Applications of electric NANO pulsing (ENP) technology

Abstract

Systems for material grain boundary engineering are described. The system is for releasing stored electric energy into polycrystalline materials within micro to nano seconds for quasi-instantaneous processing of polycrystalline materials by means of modifying grain boundary structures, microstructure and properties of the polycrystalline material.

Description

Applications of Electric Nano Pu smg (ENP) TechnologySTATEMENT REGARDING FEDERALLY SPONSORED RESEARCH OR DEVELOPMENT

[0001] This invention was made with government support under DE-SC0022244 awarded by the Department of Energy. The government has certain rights in the invention.CROSS REFERENCE TO RELATED APPLICATIONS

[0002] The current application claims the benefit of and priority under 35 U.S.C. 119(e) to U.S. Provisional Patent Application No. 63 / 613,660 filed December 21 , 2023, the disclosure of which is incorporated herein by reference in its entirety.FIELD OF THE INVENTION

[0003] This application generally refers to electropulsing technology and systems. More specifically, this application relates to electropulsing methods, techniques, and systems for modifying grain boundary structures for polycrystalline metallic materials.BACKGROUND

[0004] Grain boundary engineering (GBE) is the methodology by which grain boundary (GB) structure is characterized and material processing variables adjusted to create an optimized GB microstructure for improved material properties. GBE has been applied to numerous materials to improve their mechanical, physical, chemical, electromagnetic, optical, and biological properties. Thermomechanical treatment (TMT), doping, and magnetic / e-field assisted are several of the most widely used GBE processing methods. Thermomechanical treatment involves repeated cycles of large plastic deformation and annealing to generate grain boundaries. Doping is usually induced by high-temperature annealing, i.e., grain boundary segregation engineering (SE). SE uses segregation to optimize specific GB structure, composition, and properties to design beneficial material behaviors. Other processing methods are often used inaddition to or in conjunction with TMT and SE for GBE, such as unidirectional, rotational, and rapid solidification, magnetic field application such as magnetic annealing, electric field (E-field) application such as current-assisted sintering, additive manufacturing, and other advanced techniques. GBE has proven to be effective in improving various material properties, yet challenges persist in the finer localized control of GB interface structures and achieving desirable function for a given material.SUMMARY OF THE INVENTION

[0005] Systems and methods in accordance with some embodiments of the invention are directed electropulsing technology for grain boundary engineering and modifying grain boundary structures for polycrystalline metallic materials.

[0006] Many embodiments of the disclosure are directed to a device for generating high-current density nano pulses including: an electrical source, at least one capacitor, a pulse signal generator, and a switch configured for electric current pulsing; wherein the electric current pulsing is at a set intensity, a set duration, and a set frequency and configured for at least one of a short rise time and a short fall time; wherein the device is configured to electrically couple to a conductive material and release a current into the conductive material; and wherein the current released is configured to induce a property change in the conductive material; wherein the current density exceeds 1010A / m2and the pulse duration exceeds 100ns.

[0007] In many embodiments, the current density is approximately 1010-1012A / m2.

[0008] In many embodiments, the pulse duration is less than 1 ps.

[0009] In many embodiments, the pulsing frequency is up to 100KHz.

[0010] In many embodiments, the capacitor is configured for energy storage and charging.

[0011] In many embodiments, the capacitor is a supercapacitor.

[0012] In many embodiments, the switch is a solid state switch.

[0013] In many embodiments, the capacitor is configured for charging with a voltage of up to 1000V.

[0014] In many embodiments, the pulse signal generator outputs a square wave signal.

[0015] In many embodiments, the square wave signal output ranges from 100ns to a 5ps.

[0016] In many embodiments, the device further includes at least one constant resistor.

[0017] In many embodiments, the at least one constant resistor is a plurality of parallel constant resistors with a resistance of approximately 95m Q.

[0018] In many embodiments, the at least one constant resistor is configured for overcurrent protection.

[0019] Many embodiments of the disclosure are directed to a method of material grain boundary engineering including, charging and storing energy in a capacitor, releasing the stored energy with a pulse signal generator into an electrically coupled material; wherein the stored energy is discharged at a set intensity, a set duration, and a set frequency; producing an electric current pulse; inducing a change in the microstructure of the electrically coupled material; wherein the change in the microstructure is configured for a selected property of the material; wherein the current density exceeds 1010A / m2and the pulse duration exceeds 100ns.

[0020] In many embodiments, the material is a polycrystalline material.

[0021] In many embodiments, the releasing electric energy into the material activates a non-equilibrium structural evolution of the material.

[0022] In many embodiments, the change in the microstructure is at a nanometer spatial scale.

[0023] In many embodiments, the change in the microstructure is at a nanosecond temporal scale.

[0024] In many embodiments, the pulse signal generator outputs a square wave signal.

[0025] In many embodiments, the output ranges from 100ns to 5ps.

[0026] In many embodiments, releasing the stored energy further includes activating a switch configured for at least one of a short rise time and a short fall time of the electric current pulse.

[0027] In many embodiments, the switch is a solid state switch.

[0028] In many embodiments, the current density is approximately 1O1o-1O12A / m2.

[0029] In many embodiments, the pulse duration is less than 1 ps.

[0030] In many embodiments, the pulsing frequency is up to 100KHz.

[0031] In many embodiments, the change in the microstructure further includes inducing dislocations at multiple scales.

[0032] In many embodiments, inducing dislocations at multiple scales further includes inducing nanoscale dislocations with densities up to 1017to 1018A / m2.

[0033] In many embodiments, inducing the change in the microstructure further includes modifying a grain boundary morphology.

[0034] In many embodiments, modifying a grain boundary morphology further includes producing at least one of step-like and serrated features without affecting an adjacent grain structure.

[0035] In many embodiments, the selected property of the material is selected from the group consisting of: the diffusion rate of an element, plasticity, tensile strength, workhardening performance, yield stress, elastic modulus, and corrosion resistance.

[0036] In many embodiments, the diffusion rate of an element is configured at least one of segregation at a grain boundary and the formation of precipitates.

[0037] In many embodiments, the change in the microstructure further includes forming a nanocoating with a set structure.

[0038] In many embodiments, the nanocoating is Cr20s and the set structure is a triple hierarchical structure.

[0039] In many embodiments, the method further includes heating at a grain boundary with a temperature difference between the grain boundary and a grain interior.

[0040] In many embodiments, heating at the grain boundary leads to at least one of a localized stress field and a grain boundary transformation.

[0041] In many embodiments, the temperature difference lasts less than 100 microseconds.

[0042] In many embodiments, the method further includes heating the material to near the material melting points and maintaining a grain size and an orientation.

[0043] In many embodiments, the method includes a heating rate of 108K / s.

[0044] In many embodiments, inducing non-equilibrium structure is configured to promote at least one of dislocation generation, dislocation reconfiguration, a specific dislocation configuration.

[0045] In many embodiments, the electric current pulses are configured for a dislocation distribution density.

[0046] Additional embodiments and features are set forth in part in the description that follows, and in part will become apparent to those skilled in the art upon examination of the specification or may be learned by the practice of the disclosure. A further understanding of the nature and advantages of the present disclosure may be realized by reference to the remaining portions of the specification and the drawings, which forms a part of this disclosure.BRIEF DESCRIPTION OF THE DRAWINGS

[0047] The description will be more fully understood with reference to the following figures, which are presented as embodiments of the invention and should not be construed as a complete recitation of the scope of the invention, wherein:

[0048] Fig. 1 illustrates a comparison of electric pulsing techniques in accordance with prior art.

[0049] Fig. 2 illustrates a temperature profile across a GB in accordance with many embodiments.

[0050] Figs. 3A and 3B illustrate the ratio, for a variety of microcrystalline materials in accordance with many embodiments.

[0051] Figs. 4A through 4D illustrate a representative volume element of a microcrystalline structure in accordance with many embodiments.

[0052] Fig. 5 illustrates the main material properties for the FEA Joule heating modeling in accordance with many embodiments.

[0053] Figs. 6A through 6D Illustrate GB current density in accordance with many embodiments.

[0054] Fig. 7 Illustrates the components of an ENP assembly in accordance with many embodiments.

[0055] Figs. 8A through 8C Illustrate a circuit diagram in accordance with many embodiments.

[0056] Fig. 9 illustrates exemplary expected current density data as a function of the sample in accordance with many embodiments.

[0057] Fig. 10 illustrates exemplary current density data as a function of the sample sizes in accordance with many embodiments.

[0058] Fig. 11 illustrates a Scanning Electron Microscopy (SEM) image of a segment melting and breaking after the ENP processing in accordance with many embodiments.

[0059] Figs. 12A through 12D illustrates SEM images of the polycrystalline grain structure within the sample before and after the ENP processing in accordance with many embodiments.

[0060] Fig. 13 illustrates EDX results on the internal regions and surface areas after 100kHz ENP processing in accordance with many embodiments.

[0061] Figs. 14A. through 14C illustrates GB facets and serrations formation during the ENP processing, in accordance with many embodiments.

[0062] Fig. 15, illustrates the temperature and thermal stress changes of GBs, in accordance with many embodiments.

[0063] Fig. 16 illustrates the structure mismatch angle in accordance with many embodiments.

[0064] Figs. 17A and 17B illustrate exemplary "bus-bar" and "through-hole" in accordance with many embodiments.

[0065] Fig. 18A through 181 illustrates the temperature and microstructure of nichrome alloy in accordance with many embodiments.

[0066] Fig. 19A through 19C illustrate the grain boundary distributions and orientations in nichrome alloys in accordance with many embodiments.

[0067] Fig. 20A through 20F illustrate dislocation morphology in accordance with many embodiments.

[0068] Fig. 21 illustrates tensile testing data in accordance with many embodiments.

[0069] Figs. 22A through 22C illustrate fracture surfaces in accordance with many embodiments.

[0070] Figs. 23A through 23H illustrate the evolution of exterior surface morphology in accordance with many embodiments.

[0071] Figs. 24A through 24I illustrate the morphology of nano coating after ENP processing in accordance with many embodiments.

[0072] Figs. 25 and 26 illustrate the composition by EDS measurement of many embodiments.

[0073] Figs. 27A through 27D illustrate conventional heating treatment in accordance with many embodiments.

[0074] Figs. 28A through 28C illustrate sample morphology in accordance with many embodiments.

[0075] Figs. 29A through 29D illustrate the internal microstructures of samples in accordance with many embodiments.

[0076] Figs. 30A though 30D illustrate dislocation configuration after nano pulsing current process at 100kHz in accordance with many embodiments.

[0077] Figs. 31 A through 31 D illustrate HRTEM images of exemplary nichrome alloy samples in accordance with many embodiments.

[0078] Figs. 32A and 32B illustrate edge dislocation in accordance with many embodiments.

[0079] Figs. 33A and 33B illustrate an HRTEM image of an exemplary sample and a schematic of the extended dislocation in accordance with many embodiments.

[0080] Figs. 34A through 34D illustrate finite element modeling of the microstructure of an exemplary embodiment in accordance with many embodiments.

[0081] Figs. 35A through 35C illustrate ENP-activated precipitate in Ni-based superalloy in accordance with many embodiments.

[0082] Figs. 36A through 36H illustrate the microstructure of field-activated precipitates in Ni-based superalloy in accordance with many embodiments.

[0083] Figs. 37A through 37D illustrate properties of Ni-based superalloy field- activated precipitates in Ni-based superalloy in accordance with many embodiments.

[0084] Figs. 38A through 38D illustrate SEM-BSE images of the Ni-based superalloy after various furnace heating treatments in accordance with many embodiments.

[0085] Figs. 39A through 39C illustrate the formation of field-activated precipitates in Ni-based superalloy in accordance with many embodiments.DETAILED DESCRIPTION OF THE INVENTION

[0086] Locally controlling grain-boundary microstructure without affecting the structure and property of its adjacent grain interiors (GIs) is a difficult challenge for many material processing techniques. Traditionally, materials processing has relied on the application of heat and pressure to alter material properties. Most GBE approaches involve either considerable plastic deformation, which modifies the overall geometry and / or microstructure of the material or are achieved by GB segregation, which is limited to strong segregation systems. In most conventional methods, heat dissipates quickly but electric fields, with very short durations (nanoseconds), can spike temperatures locally, modifying grain boundary structures. When an electric current passes through a material, it generates Joule heating, which increases the temperature. This additional heat can be precisely controlled to achieve desired material properties. The electrical resistance of grain boundaries differs from that of the grains due to variations in atomic structure. This difference leads to distinct Joule heating effects. Materials processing with electropulsing technology can offer precise control over material properties through high-intensity, short- duration electric pulses and can provide new capabilities that traditional GBE methods cannot achieve.

[0087] The disclosure includes several embodiments directed toward the use of electro-nano-pulsing (ENP) processing. ENP, in accordance with various embodiments,achieves localized GBE. The ability to locally manipulate GB microstructures can enable extreme controllability and tailoring of material properties that are controlled by GBs. Accordingly, ENP can uniquely modify material properties. ENP can target heterogeneities within crystalline structures, such as dislocations, grain boundaries, and interfaces. The capabilities of ENP surpass the limitations of current electric current- assisted technologies, offering a direct method to engineer material properties by precisely controlling both thermal (localized Joule heating) and non-thermal (electric field effect) phenomena.

[0088] ENP can be applied to a wide range of advanced materials processing and manufacturing techniques. ENP can achieve quasi-instantaneous modifications at the micro and nanoscale, which improve material processing and engineered applications such as defect reconfiguration within materials, surface nanocoating technologies, and the formation of nanostructures, not achievable through conventional thermal and electric current-assisted techniques.

[0089] A comparison of high-power electric pulsing techniques, in accordance with prior art, is illustrated in Fig. 1 , highlighting the capability, techniques and technologies in the prior art. Configurations like Blumlein excel in generating high current density pulses in ultra-short durations (~40ns) without pulsing profile control and with frequency under 100Hz. The Marx generator and traditional capacitor discharge methods also fall short in providing pulsing profile customization. The discharge profiles of Blumlein, Marx, and capacitor discharge techniques typically follow an exponential decay curve. The mathematical expression for the voltage across the capacitor as a function of time (t) is given by V(t) = V exp (-tl(RC)), where Vo is the initial voltage across the capacitor at t=0. R is the resistance, and C is the capacitance of the circuit. These technologies do not offer the precision required for ENP processing. Pulses shorter than 100ns primarily affect a material’s outmost surface due to Eddy current and have restricted the prior art, such as Blumlein-based methods, to surface treatments. In contrast, high-frequency operation, substantial current density, and programmable nano pulsing profile differentiate an ENP system from other high-power electric pulsing technologies.

[0090] In accordance with many embodiments, ENP is capable of locally changing GB structures. In many such embodiments, the GB structure can be altered without affecting the grain interiors. In some embodiments, ENP can modify the GB structure of polycrystalline metallic materials. In many embodiments, an EMP apparatus can be charged to an electrical load. In many such embodiments an intense nano pulse current can induce heat localization at GBs. In some such embodiments, the nano pulse current exceeds 1O10A / m2. In many embodiments the temperature difference between the GB core and its adjacent grain interior can exceed tens of degrees Celsius. In many embodiments, the nano pulse duration is less than 100 ps. In many embodiments, GB heat localization promotes morphological changes in grain boundaries. In many embodiments electron wind force promotes morphological changes in grain boundaries. In many embodiments, morphological changes occur at GBs with a large angle during the ENP processing. In many embodiments, morphological changes occur at GBs with high- energy states during the ENP processing. In many embodiments, ENP processing results in the formation of at least one feature selected from the group consisting of atomistic facets, nanoscale serrations, and nanoscale steps. In many embodiments, ENP processing enables GB tailoring. In many such embodiments, GB tailoring further enables the synthesis of materials with finely tuned structural and functional properties.Localized Engineering Of Grain Boundary Morphology

[0091] In accordance with many embodiments, localized Joule heating is achieved at GB regions when an electric current passes through the microcrystalline microstructure due to the large difference in electrical resistivity between GBs and GIs in a polycrystalline material. The increased electrical resistance in polycrystalline materials with a greater number of GBs is evidence of the difference in electrical resistivity between GBs and GIs, in many embodiments. From the atomic structure perspective, GBs are often fully or partially disordered with a lower atomic density compared to GIs. The region of low atomic density acts as an electron repulsion barrier and scatters conduction electrons in many embodiments. A schematic diagram of a possible temperature profile across a GB is illustrated in Fig. 2, showing that the GB temperature (TGB) is significantly higher than thatof the Gl (TGI), resulting in a steep temperature gradient across the GB, in accordance with many embodiments. With the same thickness of material, the electrical resistance at a GB and within a grain interior is represented as RGB and RGI, respectively, and their ratio can be described by:Where:

[0092] Parameters PGB and pci are the respective specific resistivity of the GB and the Gl. Parameter S is the GB area or the contacting area of a grain interior with its GB. LGI is denoted as a certain thickness of grain interior material that equals the thickness or width of a GB.

[0093] The width of GB is ~1-2 nm for the majority of microcrystalline materials when the segregation of impurities at the GBs is limited. For a certain GB width of microcrystalline materials, can then be computed. Figs. 3A and 3B illustrate the ratio, for a variety of microcrystalline materials assuming 1 nm GB width. Most GBs have 10 to 1000 times greater electrical resistivity values than that of the grain interior as can be seen in Fig. 3A. Values of specific electric resistivities of grain boundaries (GB, unit: Qm2) and grain interiors (Gl, unit: Qm) used to calculate the values of ratio for a variety of microcrystalline materials are illustrated in fig 3B; the GB width is assumed to be 1 nm (=10-9m).

[0094] The Andrews method is used to determine the grain boundary electrical resistivity. The Andrews method recognizes the electrical resistance of a polycrystalline material is inversely proportional to its grain size given the same sample dimensions. The smaller the grain size, the larger the electrical resistance of the sample is. By fitting thesample electrical resistance as a function of grain size, the average value of the electrical resistivity of grain boundaries is obtained

[0095] For a microcrystalline grain structure, the total electrical resistance of a sample is calculated via. fotal grain interior + ^grain boundaryAssuming a cylindrical (radius=r, length= / ) geometry of the microcrystalline sample, the total resistance can be expressed asWhere pgi, lgi, Agi, and pgb, lgb, Agbare the respective resistivities, lengths, and cross- sectional areas for the grain interiors and grain boundaries, Agi=Agbwhen the current is passing through the electrical load, such as along the longitudinal direction of the cylindrical sample. If the average grain size is D, and the average grain boundary width is w, thenAndWhich can be rewritten as:whereThe value of pgiis very close to the reported electrical resistivity of the material, which usually has a relatively large grain size level, e.g., ^50-100pm. Therefore, once multiple values of Rtotaiare measured for a known dimension cylindrical sample that possesses different average grain sizes D, the pgband the n values can be obtained, which is an inverse relation between Rtotatand D.

[0096] A high RGB to RGI ratio allows Joule heating to maximally heat the GB regions while minimally affecting the grain interiors. The electric current induced Joule heat at GBs can be significantly higher than that within the grain interiors. If Joule heating induced temperature rises at GBs and GIs are largely different, the GBs can be manipulated without causing significant change within the grain interiors. However, because of intrinsic heat conduction, the generated Joule heat rapidly spreads throughout the entire microstructure, often in less than 1 ms, and results in negligible temperature difference between GBs and grain interiors. However, with Joule heating from intense electric current on the order of 100 ns, it is possible to achieve significant heat localization at the GBs in accordance with many embodiments.Exemplary Data

[0097] In accordance with many embodiments, polycrystalline models consisting of 1 nm-width GBs surrounding the micrometer-sized grain interior Gl components in the microstructure were used, and Joule heating finite element analysis (FEA) modeling was conducted for the polycrystalline model. FEA is one of the most efficient and reliable methods to model electric current-induced Joule heating and heat conduction in polycrystalline materials. By FEA modeling ENP processing, the distribution of electric current density and induced Joule heating in microcrystalline bulk materials can be simulated in accordance with many embodiments, and the temperature profiles across GBs and heat localization at GBs can be obtained in accordance with many embodiments.

[0098] Fig. 4 illustrates a representative volume element (RVE) of a microcrystalline structure (100 pm by 100 pm) with a 1 nm-thickness GB separating each grain interior, in accordance with many embodiments. The grain size distribution follows a normal distribution with an average grain size of about 40 pm.

[0099] The model material system was selected for the following criteria: that the material possesses a high value of i.e. , a high RGB to RGI ratio, to achieve significantly different Joule heating between GBs and GIs; that material itself, or as an essential component material, have broad structural and / or functional applications; that the material system possess a large skin depth when the electric current passing through the materialhas a duration of less than 1 microsecond, allowing a relatively homogeneous current density across the entire sample; and that the material system should have good electrical conductivity and achieve high electric current density using commercially existing power supplies. Nichrome-80 solid solution superalloy was chosen as the model material to conduct FEA modeling of the ENP processing.

[0100] Fig. 5 depicts the main material properties (at 20 °C) used for the FEA Joule heating modeling and the calculated skin depth for the chosen material system. The temperature dependencies of electrical conductivity, electrical resistivity, thermal conductivity, heating capacity, and density were taken into account in the FEA Joule heating modeling.

[0101] As shown in Fig. 4, a single nano pulse of direct electric current (DC, / ) is introduced into the microcrystalline model in which the GB width is 1 nm, then, the timedependent Joule heating FEA modeling was carried out for the Nichrome-80 alloy. Fig. 4 also presents exemplary results for temperature distribution in the microcrystalline model under a nano pulse direct electric current with a duration of 1637ns and a current density of 6.10 x lO1oA / m2in the Nichrome-80 alloy. The temperature in the grain interior regions of the alloy reaches 1400 °C after one pulse duration of 1637ns. This temperature is just below the melting point of the Nichrome-80 alloy (1403-1405 °C). Also illustrated in Fig. 4 is the significant heat localization at the GBs under an intense single nano pulse direct electric current. The temperature in GB regions is higher than that in the grain interiors, showing localized excessive heating of GBs. The temperature profile across GB showed that the temperature at the GB center is -25 °C higher than that in the grain interior when a single 1637ns electric current pulse (with a current density of 6.10 x lO1oA / m2) is introduced into the material. The temperature in the grain interiors drops quickly further from the GB center, resulting in a large temperature gradient in the vicinity of GBs. Fig. 4 also illustrates FEA Joule heating modeling for electric current pulse durations at 2000ns, 1500ns, 1000ns, 500ns, 300ns and 200ns. When decreasing the electric current pulse duration, a higher current density is needed to achieve the same level of Joule heating in the microcrystalline system. The respective current densities corresponding to 2000ns, 1500ns, 1000ns, 500ns, 300ns, and 200ns were 0.905, 1.044, 1.278, 1.808, 2.335, and2.860 times that at the case of 1673ns, and the temperature difference between GB and Gl for each pulse duration increased with a decrease of the pulse duration. For example, the temperature at GB1 is 41 .8 °C higher than that in the Gl at a pulse duration of 500ns. The GBs have lower thermal conductivity and higher heating capacity compared to the grain interiors in microcrystalline metals, which enhances the Joule heating localization at GBs, in accordance with many embodiments. Fig. 5 illustrates the current density distribution and detailed profiles across grain boundaries in accordance with an exemplary embodiment. The current density difference between GB and Gl depends on the orientation relative to the external electric current ( / ) or field direction. GBs that are more perpendicular to the field direction experience similar levels of current density as those in their adjacent grain interiors. However, GBs that are oriented more parallel relative to the field direction possess different current densities compared to their adjacent grain interiors. The GB2’s current density in region A is ~5-6 times smaller than that in its adjacent grain interiors, as illustrated in the current density profile across the GB2 in Fig. 6. The large difference in the current density between GBs and GIs creates a very steep current density gradient in the vicinity of GB2 plane.Exemplary Apparatus Embodiments

[0102] Fig. 7 depicts the components of an ENP assembly in accordance with various embodiments. The EMP assembly, according to such embodiments may comprise a high voltage DC power supply 700, an Integrated Power Module (IPM) 702, a nano pulse generator 704, and an oscilloscope 706. The IPM unit 702 is configured with Cu feed plates. The maximum voltage and single pulse current of the IPM unit 702 are 1 ,000V and 10,000A, respectively. The nanosecond (ns) pulsed electric direct current (DC) is output via the Cu feed plates, and a specimen can be electrically coupled to the feed plates. The pulsation is controlled by the nano pulse generator 704 and monitored by the oscilloscope 706. The feed plates are configured with a number of shunt resistors connected in parallel, each having a nominal resistance of 3.3Q. 33 parallel shunt resistors provide nominal 0.100Q. The ENP system can achieve an electric current of10,000A at a charging voltage of 1000V. The ENP assembly is configured to allow installation of samples with various geometry such as rod, wire, and thin films.

[0103] In many embodiments, the IPM is configured for the generation of high-intensity DC pulses. In many embodiments the, the pulse duration is as short as 100 ns. In some embodiments, the pulse duration is up to 1 ms. In many embodiments, the IPM’s charging voltage is produced from 120 AC, which is convenient to operate in regular lab conditions. In many embodiments, the IPM comprises a fast capacitor is configured for storing electrical energy. Many embodiments are configured so that with the release of electrical energy, a short pulse of large electric current is generated. In many embodiments, the current passes through the specimen (electrical load). Many embodiments further comprise a solid-state switch. In many such embodiments, the solid- state switch is configured to control the duration of the electric pulse, which in many such embodiments is further configured for ultra-fast rise / fall time. Some such embodiments are configured for a rise / fall time of <40ns of an intense electric current during pulsing. In many embodiments, electric current travels along the top plate from the IPM to the outside of the specimen, and through the specimen, along a bottom plate back to the IPM. In some embodiments the sample specimen is installed inside a through-holes on the feed plates that are integrated with the IPM device.

[0104] In many embodiments, the feed plates are configured as a pair of planar sections. The inductance of each plate (Lpiate) is given by:o is the permeability constant of free space (=4TT X 10'7«1 .2566 x lO-6H / m) and the sample and outside feed plates can be configured as a short section of a coaxial conductor, with an inductance (Lconductor) of:In accordance with an exemplary embodiment the inductance of the feed plates and conductor were calculated to be 2.6 nH and 1.1 nH, respectively. The inductance of the IPM unit was 17.3 nH provided and where the sample passes through the PCB board, 1pH / m which for a 1.5 mm-thick PCB, was 1.5 nH and, the total inductance of the IPM system was Ltotal = 17.3 nH + 2.6 nH + 1.1 nH + 1 .5 nH = 22.5 nH.

[0105] In many embodiments of the ENP system, the inductance and electric load are configured in series, as illustrated in the circuit diagram in Fig. 8. The instantaneous strength l(t) of the electric current corresponding to a pulse duration t is obtained by vZi / ! U - ^3A x where V is voltage, R is the total resistance of the shunt resistors and the sample, L is the total inductance of the ENP system, t is the pulse duration and e is the base of the natural logarithm. The calculated instantaneous strength of the electric current as a function of time up to 2000ns and the electric current flowing in the series circuit reaches maximum steady state value after a few 100ns.

[0106] In accordance with an exemplary embodiment the electric current nano pulses with durations of 200ns, 500ns and 1000ns for a the measurements show that the electric current reaches several thousands of Amps within a couple of 100ns. The cease of the pulse electric current was controlled by the solid-state switcher of the IPM unit. After the pulse was switched off, electric current drops rapidly to zero Amps. The current density can be largely increased with the reduction of the sample’s cross-sectional area.

[0107] In accordance with an exemplary ENP setup, the current density (Jn) in the sample can be calculated bywhere UtotaPs the charging voltage for the ENP system, RIPM^O.OQ Q. p, L, and r is the respective electrical resistivity, length, and radius of the sample.

[0108] Fig. 9 illustrates exemplary expected current density data as a function of the sample size for a Nichrome-80 sample at a charging voltage of Utotai = 600V using the exemplary ENP apparatus in accordance with many embodiments. The length of the sample for the calculations is set to 4 mm or 8 mm. Depending on the size of the Nichrome-80 sample, the current density is as high as ~1011A / m2, which is ~4-5 magnitudes higher than typical Spark Plasma Sintering (SPS) results. Higher currentdensity can further be achieved with a shorter sample length or a higher charging voltage for the ENP apparatus. Fig. 10 illustrates exemplary current density data as a function of the sample sizes for several metals and alloys at a charging voltage of 1000V. Greater current density can be achieved in materials that have higher electrical conductivity.Exemplary ENP Embodiment

[0109] In accordance with an exemplary embodiment, Nichrome-80, samples with a length of 8 mm and a diameter of 127 pm (i.e. , AWG 36) were ENPed with the exemplary apparatus previously discussed. The charging voltage for the IPM was 600V. The duration of each single electric pulse was set to 1000ns. The sample was run at various frequencies, including 10Hz and 100kHz. To determine the heating rates and the speed of the temperature rise in the sample and the GB modification during the ENP processing. The current density in the sample during processing was about 6.10 x lQ10A / m2at 600V charging voltage (as illustrated in Fig. 9). The sample wires melted completely and broke into two parts after only two electric current pulses were introduced into the samples. Fig. 11 illustrates an exemplary Scanning Electron Microscopy (SEM) image of a segment resulting from the melting and breaking after the ENP processing. The result aligns with the calculations that a single 1637ns pulse with a current density of 6.10 x l O1oA / m2heated the sample to about the melting point of Nichrome-80], A 2000ns pulse at the same current density would heat the sample to be above the melting point. Additionally, the heating caused by the first pulse did not dissipate completely during the first resting time (before the second pulse) in both frequency testing conditions. However, different heat loss can affect the GB behavior during the ENP process.

[0110] Fig. 12 illustrates SEM images of the polycrystalline grain structure within the sample before and after the ENP processing. There is no appreciable grain growth after the ENP processing at both 10Hz and 100 kHz nano pulsing conditions. The samples all show similar annealed morphology due to the rapid heating and cooling rates during the ENP processing. The average grain size is about 10 pm after the ENP processing. The invariant grain structure during the ENP processing suggests that any changes in the microstructure are likely to be localized, such as in the GB regions. In the thermal-onlytreated sample, the grain size is significantly larger than 20 pm. The grains grow significantly under high temperatures near the melting point. The thermal-only effective treatment time is several orders of magnitude longer than that of ENP processing.

[0111] SEM-EDX (Energy dispersive X-ray spectroscopy) was conducted to examine the sample composition and oxidation during the ENP processing. Fig. 13 illustrates the EDX results on the internal regions and surface areas in the sample and after the 100kHz ENP processing. The nearly same values of weight percentages (wt.%) of each of the constitutive elements within and on the surface of the raw sample show a homogeneous distribution of the chemical composition throughout the sample. No traceable oxygen was detected in the sample. However, considerable surface oxidation of the ENPed sample was detected whereas the internal chemical composition remained nearly unchanged compared to the raw sample. The oxides are likely to be Cr20a and SiO2 which are common oxidation products on the surface layer of Nichrome alloys. Both the 10Hz and 100 kHz ENPed samples had a similar chemical composition. Except for the very surface layer, no appreciable oxidation took place within the Nichrome-80 sample after the ENP treatment. This was likely the result of Nichrome-80’s high resistance to oxidation at elevated temperatures, and the extremely rapid ENP processing not providing enough time for oxygen to diffuse deep into the material. Consequently, any structural changes within the material during the ENP processing are not likely influenced by oxidation.

[0112] Figs 14A though 14C. Illustrate the characterization of the GB morphology in samples after the ENP processing at 10Hz and 100 kHz frequencies in accordance with an exemplary embodiment. Prior to TEM (Transmission Electron Microscopy) lamellas samples were prepared using the FIB (Focused Ion Beam). The segments near the melted region were vertically (longitudinal direction of the wire sample) embedded in resin, followed by the standard metallurgical sample preparation procedures. To preserve the original structure and morphology of the GBs, no chemical etchant was used for the samples as the etchant may cause chemical damage to the fine structure of GBs.

[0113] Step-like and serration GB morphologies are observed in the ENPed sample, which were not seen in the raw and thermal-only treated sample. Which correlates to the frequency of the current pulses during the ENP processing. Fig. 14A shows nanoscalestep-like GB morphology in the 10Hz ENPed sample. Fig. 14B shows nanoscale GB serration morphology in the 100 kHz ENPed sample. In the ENP system, a lower frequency pulsing provides a lower flux density of electrical energy (i.e., the specific rate of energy transfer) to the sample, and a higher frequency ENP processing supplies a higher flux density of electrical energy to the sample. The larger specific rate of energy transfer provides significant accumulative excess heat energy localized at GBs and facilitates the formation of nanoscale serrations at GB. The early stage of the formation of GB serrations at the atomistic scale is shown in Fig. 14C. However, the smaller flux density of electric energy does not provide sufficient excessive GB heat to change the overall GB morphology; instead, one “step” is formed at a GB (as illustrated in Fig. 14A). The dashed arrows indicate the direction of the external nano pulse current. However, the direction is only approximate. The TEM samples need to be tilted to a certain zone axis for the best imaging condition, and this tilt angle could be anywhere between 5° to 25°.Embodiments Implementing Skin Effect During ENP Processing

[0114] In many embodiments, high frequency pulses introduce skin effect to conductor samples. As a result, the electric current flows mainly at the “skin” of the conductor, between the outermost surface and a level called the skin depth. At high frequencies the skin depth becomes smaller. The skin depth is defined as the depth below the surface of the conductor at which the current density has fallen to 1 / e = 0.37 (e is the natural constant) of the surface current density. Skin depth (6) can be estimated by the following equation:where p is the electrical resistivity, f is the frequency of pulses,rand po are the relative permeability and permeability constant (=4TT X w7H / m), respectively. For the Nichrome- 80 alloy, p = 1 .07 x 10'6Q m andr= 1 (Nichrome-80 being a nonmagnetic alloy).

[0115] In an exemplary embodiment a pulse duration of 1000ns = 1 ps, which corresponds to f = 1 MHz was utilized and the estimated skin depth for the Nichrome- 80alloy is 523 pm, which is significantly larger than the radius of our sample: 63.5 pm for the AWG 36 (diameter of 127 pm) wire. Consequently, the current density in the sample is presumed to be relatively homogeneous over the cross section of the wire. Which means that all the grain interiors should possess similar levels of electric current density except that the electric current density at GBs varies depending on their relative orientations to the external electric field. More specifically, GBs that are more perpendicular to external electric field possess similar levels of electric current density as that of the grain interiors; however, GBs that are more parallel to external electric field possess distinctly different electric current density as compared to grain interiors. GBs that are more perpendicular to external electric field show greater excessive Joule heating as compared to GBs that are more parallel to external electric field. GBs that are more perpendicular or parallel to external electric field experience a large gradient of thermal stress or electron wind force stress across the GB regions and as a result, the morphology of GBs can be modified during the ENP processing.Embodiments Implementing Thermal Stress at GB Regions

[0116] In accordance with an exemplary embodiment during the ENP processing, as a result of short heating / cooling time, grain growth is hindered, as is seen from the SEM images (Fig. 12) and numerical modeling data (Fig. 4). There is significantly localized excessive GB heating resulting from the intense nano pulsing current. A current density of the nano pulse needed to be at least >1010A / m2to achieve sufficiently large temperature localization at GBs. The large temperature gradient across GB can create an elastic stress field which acts as a driving force leading to the GB changes from a high energy state to a lower energy state. According to the exemplary data at a nano pulse of 1637ns by the ENP processing, the temperature difference between GB core and the grain interior is about AT = 23.4 °C. This temperature difference generates approximately oT =51.5 MPa elastic stress at the GB region according to oT = aEAT, where a = 20 x 10-6 / K is the coefficient of thermal expansion, and E - 110 GPa is the elastic modulus of the Nichrome-80 superalloy at -1400 °C, which can be linearly extrapolated based on thelower temperature data or estimated using literature. The elastic stress is very localized in the vicinity of GBs (< a few pm thickness), and it lasts only several tens of ps.

[0117] FEA modeling (after the nano pulsing current is paused) showed that the localized heat or the excessive temperature at GB diffuses into its adjacent grain interiors in less than about 50ps. Fig. 15, illustrates the excessive temperature and thermal stress changes of GB. ~23.4 °C to ~51.5 MPa at the end of the ENP processing, to ~0.5 °C and -1.0 MPa at 50ps time lapse after the nano pulse is paused. Significant GB heat localization is maintained at 5ps after the pause of nano pulse; the GB heat localization becomes insignificant at 40ps after the pause of nano pulse. Consequently, the thermal stress at GB decays rapidly and becomes nearly the same as that in the grain interiors after several tens of ps once the nano pulse is OFF.

[0118] The mean distance, x, traveled by the diffusing atoms such as Ni or Cr in the Nichrome-80 alloy was calculated using the approximation diffusion equation: x2~ 2Dt, where D is the diffusion coefficients of the diffusing atom and t is the elapsed time since diffusion began. The self-diffusion or volume diffusion coefficients of Ni or Cr in the Ni- 20Cr alloy are on the order of ~10-8cm2 / s at about 1400 °C, and, the mean distance is calculated as x = 3.2 nm, x = 10.0 nm, and x = 31 .6 nm after 5ps, 50ps and 500ps (=0.5 ms) of diffusion time, respectively which are close to the length scales observed for the dimensions of GB facets or serrations. Diffusivities along GBs are usually several orders of magnitude higher than that of the volume diffusion. Consequently, the above calculated overall mean distances represent a lower limit of diffusion distances that Ni or Cr traveled near the GB region during the ENP processing. Due to the variations of excessiveness of the GB localized heating (caused by different inclination angles between the GB plane and the external current direction) and different GB diffusivities (due to different ordering of GB structures associated with crystal mismatches), it is possible to form GB facets or serrations at a variety of length scales, from atomistic scale to nanoscale, and to submicrometer scale, during the ENP processing, as shown in Fig. 14.

[0119] In contrast, in spark plasma sintering (SPS) processing, no direct evidence of GB heat localization was found in the SPSed samples. Typically, in SPS, the current density is <~107A / m2, and electrical stressing or SPS processing time is on the order ofseconds or minutes with each single pulse duration of a few ms and therefore, the heat localization at GB is negligible during SPS processing. Because of this, in SPS, the GB region does not possess sufficient energy to experience its own significant structure changes (faceting or serration) separated from the grain lattice structure.Embodiments Implementing Electron Eind Force Stress at GB regions

[0120] In accordance with an exemplary embodiment, because of the intense electric current density of ~6.10 x 1O1OA / m2used in the ENP processing for the Nichrome-80 samples, electrical wind force or electromigration effect can also influence the microstructural changes of GBs during processing.

[0121] The faceting / roughening transition of a grain boundary at a current density of can result in accelerated GB mobility caused by inelastic electron scattering at the GB. he electron wind force exerted on a metal ion core can be calculated by: ,.,.Where e, Z*, j, and p are the electron charge (e = -1.6 x 10-19C), effective valence, electric current density, and electrical resistivity of the material, respectively. The stress (Tew) acted on the metal lattice structure resulting from the electron wind force can be calculated by:Given the Z =-20 for Ni, j = 6.10 x l O1oA / m2during ENP processing, and p = 1.07 x 1 O’6Q m for the Nichrome-80 alloy, the Few=2.69 x W13N / atom is obtained.

[0122] The FCC crystal structure of the Nichrome-80 alloy has a lattice constant of 0.355 nm, the stress (Tew) level at the (100), (111 ), and (110) crystal planes (when these planes are perpendicular to the electron wind force direction) can be estimated to be ~3.4 MPa, ~3.9 MPa, ~2.4 MPa, respectively. This calculation is approximate because the value of the effective charge Z* used is for pure metal Ni, not for the Nichrome-80 alloy system. Doping / impurity atoms such as Cr in the Nichrome-80 alloy usually have a different effective charge value than the host atoms. However, during the ENP processingof the exemplary Nichrome-80 alloy sample, the electron wind force stress level is calculated to be a few MPa. The disparity in current density and the electron wind force or stress applied on the atoms between GB and grain interiors can be up to 5-6 times, which is within the same order of magnitude but less than the thermal stress level at GBs caused by localized heating. The chemical bonds within GB atoms are likely weaker compared to those in grain interiors, which implies that the electron wind force can have a pronounced effect on GB atoms, enhancing their diffusion. While the thermal stress on GBs remains paramount, the electron wind force can also play a consequential role for influencing GB microstructures. The electron wind force and stress values calculated are derived from an electric current density of approximately 6.10 x 1O10A / m2. Both electron wind and stress levels escalate linearly with increasing electric current density, and as shown in Fig. 9, attaining higher electric current densities (>1011A / m2) in Nichrome-80 alloy through ENP processing is feasible in accordance with the exemplary embodiment. In many such embodiments, the electron wind force and stress influencing GB atoms can be substantially more pronounced and result in more alterations in GB microstructure.

[0123] In accordance with an exemplary embodiment, the total pulsing time of the electric current nano pulses for ENP processing was only a couple of ps, and the change in the microstructure, including GB structures, can be localized from limited diffusion activities.Embodiments Implementing Atomic Structure Modification of GBs

[0124] In accordance with many embodiments, both the GB excessive heat and electron wind force can contribute to the modification of GB microstructures; however, in some embodiments, there are no observable morphological changes to all GBs in the ENPed samples. In many such embodiments this is the result of relatively different heating and electric current density amongst different GBs. In many embodiments, the structure and energy state of the GB itself is an important factor. In many embodiments with a polycrystalline microstructure, some GBs may possess higher energy states than others, and the high-energy GB can transit to lower energy states given appropriate activation. In some embodiments, when a lower energy state is activated, GB rougheningin the forms of facets, serrations, or steps can occur. In many embodiments, large angle GBs (>~10-15°) featuring pronounced crystal mismatches can exhibit elevated thermodynamic energy levels when contracted with low angle GBs (<~10°). In many embodiments, increased energy arises from the greater misorientation and crystal defects at the boundaries. The GB mismatch angle can be determined by examining the grain orientations of the GB’s adjacent grains. The local atomistic structures for the left (L) and right (R) grains can be determined by using the high-resolution TEM (HRTEM), as shown in Using the GB in Fig. 14, upon which the crystal structure and orientations of each grain can be identified and indexed. The left and right grains possess respective zone axes of (1-1-2) and (111 ), which are perpendicular to the HRTEM image plane, and the crystal mismatch at GB between the left and right grains can be measured. Fig. 16 illustrates the mismatch angle in accordance with an exemplary embodiment. The mismatch angle is about 37°, indicating a high energy state for the exemplary GB. In the thermal-only treated sample with only a dimensional gradient of temperature in it, the formation of steps or serrations at GBs was not observed. The main driving force for GB serration is the stored energy difference between adjacent grains, causing strain-induced boundary migration. Other mechanisms, such as GB sliding, can also cause the nucleation of serrated GBs. Traditional GB morphologies, such as those achieved during hot deformation processing, are typically smooth and straight. Therefore, no GB serration is expected to form in thermal-only treated (10s at 1400 °C) Nichrome-80 superalloy samples. In some such exemplary embodiments, the GB segregations or GB precipitations are imperceptible.

[0125] In many embodiments, the overall heating time for the thermal-only process is significantly longer as compared with that of the ENP process. In many such embodiments, grain growth becomes a favorable way of decreasing the overall free energy of the sample, which can be observed by the considerably larger grain size in the thermal-only treated sample illustrated in Fig. 12.

[0126] In many embodiments, serrated GBs in metals and alloys can result in significant advantages and material properties. In many embodiments, the serrated GBs enhance creep and fatigue resistance by hindering dislocation migration and crack propagation. In many embodiments, the morphology strengthens resistance tointergranular corrosion. In some embodiments, the morphology serves as a deterrent to atomic diffusion. In numerous embodiments, the morphology impacts phase transformations. In many embodiments, serrated GBs enhance the mechanical, thermal, and corrosion characteristics of metals and alloys.Quasi-lnstantaneous Materials Processing Technology Via High-Intensity Electrical Nano Pulsing

[0127] The electric field has been employed in many engineering fields to develop high-performance materials with expected functional and structural properties. However, the outcomes obtained with the electric field-assisted processing of materials still rely on long-term coupling with other electroless processes and the influence of Joule heating.

[0128] In accordance with various embodiments, ENP technology achieves quasi- instantaneous processing of material microstructures. In many embodiments, ENP processing can optimize the properties of materials. In many embodiments, ultra-high current density, ultra-short pulse duration, and ultra-high pulsing frequency occur simultaneously. Many such embodiments release massive electric energy into a target material. In many embodiments, the energy release occurs within micro or nanoseconds. In many embodiments, the ENP processing activates a non-equilibrium structural evolution of the material. In many embodiments, the non-equilibrium structural evolution is at the nanosecond temporal scale.Apparatus Configurations

[0129] Fig. 17A and Fig. 17B illustrate exemplary apparatus “bus-bar” and “through- hole” embodiments. The “bus-bar” configuration in Fig. 17A consists of a pair of block electrodes 1702 coupled with PCB 1706 with sample loading holes in them for the ENP processing of wire or foil components 1702. Component 1702 is attached to the electrode surface 1700 to form an electric connection. The component is secured by a surface 1704 to ensure good electric contact and coupling between the component 1702 and electrode 1700. In some embodiments, the surface 1704 comprises an insulating coating. In some such embodiments, the insulating coating is AI2O3.

[0130] A “through-hole” configuration in Fig. 17B comprises a pair of electrodes 1700 embedded on both sides of the PCB 1706 board connected through a hole for the ENP processing of components 1702. The component 1702 is inserted into a protection tube 1708. In many embodiments, the protection tube is insulating. In many such embodiments, the protection tube comprises AI2O3. The component 1702 is placed within a hole. In many embodiments, the component 1702 position is secured with holding elements 1710, and the ends of the component 1702 are connected to the electrodes 1700 on both sides of the PCB 1706 board with electrical coupling elements 1712.Exemplary Data

[0131] Fig. 18A through Fig. 181 illustrate the temperature and microstructure of nichrome alloy in accordance with exemplary embodiments. The distribution of the electric field and temperature within the cross-section of the treated material component determines the compatibility of ENP technology. In many embodiments, the rise of temperature at the alloy center exhibits stepped increment, as illustrated in Fig. 18A. The maximum wire temperature is about 1643 K, which is very close to the melting point of this nichrome alloy (1673 K). The corresponding temperature rate profile shows that ENP can bring nichrome alloy to an extraordinary heating rate of the magnitude of 108 K / s during the pulsing and energy release, and the cooling during the rest time is negligible, as illustrated in Fig. 18B. The temperature difference between the alloy component’s center and the edge increases over the pulse time, but can be maintained at a sufficiently low value (less than 12 K) as shown in Fig. 18C.

[0132] In many embodiments, ENP technology enables the localized modification of material structure at the micro-nano structure level without significant change of grain structure. Fig. 18D through Fig. 18F illustrate the EBSD inverse pole figures and the grain size distributions of nichrome alloy wires after conventional heating and after ENP processing. Fig.18D and Fig 18G illustrate the raw samples. Fig.18E and Fig 18H illustrate conventional heating samples, and Fig.18F and Fig 181 illustrate the ENP samples. For conventional heating, the nichrome alloy wires were loaded into a conventional furnace preheated to 1400 °C for 10s to achieve the highest possible heatingrate and the shortest possible duration time, as close as possible to the conditions in ENP processing. Annealing twins are commonly seen in this Ni-based superalloy with FCC structures subjected to thermal treatment. The comparison shows that the grain structure of the alloy matrix remains almost unchanged after ENP processing, while conventional heating results in a drastic modification. Many thermal processing technologies by heat conduction, radiation, or Joule heating require a duration time of at least dozens of seconds or longer. As illustrated in Figs. 18D, 18E, 18G, and 18H that the nichrome alloy grains undergo significant growth in size after only 10 s of conventional heating, with the mean grain size increasing by nearly three times the raw sample (from 4.8 to 13.4 pm), and a maximum size of 146.2 pm. In contrast, no obvious change in grain morphology occurs after the ENP processing, as illustrated by the similar distribution range and mean size illustrated in Fig 181). The raw nichrome wire has a pronounced grain orientation texture. The <001 > and <111 > orientations are preferred to align with (001 ) direction (axial direction of wire), while the radial direction has a random orientation distribution as illustrated in Fig. 18D. The orientation feature, at a smaller structural level, is modified after ENP processing as illustrated in Fig. 18F and has been significantly damaged after conventional heating. The grains at radius directions exhibit some preferred orientations, and the closest-packed plane {111} no longer tends to be parallel to the wire crosssection, as illustrated in Fig. 18E.

[0133] Fig. 19A through Fig. 19C illustrate the grain boundary distributions and orientations in nichrome alloys. Fig. 19A illustrates a raw sample; Fig. 19B illustrates a sample after conventional heating, and Fig. 19C illustrates a sample after ENP processing. The orientation difference at GBs also indicates that the high-angle GBs still hold the majority (> 90%) both in the raw sample (Fig. 19A) and after ENP processing (Fig. 19C ), whereas conventional heating (Fig. 19B) promotes the formation of more low angle GBs (> 10%).Quasi-lnstantaneous ENP Processing Defect Reconfiguration

[0134] In accordance with many embodiments, the massive energy released into nichrome alloy via ENP processing can bring non-equilibrium modifications to thematerial’s internal microstructure within a micro or nanosecond scale. In many embodiments, this modification enables the change to the grain insides (GIs) and hardening effects by producing dislocations. The intense electric pulsing promotes dislocation motion from electroplastic effects. In some embodiments, the wire length after ENP processing is increased due to its plastic deformation. In many embodiments, dramatic differences in dislocation morphology between raw and ENP processed samples can occur.

[0135] Fig 20A through Fig 20F illustrate dislocation morphology in accordance with an exemplary embodiment. The difference can be observed via microscopy, such as transmission electron microscopy (TEM). As shown in Fig. 20A, the raw nichrome alloy has a dislocation entanglement structure produced by plastic deformation. The preexisting dislocations could be generated by cold rolling of raw nichrome wire and reserved after its softening annealing. Conventional heating near the melting point promotes the dislocation mobility and the annihilation of opposite-signed dislocations, reducing the dislocation density significantly, as shown in Fig. 20B. ENP processing not only promotes the avalanche generation of dislocations but also alternates the dislocation configuration. In ENP processed nichrome wire with two pulses at a current density of 6.98 x 1 o10A / m2, a large number of jagged dislocations with heavily curved and randomly oriented topography can be observed by imaging at (011 ) zone axis as illustrated in Fig 20C, with selected area diffraction pattern as shown in Fig. 20D which are likely the results from of cross-slip of extended screw dislocations or the point pinning of dislocation lines during the ENP process. Similar dislocation reconfiguration is found in pulsing deformed Ti-AI alloy, which is a homogeneous wavy slip and the source of elastoplasticity. In many embodiments, ENP processed alloys can be hardened by suppressing planar slip and increasing the work hardening rate. Additionally, in ENP processed nichrome wire processed with eight pulses at a current density of 3.15 x 1 O10A / m2, a group of dislocation walls with orderly interwoven and periodically arranged topography can be observed by imaging at (011 ) zone axis as illustrated in Fig. 20E, and 20F.

[0136] The dislocation walls are the boundaries that divide the matrix into small celllike regions (cell blocks) with relatively low density. In many embodiments, this is causedby the multiplication and rearrangement of entangled dislocations with respect to the motion direction of drifting electrons in the ENP process. A similar parallel distribution of dislocations is found in quenched steel after electropulsing, which is described as a “seaweed” structure formed by electron wind force and can result in an electroplastic effect by the combined action of Joule heating and reduced residual stress. It is commonly assumed that electropulsing promotes the dislocation motion by electron wind force; however, the results indicate that different ENP parameters, including current intensity, pulsing duration, and pulsing rising rate, can produce a mechanism of dislocations- pulsing interaction, that contribute to the exhibition of various dislocation configurations. In many embodiments, ENP processing can induce specific electro-microstructure interaction by precisely controlling the input electric energy in a material system. In many such embodiments, ENP processing can provide a tool for exploring the mechanisms of electric field effects.

[0137] Fig. 21 illustrates tensile testing data in accordance with an exemplary embodiment. Fig. 22A through 22C illustrate fracture surfaces in accordance with exemplary embodiments. Tensile testing for and of nichrome alloy wires was conducted to analyze the mechanical impact of ENP processing. The experimental results show that a slight softening phenomenon with a small decrease of ultimate fracture strength is observed in this alloy after ENP processing, despite the dislocation reconfiguration produced by electropulsing. This can result from the temperature rise by Joule heating, while the heavy dislocation generation or reconfiguration locally happens in specific grains. In many embodiments, the softening effect on materials can be avoided in ENP processing by suppressing Joule heating. The yield stress did not show a significant change, which is likely due to the unaffected grain size after the ENP processing. Conventional thermal treatment of the nichrome alloy results in all-around property degradation as a result of dislocation annihilation, abnormal grain growth, and loose oxide mixture on the surface. The elastic modulus of the nichrome alloy after ENP processing changed from 315 to 140 GPa, which indicates a transformation of the anisotropic texture in a raw sample produced by cold rolling to a relaxed isotropic structure. ENP processing resulted in an almost instant release of the stored mechanical energy from previous colddeformation in the nichrome alloy wire without significantly changing its grain structure. The localization of electrical pulsing impacts predominantly at the dislocation level. In many embodiments, releasing a few high-intensity electric pulses results in significant non-thermal effects. Homogenization and uniformization of material properties in Ni- based superalloys during electropulsing treatment can occur. However, ENP processing occurs at a micro-nano seconds duration and, in many embodiments, can result in without detectable recrystallization.Quasi-lnstantaneous ENP Processing Surface Nanocoatinq

[0138] In accordance with many embodiments, the rapid heating speed of high- intensity ENP processing in an air atmosphere can activate an altered oxidation behavior of the nichrome alloy. In many embodiments, ENP processing can form a unique oxidation coating on a material surface. Fig. 23 illustrates the evolution of exterior surface morphology in accordance with an exemplary embodiment. Generally, in the oxidation of nichrome alloys, loose NiO and Cr2O3 are preferentially formed and finally transformed into NiCr2O4 spinel in the case of prolonged oxidation. Fig. 24A through Fig. 24I illustrate the morphology of nano coating after ENP processing in accordance with exemplary embodiments. In many embodiments, ENP processing results in a dense chromium oxide nanocoating. In many such embodiments, the chromium oxide has a triple hierarchical structure on an alloy exterior surface. In some embodiments, the nanocoating can be obtained after as few as eight consecutive pulses are released. In many embodiments, the rough coating surface displays a fold-like morphology at the micro-scale, as illustrated in Fig. 24A and 24B. Some embodiments comprise sub-nanoscale ravines as illustrated in Fig. 24C. Higher resolution further shows that in many embodiments, the surface coating is composed of the stacking of massive well-faceted nanocrystals (< 200 nm) as illustrated in Fig. 24D. In many embodiments, the dense hierarchical micro-nano structure on possesses high hardness. In many such embodiments, the instantaneous surface localized modification is configured for a specific surface area.

[0139] Many embodiments may find many applications in wear-resistant materials. Some embodiments may find application in gas sensors. In accordance with manyembodiments, ENP coating can be conducted in various gas environments. In accordance with many embodiments, different conductive alloys can be utilized as a substrate. In accordance with many embodiments, ENP coating can fabricate surface ceramic layers. In accordance with many embodiments, ENP coating can fabricate flexible films.

[0140] Fig. 25 and Fig. 26 illustrate the composition by EDS measurement of an exemplary embodiment. The EDS data shows that the surface oxidation coating mainly contains chromium oxide, and almost no formation of nickel-containing oxides is involved during ENP processing. The non-equilibrium structure induced by the ENP processing is unobtainable with a conventional thermal method. Fig. 27 A through Fig. 27D illustrate an example of a conventional heating treatment. With the magnification increasing from Fig. 27A to Fig. 27D. The sample was heated at 1400 °C for ten seconds, but only resulted in a roughened and loose surface layer; the hardness decreased, and tribological property degradation; the spheroidization and coarsening of crystals will also damage their electrical and gas-sensitive properties.

[0141] Fig. 28A through Fig. 28C illustrate morphology in accordance with exemplary embodiments. A raw sample is illustrated in Fig 28A. A conventionally heated sample is illustrated in Fig 28B, and an ENP sample is illustrated in Fig. 28C. The samples were mounted, and the nichrome alloys were etched by 10%HCI / 3%CuSO4 ethanol solution to enhance the visualization of nanocoating. Long-time thermal heating, even at lower temperatures, can generate surface mixed oxides and the matrix intergranular oxides.

[0142] However, as shown in Fig. 24C, a dense double-layer composite coating with a thickness of about half a micrometer on the nichrome alloy surface can be achieved by ENP processing. The formation of ultra-thin Cr2O3 coatings (~ 200 nm) likely results because there is kinetically insufficient time for crystal growth. ENP processing produces not only a denser Cr2O3 coating but also one extra ultra-thin amorphous Si-Cr-0 nanocoating beneath it as illustrated in Fig. 24H. The double-layer ultra-thin coating covers the nichrome alloy without excessive detrimental oxygen intergranular diffusion into the alloy matrix, as illustrated in Fig. 24E and Fig. 24I. High-intensity ENP processing not only promotes the formation of Si / Cr-contained oxidation coating by enhancingdiffusion but also enlarges the diffusivities of Cr and Si in varying degrees. The separated double-layered nanocoating can be obtained on localized alloy surfaces in microseconds by ENP processing; in contrast a long critical time is required by conventional thermal methods.

[0143] In many embodiments, ENP processing can induce a rapid development of protective coatings. In many embodiments, electropulsing produces a thin protective Cr2O3 layer that is strongly adhered to the alloy matrix. The folded structure of the coating layer indicates that despite the high temperature reached during ENP processing, grain boundary diffusion still is the main contribution to mass transport in many embodiments. In many embodiments, ultra-thin double-layered coating with a well-bonded and defect- free interface with alloy matrix results in anti-corrosion protection. Many embodiments result in an alloy with better room-temperature gas sensitivity and mechanical properties.High -Density Dislocation Formation

[0144] Achieving high dislocation density across the entire body of metallic materials often poses significant processing challenges. In accordance with many embodiments ENP processing can result in a tailorable high density of dislocation defects. In many embodiments, the density of dislocation is configured to optimize material performance for specific applications. In some such embodiments, metallic materials with high dislocation densities are configured for enhanced strength. In many such embodiments, the materials are optimized for hardness. In many embodiments, the material is configured to optimize for interaction between dislocations, which impedes their movement. In accordance with numerous embodiments, high dislocation density I configured to influence various physical and functional material properties. Some embodiments are configured for the presence of dislocations. Some such embodiments are further configured to scatter electrons and phonons. Many embodiments are configured for changes in the electrical and thermal transport properties of metals, alloys, and metallic-based materials. In many embodiments, dislocation defects are configured to influence low-frequency and long-wavelength phonons, which affects the thermal conductivity of materials, including metals and non-metals. In many environments thatcontrol the density and distribution of dislocations, it is possible to tune the mechanical, electrical, and thermal properties of materials. In accordance with numerous embodiment variations in dislocation density, it enables the configuration and tailoring of material properties for specific applications.

[0145] Traditionally, the most widely employed approaches for achieving high dislocation density in metallic materials have been through external deformation processes such as rolling, forging, or drawing, which induce dislocations by introducing large amounts of mechanical strain into the material. Generally, heavily cold-worked metals and alloys can achieve a dislocation density on the order of 1015 / m2. Some processes are carried out at high temperatures for significant amounts of time on the order of minutes or even for several hours, which can lead to unwanted effects such as grain growth and can alter the microstructure of the material, limiting the control over the properties of the material. Therefore, material processing methods are needed to induce high dislocation density in metallic materials.Exemplary Embodiments

[0146] In many embodiments, the frequency of the electric current pulses dictates the power of electric energy infused into the alloy. In many such embodiments, the electric current pulses influence the resultant dislocation configuration. In some embodiments, dislocation defects created through nano pulsing current processing align with the direction of the current flow. In some embodiments, ENP processing can achieve a dislocation density of 1015 / m2for microscale dislocations. Some such embodiments achieve a dislocation density comparable to heavily cold-worked metals and alloys.

[0147] In many embodiments dislocation density of 1017-1018 / m2can be achieved for nanometer or atomic scale dislocations. In some embodiments, transient stress localization at GBs during processing is coupled with a high concentration of point defects. Some such embodiments stimulate high-density dislocations. Many embodiments are configured for the formation of scale screw dislocations in the nichrome alloy.Exemplary data

[0148] Fig. 29A through Fig. 29C illustrate the internal microstructures of samples in accordance with some embodiments. Fig. 29A illustrates a raw sample of nichrome alloy wires treated by soft annealing processing at 1070°C for 3 hours, followed by air cooling, which contains a low density of dislocations. Fig. 29C and Fig. 29D show samples treated by the intense nano pulsing current at 10Hz and 100kHz, respectively, with high densities of sub-micrometer and micrometer-scale dislocations. The dislocation line defects in the 10Hz-processed sample (Fig. 29C) are relatively less dense than those in the 100kHz- processed sample (Fig. 29D). This is likely the result that at a higher frequency nano pulsing current, more electrical energy for a given amount of time is injected into the sample which results in more lattice distortions and generates a denser dislocation distribution. In contrast, Fig. 29B illustrates a thermal-only treatment sample, with a very low dislocation density, illustrating that intense nano pulsing current has a different effect on the dislocation behaviors of the sample compared with traditional thermal treatment.

[0149] Fig. 30A through Fig. 30D illustrate dislocation configuration after the nano pulsing current process at 100kHz in accordance with exemplary embodiments. Fig 30A and Fig. 30B show enlargements of regions 1 and 2 of Fig. 29D. The majority of the generated dislocations (indicated by the dashed white lines) have orientation relationships with the direction of the nano pulsing current ( / , indicated by the hollow arrows). In Region 1 (Fig. 30A), the dislocation lines are approximately oriented either parallel or perpendicular to the current flow direction and in Region 2 (Fig. 30B), most of the dislocation lines are inclined by approximately positive 50°-60° relative to the current direction, and a small amount of dislocations are inclined by about negative 30°-40° with regard to the current direction.

[0150] The densities of the sub-micrometer and micrometer scale dislocations, pdisi, of the regions in Figs. 30A and 30A (which are dark-field TEM images) were quantified based on stereological relationship methods. A TEM thin foil with thickness t is tilted to the optimum angle for the maximum observed individual dislocations. Then, a standard square grid is imposed onto the TEM image, and the number, N, of intercepted points between the grid and the dislocation lines, with mean length,is counted. The value of p<nsi = 2N!(Lt) can then be obtained.

[0151] For the exemplary samples, a 25 x 25 grid was superimposed on the TEM images of Figs. 30A and 30B. The counted intercepted points A / is equal to 481 and 636 for Figs. 30A and 30B, respectively. The TEM images have the size L of 1101nm x 1101 nm, and the FIB-prepared TEM lamella has a thickness of about t = 100nm. Therefore, the pdisi is 8.74x1015 / m2and 11.55xl 015 / m2in the TEM images in Figs. 30A and 30B, respectively. The calculated values are close to the limit level of dislocation density that can be obtained in heavily cold-worked metals and alloys, which is 1014- 1015 / m2.

[0152] The above-obtained values only quantify the visible micrometer and submicrometer scale dislocations of the local regions based on the bright or dark field TEM images. The average dislocation density in the entire sample can easily vary by one order of magnitude. However, the distribution density of these micrometer and sub-micrometer scale dislocations should be relatively uniform across the exemplary samples that are treated by the intense nano pulse current because the current density passing through the sample is relatively homogeneous as the calculated skin depth under the processing conditions is estimated to be 523pm, which is significantly greater than the radius of the wire sample: 63.5pm. The skin depth (5) was estimated by 3 = / y(7rf r o), where =1.07xW6Q m is the electrical resistivity, f = 1 MHz (at the pulse duration of 1 ps) is the frequency of pulses, pr=1 (nonmagnetic) and po=1.2568xlO’6H / m are the relative permeability and permeability constant, respectively, for the nichrome superalloy.

[0153] The uniform dislocation distribution in Fig. 30 deviates from the typical formation of cellular dislocation structures at high densities, which requires sufficient dislocation movement. This is likely the result of the extremely fast ENP process, where rapid energy deposition and quenching may not allow sufficient time for dislocation rearrangement into lower-energy configurations, effectively “freezing” them in a less relaxed state. Effectively, ENP kinetically traps dislocations in a high-energy state, preventing relaxation.

[0154] The orientation relationship between dislocations and current direction is highly correlated with the crystal orientation in region 1 and region 2. Fig. 30C shows the HRTEM image of region B in Fig. 29D, in which a GB is presented. In Fig. 30C, the crystalstructures of the upper (grain 1 ) and the lower (grain 2) grains are identified and highlighted in the insets. The crystal structures in the insets are the inverse FFT images (IFFT) of the filtered FFT (Fast Fourier Transformation) images for regions (highlighted by the white squares with rounded corners) in the HRTEM image, according to the standard HRTEM image post-processing methods. The dislocations line defects are formed according to the {111} slips systems of the FCC crystal structure (Fig. 30D) of the nichrome solid solution alloy. The Burgers vector of the dislocations follows the <110> slip directions of the crystal structure, giving rise to the typical FCC dislocation’s Burger’s vector of

[0110] ,

[0155] In many embodiments, the intense nano pulsing current induces atomic scale localized lattice shear, which results in the formation of extremely high-density atomistic length scale dislocation line defects. Fig. 31 A through 31 D illustrate HRTEM images of exemplary nichrome alloy samples. Fig. 31 A shows a raw sample. Fig 31 B shows a thermal-only sample. Fig. 31 C shows a sample treated by intense nano pulsing current at 10Hz. Fig 31 D shows a sample treated by intense nano pulsing current at 100kHz. Fig. 31 C and Fig. 31 D show the typical atomic structures of the dislocations in samples treated by intense nano pulsing current at the frequency of 10Hz and 100kHz, respectively. The dislocation takes the form of screw dislocations. The white-dotted lines indicate the dislocation lines.

[0156] The screw dislocation densities were quantified as described above for the exemplary samples based on their HRTEM images (Figs. 31 C and 31 D). Dislocation densities on the order of 1017-1018 / m2were calculated. The calculated screw dislocation density was 8.3*1017 / m2for the sample treated by the 10Hz intense nano pulse current, which is two to three orders of magnitude higher than the dislocation density of ~1014- 1015 / m2for heavily cold-worked metals and alloys.

[0157] Fig. 32A and 32B illustrate edge dislocation in accordance with exemplary embodiments. Edge dislocations were present in the exemplary samples treated by intense nano pulsing current for both frequency conditions. Similar high-level dislocation density is typically only observed in metals that were treated by long-time intense pulsing current processing or existed in a confined nanoscale volume.

[0158] No significant oxidation was observed within the central part of the exemplary wire samples. Oxidation was confined to the very top surface layer, less than a couple of micrometers in depth, as a result of the nano pulsing current process. The same samples and, subsequently, the same EDX data, as discussed with regard to the localized engineering of the grain boundary section above, were utilized in the analysis and experimentation for high-density dislocation formation. As discussed above, Fig. 13 shows the EDX quantifications for the internal and surface areas of the exemplary samples and compares the conditions before and after the nano pulsing current treatment. The weight percentages (wt.%) of each constituent element in the untreated or raw sample, as well as its surface, were almost identical, indicating uniform chemical composition throughout. No oxygen was detected in the raw sample. In contrast, significant surface oxidation was evident post-nano pulsing, with the internal chemical composition remaining largely similar to that of the raw sample. The oxides formed were likely Cr20s and SiO2, which are the commonly identified oxidation products on nichrome alloy surfaces. The minimal internal oxidation can be attributed to nichrome’s well-known high oxidation resistance at elevated temperatures and the extremely rapid nature of the nano pulsing process, which likely did not allow enough time for oxygen to deeply penetrate the material and suggest that potential structural alterations within the interior of the material due to nano pulsing are not affected by oxidation processes.

[0159] The atomic structure and formation mechanism of the screw dislocation defects results from the intense nano pulsing current processing. Fig. 33A and Fig. 33B illustrate the atomic structure of the lattice shear and screw dislocations in accordance with an exemplary embodiment. Fig 33A shows a HRTEM image of an exemplary sample and 33B depicts a schematic of the extended dislocation. The white dashed arrow indicates the approximate orientation of the electric current relative to the lattice plane. The actual external current direction is not the same as indicated but may deviate by a few degrees due to necessary sample rotation in the HRTEM imaging operation. An extended dislocation is formed of a pair of partial dislocations accompanied by stacking faults. The length of this type of screw dislocation falls at the nanoscale, ranging from a few nanometers up to tens of nanometers. The average dispersion (or mean free path)between these screw dislocation lines is also a few nanometers, resulting in an extremely high dislocation density.

[0160] In many embodiments the, the electrical resistivity at GBs is 102times higher than in grain interiors. In many such embodiments, the GB atoms experience a stress level two orders of magnitude higher during nano pulsing processing. In many embodiments, nano pulsing processing results in a stress level in the range of ~400-600 MPa, with many such embodiments lasting for 1000 ns duration of the nano pulsing processing.

[0161] The critical resolved shear stress (CRSS) of the nichrome alloy at various temperatures was calculated. For isotropic polycrystalline materials with an FCC structure, the CRSS=aymave, where oyand / 77ave=1 / 3.06 are the macroscopically observed yield stress and the average Schmid factor, respectively. The yield stress of the nichrome alloy, determined by the commonly used strain E=0.2% offset method, is ~345MPa at room temperature and results in a CRSS of 112.75 MPa at room temperature.

[0162] Applying the 0.2% offset rule at elevated temperatures (up to near the melting point), the yield stress at a given temperature T can be calculated by ay(T)=£E(T)=0.2%E(7^. With the elastic modulus E = 110GPa at 1400°C for the nichrome alloy results in oy(1400°C)=220MPa. Consequently, the CRSS at 1400°C is calculated to be ~71.90MPa for the nichrome alloy, which is an order of magnitude greater than the electron wind force stress experienced by grain interiors, yet it is several times smaller than the electron wind force stress for GBs.

[0163] In many embodiments, the intense and transient (1000 ns) stresses localized at GBs exert a significant impact on the grain interior lattices. In many such embodiments, the impact surpasses the CRSS. In some such embodiments, the stresses induce lattice shearing within the grain interiors.

[0164] Fig. 34A through Fig. 34D illustrates a finite element modeling of the microstructure of an exemplary embodiment. Fig. 34A shows the initial microcrystalline structure; and Fig. 35B through Fig. 35D show the distribution of von Mises stress at varied instances during the nano pulsing processing. The model is a 2D polycrystalline microstructure model with an average grain size of about 25 pm, analogous to themicrocrystalline structure of the exemplary nichrome alloy samples. The model, measuring 2000 pm in length and 500 pm in height, features both ends fixed to simulate the connection to electrodes in the nano pulsing experiments, while the top and bottom boundaries are free. The material parameters used for the linear elastic modeling of nichrome-80 alloy were an elastic modulus of 200 GPa and a Poisson’s ratio of 0.3. The finite element solid mechanics modeling was conducted using the COMSOL Multiphysics platform.

[0165] In the model, a stress of 500 MPa was applied on all internal grain boundaries, causing the grain interiors to experience similar stress levels. Using finite element modeling, the transient process of stress transfer within the microstructure was simulated. The stress levels within many grain interiors swiftly escalated to over 500 MPa within roughly 200 ns, as depicted in Fig. 34B and 35C. A dynamic equilibrium state of stress distribution was then achieved and maintained throughout the rest of the nano pulsing duration, as shown in Fig. 34C and Fig. 34D. The model assumes a 500 MPa electron wind force stress across all GBs. The assumed ~500 MPa stress is an approximation of the maximum stress that GBs might encounter during the nano pulsing processing. The stress level in most grain interiors was able to exceed the GB stress level after several hundred nanoseconds, and the stress intensities likely surpassed the CRSS, which indicates lattice shearing within grain interiors during nano pulsing processing.

[0166] The sample experienced a very high temperature throughout the processing, increasing the possibility of the concentration of vacancy defects in the material. However, the dislocation densities from rapid thermal stressing (expansion or contraction) are typically in the range of 1012to 1015nr2, which is significantly lower than the range observed. The discrepancy suggests additional dynamic processes during the ultraintense nano pulse electric current treatment amplify the generation of dislocations beyond what is typically induced by thermal stressing. While rapid thermal expansion / contraction contributes to dislocation generations, the dislocation densities that result from ENP processing are the result of a complex interplay of thermal, mechanical, and electrical effects facilitated by the ultra-intense nano pulse electric current.

[0167] In accordance with many embodiments, limited recrystallization occurs as a result of the short time duration of the nano pulsing current processing, despite the temperature being the melting point as the peak temperature. An infrared camera was used to record and determine the lower limit of the heating rate and temperature that was reached in the sample by the intense nano pulsing current. Within one frame of recording (30Hz, with 33.3ms per sampling), the wire sample went from heating up to melting temperature (~1400°C for the nichrome alloys), and then broke into two parts and then cooled to room temperature. The recording indicates a heating rate greater than 104K / S, but that results in little change in the grain size and the overall microstructure.Electromiqration-induced Quasi-instantaneous Precipitation Behavior

[0168] Thermally activated precipitation imparts a strengthening effect to many alloys through solution-aging treatment. A major challenge for precipitation-hardening superalloys is their microstructural instability, which results in the formation of detrimental phases and a loss of strength during prolonged thermal exposure. In accordance with many embodiments, the combination of microalloying and high-intensity electrical nano pulsing (ENP) can suppress thermal degradation. Many embodiments promote ultra-rapid field-activated NisCe precipitation in Ni-based superalloy with cerium micro-addition. In accordance with many embodiments, a single electric pulse at ultra-high intensity (1.30xl011A / m2) and ultra-short duration (330ns) is applied. In many such embodiments, the electric pulse results in improved electrochemical corrosion resistance (corrosion potential ~488.6mV) and work-hardening performance (ultimate tensile strength ~831.1MPa) in the alloy.

[0169] Precipitation-hardened superalloys are widely utilized in high-temperature environments, such as turbine blades and power engines. The thermally activated precipitation of secondary phases enhances their mechanical properties and oxidation resistance. However, improper heat treatment can lead to the coarsening of these precipitates, resulting in microstructural degradation and a decline in properties.

[0170] Microalloying, which involves the addition of small amounts of alloying elements, can facilitate grain size reduction and precipitation hardening in metallicmaterials. Rare earth (RE) micro-additions have demonstrated significant potential for grain refinement, microstructural optimization, and property enhancement in various material systems, such as texture modification in extruded Mg-Mn alloys, improved fatigue life in low-oxygen steels, reduced corrosion rates in binary Mg-RE alloys, and increased coercivity in (Nd, Dy / Tb)2Fei4B permanent magnets.

[0171] Cerium (Ce), a surface-active element, is commonly added to metallic materials. In nickel-based superalloys, its inclusion can improve oxidation resistance. Strengthening effects have been observed in nickel-based alloys containing refined CeO2 oxide dispersoids. Additionally, minor additions of cerium have been associated with enhanced tribological properties.

[0172] A challenge with rare earth microalloying additions is unintended microstructural changes, which can occur due to their high reactivity and low diffusivity. In accordance with many embodiments, ENP processing applies an instantaneous current with periodically varying amplitudes and modifies the microstructure and properties of materials by utilizing both Joule heating and unique electric field effects. In many embodiments the ENP process enhances electroplasticity in alloys by promoting dislocation reconfiguration and improves the mechanical properties. In some such embodiments, ENP processing stabilizes the alloy’s bi-lamellar microstructure, such as for Ti-6AI-4V alloy. In some environments, ENP processing can redistribute lead-rich phases, such as in Cu-Zn alloys.

[0173] In accordance with many embodiments, ENP processing can enable cerium micro-additions in superalloys without compromising the integrity of the alloy matrix. In accordance with many embodiments, releasing a high-intensity single-impulse current within nanoseconds enables electromigration and thermal accumulation. In accordance with many embodiments ENP processing enables targeted modifications of microalloying additions in precipitation-hardening superalloys.

[0174] In many embodiments, ENP processing is combined with microalloying regulation for the processing of Ni-based superalloys. In many embodiments a single electric pulse, with an ultra-high intensity such as ~1011A / m2and an ultra-short durationsuch as 330ns, triggers non-equilibrium, field-activated precipitation of cerium atoms within nanoseconds.Exemplary Data

[0175] Fig. 35A through 35C illustrate ENP-activated precipitate in Ni-based superalloy in accordance with an exemplary embodiment. Fig. 36A shows the measured current density profile of single electric pulsing in this ENP processing of Ni-based superalloy. The inset is the schematic diagram of the sample setup for the ENP processing (Fige 17A). Electrodes are embedded on PCB board. Ni-based superalloy wire is attached to the electrodes, with an AI2O3 plate pressing on it. The black arrows represent the flow path of current pulsing; Fig. 35B shows the temperature distribution on the side cross-section of the sample-electrode setup at 330ns. The enlarged view is the contact area of the sample and electrode; Fig 35C shows the temperature at the center of the alloy wire as a function of time. The inset shows the temperature distribution on the normal cross-section of alloy wire at 310ns and 330ns, respectively.

[0176] In accordance with an exemplary embodiment, nichrome superalloy wire (softening annealed; diameter 78.7pm; FCC lattice with a=0.3539 nm) with cerium microaddition was used. The Ni-based superalloy reached its melting point (1400°C) within just 330ns through ENP processing at an electric current density of 1.30xl011A / m2. The ENP processing resulted in a heating rate of about 4.17X109K / S, and the electric pulse had a rising rate of 1 ,86xl O18A / (m2s).

[0177] Fig. 36A through Fig. 36H illustrate the microstructure of field-activated precipitates in Ni-based superalloy in accordance with an exemplary embodiment. The raw Ni-based superalloy has FCC structure (a=3.54A; Fm3m) with no visible precipitates (Fig. 36A). A uniform element distribution is shown across both GBs and GIs. After ENP processing, cerium-rich precipitates appear (Fig. 36B), forming within an ultra-short time scale (330 ns). In many embodiments, the precipitates preferentially form at defect sites, such as specific GBs and dislocations, where higher electron scattering increases electrical resistance and leads to more Joule heating during interaction with the electric current. In some embodiments, ENP processing results in intergranular precipitategrowth. Fig. 36C shows TEM-EDS elemental mappings of an exemplary embodiment showing significant cerium aggregation in these precipitates, with depletion of Ni and Cr. In some embodiments, the cerium content can reach about 14wt.% at GBs or GIs and up to 22wt.% at GB junctions, which is higher than in the raw alloy. The ENP processing activates the electromigration of cerium atoms into defect regions within 330ns. The partially mismatched structure with low atomic packing density at GBs or dislocation regions is an ideal site for large cerium atoms, as it helps reduce the free energy in the Ni-based alloy. The cerium atoms are stabilized without being pushed further out of these defect areas. The precipitate size remains small (~40nm) at GB triple junctions, while regions like GBs and GIs promote precipitate growth through localized melting by Joule heating, resulting in its larger sizes.

[0178] The precipitate was determined to be NisCe intermetallic with a hexagonal lattice structure (a=b=4.87A, c / a=0.82; a=p=90°, y=120°; P6 / mm) (Fig. 36D and Fig 36E). The addition of Ce in Ni-based superalloys can also enhance oxidation resistance and strengthen the alloys. However, the formation of coarsened NisCe along GBs is often considered detrimental, as it reduces the high-temperature ductility of the alloy. The ENP processing transforms the GB structure into a semi-coherent boundary between the hexagonal NisCe lattice and the cubic nichrome lattice, as shown in Fig. 36G. The electron beam was aligned with the (011 )nichrome zone axis of the matrix grain and a set of crystal planes corresponding to the (2110)Nisce plane of NisCe grain with an interplanar spacing of 2.84A. The precipitation of NisCe intermetallic due to ENP processing results in lattice distortion of the {111 }nichrome plane, the closest-packed plane of the nichrome alloy matrix, along the (211 nichrome direction. This is likely from the release of stress concentration produced by the larger lattice mismatch. The atom array is not completely disordered at this semi-coherent boundary. Two lattices maintained a specific orientation relationship. As shown in the inset in Fig. 36G, the

[0110] Nisce direction of the hexagonal Ni5ce lattice is at an angle of about 142.5° to the (011 )nichrome direction of the cubic nichrome lattice. Fig. 37H shows a schematic diagram of the formation of interlaced alignment between (0001)Ni5Ce plane of hexagonal NisCe lattice and {011}nichrome plane of cubic nichrome lattice. The semi-coherent boundary exhibits higher interfacial energy due to elastic latticedistortion (increasing strain energy) and a complex atomic bonding state (increasing chemical energy). The rapid transition in GB regions induced by ENP processing leads to the formation of the high-energy GB structure without sufficient kinetic stabilization. In accordance with many embodiments, the GB regions in Ni-based superalloy can convert electrical energy from ENP processing into interfacial energy by forming semi-coherent NisCe intermetallic with a spatially interlaced mismatch.

[0179] Fig. 37A through Fig. 37D illustrate properties of Ni-based superalloy field- activated precipitates in Ni-based superalloy in accordance with an exemplary embodiment. Fig. 37A and Fig. 37B show Electrochemical performance, and Fig. 37C and Fig. 37D show the mechanical performance of the alloy after ENP processing at 1 ,30x1011A / m2and 330ns. Fig. 37A shows potentiodynamic polarization curves. Fig. 37B shows electrochemical impedance spectroscopy (Nyquist plot) with an equivalent circuit for data fitting. Fig. 37C shows engineering tensile curves with a schematic diagram showing the microstructure. Fig. 37D shows strain hardening rate curves calculated from the true stress-strain curves.

[0180] The potentiodynamic polarization curves of this Ni-based superalloy after ENP treatment in 3.5wt.%NaCI solution are shown in Fig. 37A. The raw alloy exhibits typical polarization behavior for passive alloys, and the curve for the treated alloy shifts to a much higher electric potential. Both alloys display the same corrosion current density (approximately 3.86x10'8A / mm2), indicating a similar corrosion rate in the electrolyte. There is a significant increase in corrosion potential, from -935.4mV in the raw alloy to 488.6mV in the alloy after ENP processing, indicating a lower tendency for metal ion formation and corrosion processes in the alloy after ENP processing.

[0181] Electrochemical impedance spectroscopy (EIS) shown in Fig. 37B further compares the corrosion performance of the specimens. The larger capacitance half-arc diameter and impedance modulus indicate better electrochemical corrosion resistance in this Ni-based alloy after ENP processing. The inset shows the high-frequency range and displays two segments in its EIS curve. The equivalent electrical circuit (EEC) for fitting the EIS data is used to fit the EIS curves of passive alloys. It includes the electrolyte resistance (Rs), the charge transfer resistance (Ret), the passive layer resistance (R1 ),and the corresponding constant phase elements (CPEct and CPE1 ). The much higher R1 value compared to the Ret value in the Ni-based superalloy suggests that its corrosion resistance primarily depends on the passive layer. Specifically, the lower Ret value in the alloy after ENP treatment indicates less resistance for charge transfer in the surface passive layer. The R1 value increases by orders of magnitude, indicating that ENP processing significantly suppresses mass transfer.

[0182] Tensile stress-strain curves for the Ni-based superalloy before and after the ENP processing at room temperature are shown in Fig. 37C. The high-intensity ENP processing, with a single pulse at 1.30*1011A / m2 and 330ns, alter the deformation behavior of the alloy. There is a hardening effect in the plastic region and a softening effect in the elastic region after ENP Processing. Compared to raw alloy, the maximum plasticity of the ENP processed alloy decreases by 8.6%, the yield strength decreases by 148.2MPa, and the ultimate tensile strength increases by 121.8MPa. The presence of discrete and fine Ni5Ce precipitates with semi-coherent interfaces at GBs resulted in enhanced room-temperature hardening and strength. The small amount of rare earth doping, localized along the GBs, strengthens the GBs against mechanical creep by altering the local bonding environment and retarding GB sliding. Precipitates with discrete morphologies are beneficial for the ductility of Ni-based alloys compared to cellular morphologies. There is also a significant decrease in the elastic modulus from 85.96GPa in the raw alloy to 43.75GPa in the treated alloy. During the electrical pulsing process, conductive materials undergo lattice expansion and stress relaxation, resulting in a decrease in the elastic modulus, and the decreased elastic modulus can persist after the ENP processing. Fig. 36G shows the interplanar spacing, is d{iooj =3.512k and d{onj =2.570A for the raw alloy, whereas d{iooj =3.571A and d{oi 1} =2.561A for the alloy after ENP process. The matrix lattice returns to its original state without significant changes in its lattice size or the equilibrium position between atoms after the electrical pulse is turned off.

[0183] Fig. 37C depicts tensile tests that show that the elastic modulus of both raw and treated alloys does not change even after thermal annealing at 900°C for 20 minutes.The decrease in elastic modulus after the ENP process is related to the physical state of metal ions and free electrons in the alloy matrix.

[0184] Fig. 37D shows the work hardening rate calculated from the corresponding true stress-strain curves. The criteria (0=OT) for the onset of necking formation are also given. The raw alloy only shows stage II (a thermal hardening range) with a much lower work hardening rate and fractures before reaching stage III (dynamic recovery range). The alloy, after ENP processing, exhibits typical work hardening behavior for Ni-based superalloys. A plateau stage (stage II) follows the initial elastic-plastic transition stage, with a gradual decrease in the work hardening rate as stress increases (stage III). The work hardening rate in stage II (9n) is much higher and reaches the expected thermal hardening rate (0H=G 2O) for Ni-based superalloys with dispersible precipitates, indicating that the ENP process can produce hardening performance in Ni-based superalloys by forming fine and dispersed NisCe precipitates in nanoseconds, which would otherwise take hours to achieve with conventional thermal treatments.

[0185] Fig. 38A through Fig. 38G illustrate the formation of field-activated precipitates in Ni-based superalloy in accordance with an exemplary embodiment. Fig. 38A through Fig. 38D show SEM-BSE images of the Ni-based superalloy after various furnace heating treatments. Fig. 38A shows the alloy after furnace heating at 1 00°C and 10s. Fig. 38B shows the alloy after furnace heating at 1400°C and 30s. Fig. 39C shows the ally after aging treatment at 950°C / 30min and 700°C / 10h. Fig. 38D shows the morphology of Ni5Ce precipitates in the alloy after ENP processing at 1.30xl 011A / m2and 330ns, followed by furnace heating at 1400°C and 10s. Fig. 38E and Fig. 38F show the Ni-based superalloy after various electropulsing treatments. Fig. 38E shows the alloy after low- intensity and long-duration electropulsing treatment (SPS) at 1.52xl 07A / m2and 14min. Fig. 38F shows the alloys after medium-intensity and medium-duration electropulsing treatment (EP) at 3.15xl 010A / m2and 8ps. Fig. 38G shows schematic diagrams illustrating the coupling relations of thermal and field effects in the electropulsing treatment.

[0186] The Ni-based superalloy was subjected to furnace heating without the presence of an electric current to illustrate that the quasi-instantaneous formation of the NisCe precipitate is an exclusive outcome of the ENP processing. When held at its meltingpoint (1400°C) for 10 seconds, the alloy exhibits significant grain coarsening without any precipitation (Fig. 38A). When extended to 30 seconds, this alloy begins to melt and has hypo eutectic structure after solidification (Fig. 38B) demonstrating that absent electropulsing, the precipitation behavior is also absent in the alloy and that the cooling process does not contribute to the precipitate formation as both the ENP and furnace treatments shared the same cooling stage behavior.

[0187] Nickel-based superalloys are highly valued for their ability to develop controlled precipitates through various aging treatments. However, some treatments can thermally induce the precipitation of NisCe intermetallic compounds within the alloy. After a solutionaging treatment, the NiCr2 Laves phase, which has a topologically close-packed (TCP) structure, is formed (Fig. 38C), which is not desirable for some applications, such as structural applications, because of its brittleness at ambient temperatures.

[0188] The precipitation of NisCe intermetallic is primarily driven by the strong electromigration inherent to ENP processing. Due to the significant difference in timescales between ENP processing and furnace heating, there could be concerns that the rapid formation of NisCe intermetallic could dissipate because of instability during extended furnace heating. However, additional furnace heating of the alloy at 1400°C for 10 seconds after ENP processing showed that NisCe precipitates (Fig. 38D) remained intact throughout the heating period.

[0189] Utilizing electric current is not exclusive to ENP processing. After SPS treatment, the alloy showed more grain growth than with furnace heating, likely due to the longer processing time, but no field-activated precipitation was observed (Fig. 38E). In SPS treatment, Joule heating overshadows the effects of electric current (Fig. 38G). While a stronger electric current enhances the field effect, Joule heating is still the main factor at an intensity of 3.15x1010A / m2and a duration of 8 ps. Slight grain growth occurs without significant changes or precipitation (Fig. 38F), suggesting grain boundary (GB) migration at the atomic scale. The strong electromigration effect, absent in SPS and furnace treatments, facilitates NisCe precipitation in ENP processing (Fig. 36B). Although coupling exists, the field effect predominates in high intensity electropulsing treatment (~1011A / m2for 330 ns) (Fig. 39G).

[0190] Fig. 39A through Fig. 39C illustrate the formation mechanism of field-activated precipitates in Ni-based superalloy in accordance with an exemplary embodiment. Fig. 39A shows TEM-EDS measurements on the elemental distributions of NisCe intermetallic after ENP processing at 1.30xl 011A / m2and 330ns. Fig. 39B is a schematic diagram showing the atomic migrations and microstructure evolution in the Ni-based superalloy during ENP processing at 1.30xl011A / m2and 330ns. The inserts show schematics of the melting point decrease at GBs due to cerium atom concentration and FEM simulation on temperature distribution around Ni5Ce intermetallic during ENP processing. Fig. 39C compares the enhanced diffusivity of various elements in Ni-based superalloy during ENP processing at 1.30xl 011A / m2and 330ns.

[0191] The eutectic transition at GBs due to localized melting is crucial for forming the NisCe intermetallic. A counter-proof approach confirms that a cerium supersaturated solid solution at GBs cannot lead to NisCe precipitation during ENP processing. Furnace heating at 1300°C for 30 seconds showed no melting while heating at 1400°C for the same duration resulted in melting, supporting the counter-proof validity. After ENP processing at a current density of 1 .30X1011A / m2for 310 ns (at 1300°C), the microstructure was similar to the raw alloy, with only minimal NisCe precipitate. Significant Ce precipitation likely occurs in the final 20 ns of ENP when localized liquid forms at GBs.

[0192] The mechanism of field-activated precipitation in a Ni-based superalloy during ENP processing is shown in Fig. 39B. Molecular dynamics simulations indicate that an external electric field enhances atomic diffusivity. In the early stages of ENP, before eutectic transitions at the GBs, atomic flux due to concentration gradients is minimal. The primary driver for mass transport, particularly the movement of cerium atoms toward the GBs, is the electromigration effect, resulting in significant atomic flux. This accumulation of cerium lowers the melting point at the GBs to 1211 °C, well below the matrix temperature of 1400°C at 330 ns. Localized Joule heating at the GBs contributes to transient liquid formation, causing superheating and enabling localized melting within nanoseconds. When the Ce-rich liquid phase, the temperature gradient increases due to electric current concentration, creating a temperature difference of about 136 K., which promotes thermal diffusion and growth of precipitates along the GBs.Quantitative analysis of electric field effect on diffusivity during ENP

[0193] In accordance with many embodiments, electromigration (JE) significantly contributes to the precipitation of NisCe intermetallic during ENP processing, while thermal diffusion (JT) becomes relevant when eutectic transitions occur at GBs. To assess the contributions, transmission electron microscopy with energy-dispersive spectroscopy (TEM-EDS) was used to analyze the concentration distributions of precipitates (Fig. 40A). Electromigration is the dominant factor for Ni5ce precipitation, accounting for about 99% of the atomic flux. Although diffusion from concentration gradients opposes NisCe formation, its impact is minimal. The temperature gradient from Joule heating produces even less atomic flux and is only notable after eutectic reactions at GBs, indicating that ENP-induced precipitation of NisCe intermetallic cannot be achieved through traditional thermal or electropulsing treatments.

[0194] The increase in atomic diffusivity in the presence of an electric field (£ ") is shown in Fig. 40C. Atomic diffusivity without an electric field (D / ) and total atomic mobility (D) were measured at approximately 1400°C, with t increase mainly attributed to electromigration effects. Electromigration causes momentum transfer from electrons to ions, generating significant lattice strain and facilitating local stress relaxation through atomic diffusion. The rapid changes in electric current create an induced vortex magnetic field (Fig. 36A), enhancing diffusivity by increasing activation entropy and atom jump frequency. Vacancy formation is crucial for this process; while atom migration via vacancies is typically more difficult in substitutional solid solutions like this alloy, extensive vacancy formation under strong electromigration during ENP processing significantly bolsters this mechanism. The diffusion coefficient can be expressed as:Here, Do is the pre-exponential factor, and Q is the diffusion activation free energy, which includes vacancy formation energy AEv and activation internal energy AE. T represents temperature, and k is the Boltzmann constant. The distance between atomic jump planes is d, while P is the probability of atomic jumps, v is the vibration frequency towardsvacancies, and z is the atomic coordination number related to vacancy concentration. Sv indicates the vibration entropy of vacancies, and AS is the diffusion activation entropy.

[0195] Diffusion activation free energy is an intrinsic parameter dependent on the alloy's lattice structure. During the recrystallization of low carbon steel, while (Q) may decrease with electric current, enhanced diffusivity is attributed to an increased (Do). Higher vacancy concentration and lattice tensile strain lead to increased (d) and (z) values, resulting in a significant rise in diffusivity (~102times), driven by increased (ASv). Nickel, the principal alloy element, shows a weaker response to diffusivity enhancement compared to cerium under electric current. This suggests that ENP is more effective in complex alloys than in pure metals. As a result, and in accordance with many embodiments, ENP processing can influence various kinetic processes in alloys with micro alloying, such as homogenization, precipitation, and corrosion.DOCTRINE OF EQUIVALENTS

[0196] This description of the invention has been presented for the purposes of illustration and description. It is not intended to be exhaustive or to limit the invention to the precise form described, and many modifications and variations are possible in light of the teaching above. The embodiments were chosen and described in order to best explain the principles of the invention and its practical applications. This description will enable others skilled in the art to best utilize and practice the invention in various embodiments and with various modifications as are suited to a particular use. The scope of the invention is defined by the following claims.

[0197] As used herein, the singular terms "a," "an," and "the" may include plural referents unless the context clearly dictates otherwise. Reference to an object in the singular is not intended to mean "one and only one" unless explicitly so stated, but rather "one or more."

[0198] As used herein, the terms "approximately" and "about" are used to describe and account for small variations. When used in conjunction with an event or circumstance, the terms can refer to instances in which the event or circumstance occurs precisely as well as instances in which the event or circumstance occurs to a close approximation. Whenused in conjunction with a numerical value, the terms can refer to a range of variation of less than or equal to ± 10% of that numerical value, such as less than or equal to ±5%, less than or equal to ±4%, less than or equal to ±3%, less than or equal to ±2%, less than or equal to ±1 %, less than or equal to ±0.5%, less than or equal to ±0.1 %, or less than or equal to ±0.05%.

[0199] Additionally, amounts, ratios, and other numerical values may sometimes be presented herein in a range format. It is to be understood that such range format is used for convenience and brevity and should be understood flexibly to include numerical values explicitly specified as limits of a range, but also to include all individual numerical values or sub-ranges encompassed within that range as if each numerical value and sub-range is explicitly specified. For example, a ratio in the range of about 1 to about 200 should be understood to include the explicitly recited limits of about 1 and about 200, but also to include individual ratios such as about 2, about 3, and about 4, and sub-ranges such as about 10 to about 50, about 20 to about 100, and so forth.

Claims

WHAT IS CLAIMED IS:1 . A device for generating high-current density nano pulses comprising: an electrical source, at least one capacitor, a pulse signal generator, and a switch configured for electric current pulsing; wherein the electric current pulsing is at a set intensity, a set duration, and a set frequency and configured for at least one of a short rise time and a short fall time; wherein the device is configured to electrically couple to a conductive material and release a current into the conductive material; and wherein the current released is configured to induce a property change in the conductive material; wherein the current density exceeds 1010A / m2and the pulse duration exceeds 100ns.

2. The method of claim 1 , wherein the current density is approximately 1010-1012A / m2.

3. The method of claim 1 , wherein the pulse duration is less than 1 ps.

4. The method of claim 1 , wherein the pulsing frequency is up to 10OKHz.

5. The device of claim 1 , wherein the capacitor is configured for energy storage and charging.

6. The device of claim 5, wherein the capacitor is a supercapacitor.

7. The device of claim 1 , wherein the switch is a solid state switch.

8. The device of claim 5, wherein the capacitor is configured for charging with a voltage of up to 1000V.

9. The device of claim 1 , wherein the pulse signal generator outputs a square wave signal.

10. The device of claim 9, wherein the square wave signal output ranges from 100ns to a 5ps.

11. The device of claim 1 , wherein the device further comprises at least one constant resistor.

12. The device of claim 11 , wherein the at least one constant resistor is a plurality of parallel constant resistors with a resistance of approximately 95mQ.

13. The device of claim 11 , wherein the at least one constant resistor is configured for over-current protection.

14. A method of material grain boundary engineering comprising, charging and storing energy in a capacitor, releasing the stored energy with a pulse signal generator into an electrically coupled material; wherein the stored energy is discharged at a set intensity, a set duration, and a set frequency; producing an electric current pulse; inducing a change in the microstructure of the electrically coupled material; wherein the change in the microstructure is configured for a selected property of the material; wherein the current density exceeds 1010A / m2and the pulse duration exceeds 100ns.

15. The method of claim 14, wherein the material is a polycrystalline material.

16. The method of claim 14, wherein the releasing electric energy into the material activates a non-equilibrium structural evolution of the material.

17. The method of 14, wherein the change in the microstructure is at a nanometer spatial scale18. The method of 14, wherein the change in the microstructure is at a nanosecond temporal scale.

19. The method of claim 14, wherein the pulse signal generator outputs a square wave signal.

20. The method of claim 19, wherein the output ranges from 100ns to 5ps.

21. The method of claim 14, wherein releasing the stored energy further comprises activating a switch configured for at least one of a short rise time and a short fall time of the electric current pulse.

22. The method of claim 14, wherein the switch is a solid state switch.

23. The method of claim 14, wherein the current density is approximately 1O1o-1O12A / m224. The method of claim 14, wherein the pulse duration is less than 1 s.

25. The method of claim 14, wherein the pulsing frequency is up to 100KHz.

26. The method of claim 14, wherein the change in the microstructure further comprises inducing dislocations at multiple scales.

27. The method of claim 26, wherein inducing dislocations at multiple scales further comprises inducing nanoscale dislocations with densities up to 1017to 1018A / m2.

28. The method of claim 14, wherein inducing the change in the microstructure further comprises modifying a grain boundary morphology.

29. The method of claim 28, wherein modifying a grain boundary morphology further comprises producing at least one of step-like and serrated features without affecting an adjacent grain structure.

30. The method of claim 14, wherein the selected property of the material is selected from the group consisting of: the diffusion rate of an element, plasticity, tensile strength, work-hardening performance, yield stress, elastic modulus, and corrosion resistance.

31. The method of claim 30, wherein the diffusion rate of an element is configured at least one of segregation at a grain boundary and the formation of precipitates.

32. The method of claim 14, wherein the change in the microstructure further comprises forming a nanocoating with a set structure.

33. The method of claim 33, wherein the nanocoating is Cr20s and the set structure is a triple hierarchical structure.

34. The method of claim 14, further comprising heating at a grain boundary with a temperature difference between the grain boundary and a grain interior.

35. The method of claim 34, wherein heating at the grain boundary leads to at least one of a localized stress field and a grain boundary transformation.

36. The method of claim 34, wherein the temperature difference lasts less than 100 microseconds.

37. The method of claim 14, further comprising heating the material to near the material melting points and maintaining a grain size and an orientation.

38. The method of claim 37, further comprising a heating rate of 108K / s.

39. The method of claim 16, wherein inducing non-equilibrium structure is configured to promote at least one of dislocation generation, dislocation reconfiguration, a specific dislocation configuration40. The method of claim 14, wherein the electric current pulses are configured for a dislocation distribution density.

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