely from chemical modifications of the space charge properties.

[0257] Electrical Conductivity Behavior: In FIGS.14A and 14B, the inverse resistance (1 / R) x temperature (T) extracted from impedance measurements (characteristic Nyquist plots measured at 350°C shown in FIGS.23A–23B of all films grown on MgO and Al2O3 are plotted as a function of inverse temperature (kBT)-1on an Arrhenius plot to enable extraction of corresponding activation energies. The conductance behavior of the films can be described, for the most part, as exhibiting single slopes. Deviations, however, appear at the lower temperatures due to expected current leakage along the film surface or the film / substrate interface, given the very high recorded resistance values (> GΩ) at the reduced temperatures. Focusing on the higher temperature behavior (> 150 ºC to 200ºC), where clear trends are observable, and single slope activation energies are present, we find that the nominally epitaxial samples show the highest conductance’s and lowest activation energies (approximately 0.67 eV), consistent with other literature reports both on thin film and bulk ceramics, with the activation energy ascribed to bulk conductivity resulting from lattice oxygen vacancy migration. The polycrystalline films may exhibit larger activation energies than those of the epitaxial films (0.67eV), consistent with grain boundary resistance dominance and space charge related phenomena.

[0258] When examining the temperature dependence of film conductance grown on MgO, as displayed in FIG.14A, one observes that films annealed at the lowest temperature (850 ºC) may exhibit the highest activation energy (1.16 eV) and lowest conductance over the entire temperature range. The values reported are consistent with other literature reports both on thin films and bulk ceramics, ascribed to blocked oxygen ion transport at the GBs. With increasing annealing temperature, the conductance (1 / R) increases, while the activation energy decreases, with the total activation energy dropping from 1.16 eV to 0.67 eV, the latter aligning with the values characteristic of the nominally epitaxial sample. Likewise, the conductance increases toward that of the nominally epitaxial films. Above 1000ºC, the conductance and activation energies reach values close to those of the nominally epitaxial films, indicating a saturation of the observed phenomena.

[0259] The trend is less straightforward in the case of films grown on Al2O3. Starting at the lowest temperature anneal (750 ºC), the films exhibit a high resistance and activation energy (approximately 1.24 eV), slightly higher than the values obtained for the film grown on MgO. With increasing annealing temperature up to 850 ºC, however, the resistance andactivation energy increase substantially, with activation energy (Ea) reaching a value of approximately 1.57 eV. Above 850 ºC, the resistance and its activation energy, however, begin to decrease, reaching a saturation point > 1050ºC. The conductance and activation energies, however, do not coincide with the values of the nominally epitaxial film, as observed for the films grown on MgO. Interestingly, the activation energy reaches a saturation value of (approximately 1.15 eV), a value often reported for GDC oxides in thin film and bulk form following high-temperature sintering.

[0260] As shown in FIG.15, if we extrapolate the film conductivities down to room temperature based on their high-temperature slopes, the changes in film conductance with annealing temperature reach factors of ≥105in both cases. However, at its peak, around 850 ºC, the conductance’s of both films also differ by a factor of 105. Using Equation 2, we obtain estimated values for the changes in space charge potential as displayed in FIG.15 (procedure outlined below in the section titled Impedance fitting and space charge model fitting and plots showing fits and table with space charge potential values provided in FIGS. 25A–25C). The space charge potential may follow the same trends as the total activation energy and resistivity with annealing temperature. For the MgO case, the potential barrier begins at 0.25 V and decreases to near zero (value 0.013V) with annealing, while for films grown on Al2O3, it begins at 0.27 V, reaches values as high as 0.45 V and finally decreases back down to 0.23 V, similar to the starting point for films on MgO. Based on these space charge values and applying Equation 4 that relates barrier height ∆ϕ(0) to GB core charge Q, the average “net” core charge density (Qcore) at the grain boundaries changes by a factor approximately 6 from 0.047 to 0.276 C / m² over the entire range of conditions examined.

[0261] Analysis of Depth Profiles by TOF-SIMS: To elucidate the role that chemical modification has on the electrical properties of the grain boundaries; secondary ion mass spectroscopy (SIMS) was performed on the samples to measure the in-diffusion profiles of the substrate cations into the grain boundaries of the supported GDC films. FIG.16A displays characteristic SIMS spectra for the precursor elements (e.g., dopants) and host elements for samples as deposited and subsequently annealed at 750 °C for 6 hours and 850 °C for 6 hours, respectively. FIG.16B shows STEM- EDS spectra images collected from a representative grain boundary in GDC grown on Al2O3 (left image) and on MgO (right image) at 1050°C. In the section titled Secondary Ions Mass Spectroscopy (SIMS) data, we provide SIMS plots for samples annealed from 650 °C to 1250 °C in the case of MgO and 750 °C to 1250 °C for the case of Al2O3. In general, the depth-independent signals of Ce+and Gd+ originate from the deposited GDC film, and Mg+and Al+each arise from the underlying MgO and Al2O3 substrates. In contrast to the as-deposited samples, where both Mg and Al signals drop sharply at the film-substrate interface, all samples that were annealed show significantly increased Mg or Al signal intensity to much deeper penetration depths. As shown in FIG.16A, the concentration profiles for Mg and Al appear to saturate by the time the samples' annealing temperature reaches 850 °C, where the number of in-diffused species presumably reaches the solubility limit of the grain boundaries of the deposited GDC films. Above 850 °C, as displayed in FIGS.26A–26B, the saturation concentration in the bulk of the films may decrease slightly with increasing annealing temperatures, whereas samples deposited on MgO may show an increase in Mg concentration near the surface. Such variation can be rationalized by noting that the decrease in grain boundary concentration may be due to grain growth that eventually lowered the apparent grain boundary solubility per surface area, while the grooving / dewetting, especially at higher temperatures, enables additional surface diffusion to occur along the film surface, consistent with previous reports. Discrepancies between Al and Mg may also relate to a fundamental difference in surface diffusivity. Detailed views of the SIMS profiles for all samples are provided in FIGS.27 and 28.

[0262] The dominant in-diffusion pathways were identified by fitting the diffusion profiles using both bulk and grain boundary diffusion expressions, detailed in the section titled Secondary Ions Mass Spectroscopy (SIMS) data, with the results plotted as a function of temperature shown in FIG.33. The activation energies for diffusion along the grain boundaries are ½ and 1 / 5 of the bulk values (Al2O31.3 eV vs 2.6 eV and MgO 0.84 eV vs 4.1 eV). We calculate the differences in diffusivities for the bulk vs grain boundaries at a common temperature, e.g., 850 °C, through extrapolation. ^^^^^^,^^^^ / ^^^^,850~10−16^^^^2 / ^^, whileequating to approximately 5 orders of magnitude enhancement in Al and Mg cation diffusivities between grain boundary versus bulk, consistent with previous reports on markedly enhanced grain boundary diffusivity of foreign cation elements.

[0263] Analysis of Grain Boundary Segregation by STEM-EDS: Scanning transmission electron microscopy (STEM) energy dispersive x-ray spectroscopy (EDS) corroborates the SIMS measurements, showing in-diffusion from the substrates into the GDC films as in FIG.16B for samples annealed on alumina and magnesia at 1050 °C. Additional STEM EDS line plots are presented in FIGS.36A–36D, 37A–37D, and 38A–38D for samples on alumina annealed at 850 °C and 1050 °C, and in FIGS.39A–39D and 40A–40Dfor samples on MgO annealed at 1050 °C and 1250 °C). In both cases, the substrate cation may be found to be enriched at the grain boundary and more pronounced at higher annealing temperatures.

[0264] In both cases, there is additionally evidence of Gd segregation to grain boundaries—this may also lower the measured space charge potentials. Further, Si contamination was also found at the grain boundary (FIGS.37A–38D, 39A–39D, and 40A– 40D ) for samples annealed at higher temperatures (e.g., > 1000°C), though no significant presence was found when analyzed by SIMS (see FIGS.34 and 35). The observation of Si contamination may correlate with higher temperature anneal, as little evidence of Si was present in GDC3 samples grown on Al2O3 annealed at 850 °C as shown in FIGS.36A–36D. In general, no other significant traces of potential contaminants (Ca, Na, H, Sr) could be found in either the SIMS or TEM results (see FIGS.34 and 35). Certain boundaries show little or no segregation and appear to be low-angle grain boundaries (see FIGS.37A–37D). This highlights the coupling between grain boundary structure and the resultant electrochemical properties.

[0265] Discussion

[0266] These results demonstrate the ability to systematically and markedly vary the GB resistance in a model polycrystalline oxygen ion conducting solid electrolyte thin film (Gd0.03Ce0.97O1.985). This may lead to changes in room temperature extrapolated values by up to approximately 12 orders of magnitude by controlling the net core charge density (positive – in this case) of the grain boundaries - as illustrated in FIG.15. This may be achieved by growing GDC3 thin films on both Al2O3 and MgO substrates and annealing them over a broad temperature range (750 ºC to 1250ºC). Potential contributions from structural variations may be ruled out, as the same evolution was observed on both substrates regarding grain size, morphology, and grain orientation. Strain and secondary phases were furthermore ruled out by XRD and TEM studies. Since we were able to confirm from SIMS profile fitting that both Al3+and Mg2+ions diffuse approximately 105times faster along the GBs than in the GDC grains, we were able to selectively and controllably modulate the grain boundary chemistry, and therefore the GB core charge, while maintaining the “bulk” grain chemistry constant.

[0267] GDC3 polycrystalline films, as grown, may already exhibit a built-in space charge potential of approximately 0.25 V, consistent with a positive core charge as described in the section titled Space Charge Model and as previously reported. Our electrical, SIMS,and TEM results are all consistent with the assumption that in-diffusion of Al3+results in a net increase in positive charge, whereas in-diffusion of Mg2+results in a net decrease in positive charge at the GB core resulting in an overall change in net positive charge by a factor of approximately 6 (0.047 to 0.276 C / m²). These results align with the expectation that the much smaller Al3+ion (0.535 Å – 6 fold coordination) may enter interstitially into sites within the grain boundary regions, leading to a net positive defect ^^^^i…, while the Mg2+ion (0.89 Å – 8-fold coordination) with a net negative charge ^^^^ce′′, and with comparable size to the Ce4+ion (0.97 Å– 8 − fold coordination), may enter substitutionally.

[0268] While this explanation is consistent with the observed trends upon heating up to 850 ºC, the reversal in trends for the Al in-diffused films above that temperature may require further explanation. It is instructive to examine the STEM-EDS sppectra that show that the net negatively charged bulk dopant ^^^^ce′may begin to significantly segregate to the grain boundaries above 850 ºC (see FIGS.31A–40D). This may be expected to result in: i) the addition of negative charges within the grain boundary core and thereby ii) a decrease in the space charge width. In the case of the MgO substrate, the redistribution of both ^^^^^^^^′′ and ^^^^^^^^ to the grain boundary core together may contribute to the decrease in space charge potential, while in the case of Al2O3, ^^^^’^^^^ and ^^^^i…may compete.

[0269] This may explain why Mg in-diffused films may exhibit a monotonic decrease in GB resistance, while Al in-diffused films may exhibit an initial decrease in conductivity at reduced temperatures and then an increase at higher temperature anneals. Without being bound to any particular theory, at the lower temperatures, only Al diffusion along the grain boundaries may actively contribute to increased positive GB core charge. Then at higher temperatures during which the bulk dopant ^^^^ce′may become sufficiently mobile even within the grain, including within the space charge region, to enable it to diffuse the short distance to the GB core, at which point its negative charge may begin to compensate for the extra positive charge introduced during the Al ion in diffusion at lower temperatures. Similar effects of bulk impurity redistribution in the vicinity of grain boundaries in solid electrolytes are often reported.

[0270] The built-in non-uniform dopant profile in the grain boundary vicinity due to segregation may serve to influence the electrical behavior of the grain boundary following a restricted equilibrium scenario which may describe an intermediate scenario between the Mott-Schottky and Gouy-Chapman approximations (see section titled High Temperature Restricted Equilibrium Scenario additional details). This case may be evaluated vianumerical simulations with knowledge of the dopant frozen-in profile in response to the thermal history but may be impractical for most experimental studies. Si appeared to segregate at elevated temperatures (>1000°C), however its impact on our observed trends appears minimal. The fact that it is present in both sets of films (on Al2O3 and MgO) that still exhibit significant conductance differences at higher temperatures may rule out any significant role played by Si in our observations (see detailed discussion in the section titled STEM-EDS dataset collected from grain boundaries of GDC3 grown on Al2O3 and MgO).

[0271] The findings disclosed herein are summarized in FIG.17A, which illustrates the increased degree of penetration along the GBs 1312a / 1312b by Al and Mg ions (e.g., precursor elements 1370b and 1370, respectively) by selective in-diffusion from a MgO substrate 1320a or a Al2O3 substrate 1320b with increasing anneal temperatures, as well as increasing grain growth and segregation of the bulk Gd dopant 1780 with increasing temperature anneals for a given time, t. For temperatures of 600 °C and below, partial selective penetration of substrate cations (e.g., precursor elements 1370b and / or 1370) along the GBs 1312a / 1312b is obtained. For intermediate temperatures of about 800°C, full relatively uniform penetration is achieved. At even higher temperatures on the order of 1000°C, grain growth occurs accompanied by Gd dopant 1780 segregation to the grain boundaries 1312a / 1312b as well as surface diffusion of the substrate derived cations (e.g., precursor elements 1370b and / or 1370) on the surface of the film 1310a / 1310b.

[0272] FIG.17B describes the relative impact of the segregation of the Gd bulk dopant 1780, resulting in its accumulation in the space charge region and ultimately insertion of Gd into the core, causing a reduction of space charge potential and a decrease in space charge width (following a restricted equilibrium scenario) of the GB 1312a / 1312b, which may either compete or support the initial of Al and Mg on the grain boundary resistance. This demonstrates that by careful control of annealing time and temperature, it may become possible to enrich the grain boundaries 1312a / 1312b with selected elements (e.g., precursor elements 1370b and / or 1370), while reducing the segregation of the bulk dopant 1780.

[0273] In summary, one can systematically decrease the grain boundary potential to near zero for films grown on MgO, allowing GDC films to exhibit conductivities characteristic of grain boundary-free, epitaxial-like samples. This has implications for oxide based solid-state electrochemical devices (e.g., radiation detector device 100 / 150 / 700 / 800 / 900 / 1250) since such devices, regardless of conducting ion, rely on polycrystalline solid electrolytes or mixed ionic electronic conducting electrodes ormembranes. Blocking of ion transport at grain boundaries is a common observation and, at a minimum, leads to increased ohmic losses and thus decreases device efficiency. Ion blockage at grain boundaries has furthermore been reported to contribute to device degradation, such as microcracking in solid oxide electrolysis cells and dendrite formation in high-energy density sodium or lithium batteries.

[0274] Grain boundary barriers to ion migration may be substantially increased, e.g. from about 0.25 V to about 0.45 V, when annealing at temperatures < 900 ºC on Al2O3 substrates. This may lead to a large activation energies for grain boundary resistance in the 3 at% GDC system (approximately 1.6 eV). While to date, there have been few applications where large grain boundary barriers in solid electrolytes have been of interest, such ion blocking barriers may be effectively utilized to detect UV rays, gamma rays X-rays and / or neutrons. As with other light or radiation detectors, a figure of merit is a high dark resistance to achieve enhanced sensitivity (e.g., a dark resistance of about >1010ohm at room temperature). In the case of the ionic-based radiating detectors disclosed herein (e.g., radiation detector device 100 / 150 / 700 / 800 / 900 / 1250), this may be achieved not only by having large energy band gaps, but also by having grain boundary barriers as large as possible, achieved through the methods described herein.

[0275] The influence of grain boundary structure, e.g., grain boundary misorientation angles or complexions, on diffusivity of the elements and their solubility within the interface and their impact on the segregation of mobile charged species may need to be developed further. For example, the highest space charge potentials that one can achieve in each system may be likely to relate to the solubility limits of the precursor elements at these interfaces, as well as the competition between the relative segregation of the various charged mobile defects, bulk dopants, and / or grain boundary impurities to the boundaries. How this may change as a function of grain boundary structure and grain boundary chemistry and how to identify and control this structure, and how the segregation processes depend on the existing presence of charged impurities may in turn be beneficial in defining upper limits for space charge values.

[0276] Demonstrated herein is the feasibility of systematically engineering grain boundary resistances in ion-conducting solid electrolytes (e.g., the ion-conducting polycrystalline materials disclosed herein) over many orders of magnitude by careful incorporation of select cationic species at the interfaces. By employing targeted additives, it may be possible to either reduce the built-in space charge potential to near zero orsignificantly increase it, depending on the application. As demonstrated herein, a consideration when attempting to increase the space charge barrier potential via selective addition of charged defect species is to limit bulk dopant redistribution by careful control of annealing temperatures, while it is a favorable effect in the case in which one wishes to decrease the barriers. This approach exemplifies the potential of extrinsic control of space charge properties through point defect engineering, offering parallels to the development of semiconductors where the ability to control their electrical properties by intentional doping revolutionized the field.

[0277] Relying on the in-diffusion of elemental impurities along grain boundaries at intermediate temperatures, the methods described herein, enable the precise manipulation of space charge properties while maintaining the bulk properties unchanged. One example focuses on a single oxygen ion solid electrolyte (Gd-doped CeO₂), but these findings may have broad applicability to other solid state ionic systems, including, but not limited to, protons, alkali metal (e.g., lithium and / or sodium), silver, halide ion conducting materials, and / or any of ion-conducting polycrystalline materials disclosed herein. For example, this strategy may be extended to other ion-conducting systems of current interest for energy device applications, such as oxygen ion conducting yttria-stabilized zirconia (YSZ), lithium- ion conducting Li7La3Zr2O12 (LLZO), and / or proton-conducting BaZr0.8Y0.2O3-δ (BZY). By employing lower annealing temperatures for extended durations, it may become possible to reduce unwanted bulk dopant diffusion and / or microstructural evolution. Understanding the influence of grain boundary structure—including misorientation angles and / or complexions—on dopant diffusivity, solubility, and / or space charge properties may also play a role in the infusion of selected precursor elements along the grain boundaries. These additional considerations may help determine the upper limits of achievable space charge potentials and tailor grain boundary properties for specific applications.

[0278] This example provides a demonstration of the ability to eliminate grain boundary resistive losses, paving the way for reduced ohmic losses and / or improved performance in devices like electrolyzers and / or solid-state batteries. It also highlights the potential to intentionally increase barrier heights, enabling high dark resistances for applications such as radiation detectors (e.g., radiation detector device 100 / 150 / 700 / 800 / 900 / 1250) with enhanced sensitivity.

[0279] Methods

[0280] Material Selection: Gd doped ceria (GDC) was selected as a model material, as it is one of the most highly oxygen ion conducting solid electrolytes. Grain boundary blocking of ion transport in this system has been well characterized and confirmed to be dominated by space charge depletion of oxygen vacancies. A review of the theoretical basis describing the space charge properties of grain boundaries is provided succinctly the section titled Space Charge Model. Because of their relatively larger grain boundary resistance contributions, a lower dopant level (3 at% Gd, in the following referred to as 3GDC) was chosen than what is typically selected to achieve maximum oxygen ion conductivity (e.g., 10–20 at% Gd).

[0281] Substrate Selection: MgO and Al2O3 substrates were selected based on several characteristics: 1) insulating nature enabling electrical conductivity measurements of the overlaying films, 2) binary composition enabling a simple up-diffusion process that is easier to interpret, and 3) the charge each elements (Mg and Al) is expected to impart to the grain boundary core of CeO2. Magnesium ions (Mg²⁺), with a larger ionic radius, may be expected to incorporate substitutionally onto Ce4+sites, thereby creating a doubly charged defect with net negative charge of 2-, which in conventional defect notation is given by ^^^^ce′′. This may be expected to compensate for the net positive core charge commonly found in ceria-based solid electrolyte grain boundaries, thereby lowering the built-in space charge potential. Conversely, the smaller aluminum ions (Al³⁺), was anticipated to incorporate interstitially, creating defects with a net positive charge of 3+, which in conventional defect notation is given by ^^^^i3∙, leading instead to an increase in space charge potential. However substitutional and interstitial sites may not be able to be explicitly assigned in the grain boundary core due to ill-defined lattice structure, it still may provide a helpful guide for evaluating the relative charge cations might possess within the interface. Since the ion- conducting polycrystalline material are micro single crystals, the materials may have well defined lattice sites and thus it may be possible to distinguish between a normal lattice site and a lattice site with interstitials. Grain boundaries, in contrast, may have a much wider range of structures and may have ill-defined lattice sites. Additionally, many types of grain boundary structures may co-exist in a given ion-conducting polycrystalline material. Regardless of if the precursor elements enter substitutionally or interstitially, the precursor elements may possess a net charge that may be determined by the selection criteria explained herein. The selection criteria for precursor elements disclosed herein may be used to determine whether a grain boundary precursor element may end up giving a net negativecharge (which may be more likely substitutional) versus a net positive charge (which may be more likely interstitial) to the grain boundary.

[0282] Pulse Laser Deposition Target: The GDC powder used to prepare the Pulse Laser Deposition (PLD) target was synthesized through a co-precipitation route. A solution of gadolinium and cerium in stoichiometric proportion (3 / 97) was prepared by dissolving cerium nitrate, Ce(NO3)3:6H2O, and gadolinium nitrate, Gd(NO3)3:6H2O (Strem chemicals) (99.99%), in distilled water to reach a concentration of 0.1 molL-1. The precipitation solution was prepared by dissolving ammonium carbonate in distilled water to reach a concentration of 0.5 mol-1. The ammonium carbonate solution was calculated to have a molar excess of 2.5 compared to the total amount of cations in the nitrate solution. The ammonium carbonate solution was quickly poured into the vigorously stirred nitrate to trigger precipitation. The resulting residue was filtrated and washed four times in distilled water in a Buchner filter connected to a vacuum pump and dried at 100 °C overnight, followed by calcination at 600 °C for 1 hour to obtain the GDC powder crystallized in the fluorite structure. The powder was then pressed into a 30 mm diameter disk with a uniaxial press (1 Tcm-2) and sintered at 1500 °C for 6 hours, followed by cooling at 1 °C min-1to limit crack formation, resulting in a pellet with about 97% density.

[0283] PLD thin film: Both polycrystalline and nominally epitaxial samples were grown by PLD (Surface, Germany) using a KrF excimer laser with a 248 nm wavelength (Coherent, USA). The polycrystalline samples were directly grown on (001) MgO and (0001) Al2O3 single crystal substrates with a laser repetition rate of 10 Hz. The nominally epitaxial samples were grown following previously established protocols > 600ºC. Films were grown directly onto (0001) Al2O3 substrates, whereas a double buffer layer system of BaZrO3 and SrTiO3 (about 2-5 nm each) was initially grown onto (001) MgO substrate, to act as a seed layer. After a base pressure of 7·10- 6mbar was reached, pure oxygen was continuously leaked into the PLD chamber, keeping the total pressure at 0.013 mbar during film growth and cooling. During deposition of the polycrystalline films, the substrate temperature was held at 300 °C (heating rate 10 °C / min and cooling rate 10 °C / min). The laser energy was set to 100 mJ, resulting in a power density of about 1 Jcm-2. The target substrate distance was 7.5 cm. The resultant GDC films formed continuous pinhole-free layers. XRD was performed by a Bruker cobalt source D8 with a General Area Detector Diffraction System (GADDS). This system used a conventional 1.6 kW sealed tube cobalt anode. Incident- side optics included a variety of double-pinhole collimators and mono-capillary devices, which were used to adjustthe beam size, intensity, and divergence. The beam diameter for this instrument can range from 0.05 to 0.8 mm depending on the choice of collimator. The goniometer was a χ-cradle type, with full ^^ axis rotation and x-y-z translation. Including ω, this gave six positioning axes for the sample (not including the detector axis, 2θ). The Vantec-2000 detector had a very wide dynamic range and a maximum frame resolution of 2048 x 2048 pixels. Two- dimensional detectors facilitated the study of grainy and textured materials because the detector captured a slice of the Ewald diffraction sphere instead of a single point. Images of the cross-sections of the device were acquired at KAIST on a high-resolution scanning electron microscope (SEM, Hitachi SU8230). Samples were prepared by cutting the substrate with a low-speed precision diamond cutter to ensure a perpendicular cutting angle.

[0284] Annealing: After deposition, the samples, they were placed in Al2O3 crucibles and annealed in a Al2O3 tubular furnace at various temperatures (from 650 ºC to 1250 ºC) with 5°C / min ramp rates for 6 hours in air to ensure complete oxidation of the lattice and to allow for microstructural evolution and chemical diffusion. Approximately 7 mm long inter- digitated Pt or ITO electrodes were deposited on top of the film with the other dimensions outlined in Table 1. The Pt electrodes were deposited at room temperature by DC sputtering while ITO was deposited by PLD using similar conditions to those for the GDC films. By employing substrate-assisted in-diffusion rather than top-down metallic sources, we may preserve film integrity and reduce extraneous effects on conductivity measurements.

[0285] Electrochemical Impedance Spectroscopy (EIS): All electrochemical measurements were performed in a Linkam stage HS600 with a quartz window, allowing for heating and illumination of the sample. The chamber was flooded with 50 sccm of synthetic air to ensure oxidizing conditions throughout the measurements. All EIS measurements were performed with an MFIA Impedance Analyzer (Zurich Instruments). A 100mV amplitude was necessary for a sufficiently high current response in the highly resistive samples. The frequency range was from 0.01 Hz to 1 MHz. Each EIS spectra was repeated at least once.

[0286] SEM and TEM characterization: SEM characterization was performed using a Zeiss Merlin SEM. Samples were carbon-coated prior to imaging. Grain size measurements were performed using the line intercept method, with at least 100 intercepts at each temperature. Grain growth activation energies were found by assuming ideal grain growth (grain growth exponent = 0.5), and an initial grain size of 30 nm. TEM samples were prepared using two methods: cross-section wedge polishing and FIB liftout. In either case, final polishing was performed by low-energy broad-beam Ar+ion milling. STEMcharacterization was performed on a Thermo Fisher Scientific Themis Z S / TEM at 200 keV. EDS quantification was performed using the Velox software. For line traces, a 5 px mean pre-filter was applied prior to quantification. For spectrum image maps, a 5 px standard deviation Gaussian pre-filter was applied. Empirical background subtraction was performed prior to quantification.

[0287] In diffusion measurements by SIMS: Diffusion profiles of Mg and Al following annealing of the 3 at% Gd doped ceria films grown on MgO and Al2O3 substrates respectively were obtained through SIMS analysis measured by TOF-SIMS5 instrument (ION-TOF GmbH, Germany). Positive detection mode was used for the detection of Ce+, Gd+, Al+, and Mg+. Continuous bombardment of focused Bi+ion beam was made over a 100 μm × 100 μm region during the measurement with a 30 keV Bi ion gun. Released secondary ions were analyzed using a time-of-flight-detector. The sputter beam was set at 2 keV for O2+ ion-sputtering made over a region of 300 μm × 300 μm throughout the analysis. The analysis was conducted with a cycle time of 100 μs.

[0288] Space Charge Model

[0289] The conductivity of a grain boundary across the space charge zone can be defined as Equation 1: , (1)which describes how the space charge zone conductivity relates to the bulk conductivity (^^^^^^^^^^) and the depletion of ionic carriers within the depletion zone following the spatial distribution along x (described by space charge width - λ), of the space charge potential in the vicinity of the interface. Equation 1 shows that as the space charge potential approaches zero, the grain boundary conductivity approaches the bulk conductivity.

[0290] The potential distribution is found by solving Poisson’s equation, which describes how the system responds to the presence of a net core charge Qcore, but no exact analytical solution can be obtained. One can, however, derive an approximate expression in the case of sufficiently large space charge potentials3). For the sake ofsimplicity, we can also consider a simplified dilute limit picture and apply the Mott-Schottky case to describe the space charge zone, whereby we assume the precursor element profile to be flat and frozen-in during processing, and only the oxygen vacancies can redistribute within the proximity of the grain boundary. As previously derived in this simplified scenario,the grain boundary resistance and its relation to the space charge potential in the grain boundary core (∆^^(0)) can be obtained by integrating Equation 1, yielding Equation 2:

[0291] The effective grain boundary space charge width, defined as ^^^^^^ = 2λ, generally thought of as the electrical grain boundary width, where we typically neglect the grain boundary core contributions is defined according to λ, being the Mott-Schottky space charge width using Equation 3:which relates to the Debye length LD (first term in Equation 3 in parenthesis) and the space charge potential. The space charge potential, in turn, is defined according to the balance between the core charge at the grain boundary and the amount of opposite compensating charges in the space charge zone. In the Mott Schottky case, where the bulk dopant is assumed to have a flat concentration profile, the space charge potential in the core can be equated to Equation 4: Δϕ(4)

[0292] where Qcore (Core Charge in C / cm2) and ^^ The dielectric constant. This Equation 4 shows how the potential barrier heights stem from a charge balance strongly dictated by the core charge (Qcore) and secondarily influenced by the bulk dopant concentration. Equation 4 explains both the effectiveness of our chemical modification of the space charge potential, owing to the squared term in the numerator, and why we decided on a lightly doped Gd doping level (3at%) as the concentrations of the bulk dopant. [^^^^ce′] is in the dominator, with smaller values leading to larger space charge potentials.

[0293] It may also be possible to derive an approximate solution for another scenario, called the Gouy-Chapman case, that considers the bulk dopant to also be mobile and able to redistribute in the space charge zone throughout the measurement conditions. In that case, the space charge width is smaller and equates to the Debye length (e.g., λ = LD =The Gouy-Chapman case, however, is only valid when the bulk is mobile during the measurement conditions, which is only valid at elevated temperatures (> 1000 °C) and may generally not considered for lower temperature conductivity measurements, where the bulk dopant is effectively immobile during the measurement time,warranting the use of the Mott-Schottky approximation. This equation is a first-order approximation that can lead to errors in absolute values, as non-uniform concentrations of the bulk dopant profile in the grain boundary vicinity arise during high-temperature processing conditions and end up subsequently frozen-in upon cooling (330%). This leads to a so-called restricted equilibrium scenario (See FIG.17B and FIGS.21A–21D), which is out of equilibrium and lies between the Mott-Schottky and Gouy-Chapman cases. This may only be evaluated via numerical simulations with knowledge of the dopant frozen-in profile, impractical for most experimental studies.

[0294] Furthermore, for this example, we use epitaxial films as a measure of the bulk resistance to fit the space charge potential, assuming it does not substantially change with the annealing protocol through the sintering temperature window. This is to be acceptable in our case as we are more interested in general trends, and the results obtained from the resistivity ratio can be self- consistently compared.

[0295] While most studies of GB resistivity in ionic conductors derive a single effective barrier potential, GBs in given solids differ depending on their misorientation angles between adjacent grains, impacting impurity and defect segregation, and leading ultimately to a range of GB potentials that may not be easily distinguished from simple electrical measurements. This may be exacerbated by the fact that electrical measurement techniques measure the weighted average of all current paths across the numerous grain boundaries, dominated by the paths of least resistance. This implies that approximate analytical solutions for fitting grain boundaries space charge potentials based on electrical measurements are mainly used to consider general trends and for the sake of discussion, as we do in this example.

[0296] Electrode Geometry

[0297] Table 1: List of electrode dimensions deposited on GDC films grown on Al2O3 and MgO and annealed at various temperatures.

[0298] Structural Characterization

[0299] The X-ray diffraction patterns of both GDC polycrystalline films 1310a and 1310b grown on Al2O3 substrate 1320b and MgO substrate 1320a are consistent with the expected cubic fluorite structure exhibiting grains with (111) and (200) dominant orientations (see FIG.18C). The XRD patterns of the underlying substrate, Al2O3, and MgO are present at 41.5º 2θ and 42.9º 2θ.

[0300] Structural Evolution of GDC3 films on MgO as a function of temperature (XRD / SEM)

[0301] We display the 2 theta scans for the GDC films grown on Al2O3substrates and subsequently annealed from 650ºC up to 1250ºC in FIGS.20A and 20B. Similar to the as- grown films, no evidence of secondary phases is observed. This is additionally corroborated from STEM images shown in the section titled Impedance fitting and space charge model fitting of select boundaries at various key temperatures, where no evidence of secondary phases can be observed. While no clear trend in change in 2 theta position may be observed for the dominant film peaks, with increasing annealing temperature, the peak height of the (111) decreases relative to the (200) peak orientation, indicating a gradual increase in preferential film orientation texture towards the (200) orientation, which is fully achieved by 1000 °C, at which point the (111) peak is no longer visible. Moreover, these rocking curves displayed in FIGS.20C, performed on the dominant (200) peak, shows an apparent decrease in its full-width half max, implying a decrease in wide angle in favor of lower angle grain boundaries, consistent with grain growth and improvement in film quality. We observe similar trends for films grown on MgO and annealed from 650 ºC to 1250 ºC, as displayed in FIGS.21A–21D, though the disappearance of (111) may occur at lower temperatures (e.g., about 750°C or below).

[0302] Scanning electron microscope (SEM) measurements of the surface of the films were performed to characterize the microstructure and systematically measure the grain size of the synthesized films. In FIG.20D, representative SEM micrographs are shown forGDC films grown on Al2O3 substrates, while similar plots are present in FIG.21D for GDC films grown on MgO substrates. For both cases, similar microstructural evolution was observed. As deposited, and at low annealing temperatures, films are equiaxed with fine grain structure (d=30 nm). A columnar grain structure is observed at intermediate temperatures, e.g., 950 °C. Additionally, abnormal grain growth occurs, may be due to significant surface energy differences between grains of preferred orientation and grains of similar orientation merging during growth. At extreme temperatures, e.g., 1150°C, the average grain size may become much greater than the film thickness, and dewetting may occur, leaving roughly spheroidized grains.

[0303] By applying an Arrhenius fit to the grain sizes measured at different temperatures, an activation energy for grain growth of 1.55 eV (Al2O3) and 1.32 eV (MgO) is found in FIGS.22A and 22B. Details of the grain size measurements and fitting process are provided in the methods section. Similar microstructural evolution trends and activation energies in nanocrystalline ceria thin films have previously been reported, when columnar thin films of 26 at% and 10 at% Gd doped CeO2 were grown on Al2O3 substrates by spray pyrolysis and PLD. These exhibited mixed polycrystalline films characterized by (111) / (100) orientations, which, when exposed to increased thermal anneals (from 600 °C to 1200 °C), ultimately also led to single grain orientation. However, in those cases, the (111) orientation ended up dominating. Contrary to our initial films, we observed a predominance in (111) orientation in the as-deposited state, whereas we observe a predominance of (100), which suggests that the end results may arise from a competition in grain boundary interface mobility. Presumably, the (100) / (111) may be the slowest, and our final results depend on the initial distribution of boundary orientations. Additional differences in grain boundary mobility may arise from the in-diffusion of Al versus Mg.

[0304] Further evidence that the trends in changing microstructure may not cause the major observations in conductivity observed in this example stems from the case of the GDC3 film grown on Al2O3 annealed at 1150 °C, where an unexpected decrease in conductance relative to the sample annealed at 1050 °C is observed, as displayed in FIG. 14B. This may occur even though the activation energy continues to decrease. Without being bound to a particular theory, the change in conductance may coincide with the decrease in grain size that can be observed in FIG.20D and FIG.22A. The abnormal morphology observed for the sample annealed at 1050°C may have reached a point of instability and broken up due to surface tension. The decreased grain size may in turn correlate with anincrease in total number of grain boundary barriers along the current path adding additional series resistance, consistent with the decrease in conductance. This supports the idea that the change in activation energy may not be caused by changes in microstructure. The activation energy trend reported in FIG.15 remains consistent with the thermal annealing trend.

[0305] Impedance spectra for all GDC3 films on MgO

[0306] Exemplary measured electrochemical impedance spectra (EIS) for thin films grown on Al2O3 and MgO measured at 350°C are shown in FIGS.23A and 23B. A single large semicircle may be observed at higher frequencies, with a small tail at lower frequencies, consistent with our previous measurements on similar thin film samples. The main observed semicircle for GDC is ascribed to the combination of bulk and grain boundary contributions. The bulk contribution is not visible because it is shielded by the stray capacity that forms between the substrate and the electrodes. The low-frequency arc may be expected to be associated with the electrode contribution. The characteristic frequency and time constant τ (=2πfchar)−1of the main semicircle may depend mainly on the sample’s grain boundary resistance and electrode geometry. Note for polycrystalline GDC3 films grown on Al2O3, the lowest resistance measured at 350 °C, observed in the inset of FIGS.23A and 23B , is equal to approximately 4×106Ω, which is 103times more resistive than the most resistive film annealed at 850 °C, possessing a value of approximately 7×109Ω. The epitaxial value is even smaller and is barely visible in the inset near the origin, with a resistance value of approximately 0.5×106Ω. On the other hand, for the case of films on MgO, the lowest values are recorded for films annealed at 1050 °C, equal to approximately 0.7×106Ω, vs the highest resistance for the film grown at 850 °C with a value of 1×108Ω.

[0307] As shown in FIG.24, the nominally epitaxial films grown on MgO and Al2O3 show similar activation energies but deviate in absolute resistance by a factor of about 5. Variation in measured conductance may arise from differences in film thickness, electrode geometry, digit spacing, and / or built-in strain owing to the differing near epitaxial growths on the two structurally different substrates. These near epitaxial films may serve as bulk conductance references against those of the respective polycrystalline films.

[0308] Impedance Fitting and Space Charge Model Fitting

[0309] Capacitance Values: The spectra are fit using distributed elements composed of a R / / CPE, where CPE is a constant phase element from which capacitance values were calculated based on the expression C=(R(1-n)Q)(1 / n), where Q is the constant phase element capacitance, R the resistance and n the non-ideality factor obtained from the fittingprocedure. The measured capacitances obtained for all films, including the epitaxial ones, are all about equal to approximately 1-5×10-12F. Using Table 1, we can calculate the expectedsubstrates capacitance assuming a simple parallel plate capacitor, described by: ^^ =^^^^ ^^^^^^0( ^^ ), where L is the length of digits, N the number of pairs of digits, d the distance between digits, and εrε0 the dielectric constant of the materials. From this relation, we obtain capacitance values ranging from about 7×10-13to about 5×10-12F, which align well with the measured values.

[0310] Space charge Model Fitting: As discussed in the Space Charge Model section, we can use Equation 2 to fit our impedance results to obtain an approximate value for the space charge potentials of each film. As discussed in the section titled Impedance spectra for all GDC3 films on MgO, at higher frequencies we may only observe a single large semicircle in each impedance spectra for all films, which may be ascribed to the combination of bulk and grain boundary contributions. The bulk contribution may not be visible because it may be shielded by the stray capacity that forms between substrate and electrodes. To use Equation 2, we therefore use the impedance value of the epitaxial samples as a first order approximation of the bulk resistance. This should be a valid approximation as the bulk resistance of the polycrystalline films may not be expected to vary with annealing temperature (see section titled Structural Evolution of GDC3 films on MgO as a function of temperature (XRD / SEM), indicating that there may be no obvious change in strain state of the film that would induce a change in bulk conductance). Fitting procedure and results are displayed in FIGS.25A–25C.

[0311] As displayed in FIGS.26A–26B, for both Al and Mg, the concentration of precursor elements within the GBs exhibits a saturated behavior when annealed at 850 °C for 6 hours, where the concentration of in-diffused species presumably reached the solubility limit within the grain boundaries of the deposited GDC films. For samples deposited on Al2O3, the saturation concentration slightly decreased with increasing annealing temperatures above 950 °C. This may be rationalized by noting the reduction in grain boundary concentration due to grain growth during high-temperature heat treatments that eventually lowered the apparent grain boundary solubility per surface area. Meanwhile, samples deposited on MgO showed increased Mg concentration at the surface following annealing at temperatures above 950 °C. The precursor element concentration increased continuously with increases in annealing temperature, suggesting additional diffusion may occur along thefilm surface, potentially exacerbated by the increased amount of surface area due to grooving and / or dewetting of the film.

[0312] For annealing profiles below 950 °C, grain boundary diffusion coefficients were fitted by applying a 1D diffusion equation (Equation 5) to the precursor element depth profiles, which were normalized (see FIGS.29 and 30) with respect to the host Ce+signals. I(x) = A ∙ erfc(5) 2√DGB ∙ 21600

[0313] Here, I(x) is the normalized intensity at depth x, A in a proportional constant, and DGB is the diffusion coefficient of the precursor element. We adopted constant source grain boundary conditions for the precursor elements, and the background signal (L) was set by the precursor element concentration in the as- deposited samples. The diffusion behavior of precursor elements within the grains was also investigated by fitting the near-interface region depth profiles above 1000ºC (see FIGS.31 and 32) using a modified solution for diffusion from a slab source as shown in Equation 6. x + h (x − h) I(x) = a [erf ( b D t + 4σ2) − erf ( √4D t + 4σ2)] + I (6) √4 b b

[0314] I(x) is the normalized intensity of precursor elements at depth x, a is the proportional constant, h is the thickness of the film, σ is the SIMS mixing parameter, Ib is the normalized background intensity, and Db is the bulk diffusion coefficient. The fitted results for grain boundary and bulk diffusivities obtained from SIMS profiles for both Al and Mg are shown in FIG.33. Below we provide the fitted equations for each diffusion coefficient alongside the calculated diffusivities extrapolated to 850 ºC for comparison. The activation energies for diffusion along the grain boundaries are ½ and 1 / 5 of the bulk values (for Al2O3 1.3 eV vs 2.6 eV and MgO 0.84 eV vs 4.1 eV) and the differences in diffusivities for the bulk vs grain boundaries at a common temperature, e.g.850 °C, through extrapolation is about 5 orders of magnitude (^^^^^^,^^^^ / ^^^^,850 is approximately 10−16^^^^2 / ^^, while ^^^^^^^^^^,^^^^ / ^^^^,850 is approximately 10−21^^^^2 / ^^).

[0315] Equations for Grain Boundary Diffusivity and Bulk Diffusivity (cm2 .s-1)^)

[0316] Diffusivity (cm2 .s-1) at 850 °C ^^^^^^,^^^^,850 = 1.11 × 10−15(11)^^^^^^,^^^^,850 = 4.08 × 10−16(13) ^^^^^^^^^^,^^^^,850 = 5.85 × 10−21(14)

[0317] STEM-EDS Dataset Collected from Grain Boundaries of GDC3 Grown on Al2O3and MgO

[0318] Note on Si segregation at grain boundaries: From FIGS.38A–40D we observe Si segregation to the grain boundaries, but only at elevated temperatures (e.g., above > 1000 °C). Its impact on GB resistance remains debatable. Some literature suggests interstitial incorporation due to its small size (radius 0.4 Å – 6 fold coordination), imparting a positive core charge and increasing the boundary resistance due to an increase in space charge potential, while others indicate the possible formation of insulating siliceous phases, also increasing boundary resistance, but gettering impurities leading to decreased potential barriers. As shown in FIGS.14A and 14B, increases in resistance at higher temperatures may not be observed, while decreases in space charge potential may be observed. This would be indicative of the second scenario (e.g., gettering interface impurities). At the same time, given that the resistance did not increase, this suggests that any “siliceous phases” formed may be present only in non-continuous patchy form. Furthermore, clear differences in resistance between the series of thin films on Al2O3 versus those on MgO remain, even at higher temperatures. This indicates that if Si is playing a gettering role, it may be only partial and may not significantly impact the observed conductance trends which we assign to the presence of Mg and Al.

[0319] High Temperature Restricted Equilibrium Scenario

[0320] FIG.41 shows a physical model explaining spatial distributions of space charge potential, bulk dopant 4180 concentration, majority oxygen vacancy, and minority electron defect concentrations in response to combined substrate cation precursor elements 4170a / 4170b up-diffusion and bulk dopant 4180 Gd segregation at the GB 4112 of polycrystalline films 4110a / 4110b. As samples are annealed at higher temperatures, the grain boundaries 4112 may become saturated with precursor element 4170b Al (right) andprecursor element 4170a Mg (left), while ultimately, bulk dopant 4180 Gd segregation at the grain boundary 4112 may further reduce the space charge potential for samples with MgO (right), while counteracting the increase observed with GB decoration by Al in samples with Al2O3 (left). The concentration profile and space charge width may follow a restricted equilibrium scenario instead of simple mott Schottky approximation.

[0321] FIG.41 describes the expected impact of competition between substrate precursor element 4170a / 4170b up-diffusion, at lower temperatures, and bulk dopant 4180 segregation to the grain boundary 4112 core occurring at higher temperatures. This may lead to a modified restricted equilibrium scenario where the space charge width becomes smaller,proportional to the Debye length (^^^^ =and less sensitive to the space charge potential itself.

[0322] Conclusion

[0323] While various inventive embodiments have been described and illustrated herein, those of ordinary skill in the art will readily envision a variety of other means and / or structures for performing the function and / or obtaining the results and / or one or more of the advantages described herein, and each of such variations and / or modifications is deemed to be within the scope of the inventive embodiments described herein. More generally, those skilled in the art will readily appreciate that all parameters, dimensions, materials, and configurations described herein are meant to be exemplary and that the actual parameters, dimensions, materials, and / or configurations will depend upon the specific application or applications for which the inventive teachings is / are used. Those skilled in the art will recognize, or be able to ascertain using no more than routine experimentation, many equivalents to the specific inventive embodiments described herein. It is, therefore, to be understood that the foregoing embodiments are presented by way of example only and that, within the scope of the appended claims and equivalents thereto, inventive embodiments may be practiced otherwise than as specifically described and claimed. Inventive embodiments of the present disclosure are directed to each individual feature, system, article, material, kit, and / or method described herein. In addition, any combination of two or more such features, systems, articles, materials, kits, and / or methods, if such features, systems, articles, materials, kits, and / or methods are not mutually inconsistent, is included within the inventive scope of the present disclosure.

[0324] Also, various inventive concepts may be embodied as one or more methods, of which an example has been provided. The acts performed as part of the method may be ordered in any suitable way. Accordingly, embodiments may be constructed in which acts are performed in an order different than illustrated, which may include performing some acts simultaneously, even though shown as sequential acts in illustrative embodiments.

[0325] All definitions, as defined and used herein, should be understood to control over dictionary definitions, definitions in documents incorporated by reference, and / or ordinary meanings of the defined terms.

[0326] The indefinite articles “a” and “an,” as used herein in the specification and in the claims, unless clearly indicated to the contrary, should be understood to mean “at least one.”

[0327] The phrase “and / or,” as used herein in the specification and in the claims, should be understood to mean “either or both” of the elements so conjoined, i.e., elements that are conjunctively present in some cases and disjunctively present in other cases. Multiple elements listed with “and / or” should be construed in the same fashion, i.e., “one or more” of the elements so conjoined. Other elements may optionally be present other than the elements specifically identified by the “and / or” clause, whether related or unrelated to those elements specifically identified. Thus, as a non-limiting example, a reference to “A and / or B”, when used in conjunction with open-ended language such as “comprising” can refer, in one embodiment, to A only (optionally including elements other than B); in another embodiment, to B only (optionally including elements other than A); in yet another embodiment, to both A and B (optionally including other elements); etc.

[0328] As used herein in the specification and in the claims, “or” should be understood to have the same meaning as “and / or” as defined above. For example, when separating items in a list, “or” or “and / or” shall be interpreted as being inclusive, i.e., the inclusion of at least one, but also including more than one, of a number or list of elements, and, optionally, additional unlisted items. Only terms clearly indicated to the contrary, such as “only one of” or “exactly one of,” or, when used in the claims, “consisting of,” will refer to the inclusion of exactly one element of a number or list of elements. In general, the term “or” as used herein shall only be interpreted as indicating exclusive alternatives (i.e. “one or the other but not both”) when preceded by terms of exclusivity, such as “either,” “one of,” “only one of,” or “exactly one of.” “Consisting essentially of,” when used in the claims, shall have its ordinary meaning as used in the field of patent law.

[0329] As used herein in the specification and in the claims, the phrase “at least one,” in reference to a list of one or more elements, should be understood to mean at least one element selected from any one or more of the elements in the list of elements, but not necessarily including at least one of each and every element specifically listed within the list of elements and not excluding any combinations of elements in the list of elements. This definition also allows that elements may optionally be present other than the elements specifically identified within the list of elements to which the phrase “at least one” refers, whether related or unrelated to those elements specifically identified. Thus, as a non-limiting example, “at least one of A and B” (or, equivalently, “at least one of A or B,” or, equivalently “at least one of A and / or B”) can refer, in one embodiment, to at least one, optionally including more than one, A, with no B present (and optionally including elements other than B); in another embodiment, to at least one, optionally including more than one, B, with no A present (and optionally including elements other than A); in yet another embodiment, to at least one, optionally including more than one, A, and at least one, optionally including more than one, B (and optionally including other elements); etc.

[0330] In the claims, as well as in the specification above, all transitional phrases such as “comprising,” “including,” “carrying,” “having,” “containing,” “involving,” “holding,” “composed of,” and the like are to be understood to be open-ended, i.e., to mean including but not limited to. Only the transitional phrases “consisting of” and “consisting essentially of” shall be closed or semi-closed transitional phrases, respectively, as set forth in the United States Patent Office Manual of Patent Examining Procedures, Section 2111.03.

Claims

CLAIMS What is claimed is:

1. A detector for radiation, the detector comprising: an ion-conducting polycrystalline material to absorb the radiation, wherein the ion- conducting polycrystalline material comprises grain boundaries and a precursor element diffused within the grain boundaries; a pair of electrodes, electrically coupled to the ion-conducting polycrystalline material, to apply a voltage across the ion-conducting polycrystalline material; and a sensor, electrically coupled to the pair of electrodes, to measure a change in conductance of the ion-conducting polycrystalline material caused by absorption of the radiation.

2. The detector of claim 1, wherein the grain boundaries are positively charged grain boundaries and wherein the precursor element is a positively charged precursor element.

3. The detector of claim 2, wherein the positively charged precursor element increases the positive charge of the grain boundaries.

4. The detector of claim 1, wherein the grain boundaries are negatively charged grain boundaries and wherein the precursor element is a negatively charged precursor element.

5. The detector of claim 4, wherein negatively charged precursor element increases the negative charge of the grain boundaries.

6. The detector of claim 1, wherein the precursor element increases a space-charge potential barrier of the grain boundaries.

7. The detector of claim 1, wherein the precursor element is substitutionally inserted within the grain boundaries.

8. The detector of claim 1, wherein the precursor element is interstitially inserted within the grain boundaries.

9. The detector of claim 1, wherein the precursor element has an ionic radius at least 20% smaller than an ionic radius of a host element of the ion-conducting polycrystalline material.

10. The detector of claim 1, having a room temperature dark resistance >1010ohm.

11. The detector of claim 1, having a sensitivity ∆R / R of about 103.

12. A method of diffusing an ion-conducting polycrystalline material with a precursor element, the method comprising: selecting a precursor element based on a charge of an ionic carrier of the ion- conducting polycrystalline material; contacting a film the ion-conducting polycrystalline material with a source of the precursor element; and annealing the source of the precursor element and the film of the ion-conducting polycrystalline material at a temperature between 600 ℃ and 900 ℃ for 2 hours to 48 hours.

13. The method of claim 12, wherein selecting the precursor element comprises: comparing an ionic radius of the precursor element to an ionic radius of the ionic carrier of the ion-conducting polycrystalline material; and selecting the precursor element with an ionic radius at least 20% smaller than the ionic radius of the ionic carrier of the ion-conducting polycrystalline material.

14. The method of claim 12, wherein the source of the precursor element is a film of the precursor element.

15. The method of claim 14, wherein the film of the precursor element is deposited below the film of the ion-conducting polycrystalline material, the method further comprising: diffusing, by the precursor element, from the film of the precursor element into grain boundaries of the ion-conducting polycrystalline material.

16. The method of claim 14, wherein the film of the precursor element is deposited above the film of the ion-conducting polycrystalline material, the method further comprising: diffusing, by the precursor element, from the film of the precursor element into grain boundaries of the ion-conducting polycrystalline material; andafter annealing the film of the precursor element and the film of the ion-conducting polycrystalline material, etching the film of the precursor element to remove the film of the precursor element.

17. The method of claim 12, wherein the source of the precursor element is a substrate comprising the precursor element and wherein contacting the film the ion-conducting polycrystalline material with the source of the precursor element further comprises depositing the film of the ion-conducting polycrystalline material on a surface of the substrate.

18. A detector for radiation comprising the ion-conducting polycrystalline material of claim 12.

19. A method of diffusing an ion-conducting polycrystalline material with a precursor element, the method comprising: mixing particles of the ion-conducting polycrystalline material with particles of the precursor element to form mixed particles; drying the mixed particles at 80 ℃ to 120 ℃ to form coated particles comprising the ion-conducting polycrystalline material coated with the precursor element; calcinating the coated particles at about 400 ℃ to about 600 ℃ for about 2 hours to partially decompose the precursor element coating of the coated particles; pressing the coated particles to form a consolidated body; and sintering the consolidated body at 600 ºC to 1000 ℃ for 2 hours.

20. A detector for radiation comprising the ion-conducting polycrystalline material of claim 19.

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