Iron based alloy

Abstract

To provide an iron based alloy.SOLUTION: The alloy is represented by formula (Fe1-xCox)100-y-z-aByCuzMa, where x=0.1-0.4, y=10-16, z=0-1, a=0-8, and M=Nb, Mo, Ta, W, Ni, or Sn. The alloy has crystalline grains with an average size of 30 nm or less.SELECTED DRAWING: Figure 5

Description

[Technical field]

[0001] The present invention relates generally to alloys and methods of making same, and more particularly to Fe-based alloys and methods of making same. [Background technology]

[0002] Nanocrystalline Fe-based alloys can have soft magnetic properties and are typically produced by crystallization of rapidly quenched amorphous precursors. These alloys have a two-phase microstructure consisting of Fe-rich grains embedded in an amorphous matrix containing glass-forming elements.

[0003] Such materials have soft magnetic properties that make them desirable in applications requiring enhancement and / or channeling of magnetic flux generated by an electric current. For example, such materials have advantageously low coercivity (H c ), low or near-zero saturation magnetostriction, and very low core loss. However, compared to conventional Fe-Si steels, for example, the saturation magnetization (J s ) is the J of Fe-Si steel. s (i.e., about 2 T), large-scale production and application of these materials is limited, and devices made using these alloys have limited specific power densities, making them undesirable for weight-sensitive applications such as those found in the aerospace industry.

[0004] In recent years, there has been a continuous effort to develop alternative Fe-based alloy compositions to replace conventional Fe-Si steels for soft magnetic applications. However, these alloys lack the ability to provide a sufficiently high J s or H c High J only at the expense of s , but still not as high as desired.

[0005] Therefore, there is scope for developing Fe-based alloys with improved soft magnetic properties over existing alloys. Summary of the Invention [Means for solving the problem]

[0006] The present invention relates to a compound of formula (Fe 1-x Co x ) 100-y-z-a B y Cu z M a The present invention provides an alloy represented by the formula: (wherein x=0.1-0.4, y=10-16, z=0-1, a=0-8, and M=Nb, Mo, Ta, W, Ni, or Sn), which has crystal grains with an average grain size of 30 nm or less.

[0007] The specific composition and microstructure of the alloy of the present invention surprisingly results in a higher magnetic saturation (J s ) and low magnetic coercivity (H c ) can be advantageously combined.

[0008] As will be understood by those skilled in the art, the term "magnetic saturation" is used herein to refer to the magnetic state an alloy reaches when increasing applied external magnetic fields no longer increase the magnetization of the material. Additionally, the term "magnetic coercivity" is used herein according to its conventional meaning, i.e., as a measure of the ability of an alloy to withstand an external magnetic field without becoming demagnetized.

[0009] The alloys of the present invention advantageously have a high J s value (over 1.98T, etc.) and low H c (such as less than 25 A / m, e.g., less than 10 A / m). In some embodiments, the J s is higher than 2T. Typically, H is less than 25A / m c The values ​​are highly desirable for commercial applications. This makes the alloys of the present invention a preferred replacement for conventional Fe-Si steels for soft magnetic applications. Thus, the alloys of the present invention can function as soft magnetic alloys and are particularly suitable for use in applications requiring enhancement and / or channeling of magnetic flux generated by an electric current.

[0010] By functioning as a "magnetically soft" alloy, the alloy of the present invention is susceptible to the influence of a magnetic field, but the ferromagnetic properties of the alloy are only observed after an external magnetic field is applied. Thus, the alloy of the present invention is considered a magnetically soft alloy. 1-x Co x ) 87-y-z-a B y Cu z M a It can also be said that the present invention provides a soft magnetic alloy represented by the formula (wherein x=0.1 to 0.4, y=10 to 16, z=0 to 1, a=0 to 8, and M=Nb, Mo, Ta, W, Ni, or Sn), which has crystal grains with an average grain size of 30 nm or less.

[0011] The present invention further provides a method for producing an alloy, the method comprising the steps of: (i) forming an alloy of formula (Fe 1-x Co x ) 100-y-z-a B y Cu z M a (wherein x=0.1-0.4, y=10-16, z=0-1, a=0-8, and M=Nb, Mo, Ta, W, Ni, or Sn); and (ii) heating the amorphous alloy at a heating rate of at least 200° C. / sec.

[0012] By heating the alloy compositions described herein at a heating rate of at least 200° C. / sec, the method of the present invention advantageously improves soft magnetic properties (i.e., H c ) without significantly impairing the s The method of the present invention is particularly useful for producing alloys having a high Co content (high J s This invention is advantageous over conventional methods in that it allows the synthesis of alloys having coercivity levels significantly lower than those traditionally associated with alloys having Co contents greater than 8% (atomic).

[0013] Further aspects and / or embodiments of the present invention are outlined below.

[0014] DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS Hereinafter, embodiments of the present invention will be described with reference to the following drawings, but are not limited thereto. [Brief description of the drawings]

[0015] [Figure 1] 1 shows (a) a schematic diagram of the temperature evolution of the alloy during heating, and (b) the measured magnetic hysteresis curves for an embodiment of the (Fe0.8Co0.2)87B13 alloy obtained by rapid transverse magnetic annealing (TFA) and field-free annealing (NFA). [Diagram 2] 1 shows an example of an annealing configuration using a preheated (a) block or (b) roll according to a procedure of an embodiment. [Diagram 3] 1 shows measured core loss for an embodiment of an (Fe0.8Co0.2)87B13 alloy obtained by rapid transverse magnetic annealing (TFA) and field-free annealing (NFA). [Figure 4] (a) X-ray diffraction (XRD) pattern obtained from the (Fe1-xCox)87B13 as-cast material and (b) X-ray diffraction (XRD) pattern obtained after annealing are shown. [Diagram 5] The direct current (DC) coercivity (Hc), average grain size (D), and saturation magnetic polarization (Js) are shown versus heating rate for (Fe0.75Co0.25)87B13. [Figure 6] 1 shows DC coercivity versus annealing temperature for (Fe1-xCox)87B13. [Figure 7] For (Fe1-xCox)87B13, the DC coercive force (Hc), average grain size (D), and saturation magnetic polarization (Js) are shown versus Co content. [Figure 8] (a) XRD patterns of the (Fe0.8Co0.2)87-zB13Cuz as-cast samples (z = 0, 0.5, 1, and z = 1.5 for comparison), and (b) XRD patterns of the annealed (Fe1-xCox)86B13Cu1 samples (x = 0 to 0.3). [Figure 9]The DC coercivity versus annealing temperature is shown for the (Fe0.8Co0.2)87-zB13Cuz samples (z=0, 0.5, 1, and z=1.5 for comparison). [Figure 10] The DC coercivity (Hc), average grain size (D), and saturation magnetic polarization (Js) versus Cu content are shown for the (Fe0.8Co0.2)87-zB13Cuz samples (z=0, 0.5, 1, and z=1.5 for comparison). [Figure 11] The DC BH hysteresis curves of (Fe0.5Co0.5)87B13 and the listed grain sizes after ultrafast annealing at 460°C (733K) to 540°C (813K) for 0.5 seconds are shown. [Figure 12] The relationship between coercivity and average grain size is shown for (Fe1-xCox)87B13, (Fe0.8Co0.2)87-zB13Cuz, and (Fe1-xCox)86B13Cu1 annealed at a heating rate of 10,000°C / s, and for (Fe0.75Co0.25)87B13 annealed at heating rates ranging from 3.7 to 10,000°C / s. [Figure 13] The Js versus Co content x is shown for annealed (Fe1-xCox)87B13 as-cast, annealed (Fe1-xCox)86B13Cu1, and crystalline Fe1-xCox. [Figure 14] The complex permeability versus applied magnetic field taken at 1000 Hz (frequency of the magnetic field used during the measurements) is shown for the transverse magnetic field annealed (TFA), longitudinal magnetic field annealed (LFA) and no external magnetic field annealed (NFA) samples. [Figure 15] 1 shows the coercivity as a function of annealing temperature for the (Fe0.8Co0.2)87B13 embodiment alloy. [Figure 16] 1 shows DC hysteresis loops measured for an (Fe0.8Co0.2)87B13 embodiment alloy after annealing at the optimum annealing temperature. [Figure 17]1 shows the effect of annealing and cooling on the magnetic polarization properties of an (Fe0.8Co0.2)87B13 embodiment alloy obtained by annealing in the presence of a transverse magnetic field followed by cooling with or without the magnetic field. [Figure 18] Measured core losses at 50 Hz, 400 Hz, and 1000 Hz for a comparative sample of 3 wt. % iron-silicon steel are shown in comparison to an alloy of the (Fe0.8Co0.2)86B13Cu1 embodiment. DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0016] The present invention relates to a compound of formula (Fe 1-x Co x ) 100-y-z-a B y Cu z M a (wherein x=0.1-0.4, y=10-16, z=0-1, a=0-8, and M=Nb, Mo, Ta, W, Ni, or Sn). Unless otherwise indicated, elemental ranges and compositional values ​​used herein are intended to refer to atomic percent.

[0017] Due to their specific compositions, the alloys of the present invention have a dual-phase microstructure characterized by a crystalline phase consisting of body-centered cubic (bcc) Fe-Co grains, or bcc Fe-Co-Ni grains when Ni is present, embedded in an amorphous phase containing high concentrations of non-ferromagnetic elements such as B, Cu, Nb, Mo, Ta, W, and Sn, when present according to the formula definitions above.

[0018] When x is in the range of about 0.1 to about 0.4, the alloy of the present invention has a magnetic saturation J exceeding 1.98 T. s has sufficient cobalt to be advantageously obtained. s Such values ​​make the alloy of the present invention competitive with conventional soft magnetic alloys based on Fe-Si steels, for example. sIt has been found that the magnetic saturation of the alloy is less than 1.98 T, making the alloy less preferred for practical purposes. In some embodiments, the magnetic saturation of the alloy is advantageously at least 2 T.

[0019] In some embodiments, x ranges from about 0.2 to about 0.3. In these embodiments, the alloy has a J s Contains sufficient cobalt to ensure

[0020] The alloy of the present invention contains boron in an atomic content range of about 10% to about 16% (i.e., y=10-16). This range ensures the stability of the amorphous phase and is favorable for the large magnetocrystalline anisotropy of H c In particular, the presence of at least 10% boron in the alloy increases the stability of the amorphous phase, and less than 16% boron minimizes unwanted Fe-B compounds after heating.

[0021] In some embodiments, y is at least 11. For example, y can be at least 12. In these embodiments, the glass-forming ability of the amorphous phase upon casting is improved (i.e., the production of an amorphous phase that is free of crystalline phase is improved).

[0022] In some embodiments, y is at least 15% or less, such as 14% or less, such that such a concentration advantageously allows for the absence of unwanted Fe-B compounds in the alloy and improves the magnetic saturation of the alloy (i.e., magnetic saturation increases as y decreases).

[0023] The alloy may further include copper. Specifically, the alloy of the present invention includes 0-1% (i.e., z=0-1) of copper in atomic concentration. The copper in the alloy composition may contribute to the refinement of the grains constituting the crystalline phase of the alloy. This may be advantageous, for example, during the synthesis of the alloy. This is because copper is believed to provide heterogeneous nucleation sites for the crystalline phase. Even at low concentrations of copper (e.g., z=0.2 or z=0.5), the refinement of the grains of the crystalline phase has been confirmed. On the other hand, an excessive amount of copper (e.g., more than 1%) may firstly hinder the formation of the amorphous phase, making the alloy too brittle for practical use and reducing the soft magnetic properties. Thus, in some embodiments, z is in the range of 0.2-1, 0.2-0.7, or 0.2-0.5.

[0024] The alloy of the present invention may further comprise an element M selected from Nb, Mo, Ta, W, Ni, and Sn. Specifically, the alloy has an atomic content of Nb, Mo, Ta, W, Ni, or Sn of 0-8% (i.e., a=0-8). The presence of element M increases the H content of the alloy. c For example, any of these elements during the synthesis of the alloys can suppress grain growth of crystalline phases, resulting in the formation of H c Moreover, the presence of element M ensures a greater stabilization of the amorphous phase over a wider temperature range than alloys without element M. On the other hand, the presence of an excess of element M of more than 8% in the alloy reduces the J of the alloy due to the corresponding reduction in the Fe and Co contents in the alloy. s This could have a negative impact on

[0025] Thus, in some embodiments, a ranges from 0 to 7.5, 0 to 5, 0 to 2.5, or 0 to 1.

[0026] In some embodiments, both z and a are 0.

[0027] The alloys of the present invention have grains with an average grain size of 30 nm or less. For a particular alloy, the "average grain size" of the grains is the average grain size determined by Scherrer's formula from the X-ray diffraction (XRD) pattern of the alloy, with reference to the line broadening of the Fe(110) bcc reflection, according to procedures that would be known to one of ordinary skill in the art.

[0028] According to the XRD pattern measured for the alloy of the embodiment, the grains have a body-centered cubic (bcc) crystal structure. Without wishing to be limited by theory, the composition of the grains is approximately Fe 1-x Co x where x is the standard composition. The elements B, Cu (if present), Nb (if present), Mo (if present), Ta (if present), W (if present), and Sn (if present) are generally not considered to be included in the grains as they are generally considered to be excluded into the residual amorphous phase upon crystallization. The only exception is Ni, if present. Thus, for Ni-containing alloys, the grains are considered to contain the same proportions of Fe, Co, and Ni as expressed in the standard composition.

[0029] The average grain size of the crystal grains is not limited as long as it is less than 30 nm. In some embodiments, the alloy includes crystal grains having an average grain size of about 20 nm or less, about 15 nm or less, about 10 nm or less, or about 5 nm or less. For example, the alloy may include crystal grains having an average grain size of about 10 to about 30 nm.

[0030] Due to the particular combination of cobalt content (x=0.1-0.4) and grain size less than 30 nm, the alloy of the present invention advantageously has a magnetic saturation J while maintaining a coercivity less than 25 A / m, e.g., less than 10 A / m. s This indicates that alloys with a cobalt content of more than 8 atomic percent have a high J s However, due to magnetically induced anisotropy, c This is surprising from the perspective of the conventional understanding that

[0031] Without wishing to be limited by theory, it is believed that the crystalline phase of the alloys of the present invention is advantageously characterized by low values ​​of magnetically induced anisotropy associated with cobalt, which allows the alloys of the present invention to have higher cobalt contents while maintaining a magnetically softer magnetic property than conventional Fe-Co alloys, thereby improving the J s At least 1.98T,H c The alloy has a dielectric constant of about 25 A / m or less, for example about 10 A / m or less.

[0032] It is believed that the particular microstructure of the alloys of the present invention results in a globally randomized magnetocrystalline anisotropy, which acts to average out the local magnetocrystalline anisotropy of the grains. Specifically, each grain may have a well-defined magnetic axis, but the randomized spatial orientation of all grains may be such that the overall magnetic anisotropy of the resulting alloy is minimized. As a result, the effect of large intrinsic magnetocrystalline anisotropy on the coercivity is minimized. The effectiveness of this averaging process is diminished by the presence of coherent magnetically induced anisotropy in the alloy. In principle, the degree of magnetically induced anisotropy can be quantified with reference to certain parameters, one useful parameter being the overall uniaxial anisotropy coefficient of the alloy (K u As would be known to one skilled in the art, such parameters can measure the directional dependence of the magnetic properties of the alloy.

[0033] In this context, the anisotropy coefficient associated with the alloys of the present invention can be significantly smaller than that of conventional soft magnetic alloys. For example, the uniaxial anisotropy coefficient (K u ) is about 200 J / m 3 In some embodiments, the anisotropy coefficient (K u ) is about 100 J / m 3 Less than 50 J / m 3 Less than 25 J / m 3 Less than or about 10 J / m 3 is less than.

[0034] As will be appreciated by those skilled in the art, the alloys of the present invention may further include unavoidable impurities. As used herein, the term "unavoidable impurities" refers to elements other than the elements of the alloys of the present invention that are necessarily present in the alloy obtained by a particular synthesis, for example because they are originally present in the alloy precursors. Such impurities include S, O, Si, Al, C, and N, among others.

[0035] The present invention also relates to a compound of the formula (Fe 1-x Co x ) 100-y-z-a B y Cu z M a (wherein x=0.1-0.4, y=10-16, z=0-1, a=0-8, and M=Nb, Mo, Ta, W, Ni, or Sn). The alloy being "amorphous" means that at least 80% of the volume of the alloy is in a non-crystalline state.

[0036] The amorphous alloy may be prepared by any procedure known to those skilled in the art that allows for obtaining an amorphous alloy having a particular composition, for example, the amorphous alloy may be produced by quenching a molten alloy.

[0037] In a typical procedure, the molten alloy is first synthesized. For example, the molten alloy may be produced by melting the alloy components (also referred to as "alloy precursors" in the specification). The alloy precursors may be melted individually and then mixed to form the molten alloy. Alternatively, at least one of the alloy precursors is melted (typically the main element of the alloy), and the other elements are added thereto and completely melted. As a further alternative, the solid alloy precursors (e.g., in particulate, powder, or ingot form) are first mixed together, heated to a temperature sufficient to melt the elements, and the molten elements are blended to produce the molten alloy. The alloy precursors are heated to a melting temperature sufficient to liquefy them in their entirety. The melting temperature is preferably 50°C, 100°C, or 300°C (or more) higher than the temperature at which the alloy precursors become liquid. The atmosphere in which the molten alloy is discharged is not particularly limited, but an inert gas atmosphere is preferable from the viewpoint of reducing the inclusion of oxides and the like in the amorphous alloy.

[0038] The molten alloy may then be held at the melting temperature for a time sufficient to ensure homogenization of the molten alloy. Thus, the actual melting temperature and time of the molten state are not limited, provided that the temperature and time ensure complete homogenization of the alloy precursor. In some embodiments, the molten alloy may be homogenized by heating and holding at a temperature of about 300° C. to 2,000° C. for at least 10 minutes.

[0039] In some embodiments, one or more of the alloy precursors are heated separately. For example, each alloy precursor may be liquefied or partially liquefied and then mixed together to form a molten alloy. In many embodiments, one or more of the alloy precursors are heated to various temperatures and then mixed.

[0040] The alloy precursor may be heated to obtain a molten alloy according to any suitable procedure known to those skilled in the art. For example, the molten alloy may be prepared by resistance melting, arc melting, induction melting, or a combination thereof. In resistance melting, electrical resistance is used as a heat source. In arc melting, heating is performed by an electric arc used as a heat source. In induction heating, heating is performed by electromagnetic induction through heat generated in the target by eddy currents at high frequency.

[0041] The molten alloy may then be quenched according to any procedure that ensures the formation of an amorphous alloy, for example, by melt spinning, centrifugal spinning, or solution quenching at a cooling rate fast enough to ensure the formation of an amorphous alloy.

[0042] In some embodiments, the amorphous alloy is produced by melt spinning, e.g., by dropping the molten alloy onto a rotating chill roll in a planar flow casting procedure. The procedure may be performed under inert conditions, e.g., under argon. The chill roll may rotate at any rotational speed that facilitates quenching of the molten alloy to produce the amorphous alloy. For example, the chill roll may rotate at a peripheral speed of about 15 m / s or more, about 30 m / s or more, or about 40 m / s or more. In some embodiments, the chill roll rotates at a peripheral speed of 55 m / s or less, 70 m / s or less, or 80 m / s or less. Those skilled in the art will be able to devise suitable rotational speeds that facilitate quenching of the molten alloy to produce the amorphous alloy.

[0043] Depending on the quenching method, the amorphous alloy may be obtained in the form of a ribbon, flakes, granules, or bulk. For example, when the amorphous alloy is produced by melt spinning, the alloy is obtained in the form of a ribbon. The ribbon may have a size according to the melting and spinning conditions. The thickness of the ribbon may be in the range of about 5 μm to about 45 μm, for example, about 10 μm to about 15 μm. The width of the ribbon may be in the range of about 0.5 mm to about 220 mm, for example, about 1 mm to about 200 mm, about 1 mm to about 150 mm, about 1 mm to about 100 mm, about 1 mm to about 50 mm, about 1 mm to about 25 mm, or about 1 mm to about 12 mm.

[0044] Certain elements in the alloy composition may play a role in determining the microstructure and composition of the alloy when the melt is quenched. For example, the presence of at least 10 atomic % B (boron) in the alloy composition (y≧10) favors the formation of an amorphous form of the alloy and promotes the stability of the amorphous phase. At the same time, as described herein, having 16% or less boron (y≦16) minimizes the formation of unwanted hard magnetic Fe-B compounds during annealing. Also, having a B content of 16 atomic % or less in the amorphous alloy avoids the formation of Fe-B compounds when the amorphous phase crystallizes.

[0045] Thus, in some embodiments, y is at least 11. For example, y can be at least 12. In some embodiments, y is at least 15% or less, such as at most 14%. Such concentrations advantageously ensure that unwanted Fe-B compounds are not present in the alloy.

[0046] The method of the present invention also requires that the amorphous alloy then be heated at a heating rate of at least 200° C. / sec. In the present invention, "heating rate" is understood to be the rate at which a particular amorphous alloy is heated, as measured by a non-insulated K-type thermocouple with a tip diameter of 0.1 mm in intimate thermal contact with the alloy.

[0047] In a typical procedure, the heating rate may be determined in terms of the rate of temperature increase measured with reference to the start and end temperatures in a single-step process. The start temperature may be room temperature (e.g., about 22°C) and the end temperature may be 95% of the difference between the start temperature and the temperature of the preheated surface used in the annealing process. Figure 1(a) shows a schematic diagram of the temperature profile associated with this type of measurement, along with a description of the relevant reference parameters. In the example procedure, the thermocouple tip is quickly (i.e., in less than 0.1 seconds) contacted to two parallel preheated surfaces with sufficient force to ensure good contact (i.e., thermocouple surface pressure of about 1 GPa). The temperature of the preheated surfaces is measured with a secondary thermocouple embedded in the heated surfaces within 1 mm of the contact area, and is measured after allowing the temperature reading to stabilize for at least 10 seconds to provide an accurate indication of the surface temperature. The mass of the heated surface should be large enough so that its measured temperature change does not exceed 5°C / s throughout the annealing process.

[0048] Annealing the amorphous alloy at a heating rate of at least 200° C. / sec promotes the formation of a fine crystalline phase of bcc Fe-Co, or Fe-Co-Ni if present, embedded in the amorphous phase, with the average grain size advantageously being less than 30 nm. In general, the faster the heating rate, the smaller the average grain size. As a result, the faster the heating rate, the more advantageously the H c In particular, a heating rate of at least 200° C. / sec advantageously provides precise control of the alloy microstructure (i.e., grains with a size of less than 30 nm) and provides a H c This leads to a significant decrease in s It was found that it was possible to ensure (i.e., exceed 1.98T).

[0049] Thus, in this context, a faster heating rate is beneficial in reducing the overall magnetically induced anisotropy of the alloy, and as described herein, cIt will be appreciated that this promotes a reduction in the H value. Thus, the method of the present invention allows for the control and minimization of magnetically induced anisotropy during annealing. c High Co content (and therefore high J) without compromising s Advantageously, it is possible to synthesize Fe-Co alloys having the above properties.

[0050] Thus, in some embodiments, the heating rate is greater than 200° C. / sec. For example, the amorphous alloy may be heated at a heating rate of at least about 250° C. / sec, at least about 500° C. / sec, at least about 750° C. / sec, at least about 1,000° C. / sec, at least about 1,500° C. / sec, at least about 2,000° C. / sec, at least about 5,000° C. / sec, at least about 7,500° C. / sec, at least about 10,000° C. / sec, or at least about 15,000° C. / sec.

[0051] It will be appreciated that, for methods involving heating the amorphous alloy at a heating rate of at least 200° C. / sec, the heating procedure may consist entirely of a heating step at that rate, or may involve heating at that rate as part of a multi-step heating procedure, in either case where the rapid heating rate is implemented during the majority (i.e., greater than 50%) of the crystallization process.

[0052] Any annealing procedure capable of heating the amorphous alloy at the rates disclosed herein may be suitable for use in the methods of the present invention.

[0053] For example, the amorphous alloy may be contacted with a heating element that has been preheated to an elevated temperature. In that regard, the heating element may be preheated to any temperature that will heat the amorphous alloy at a heating rate of at least 200°C / sec when the amorphous alloy is thermally contacted with the heating element. For example, the heating element may be preheated to at least about 500°C, at least about 750°C, or at least about 1,000°C. In some embodiments, the heating element is preheated to about 500°C.

[0054] Contacting the amorphous alloy with a heating element that has been preheated as described herein can be accomplished by any means known to those skilled in the art that would be suitable for the intended purpose.

[0055] For example, the amorphous alloy may be contacted with a preheated heating element in the form of a heating block. This may be done, for example, by a device capable of clamping the amorphous alloy between preheated blocks. The blocks may be made of any material capable of being preheated to a desired heating temperature and ensuring rapid heat transfer to the alloy. Suitable materials for the blocks are therefore metals (e.g., copper, titanium), alloys (e.g., steel, aluminum alloys), and ceramic materials (e.g., alumina). The clamping may be performed by applying a clamping force capable of ensuring uniform distribution of heat throughout the alloy. In some embodiments, the heating of the amorphous alloy is performed by clamping the alloy between preheated blocks at a pressure of at least about 3 kPa, e.g., at least 30 kPa, or at least 100 kPa. In one embodiment, the clamping force is 133 kPa. An example of such a configuration is shown in FIG. 2(a) in which a ribbon of amorphous alloy is clamped between preheated heating blocks.

[0056] According to another configuration, the amorphous alloy may be contacted with a heating element in a hot rolling configuration. These configurations are particularly preferred in that they allow the amorphous alloy to be annealed continuously. In these examples, the heating element may be in the form of two rolls preheated to the desired temperature and in contact with each other such that the rotation of one roll corresponds to the counter-rotation of the other roll. According to such an arrangement, the amorphous alloy in the form of a thin strip passes between the rotating rolls. Each roll may be made of any material that can be preheated to the desired heating temperature and ensures fast heat transfer to the alloy. Thus, suitable materials in this regard include metals (e.g., copper, titanium), alloys (e.g., steel, aluminum alloys), and ceramic materials (e.g., alumina). The rolls may be pressed together to obtain a clamping pressure of at least about 3 kPa, for example at least 30 kPa, or at least 100 kPa. In one embodiment, the rolls are pressed together to obtain a clamping force of 133 kPa.

[0057] An example of a hot rolling configuration suitable for use in the present invention is shown in FIG. 2(b). The figure shows a configuration based on a pair of preheated rollers through which the amorphous alloy ribbon is passed. The rollers may be preheated to any suitable temperature described herein, and the temperature of each roll may be independently adjusted to achieve the desired alloy structure. As the rolls rotate, the amorphous alloy ribbon is unwound from a pay-out reel and passes between the rolls, which may be pressed against each other with pressure as described herein. In the configuration shown in the figure, the ribbon is brought into tangential contact with one of the rolls along half of the circumference of the roll. However, the degree of contact between the ribbon and the rolls can be varied to allow the ribbon to be heated to the desired degree. When the ribbon leaves the contact point between the rolls, the temperature of the ribbon has risen to a level sufficient to begin to crystallize. By maintaining contact with one of the rolls as the rolls rotate, heat generated during crystallization can be removed. The ribbon then leaves the roll surface and is cooled (either by natural convection, forced convection, chilled blocks, or liquid cooling bath) before being wound onto the take-up reel. In certain configurations, a servo motor may be attached to one of the rolls to provide a controlled speed of rotation. The speed of rotation may be adjusted to control the annealing time of the ribbon. Also, servo motors attached to the payout and take-up reels may be used to provide a constant torque, and therefore tension, to the ribbon. Furthermore, an encoder may be attached to the servo motor to monitor and record the difference in the total number of rotations of the two mandrels, so that the tensile strain on the ribbon during the annealing process may be evaluated and controlled. In such cases, it is particularly advantageous to produce alloy ribbons with a minimum thickness, typically less than 18 μm. This will ensure that the formation of undesirable eddy currents is limited when the ribbon is formed into a lamination core and exposed to an alternating magnetic field. As a result, alloy production systems can be designed to have servo motors with higher efficiency (i.e., lower power loss), which is economically beneficial.

[0058] Additionally, annealing procedures that may be suitable to obtain a heating rate of at least 200° C. / sec include liquid bath annealing and hot air annealing.

[0059] In liquid bath annealing, the amorphous alloy is immersed in a liquid bath held at an elevated temperature. The bath acts as a heating element and can be held at a preheat temperature as described herein. The amorphous alloy may be immersed for any suitable time (e.g., on the order of seconds to minutes, such as 0.5 to 5 seconds) to obtain the desired structure. The bath may be made of any material that will become molten at the required bath temperature. Suitable materials in this regard include molten Pb-Sn solder, molten gallium, molten aluminum-gallium alloy, and molten salts.

[0060] In hot air annealing, the amorphous alloy (e.g., in the form of a ribbon) is rapidly heated by passing it over a stream of hot air that acts as a heating element. In some configurations, the alloy may be in the form of a ribbon that is drawn off a first spool and taken up by a second spool. In such cases, the tension in the ribbon during annealing can be adjusted by controlling the torque and / or speed of the spool (e.g., with a servo motor).

[0061] Regardless of how the amorphous alloy is heated, the actual heating rate of the amorphous alloy can be controlled by interposing one or more insulating layers between the heating element and the amorphous alloy sample. Such layers can be made, for example, of a material having the same or lower thermal conductivity than the material of the heating element. For example, the heating rate can be controlled by interposing one or more layers of metal (e.g., iron, titanium), alloy (e.g., steel, aluminum alloy), or ceramic material (e.g., alumina) between the heating element and the amorphous alloy sample.

[0062] In the method of the invention, the amorphous alloy may be heated at any annealing temperature suitable to provide an alloy having a microstructure characterized by a crystalline phase consisting primarily of bcc Fe grains with Co and, if present, Ni embedded in an amorphous phase. Without wishing to be limited by theory, it is believed that the microstructure of the amorphous alloy changes during heating according to a two-step crystallization mechanism in the following sequence: (amorphous) → (bcc Fe with Co or Ni, if present) + (amorphous phase) → (bcc Fe with Co or Ni, if present) + (hard magnetic compound such as Fe-B).

[0063] Thus, an appropriate annealing temperature for a particular heating rate may be determined to ensure that the formation of hard magnetic compounds is minimized or not formed, i.e., that the coercivity is minimized. Generally, a crystalline phase is formed when the annealing temperature is equal to or greater than the crystallization onset temperature. In that respect, if the annealing temperature exceeds the crystallization onset temperature of the Fe-B compound, strong magneto-crystalline anisotropy associated with the formation of hard magnetic Fe-B compounds may be induced. Thus, the annealing temperature may be determined to not reach or exceed the crystallization onset temperature of the Fe-B compound. For example, the annealing temperature of the amorphous alloy may be slightly lower (e.g., 5-20°C lower) than the temperature at which the Fe-B compound begins to form.

[0064] Thus, in some embodiments, the annealing temperature ranges from about 350° C. to about 650° C., from about 400° C. to about 650° C., from about 450° C. to about 600° C., from about 450° C. to about 550° C., or from about 450° C. to about 500° C. For example, the annealing temperature may be about 490° C., about 500° C., about 510° C., or about 520° C.

[0065] When selecting a suitable annealing temperature for the purposes of the present invention, one or more other factors may need to be taken into account. For example, the crystallization reaction associated with the formation of crystalline phases in the alloy may involve the release of significant latent heat that may itself contribute to the heating of the alloy. In this regard, the skilled person will take such additional contributions into account when devising a heating procedure. For example, the skilled person will take suitable precautions to suppress or remove excess latent heat of crystallization during annealing (e.g., using a preheating surface with suitable mass and thermal conductivity that will allow the removal of latent heat during the formation of crystalline phases).

[0066] In the method of the present invention, the amorphous alloy may be maintained at a predetermined annealing temperature for a time required to provide an alloy having a microstructure characterized by a crystalline phase consisting primarily of bcc Fe grains containing Co and, if present, Ni, embedded in the amorphous phase. Suitable annealing times are, for example, from about 0 to about 80 seconds, from about 0.1 to about 80 seconds, from about 0.1 to about 60 seconds, from about 0.1 to about 30 seconds, from about 0.1 to about 15 seconds, from about 0.1 to about 10 seconds, from about 0.1 to about 5 seconds, from about 0.1 to about 1 second, or from about 0.1 to about 0.5 seconds.

[0067] In some embodiments, the amorphous alloy is also subjected to an external force, such as tensile and / or compressive stress, while being heated. The application of tensile and / or compressive stress during annealing induces elastic strain in the structure of the crystals formed during annealing, which helps control the directionality of the magnetically induced anisotropy formed during annealing of the alloy.

[0068] The amorphous alloy may be subjected to tensile and / or compressive stress during heating by any means known to those skilled in the art. For example, if heating is performed by placing the amorphous alloy between heating elements and placing the alloy in thermal contact with the heating elements, the heating elements may be pressed together to apply compressive stress to the alloy. Additionally or alternatively, the amorphous alloy may be subjected to tensile stress by pulling the alloy at both ends while in contact with the heating elements. This may be done by any means known to those skilled in the art. For example, the alloy may be clamped at both ends and mechanically stretched. Alternatively, if the heating elements are in the form of heated rolls, the tension of the alloy may be adjusted as described herein.

[0069] In some embodiments, heating the amorphous alloy includes exposing the alloy to a magnetic field, which can further control the directionality of the magnetically induced anisotropy formed during annealing of the alloy. In particular, exposing the alloy to a magnetic field during annealing can maximize the effectiveness of the randomization of the magnetocrystalline anisotropy and promote the averaging of the local magnetocrystalline anisotropy of the grains as they are formed. As a result, the H of the resulting alloy can be controlled. c can be further minimized.

[0070] The strength of the magnetic field may be any strength that would be suitable to align the magnetization of the material during grain formation and / or during the cooling process after annealing is complete. In some embodiments, the strength of the magnetic field is at least about 0.3 kA / m. For example, the strength of the magnetic field may be at least about 1 kA / m, at least about 3 kA / m, at least about 10 kA / m, at least about 30 kA / m, or at least about 300 kA / m. In some embodiments, the strength of the magnetic field is about 1000 kA / m.

[0071] In some embodiments, the magnetic field is rotated or has a varying orientation and / or magnitude relative to the alloy material. Employing a magnetic field that is rotated or has a varying orientation and / or magnitude relative to the alloy material can result in an alloy with an essentially isotropic magnetization distribution. This significantly suppresses magnetically induced anisotropy, thereby improving the soft magnetic properties of the alloy (i.e., lower H c ) can be dramatically improved.

[0072] Any means that would allow the amorphous alloy to be annealed in the presence of a magnetic field whose orientation and / or magnitude is changing relative to the alloy material would be suitable for the purposes of the present invention. For example, a rotating magnetic field may be obtained by rotating a magnetic source around the alloy during annealing. Alternatively, the alloy may be fixed to a suitable rotating support and rotated in a fixed magnetic field during annealing. Alternatively, a magnetic field of alternating magnitude (i.e., the magnitude of the applied magnetic field can change with time) may be applied to the sample material in multiple fixed orientations across three dimensions.

[0073] The orientation or magnitude of the magnetic field relative to the alloy may be varied at any rate suitable for randomizing magnetically induced anisotropy in the alloy. In some embodiments, the rate at which the orientation or magnitude of the magnetic field is varied is at least about 1 Hz, at least about 30 Hz, at least about 100 Hz, at least about 300 Hz, at least about 1,000 Hz, or at least about 3,000 Hz. For example, the rate at which the orientation or magnitude of the magnetic field is varied is between about 1,000 Hz and about 3,000 Hz.

[0074] In some embodiments, the magnetic field is a transverse magnetic field. In that regard, FIG. 1 shows the (Fe 0.8 Co 0.2 ) 87 B 131 shows the magnetic hysteresis curves measured for the alloy, which are plotted against the hysteresis curves of the corresponding alloy samples annealed in the presence of a transverse magnetic field (TFA curve) compared to the hysteresis curves of the corresponding samples annealed in the absence of a magnetic field (NFA curve).

[0075] In some embodiments, the magnetic field is a longitudinal magnetic field. In those cases, the magnetic field is such that the magnetic field lines run substantially parallel to a major axis of the alloy. In these embodiments, the alloy sample may be a longitudinal field annealed (LFA) sample.

[0076] Moreover, an advantage of heating the amorphous alloys in the presence of a magnetic field is that they result in alloys with lower core loss compared to the corresponding alloys annealed without an applied magnetic field. In that regard, FIG. 3 shows the results of 0.5 second rapid annealing at 490°C (Fe 0.8 Co 0.2 ) 87 B 13 The core losses at 50 Hz, 400 Hz, and 1,000 Hz for the alloys are shown. The lower magnetic core loss observed in the TFA samples is believed to indicate a lower magnetic permeability (i.e., the slope of the curve in the range of 0 A / m to 400 A / m in Figure 1), suppressing the formation of eddy currents in the TFA samples.

[0077] After heating, the alloy may be cooled. Cooling may be accomplished by any means known to one of skill in the art. For example, cooling may be accomplished by natural or forced convection. In some embodiments, the alloy is cooled by exposing it to ambient conditions and allowing it to cool naturally to room temperature. In some embodiments, the alloy is cooled by placing it in thermal contact with a cooler surface or element. For example, the alloy may be placed in thermal contact with a cooling block, a bath of cold liquid, or a stream of cold air. One of skill in the art will be able to devise suitable cooling procedures in this regard.

[0078] Typically, the alloy may be cooled at any cooling rate that promotes the maintenance of the crystalline structure of the alloy obtained during heating. For example, the alloy may be cooled at a cooling rate of at least about 1° C. / sec, at least about 10° C. / sec, at least about 50° C. / sec, or at least about 100° C. / sec. In some embodiments, the alloy is cooled at a cooling rate of at least about 100° C. / sec. One of ordinary skill in the art would know how to monitor the cooling rate according to the procedures described herein for heating rates.

[0079] In some embodiments, the alloy is heated and then cooled in the presence of a magnetic field as described herein. For example, the alloy is heated and then cooled, e.g., to room temperature, in the presence of the same magnetic field used during the heating step. Advantageously, it has been observed that when the alloy is cooled in the presence of a magnetic field, the soft magnetic properties of the alloy can be further improved.

[0080] As used herein, "room temperature" refers to an ambient temperature which may be, for example, from 10° C. to 40° C., more typically from 15° C. to 30° C. For example, room temperature may be a temperature from 20° C. to 25° C.

[0081] Certain compositional features of the alloy may play a role in the crystallization kinetics of the alloy during heating. For example, the presence of Cu in the alloy may be effective in reducing the average grain size of the alloy. Without wishing to be limited by theory, it is understood that Cu acts as a heterogeneous nucleation site during heating of the amorphous alloy. Specifically, the addition of Cu to Fe-based nanocrystalline soft magnetic alloys may result in the formation of Cu-rich clusters before the onset of crystallization. Such Cu-rich clusters can act as heterogeneous nucleation sites and aid in grain refinement. It is also believed that an increase in Cu content reduces the onset temperature of Cu clustering, which results in a higher number density of Cu clusters before the onset of crystallization, thus improving grain refinement. In general, low concentrations of copper (e.g., z=0.2 or z=0.5) can have a significant effect on the grain refinement of the crystalline phase, while excessive amounts of copper (e.g., greater than 1%) may make the alloy too brittle for practical applications or may prevent the formation of the amorphous phase in the first place. Thus, in some embodiments, z ranges from 0.2 to 1, from 0.2 to 0.7, or from 0.2 to 0.5.

[0082] The alloy of the present invention may also contain an element M selected from Nb, Mo, Ta, W, Ni, and Sn. Specifically, the alloy contains 0-8 atomic % of Nb, Mo, Ta, W, Ni, or Sn (i.e., a=0-8). It has been found that the role of the additional element M is related to the grain refinement and / or stabilization of the amorphous matrix phase during heating of the amorphous alloy. As a result, the presence of element M increases the H content of the alloy. c For example, both of these elements can suppress grain growth of crystalline phases during the synthesis of the alloy, resulting in the H c Moreover, the presence of element M ensures a greater stabilization of the amorphous matrix phase over a wider temperature range than alloys without M. On the other hand, the presence of an excess of element M, exceeding 8%, in the alloy reduces the J of the alloy due to the corresponding reduction in the Fe and Co contents in the alloy. sThus, in some embodiments, a ranges from 0 to 7.5, 0 to 5, 0 to 2.5, or 0 to 1. In some embodiments, z and a are both 0. EXAMPLES

[0083] Example 1 (Fe 1-x Co x ) 87 B 13 Precursor amorphous ribbons with a nominal composition of (where x = 0-0.5) were produced by melt spinning (planar flow casting) in an Ar atmosphere. Ribbons with thicknesses of about 10-15 μm and widths of 1-12 mm were obtained. The ribbons were placed in 20 μm thick Cu foil packets and subjected to ultra-rapid annealing in an Ar atmosphere. These packets were then compressed between two preheated Cu blocks (length 150 mm, width 50 mm) with a force of 950 N for 0.5 s using a pneumatic cylinder and an automatic timing mechanism.

[0084] The average particle size (D) was evaluated by X-ray diffraction (XRD) using a CoKα source using the Scherrer equation. The density was evaluated using a He gas pycnometer. The saturation magnetic polarization (J s =μ 0 The magnetization strength (Ms) was evaluated using a vibrating sample magnetometer (VSM) BHV-35H (Riken Denshi Co., Ltd.) at 0.8 MAm and 22°C (295 K). c The evaluation was carried out at 295K using a hysteresis loop tracer BHS-40DC (manufactured by Riken Denshi Co., Ltd.).

[0085] Figure 4(a) shows the (Fe 1-x Co x ) 87 B 134(b) shows XRD patterns obtained from a selection of amorphous ribbon castings having compositions of 0.5 to 1.0%. The patterns were obtained from the side of the ribbon not in contact with the casting wheel. Since no discernible crystalline reflection peaks are observed at x=0-0.3, the ribbons are considered amorphous over the length scale detectable by XRD. At x=0.4 and 0.5, a crystalline reflection peak identified as bcc Fe is observed at approximately 52.8°. However, the intensity of this crystalline reflection peak is low compared to the broad amorphous background, suggesting that the volume fraction of bcc Fe in the as-cast product is less than 20%. FIG. 4(b) shows the XRD patterns obtained after the ultra-rapid annealing process. The patterns show crystalline reflection peaks identified as belonging to bcc Fe.

[0086] Figure 5 shows the relationship between the heating rate (α) and the (Fe 0.75 Co 0.25 ) 87 B 13 H c , D, and J determined by XRD s For each heating rate (α) used, the H c The annealing time was chosen to minimize H. The heating rate was varied by placing an insulating material between the sample and the preheated copper block. As the heating rate increased from 3.7 to about 10,000 °C / s, H c It can be seen that the current density decreases from about 70 A / m to 10 A / m, but J s remains above 2T in all conditions. c This decrease in is believed to be associated with a corresponding decrease in D from 24.3 to 19.7 nm, revealing that an ultrafast annealing process can be utilized to maximize the soft magnetic properties of this alloy system.

[0087] The data in FIG. 5 confirms that the coercivity and grain size decrease with increasing heating rate. From the trend line (dashed line) shown in the plot of FIG. 5, it can be seen that a heating rate of 200° C. / sec or higher is advantageous for obtaining grain sizes less than 30 nm (22 nm or less in this example). This is due to the fact that the coercivity (H c ), while magnetization saturation (J s ) can be sustained at values ​​in excess of 1.98 T. As discussed herein, low coercivities of 25 A / m or less would typically be required for commercial applications. Overall, the data confirms the significant advantages gained by heating the alloy at rates of 200° C. / sec or greater.

[0088] FIG. 6 shows the annealing temperature (T a ) for H c The optimal annealing temperature (T op ) can be identified for each alloy as the temperature at which the coercivity is minimal. op is observed near about 490°C (763K) for x = 0 and 0.2, and near about 500°C (773K), 510°C (783K), and 520°C (793K) for x = 0.3, 0.4, and 0.5, respectively.

[0089] In Fig. 7, the heating rate is about 10,000°C / s and the holding time is 0.5 s. op After annealing at 0.5 1-x Co x ) 87 B 13 H c , D, and J s The x value (relative to the Co content) is less than 0.25, which indicates moderate H c The only increase in the H is 6.4 A / m at x=0 and 10.2 A / m at x=0.25. c A sharp increase in H with Co content is observed, with a peak at 24 A / m at x=0.5. cThe increase in can be partly attributed to the coarsening of the microstructure, since D increases by about 1.3 nm for every 0.1 increase in x. However, the increase in H c The sudden increase in J is not reflected in the gradual change in D seen in Figure 7. s is observed to increase, with a maximum value of 2.04 T seen at x = 0.25, which is just as good as the measured value of 2.0 T for Fe-3 wt. % Si.

[0090] Example 2 Nanocrystalline (Fe 0.8 Co 0.2 ) 87-z B 13 Cu z The effect of Cu addition on the z = 0 to 1.5 alloys is also investigated. In this context, a sample with z = 1.5 was prepared for comparison. (Fe 0.8 Co 0.2 ) 87-z B 13 Cu z and (Fe 1-x Co x ) 86 B 13 Cu 1 Precursor amorphous ribbons with standard compositions of z=0-1.5 (z=1.5 sample is for comparison) and x=0-0.3 were produced by melt spinning (planar flow casting) in Ar atmosphere. Ribbons with thicknesses of about 10-15 μm and widths of 1-12 mm were obtained. The ribbons were placed in 20 μm thick Cu foil packets and subjected to ultra-rapid annealing in Ar atmosphere. These packets were then compressed between two preheated Cu blocks (length 150 mm, width 50 mm) with a force of 950 N for 0.5 s using a pneumatic cylinder and an automatic timing mechanism.

[0091] The grain size was evaluated by X-ray diffraction (XRD) using a CoKα source using the Scherrer equation. The density values ​​reported in this study were evaluated using a He gas pycnometer. The saturation magnetic polarization (J s =μ 0 The magnetization was evaluated at 0.8 MAm and 295 K using a vibrating sample magnetometer (VSM) BHV-35H (Riken Denshi Co., Ltd.).c The evaluation was carried out at 295K using a hysteresis loop tracer BHS-40DC (manufactured by Riken Denshi Co., Ltd.).

[0092] FIG. 8(a) shows the results of quenched (i.e., before annealing) amorphous (Fe 0.8 Co 0.2 ) 87-z B 13 Cu z The XRD patterns of the samples (z = 0, 0.5, 1, 1.5 (the last one is for comparison)) are shown in Fig. 8(b). 1-x Co x ) 86 B 13 Cu 1 The XRD patterns measured for (x=0 to 0.3) are shown.

[0093] FIG. 9 shows the (Fe 0.8 Co 0.2 ) 87-z B 13 Cu z (where z=0, 0.5, 1, and 1.5 (the last one is for comparison)) a H for c When thin ribbon castings of these compositions were examined by XRD (see FIG. 8), no discernible crystalline reflection peaks were found for the alloys with z=0-1 (e.g., z=0.5 and 1.0). In those cases, the data show broad reflections indicative of an amorphous alloy phase. However, some crystallization was observed for the alloy with z=1.5. From FIG. 9, it can be seen that the addition of Cu increased the T op The temperature drops by about 10°C, and H c It can be seen that the coercivity becomes more sensitive to changes in annealing temperature. Considering the general trend of the coercivity data in the plot of Figure 9, it can be seen that as the Cu content decreases from 1% to 0% (i.e., z=1 to z=0), the annealing temperature window that allows very low coercivity (i.e., less than 15 A / m) gradually expands. Overall, compared to alloys with Cu contents greater than 1% (e.g., 1.5%), alloys with z=0-1 have a wider annealing temperature window that can be employed to obtain the advantageous combination of very high magnetic saturation and very low coercivity.

[0094] Figure 10 shows the T op Annealing was performed at 1000 K (Fe 0.8 Co 0.2 ) 87-z B 13 Cu z H vs. Cu content c , D, and J s The only phase identified by XRD in the annealed sample was that belonging to bcc Fe. The addition of Cu decreased D, and z = 0 and z = 1.5 showed average grain sizes of 20.6 to 16.8 nm, respectively.

[0095] From Fig. 10, when 0.5 atomic % Cu is added, H c The H c It can be seen that the coercivity remains below 10 A / m. Overall, the data in FIG. 10 confirms that as the Cu(z) content increases from z=0 to z=1.0, the coercivity decreases favorably (i.e., by about 2.4 A / m, from 9.3 to 6.9 A / m, respectively). However, as the Cu(z) content increases beyond 1% (i.e., z is greater than 1), the coercivity begins to increase (to 8 A / m for z=1.5).

[0096] Saturation magnetization (J s ) data, it is also confirmed that the value decreases slightly with increasing Cu content, decreasing by about 0.01 T on average per atomic % of Cu. s Although the decrease in J is small, the data show that by controlling the Cu content below 1%, sIt has been shown that alloys with a z value of more than 1.98 T, e.g., 2 T or more, can be obtained. This further reinforces the discussion made herein regarding controlling the Cu content to less than 1%. In that regard, it is also recommended to limit the Cu content to 1% or less (i.e., z=0-1) to ensure adequate mechanical properties of the alloy and the formation of amorphous phase. This is particularly relevant when the alloy is produced in ribbon form, where the poor mechanical properties of alloys with z greater than 1 (i.e., Cu content greater than 1%) may prevent the formation of ribbons with thicknesses significantly smaller than 20 μm, e.g., less than 15 μm.

[0097] When Co is added, the Cu clustering onset temperature (T clust For example, when 20% of Fe is replaced by Co, T clust increases to a value equal to the crystallization onset temperature. Because Cu clustering must occur before significant crystallization begins to aid grain refinement, substituting Co for Fe may reduce the effectiveness of Cu as a nucleating agent. Also, increasing Cu content may improve grain refinement by decreasing the Cu clustering onset temperature, resulting in a higher number density of Cu clusters before the onset of crystallization.

[0098] From the data presented herein, (Fe 0.8 Co 0.2 ) 87 B 13 It can be seen that the grain size is obviously reduced by adding Cu to the alloy. According to the trend seen in Fig. 10, even a small amount of Cu, 0.5 atomic %, can be effective in grain refinement. This indicates that the grain size of the alloy op Even a small amount of Cu added during rapid annealing at 1000 K resulted in a T clust This may suggest that the onset temperature is lower than the onset temperature of crystallization. This effect may be due to the relatively high annealing temperature made possible by the ultrafast annealing technique. Furthermore, the decrease in D with higher Cu addition may suggest a higher number density of Cu clusters before the onset of crystallization.

[0099] Thus, in a general sense, the data supports the view that Cu is effective in reducing the average grain size and can provide some improvement in the soft magnetic properties of the sample alloys. Nevertheless, care should be taken to ensure that the Cu content does not impair the mechanical stability or the formation of amorphous phases in the alloys. In that regard, as discussed herein, it is recommended to limit the Cu content to 1% or less (i.e., z=0-1) to ensure adequate mechanical properties and the formation of amorphous phases in the alloys. The data also indicate that any discontinuity between grain size and soft magnetic properties, as shown in FIG. 11, may be due to the formation of a significant magnetically induced anisotropy upon the addition of Co, which adversely affects the exchange softening process.

[0100] Figure 11 shows the (Fe 0.5 Co 0.5 ) 87 B 13 The DC BH hysteresis curves of the annealed specimens and the respective grain sizes are shown in Figure 4. The samples used to obtain the BH curves in Figure 4 were approximately 100 mm long and 1 mm wide, and were measured using an open magnetic circuit in a 0.5 m long solenoid with an air-core compensation pick-up coil. D was also evaluated and is also shown in Figure 4. It is clear that D decreases with increasing annealing temperature. This improvement in grain refinement with annealing temperature is due to the observed H c However, it can also be seen from Figure 4 that the BH curve for the sample annealed at 480°C (753K) clearly shows signs of a Barkhausen jump. This, together with the squareness of the BH curve (high residual / saturation ratio), suggests that significant induced anisotropy may be present in this material. Furthermore, for samples annealed at temperatures above 480°C (753K), there is no sign of a Barkhausen jump and the squareness of the BH curve is observed.

[0101] Example 3 FIG. 12 shows the results of annealing at a heating rate of 10,000°C / s (Fe1-x Co x ) 87 B 13 , (Fe 0.8 Co 0.2 ) 87-z B 13 Cu z , and (Fe 1-x Co x ) 86 B 13 Cu 1 About H c The relationship between D and the heating rate is shown in Fig. 5. The heating rates of the (Fe 0.75 Co 0.25 ) 87 B 13 H c and D are also included.

[0102] When the grain size is larger than 20 nm, the coercivity is D 6 For small grain sizes, this dependence is well described by D 3 Close to. D 6 Dependency and D 3 Both dependencies are predicted by Herzer's random anisotropy mode. D 3 It has been shown that the dependence occurs when the exchange length is controlled by an induced anisotropy that is coherent over length scales larger than the exchange length. Thus, the magnetically induced anisotropy (K u ) is considered to correspond to the square of the Co content of the test sample. 0.8 Co 0.2 ) 87-x B 13 Cu z H for change in D c The relative insensitivity of K to these materials is quite large. u This further supports the view that the phenomenon is due to the presence of

[0103] Furthermore, from FIG. 12, D 3 It can also be seen that the data are quite scattered across regions. This scattering can be understood to reflect the different Ku levels present in each composition. c D6 From D 3 The switch to grain size dependence is due to the random magnetocrystalline anisotropy K u It is known that this occurs when the ratio of K to K is about 1:2. u A change of about an order of magnitude is expected to cause a switch in particle size dependence at various particle sizes, resulting in a scattering of the data.

[0104] As already seen in Fig. 7, the nanocrystals (Fe 1-x Co x ) 87 B 13 So, even though D changes gradually, H c increases sharply at x = 0.2. From Fig. 12, (Fe 1-x Co x ) 87 B 13 In this kind of H c The increase in D 6 From D 3 It can be seen that this corresponds to the transition to dependence. c The sharp increase in K was caused by the addition of Co. u Therefore, randomizing the magnetically induced anisotropy by rotating magnetic field annealing would be effective in improving the soft magnetic properties of the test samples.

[0105] FIG. 13 shows the as-cast nanocrystalline (Fe 1-x Co x ) 87 B 13 and nanocrystalline (Fe 1-x Co x ) 86 B 13 Cu 1 J s The J of non-oriented Fe-Si steels containing 3 wt% and 6.5 wt% Si is shown. s Typical values ​​of are also shown. 1-x Co x ) 87 B 13 At (x=0.2, 0.25, and 0.3), J exceeds 2T. sThis is comparable to Fe-3 wt.% Si steel.

[0106] Co-free composition Fe 87 B 13 When Co content of 0.1 was added to J in the as-cast state, s A single maximum increase in J for such a thin strip casting is observed. s The increase in the Curie temperature (T c ) from about 220° C. (497 K) to values ​​above the primary crystallization onset temperature of about 370° C. (643 K).

[0107] J of crystalline Fe-Co s It is well-established that the peak is located near the Co content x = 0.35. However, the annealed (Fe 1-x Co x ) 87 B 13 For the as-cast sample, the peaks are centered around x = 0.2 and 0.25, respectively.

[0108] Such a J s The difference in peak position is J s =V f cry J s cry +(1-V f cry )J s amo (In the formula, V f cry where J is the crystalline volume fraction) s amo ) and crystalline phase (J s cry ) can be understood to reflect the local volume-weighted average contribution from the

[0109] Assuming a uniform partitioning of Co into both phases after nanocrystallization, the equilibrium volume fraction of the crystalline phase can be evaluated by mass balance. The composition of the remaining amorphous phase after annealing is Fe 3Assuming a composition approaching that of B, the crystalline volume fraction is expected to be about 50%. Thus, the B-rich residual amorphous phase and the crystalline Fe–Co phase have J-phases similar to their bulk counterparts. s If the Co dependence of the two-phase nanocrystalline material has a Co content between the amorphous (x=0.2) and Fe-Co crystalline (x=0.35) phases, then the two-phase nanocrystalline material will have a J s It is expected to have a peak.

[0110] Table 1 shows the results of rapid annealing (Fe 0.8 Co 0.2 ) 87 B 13 and (Fe 0.8 Co 0.2 ) 86 B 13 Cu 1 H c , J s The properties of the soft magnetic materials are summarized in Table 1. ... 73.5 Cu 1 Nb 3 S 15.5 B 7 (Finemet), which has a high J that is superior to conventional soft magnetic materials, including Fe-based amorphous non-oriented (NO) Fe-Si steels. s (over 2T) and low H c It can be seen that a combination of (less than 10 A / m) can be obtained.

[0111] Table 1: Properties of the (Fe-Co)-B-(Cu) compositions investigated in this study and literature values ​​for nanocrystalline, amorphous, and crystalline materials.

[0112] [Table 1]

[0113] Example 4 FIG. 14 shows the complex permeability versus applied magnetic field obtained at 1000 Hz (the frequency of the magnetic field used during the measurements) for a transverse magnetic field annealed (TFA) sample, a longitudinal magnetic field annealed (LFA) sample, and a sample annealed without an external magnetic field (NFA). The compositions of the samples are (Fe 0.8 Co 0.2 ) 87 B 13 and all three conditions were annealed at 490°C for 0.5 seconds with a heating rate of 10,000°C / s (10,000K / s).

[0114] TFA was performed by placing the sample between two preheated copper blocks in the presence of an applied magnetic field of approximately 24,000 A / m oriented transversely to the measurement direction. LFA was performed by placing the sample between two preheated copper blocks in the presence of an applied magnetic field of approximately 3,000 A / m oriented longitudinally to the measurement direction.

[0115] In FIG. 14, it can be seen that the complex permeability is maximized at about 40 A / m for all three annealing methods. The LFA sample is seen to have the highest peak value of complex permeability (about 30,000) and the TFA sample is seen to have the lowest peak value (about 7,000). This reduction in complex permeability of the TFA sample is due to the formation of directional magnetically induced anisotropy. This directional magnetically induced anisotropy is perpendicular to the measurement direction of the TFA sample and therefore acts to reduce the complex permeability compared to the NFA sample. The LFA sample shows the opposite effect, where a magnetically induced anisotropy is induced parallel to the measurement direction, increasing the relative complex permeability of the sample.

[0116] It is well accepted that high permeability is associated with rapid re-arrangement of magnetic domains in soft magnetic materials. It is also well known that this rapid change in domain structure is associated with the formation of larger eddy currents than the slowly rotating domain structures that are typical for low permeability materials. Therefore, the reduction in core loss seen in the TFA sample versus the NFA sample in Figure 3 and Table 2 is the result of reduced eddy current losses due to the reduction in the permeability of the material made possible by the transverse magnetic field annealing process.

[0117] Also, from Table 2, the rapid annealed (Fe 0.8 Co 0.2 ) 87 B 13 It is observed that the core loss of the samples is significantly lower compared to Fe-3 wt.% Si steel, with or without an applied magnetic field.

[0118] Table 2: Rapid annealing at 50 Hz, 400 Hz, and 1000 Hz with a maximum magnetization of 1.5 T (Fe 0.8 Co 0.2 ) 87 B 13 AC core loss

[0119] [Table 2]

[0120] Example 5 Nanocrystalline (Fe 1-x Co x ) 87-y-a-z B y Cu z M a (wherein x = 0.1 to 0.4, y = 13 to 14, z = 0 to 1, and a = 0 to 8) The effect of adding M was also investigated. Precursor amorphous ribbons having the same standard compositions as those shown in Table 3 below were produced by melt spinning (planar flow casting method) in an Ar atmosphere.

[0121] Ribbons with thicknesses of approximately 10–15 μm and widths of 1–12 mm were obtained. The ribbons were placed in 20 μm thick Cu foil packets and subjected to ultra-rapid annealing in an Ar atmosphere. These packets were then compressed between two preheated Cu blocks (length 150 mm, width 50 mm) with a force of 950 N for 0.5 s using a pneumatic cylinder and an automatic timing mechanism.

[0122] XRD using a CoKα source confirmed the formation of an amorphous phase after the casting process, with a volume fraction of 80% or more. XRD also confirmed the formation of bcc Fe-Co or Fe-Co-Ni (when Ni is present) crystalline phases embedded in the residual amorphous phase. The saturation magnetic polarization (J s =μ 0 The magnetization was evaluated at 0.8 MA / m and 295 K using a vibrating sample magnetometer (VSM) BHV-35H (Riken Denshi Co., Ltd.). c The evaluation was carried out at 295K using a hysteresis loop tracer BHS-40DC (manufactured by Riken Denshi Co., Ltd.).

[0123] Table 3 shows the (Fe 1-x Co x ) 100-y-a-z B y Cu z M a H of various rapidly annealed nanocrystalline soft magnetic materials with compositions c and J s Shows.

[0124] Table 3: The iron-containing compounds investigated in this study (Fe 1-x Co x ) 100-y-a-z B y Cu z M a Properties of the composition (wherein M=Nb, Mo, Ta, W, Ni, or Sn)

[0125] [Table 3]

[0126] The addition of element M is mainly to improve glass-forming ability. c However, the addition of all elements M reduces J s It is also observed that the y and z elements are added to replace the ferromagnetic Fe and Co.

[0127] Example 6 Figures 15 to 17 show (Fe 0.8 Co 0.2 ) 87 B 13 Another magnetic characterization of the sample is shown.

[0128] FIG. 15 shows the (Fe 0.8 Co 0.2 ) 87 B 13 The figure shows the coercivity versus annealing temperature measured for the samples. The samples were rapidly annealed by clamping them between preheated copper blocks for 0.5 seconds. The figure also shows that for a minimum coercivity of 3.4 A / m, the optimal annealing temperature (T op ) is shown.

[0129] Figure 16 shows the results obtained at the optimal annealing temperature (Fe 0.8 Co 0.2 ) 87 B 13 The direct current (DC) hysteresis loop measured for the sample is shown. The coercivity was 3.4 A / m. Independent measurements by VSM showed that the saturation polarization of the sample was 2.02 T.

[0130] FIG. 17 shows that in the presence of an applied transverse magnetic field (Fe 0.8 Co 0.2 ) 87 B 13 The figure shows the effect of rapid annealing of (Fe 2+ 0.01 μm) after rapid annealing at 753 K (i.e., 480 °C) using a preheated copper block and subsequent cooling, either with or without the applied magnetic field. 0.8 Co 0.2 ) 87 B 13 The effect of magnetic field annealing on the shape of the DC hysteresis loop of the annealed ribbon can be understood. It is observed that the effectiveness of the magnetic field annealing method is reduced when the annealed ribbon is cooled without the influence of the magnetic field, compared to cooling the ribbon under the influence of the magnetic field. Therefore, to obtain optimal magnetic properties, if magnetic field annealing is employed, a magnetic field should be present during all stages of the annealing. The relevant parameters are outlined in the table below.

[0131] Table 4: Relevant parameters for the data shown in Figure 17

[0132] [Table 4]

[0133] Example 7 In Fig. 18, (Fe 0.8 Co 0.2 ) 86 B 13 Cu 1 The magnetic characterization of the samples is shown. In particular, the data shows the magnetic properties of the rapidly annealed (Fe 0.8 Co 0.2 ) 86 B 13 Cu 1 The core losses measured at 50 Hz, 400 Hz, and 1000 Hz for the 3 wt.% iron-silicon steel compared to the sample. The data show that at all frequencies and magnetization levels tested, the (Fe 0.8 Co 0.2 ) 86 B 13 Cu 1 It can be seen that the core loss of is significantly lower than that of iron-silicon steel.

[0134] Throughout this specification and the claims which follow, unless the context requires otherwise, the word "comprise" and variations such as "comprises" and "comprising" should be understood to mean the inclusion of a stated integer or step or group of integers or steps, but not to exclude any other integers or steps or group of integers or steps.

[0135] Reference in this specification to any prior publication (or information derived therefrom) or to any known matter is not, and should not be considered, as an admission or acknowledgement, or in any way an indication that the prior publication (or information derived therefrom) or known matter is part of the common general knowledge in the field to which this specification pertains.

Claims

1. Formula (Fe １－ｘ Co ｘ ) １００－ｙ－ｚ－ａ B ｙ Cu ｚ M ａ (In the formula, x = 0.1 to 0.4, y=10 to 16, z=0 to 1, a=0 to 8, and M=Nb, Mo, Ta, W, Ni, or Sn), An alloy having crystal grains with an average grain size of 30 nm or less.

2. 2. The alloy of claim 1, wherein x ranges from about 0.2 to about 0.

3.

3. The alloy of claim 1 or 2, wherein z is in the range of about 0.2 to 1.

4. The alloy of any one of claims 1 to 3, wherein z and a are both 0.

5. Magnetization saturation (J ｓ 5. The alloy of claim 1, wherein ZnO is at least 2T.

6. The alloy according to any one of claims 1 to 5, wherein the average grain size of the crystal grains is between 10 nm and 30 nm.

7. 1. A method for producing an alloy, the method comprising: Formula (Fe １－ｘ Co ｘ ) １００－ｙ－ｚ－ａ B ｙ Cu ｚ M ａ (In the formula, x = 0.1 to 0.4, y=10 to 16, z=0 to 1, a=0 to 8, and Preparing an amorphous alloy represented by M=Nb, Mo, Ta, W, Ni, or Sn; and heating the amorphous alloy at a heating rate of at least 200° C. / sec. The method according to claim 1,

8. The method of claim 7 , wherein heating the amorphous alloy comprises exposing the alloy to a magnetic field.

9. 9. The method of claim 7 or 8, wherein heating the amorphous alloy comprises exposing the alloy to a rotating magnetic field in the range of at least 0.3 kA / m.

10. 10. The method of any one of claims 7 to 9, wherein the step of heating the amorphous alloy comprises exposing the alloy to a magnetic field whose orientation and / or magnitude varies in the range of about 1 Hz to about 3,000 Hz.

11. A method according to any one of claims 8 to 10, wherein after the heating step, the alloy is cooled in the presence of the magnetic field.

12. The method of any one of claims 7 to 11, wherein the amorphous alloy is heated to an annealing temperature in the range of about 350°C to about 650°C.

13. The method of any one of claims 7 to 12, wherein the amorphous alloy is heated to a predetermined annealing temperature and held at the annealing temperature for about 0 to about 80 seconds.

14. The method of any one of claims 7 to 13, wherein the amorphous alloy is in the form of a thin ribbon having a thickness ranging from about 5 μm to about 15 μm.

15. The method according to any one of claims 7 to 14, wherein the step of heating the amorphous alloy is carried out by sandwiching the alloy between preheated blocks at a pressure of at least about 3 kPa.

16. A method according to any one of claims 7 to 14, wherein the step of heating the amorphous alloy is carried out by passing the alloy between preheated rolls.

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