Patent Description:
In one aspect, the present invention relates to manganese-aluminum (MnAI) alloys that have a high content of τ phase or ε phase, and methods for their production.

In one aspect, the present invention relates to MnAI alloys having a high content of τ phase and improved magnetic properties, and methods for their production.

In another aspect, the present invention relates to methods for preparing MnAI alloys, and to methods for producing MnAI alloys, comprising the ε phase in high purity.

In yet another aspect, the invention relates to methods for converting the ε phase of MnAI alloys into the τ phase.

Magnetic materials are used in many industrial applications and devices, such as speaker systems, earphones, direct current electric motors in battery (primary or secondary cell) powered devices, wind generators etc., just to mention a few.

Such magnetic materials are however expensive, and furthermore generally use rare earth metals. Accordingly, these conventional materials lead to both high costs and high environmental burden. There is thus a desire for a replacement of these rare earth metal based magnetic materials.

The low cost, rare earth free τ-phase of MnAI alloys has magnetic properties, and is currently considered to have potential to replace currently used magnets, such as bonded Nd<NUM>Fe<NUM>B rare earth permanent magnets. MnAI based magnetic materials (~<NUM>-<NUM> at. % Mn) with the L1°-type structure (τ-phase) have in particular great potential to become a high performance permanent magnet material at low cost, if a suitable processing route could be developed.

The room temperature stable phases in the MnAI phase diagram include γ2 (Al8Mn5), β (Mn) and the ferromagnetic τ-phase. The most common approach for synthesizing τ-phase MnAI alloys is by rapid quenching the high-temperature ε phase, followed by a heat treatment between <NUM> and <NUM>. However, the τ-phase is metastable (see <FIG>), and it is difficult to obtain MnAI alloys containing the τ-phase in high purity. So far, various heat treatments have been employed in order to obtain and preserve the τ-phase.

Oftentimes the τ-phase decomposes spontaneously and in particular upon further processing, and it is difficult to obtain pure τ-phase of high quality. This is in particular a problem if one wants to obtain powders of the alloy suitable for a permanent magnet alignment process, as this generally requires the powders to have a regular round shape and simultaneously high τ-phase purity. The powders should also be very fine, and most preferably there is only one grain (crystal) per particle, as this improves alignability and hence magnetic remanence. Yet, currently used processing methods for producing a powder form of the alloy (such as ball milling) distort the structure and partially destroy the τ-phase.

As derivable from <FIG>, the τ-phase can only exists between about <NUM> to <NUM> atom% being Mn, the remainder being Al and unavoidable impurities. Further, as indicated by the hyperbolic dotted line in <FIG>, the likelihood to form the τ-phase in binary compositions of Mn and Al is greatest at <NUM> - <NUM> atom% Mn. At the same time, the theoretical magnetic moment of the material is higher for compositions containing less Mn.

It has also been reported that the presence of dopants such as carbon can stabilize the τ-phase and facilitates the formation of an alloy resulting in an improved saturation magnetization, Ms. Further, Ms has also been reported to increase with carbon. These advantages are in part outbalanced by disadvantages, such as that the Curie temperature and anisotropy field are lower than those of undoped alloys.

In view of the above, most of the research up do date is devoted to MnAI alloys having a Mn content of <NUM> atom % or less, and containing no or only very little dopant, in order to obtain a material that maintains a τ-phase that has high magnetic moment as well as highest possible Curie temperature.

As one example, <NPL>) the preparation of τ-phase MnAI alloys having a composition Mn<NUM>Al<NUM>. Reported is here also on the attempt to prepare τ-phase by strip-casting compositions of Mn<NUM>Al<NUM> and Mn<NUM>Al<NUM>. While the τ-phase could be successfully synthesized for Mn<NUM>Al<NUM> and Mn<NUM>Al<NUM>, it is reported that the strip-casted Mn<NUM>Al<NUM> tends to form other samples, i.e. γ2 and β. All materials showed increasing amounts of γ2 and β after grinding. The material having the highest saturation magnetization (Ms) was Mn<NUM>Al<NUM> strip, i.e. prior to grinding, exhibiting <NUM> emu/g (at <NUM> Tesla).

Fang et al reported that the τ-phase can be obtained by directly casting Mn<NUM>Al<NUM> and (Mng<NUM>Al<NUM>)<NUM>C<NUM> (see <NPL>) using a drop synthesis method involving a rapid cooling from <NUM> of the melt to room temperature in <NUM> minutes. Here, the highest Ms of <NUM> emu/g (at 9T) was obtained in the system of (Mn<NUM>Al<NUM>)<NUM>C<NUM>.

The document further teaches that the crystalline order in the MnAI system is crucial for the magnetization, and that surplus Mn yields a "built in" disorder since on average <NUM>% of the Mn atoms occupy the Al sites and interact antiferromagnetically with the Mn on the (<NUM>,<NUM>,<NUM>) sites. This is illustrated in <FIG>. Here, the lines for "<NUM>%, "<NUM>%" and "<NUM>%" illustrate how the magnetic moment is influenced by the amount of misplaced Mn, i.e. where <NUM>, <NUM> or <NUM>% of the Al sites are replaced by Mn atoms, as compared to the ideal case. The ideal case for a <NUM>%Mn/<NUM>%AI alloy is that <NUM>% of the Al-sites are occupied by Mn atoms. Thus, for this composition <NUM>% mismatch means that <NUM>% Al-sites are taken by Mn.

Fang further suggests that the stability region of the β phase extends further than previously expected, extending up to <NUM> at% Al (corresponding to <NUM>% Mn), in particular for materials that are processed subsequently by mechanical treatments such as high energy ball milling.

A comparison of τ phase MnAI prepared via different routes was provided by <NPL>, also using materials having the stoichiometric composition Mn<NUM>Al<NUM>. The materials were either gas atomized from the melt under argon, casting into ribbons using melt spin casting under argon, or pouring the molten alloy onto the face of a cooled copper dish rotating at high speed under argon, thereby forming alloy droplets.

The gas atomized material contained about <NUM>% of the ε phase, and β and γ2 were also present. The document further states that annealing of milled powder (i.e. the material obtained after gas atomization of the melt and additional milling thereof using a water-cooled Union Process 1SD Svegari attritor) transforms the ε phase to the τ phase, yet the equilibrium γ2 and β phase are still present in essentially the same amounts as before annealing. It is further described that in milled materials the cooling rates of powders appears to be so high that the formation of the high temperature phase does not occur, and that β and γ2 form directly from the melt. It is concluded that the τ phase would not form this way, since it transforms from the ε phase and not directly from the melt.

(<NUM>) presents a neutron diffraction study of the [tau]-phase of carbon-doped Mn-Al alloy (<NPL>.

(<NUM>) investigates the effects of ambient aging upon upon the magnetic and structural properties of mechanical alloy Mn<NUM>Al<NUM>-xCx (x = <NUM>, <NUM> and <NUM>) alloy powders (<NPL>).

et al (<NUM>) studies structural and magnetic properties of nanostructured Mn-Al-C magnetic materials (<NPL>).

It is an object of the present invention to provide a material having a high content of ε phase. This material is believed to be a good precursor for preparing the τ phase, as the τ phase forms from the ε phase.

It is a further object of the present invention to provide a method for preparing a MnAI alloy having a high content of the ε phase, i.e. to prepare a MnAI alloy having a high purity of the ε phase.

It is a further object of the present invention to provide a MnAI alloy having improved magnetic properties. In a related aspect, it is an object of the present invention to provide a MnAI alloy having a high content (or high purity) and/or quality of the τ phase, in particular (but not limited to) MnAI alloy materials in powder form. Here, the term "quality" in particular includes the aspects of homogeneity of the microstructure of the τ phase as well as a low mismatch of the site occupancy (i.e. ideal or close-to-ideal site occupancy by Mn and Al in the crystal lattice).

The present invention also aims at providing a method that is capable of providing such MnAI alloys, in particular a method that is industrially useful.

It is a further object of the present invention to provide an improved method for converting the ε phase into the τ phase.

It is yet another object of the invention to provide a MnAI alloy composition able to form τ or ε phase, which is relatively robust towards further processing steps and/or towards variations in the synthesis conditions such as cooling rate from the melt, and still leads to high purity of ε phase, and subsequently τ phase formed therefrom. Herein, the τ phase shows a good homogeneous microstructure, thereby allowing to obtain a material with improved magnetic properties (in particular Ms).

It is yet another object of the present invention to provide an improved synthesis route that includes a proper composition of the MnAI alloy as well as a preparation process including a suitable quenching (cooling) rate from the melt, ensuring a microstructure suitable for preparing high purity ε and subsequently τ phase. The resulting material is in particular suitable for later micronization steps, which are known to show the tendency to deteriorate the τ phase.

It is a further object of the invention to provide an MnAI alloy, preferably in particle form, showing both high magnetic saturation (Ms) and high coercitivity (Hc).

It is a further object of the invention to provide a method able to homogenize the microstructure of an MnAI alloy.

It is a further object of the invention to provide for a micronization process that can reduce the particle size and reduce grain size, and which allows the microstructure to become randomized and homogenized while also allowing for suitable magnetic properties.

In yet another aspect, it is an object of the present invention to provide means for obtaining an MnAI alloy containing τ phase and a having high magnetic saturation before a micronization process.

All aspects of the present invention aim at providing materials and/or methods for their production that either have favorable magnetic properties themselves, or which are favorable precursor materials for preparing materials having favorable magnetic properties.

The present invention solves one or more of the above objects specifically by the following aspects:.

Further aspects and embodiments of the present invention will become apparent from the following detailed description.

In the present invention, all physical parameters are measured at room temperature (<NUM>) and at atmospheric pressure (<NUM><NUM>Pa), unless indicated differently.

As used herein, the indefinite article "a" indicates one as well as more than one and does not necessarily limit its reference noun to the singular.

The term "about" means that the amount or value in question may be the specific value designated or some other value in its neighborhood, generally within a range of ±<NUM>% of the indicated value. As such, for instance the phrase "about <NUM>" denotes a range of <NUM> ±<NUM>.

The term and/or means that either all or only one of the elements indicated is present. For instance, "a and/or b" denotes "only a", or "only b", or "a and b together". In the case of "only a" the term also covers the possibility that b is absent, i.e. "only a, but not b".

The term "comprising" as used herein is intended to be non-exclusive and open-ended. A composition comprising certain components thus may comprise other components besides the ones listed. However, the term also includes the more restrictive meanings "consisting of" and "consisting essentially of". The term "consisting essentially of" allows for the presence of up to and including <NUM> weight%, preferably up to and including <NUM>% of materials other than those listed for the respective composition, which other materials may also be completely absent.

Whenever ranges are indicated in the present invention, be it by the use of a hyphen as in "<NUM> - <NUM>", by using the word "to" instead of a hyphen, or by using the word "between", such as in "between <NUM> and <NUM>", the respective end values of the range are included, so that e.g. the range of "<NUM> - <NUM>" includes the values of <NUM> and <NUM>. These expressions however also cover the entire set of values within that range, so that the terms are used as abbreviation of the term "<NUM> or more, but <NUM> or less".

The percentages relating to alloy compositions are atom-%, unless specified differently. All other percentages are by weight, unless indicated differently.

A particle size is typically defined as D50, unless stated otherwise. D50 defines that <NUM>% by weight of the particles have the indicated size or smaller, as determined by a laser light scattering method.

During the manufacture of the alloy of the present invention and in the method for producing an alloy of the present invention, certain heat treatments within a specified temperature range for a time within a specified range may be conducted. An example is a case wherein "a heat treatment at a temperature of <NUM> to <NUM> for a time of <NUM> to <NUM> hours" means that the alloy is kept at a temperature within the range of <NUM> to <NUM> for a time of <NUM> to <NUM> hours. The time used for heating up to the lower limit (e.g. <NUM>) and cooling below the lower limit is not taken into account for the specified duration of the heat treatment. Further, the upper limit is typically not exceeded in such heat treatment steps.

The numerical values of characteristics and properties of the materials employed or produced in the present invention relate to the values obtained by using the instruments and conditions employed in the Examples. In case of differences or discrepancies, the following conditions prevail:
The specified saturation magnetization values are in Emu/g as measured at an applied field of 9T (using a Physical Property Measurement System, PPMS, as available from Quantum Design, Inc. at room temperature <NUM>° and <NUM> atm), unless specified differently.

For saturation magnetization data measured at an applied field of <NUM>. 8T (using a Vibrating Sample Magnetometer, VSM), the values have for comparison purposes been converted to correspond to a 9T-measurement, based on a conversion factor obtained from previous measurements where both instruments have been used on the same sample.

The coercive force is specified in kA/m assuming a material density of <NUM>/m<NUM>. This density may also be assumed for any other desired conversion, such as for converting a volume-based parameter into a weight-based parameter.

Whenever in the present invention reference is made to the content of a crystalline phase in %, the respective value refers to the value obtained by an XRPD analysis with Rietveld refinement as described in example section <NUM>.

In the present invention, the alloys are represented by e.g. formula (I) with the composition (MnxAly)Cz, wherein x = <NUM> to <NUM>, y = <NUM> to <NUM>, x + y = <NUM>, and z = <NUM> to <NUM>. This is equivalent to the representation (Mn<NUM>. Hence, there are z atoms of C relative to <NUM> atoms of the total of Al and Mn. The same applies to formula (II).

As outlined in the passages above, most of the attempts for preparing MnAI alloys comprising a high amount of τ phase have focused on materials that have a Mn content of <NUM>% or less. Further, while there have been attempts to stabilize the τ phase by addition of carbon as a dopant, these attempts have also been made with MnAI alloys having such a Mn content. Little attention has so far been paid to producing a MnAI alloy showing a high content of Mn (e.g. <NUM> at. -% or more, such as between <NUM> and <NUM> at. -%, relative to the total of Mn and AI).

The present invention is in one aspect based on the finding that a key step in the preparation of an MnAI alloy having a high content of τ phase is the preparation of an MnAI alloy having a high content of ε phase. The present invention furthermore provides a method for preparing such MnAI alloys having a high content of ε phase (e.g. <NUM>% or more, such as <NUM>% or more, <NUM>% or more, or <NUM>% or more) by employing specific process conditions other than those used or suggested in the prior art, as will be explained in more detail below.

The present invention is in another aspect based on the surprising finding that MnAI alloys having a high content of τ phase can be prepared from compositions of formula (I), i.e. compositions consisting of relatively high amounts (<NUM> - <NUM> at% of the total of Al and Mn) of manganese and low amounts (<NUM> - <NUM> at% of the total of Al and Mn) of aluminum, in combination with a specific amount of carbon of <NUM> to <NUM> at-%. Compositions satisfying formula (I) can surprisingly be processed such as to contain a high content of τ phase, and may thus provide materials with improved magnetic properties. The carbon content may preferentially be chosen such as to be linked to the amount of Mn, such that the molar ratio of Mn/C is in the range of <NUM> to <NUM>, more preferably <NUM> to <NUM>. With these preferably and more preferably carbon contents, the alloy is experimentally saturated with carbon dissolved in the alloy, so that the τ phase can be stabilized most effectively.

The present invention is in a further aspect based on the surprising finding that particles of an MnAI alloy having favorable magnetic properties can be prepared by a process involving specific milling conditions at low temperature and/or specific heating conditions (also referred to as "flash heat treatment" in the following).

The features and aspects of the present invention will be described separately in the following more detailed description:.

In a first aspect, the present invention relates to alloys of formula (I) as defined below:.

Herein, the term "unavoidable impurities" includes any element other than Al, Mn and C, and typically denotes impurities that originate from the raw materials used for producing the powder or that are introduced during the manufacturing process. The total content of such impurities is typically <NUM> % by mass or less, preferably <NUM> % by mass or less, and further preferably <NUM> % by mass or less, relative to the total mass of the alloy.

Compared to prior art alloys, the alloy of formula (I) combines a relatively high manganese content with a carefully defined carbon content. Without wishing to be bound by theory, it is believed that the interstitial positions in the ferromagnetic τ phase are occupied by carbon, making the τ phase more thermodynamically stable or favorable. Further, again without wishing to be bound by theory, the combination of carbon with manganese in an amount of <NUM> -<NUM> at. -% (relative to the total of Al and Mn) allows obtaining a more homogeneous microstructure of the MnAI alloy in both ε and τ phase (see <FIG>). This homogeneous microstructure may at least be partially be preserved during possible further processing steps, so that the alloy of formula (I) provides an improved precursor material for final applications.

Again without wishing to be bound by theory, addition of carbon at the octahedral interstitial sites (½, ½, <NUM>) is believed to be an effective way of stabilizing the tetragonal structure of the τ phase with an elongation along the c-axis. While carbon reduces the Curie temperature, it increases the saturation magnetization (Ms) with a larger resultant magnetic moment. The increase in stability by carbon doping is believed to occur because the interstitial atoms inhibit the diffusion of the Mn and Al atoms.

This finding is truly surprising, as it was generally considered that carbide precipitates will form at higher carbon concentrations that will act as nucleation sites for the equilibrium phases when the C addition exceeds the solubility limit. It has now surprisingly been found that at higher manganese contents it is possible to dissolve even more carbon, and thus, it is possible to increase the stability of the τ phase. Furthermore, the workability is also improved due to the small carbon atoms, which are believed to relieve internal lattice stresses.

Suitable magnetic properties in a material having a high τ phase content have been obtained when the value of z, representing the number of carbon atoms relative to <NUM> (Al+Mn) atoms, is <NUM> or higher, preferably <NUM> or higher, more preferably <NUM> or higher. At the same time, the maximum value of z is <NUM> or less, preferably <NUM> or less, more preferably <NUM> or less.

The value of x, representing the number of manganese atoms relative to <NUM> (Al+Mn) atoms, is <NUM> or higher, preferably <NUM> or higher, more preferably <NUM> or higher. At the same time, the value of x is <NUM> or less, preferably <NUM> or less, more preferably <NUM> or less.

The value of y, representing the number of aluminum atoms relative to <NUM> (Al+Mn) atoms, is chosen such that the sum of x + y = <NUM>. y is <NUM> or less, preferably <NUM> or less, mor preferably <NUM> or less, and is at the same time <NUM> or higher, <NUM> or higher, more preferably <NUM> or higher.

The value of z, representing the number of carbon atoms relative to <NUM> (Al + Mn) atoms, is <NUM> to <NUM>. The lower limit of z is thus <NUM> or higher, such as <NUM> or higher, <NUM> or higher, <NUM> or higher or <NUM> or higher. The upper limit is <NUM> or lower, but can also be <NUM> or lower, <NUM> or lower, or <NUM> or lower. These lower and upper limits can be combined in any way, and in preferred embodiments the value of z is <NUM> to <NUM>, more preferably <NUM> to <NUM>.

The values of x, y and z can be adjusted by appropriately mixing suitable amounts of the starting compounds for the alloy production, which are typically elementary manganese, elementary aluminum and graphite. In the final alloy, the values of x, y and z can be determined by analyzing the relative content of the metals by a suitable technique, such as ICP-AES (Inductively Coupled Plasma - Atomic Emission Spectroscopy and LECO elemental analysis for C, N, P and S, or Spark Optical Emission Spectroscopy.

Given the above relationship between the amount of Mn and C, the ratio of x to z (x/z) is preferably <NUM> or higher, more preferably <NUM> or higher or <NUM> or higher, and still more preferably <NUM> or higher, and preferably simultaneously <NUM> or less, more preferably <NUM> or less or <NUM> or less, and still more preferably <NUM> or less. By adjusting the relative amounts of Mn and C such that x/z falls within these ranges, maximum stabilization of the τ phase and at the same time good magnetic properties can be maintained. Without wishing to be bound by theory, it is believed that with a lower Mn/C-ratio, such as lower than <NUM>, the risk of excess carbon being present that might form e.g. Mn<NUM>AlC upon heat treatment may not be completely excluded. This risk is further minimized if the ratio of Mn/C, i.e. (x/z), is adjusted to satisfy the above-mentioned preferable and more preferable ranges.

While it is possible to form the τ phase directly from the melt, it is common to first form the ε phase, followed by formation of the τ phase by suitable heat treatment.

Hence, in a preferred aspect, the alloy of formula (I) has a content of the ε phase of <NUM>% or more, more preferably <NUM>% or more, further preferably <NUM>% or more, determined by a the XRPD method using Rietveld refinement as described in Example <NUM>.

While the method for preparing the alloy of formula (I) is not specifically limited, the alloy of formula (I) can suitably and preferably be prepared by the method described below for the alloys of formula (II). In this case, particularly preferred characteristics can be obtained, yet the method is also applicable to other carbon-containing alloys of formula (II). Formula (I) is thus a preferred sub-set of the alloys of formula (II).

As is well known, the synthesis parameters, i.e. melt temperature, quenching- or cooling rates of the melt, have an impact on the homogeneity of the microstructure, and suitable parameters can generally be chosen by a skilled person based on common general knowledge and/or be turning to the preparation conditions described in the examples. It is therefore preferred to choose synthesis conditions and subsequent process conditions for the material in such way that inhomogeneity is as far as possible eliminated.

The melt temperature is typically in the range of <NUM> or higher, preferably <NUM> or higher, more preferably <NUM> or higher, such as <NUM>. The melt should be homogeneous, so that the melt is held for some time (e.g. <NUM> - <NUM> minutes or even longer) in the molten state at high temperature in order to allow a good mixing and diffusion of the components.

Subsequently, the melt is processed to form a solidified material. Depending on the cooling rate either ε- or τ-phase can be formed directly from the melt. Normally, higher cooling rates than about <NUM>/s will result in ε-phase, while slower cooling rates at below ~<NUM> could result in the magnetic τ-phase directly.

Hence, where the material obtained directly from the melt is desired to have a large content of the ε phase, rapid cooling from the melt is preferred. A high cooling rate is preferred not only to form the ε phase in high purity, but also to ensure a homogeneous composition as well as microstructure (minimize segregates) inside the solidified material. In order to avoid formation of the more thermodynamic stable phases (in particular β and γ2) as well as to obtain an even distribution of the raw materials, the average cooling rate between the melt temperature and approximately <NUM> is preferably <NUM><NUM>°C/s or higher, more preferably <NUM><NUM>°C/s or higher, still more preferably <NUM><NUM> C/s or higher. Furthermore, in order to favor a direct formation of an alloy having high content of the τ phase, and suppress formation of e.g. β and γ2-phases, the average cooling rate below approximately <NUM>, and above approximately <NUM> is between <NUM>/min and <NUM>/min, preferably between <NUM>/min and <NUM>/min, more preferably about <NUM>/min.

If the quenching rate and/or the mixing of the components in the melt prior to atomization has been insufficient, the material may exhibit regions of different composition caused by microstructural variations with different crystallinity, and may then also show a relatively large difference in composition between the regions/segregates. The regions of lower crystallinity exhibit a significantly lower Mn/Al-ratio and, thus also lower carbon content due to the decreased solubility of carbon in the presence of less manganese. Larger differences in composition between these region/segregates has been found to be undesirable, as it may lead to cause inhomogeneities on a microstructural level even in materials that have undergone further processing steps.

Further, the cooling rate in a "natural cooling process" (such as described e.g. by <NPL>) depends on various factors, such as shape and volume of the alloy cooled (e.g. in ingot form), the temperature of the atmosphere, etc. In this article, the cooling process was described to exhibit a cooling rate from <NUM> to room temperature in about <NUM> minutes, which equates to an average cooling rate of about <NUM>/s.

This shows that a natural cooling process, of e.g. an alloy cast from the melt to form an ingot, is typically too slow to lead to the formation of a pure ε phase of homogeneous microstructure, but still too fast to allow the direct formation of τ phase. If a material having a high content of ε phase is desired, additional measures need to be taken in order to achieve higher cooling rates. These include in particular measures for forming droplets from the melt in order to increase surface area relative to the mass of the melt, thereby increasing cooling rates, and forced cooling, e.g. bringing the molten alloy (preferably in droplet form) into contact with a cooling medium such as a cooling substrate, a cooling gas or a cooling liquid, e.g. water. Suitable methods thus include gas or liquid atomization, melt spinning or spin casting. Gas atomization is preferred, as it is scalable and produces particles having an almost round or spherical shape that allow good orientation in a magnetic field, and which require little or no further mechanical processing such as milling in order to obtain a powder material suitable for further processing and final application after conversion of the ε phase to the ultimately desired τ phase. Further, gas atomized particles show good homogeneity and little crystal defects, as gas atomization is believed to exert little stress on the internal structure of the alloy during cooling, thereby facilitating the formation of a relatively homogeneous crystal structure with large single crystalline domains.

The cooling of the alloy (e.g. during gas atomization) from the melt is preferably conducted in inert gas atmosphere, such as argon or nitrogen, in order to avoid the formation of oxides and to maintain a homogeneous structure. Cooling the liquid droplets into a tank of water can also be applied to reach even higher cooling rates, however, with the expense of higher oxygen contents.

By providing a melt of the MnAI alloy of formula (I) and performing a cooling process under inert gas atmosphere or water employing a cooling rate rapidly enough to avoid substantial formation of impurity phases (e.g. β, γ2) or the τ phase, a material that is homogeneous and rich in ε phase can be obtained. This material, such as obtained by gas atomization of the melt and cooling under inert gas or water conditions, is an excellent precursor material for the further processing steps, in particular the flash heat treatment step for transforming the ε phase into the τ phase. Care should however be taken to avoid extreme cooling rates of e.g. higher than <NUM><NUM> or <NUM><NUM> °C/s, as this might disturb the crystallization process and may lead to a lower purity of the desired ε phase.

It has further been found that the alloy composition of formula (I) is less sensitive to the synthesis conditions, in particular as regards the quench rate from the melt, and still allows obtaining a homogeneous microstructure under conditions that lead to less homogeneous/more heterogeneous microstructure for alloys that do not conform to formula (I), due to a lower Mn content. This is illustrated in the Figures, in particular in <FIG>, which will also be explained in more detail in the following description of step a. It is however noteworthy that the following effects are believed to occur independently of the specific step of gas atomization (step a), i.e. also for alloys of formula (I) that are produced without a step of gas atomization from the melt, but by different techniques.

<FIG> shows on the top the diffractogram (XRPD) of an alloy having the composition Mn<NUM>Al<NUM>C<NUM>, and in the bottom of an alloy of the composition Mn<NUM>Al<NUM>C<NUM>, both directly after gas atomization under identical conditions. As can be seen from the XRPD data, the sample at the top (falling within the scope of formula (II), but not formula (I)) shows a reasonably good purity, yet some impurities (γ2, β) are also present as derivable from the minor peaks around <NUM> - <NUM>° in 2θ. Such peaks are also present in the lower figure showing the diffractogram of the alloy having a composition of formula (I), yet to a lesser extent.

Interesting to note is also the difference in structural homogeneity, as represented by the SEM pictures shown in <FIG>). Here, both samples show a microstructure composed of grains (dark in <FIG>)) and a matrix phase (bright in <FIG>)). Both of the grains and the matrix are present mainly in the ε phase, yet differ in respect of their composition. The grains are formed of a phase ε<NUM> having a high Mn and C content, while the matrix is formed of a phase ε<NUM> having a high Al content. The matrix areas show nanocrystallinity, i.e., have a lower ability to crystallize.

Interestingly, in the alloy conforming to formula (I), the difference in composition between the matrix and the grains is smaller as compared to the alloy conforming to formula (II) (but not formula (I)). The sample conforming to formula (I) has thus a more homogeneous microstructure. Without wishing to be bound by theory, it is believed that this difference in the compositional difference between the ε<NUM> and ε<NUM> phase is caused by the combined presence of carbon and relatively high amounts of Mn, such that the Mn/C ratio is within the preferred and more preferred ranges described above. Apparently, the presence of carbon in the required amount helps avoiding the formation of segregates, leading to a more homogeneous structure.

This homogeneity of the microstructure of the as-synthesized ε phase also influences the microstructure of the subsequently formed τ phase. This is illustrated in <FIG> and d). The sample conforming to formula (II) (but not formula (I)) having the greater inhomogeneity of the microstructure produces, upon heating at <NUM> in vacuum for <NUM> minutes, a microstructure showing high inhomogeneity consisting mainly of τ, γ2, and β (see <FIG>). Conversely, the sample conforming to formula (I) leads to a much more homogeneous microstructure of predominately τ-phase (see <FIG>), which in turn allows obtaining a higher Ms.

This difference in structure is also observable by a DSC analysis. <FIG> shows the DSC analysis of a sample conforming to formula (II) ((Mn<NUM>Al<NUM>)<NUM>C<NUM>), but not formula (I). The DSC shows an exothermic peak around <NUM> that represents the formation of the τ phase. Upon further heating, there are two endothermic peaks around <NUM> and <NUM>, which represent the re-formation of the ε phase from the τ-phase. The occurrence of two different peaks in the DSC shows that the τ-phase formed actually consists of at least two different phases, one converting back to the ε-phase at around <NUM> and the other one converting back to the ε phase at around <NUM>. This is in good agreement with the structure observed in <FIG>.

The DSC graph of a sample conforming to formula (I) ((Mn<NUM>Al<NUM>)<NUM>C<NUM>) is shown in <FIG> b. Here, there is one exothermic peak around <NUM> representing the formation of the τ phase, and one peak around <NUM> representing the re-formation of the ε phase. Notably, there is no peak at higher temperatures as in <FIG>, showing that the microstructure is more homogeneous and essentially consists of one τ-phase only. This is good agreement with the structure observed in <FIG>.

It follows that the alloy composition of formula (I) having the right balance between Mn and C (i.e. having the suitable Mn/C ratio) allows obtaining a more homogeneous ε phase and a more homogeneous τ phase as compared to alloys having lower Mn content and/or in an incorrect carbon amount, i.e. with an unsuitable Mn/C ratio. The alloy of formula (I) is thus in one embodiment characterized by a DSC graph showing one endothermic peak of <NUM> J/g or more, preferably <NUM> J/g or more and further preferably <NUM> J/g or more, such as <NUM> J/g or more, in the temperature range of <NUM> - <NUM> (or <NUM> - <NUM>), and either no endothermic peak in the temperature range of <NUM> - <NUM> or an endothermic peak in the in the temperature range of <NUM> - <NUM> of <NUM> J/g or less, preferably <NUM> J/g or less, and more preferably <NUM> J/g or less, determined by a DSC method using a heating rate from room temperature to <NUM> C° at <NUM>/min. Thereby, a more homogeneous τ-phase can be formed. While it is also possible to obtain the ε phase outside the composition of formula (I), this ε shows the tendency to not only form the τ-phase upon proper heat treatment, but to also form the thermodynamically more stable β- and <NUM>-phases. Within the scope of formula (I), thus an ε-phase can be obtained that is more suitable for obtaining a more pure and more homogeneous τ-phase.

An MnAI alloy of formula (I) can then, after formation of the ε and/or the τ phase, be processed further. For instance, the τ phase MnAI alloy of formula (I) (e.g. obtained by first forming the ε phase and then converting the ε phase into the τ phase by a suitable heat treatment, as described for instance below for step b. , or by other suitable treatment) can be subjected to a milling process. This milling process can be the low temperature cryo milling step c. , but can also be another suitable milling process known in the art, such as ball milling at another temperature, such room temperature (<NUM>).

In the method of the present invention, an alloy of formula (II) is used. This formula is given below:.

wherein
x' = <NUM> to <NUM>; y' = <NUM> to <NUM> x' + y' = <NUM>, and z' = <NUM> to <NUM>; the alloy consisting of Mn, Al, optionally C and optionally unavoidable impurities. Herein, the term "unavoidable impurities" includes any element other than Al, Mn and C, and typically denotes impurities that originate from the raw materials used for producing the powder or that are introduced during the manufacturing process. The total content of such impurities is typically <NUM> % by mass or less, preferably <NUM> % by mass or less, and further preferably <NUM> % by mass or less, relative to the total mass of the alloy.

Notably, formula (I) is a subset of formula (II), i.e. formula (I) is completely encompassed by formula (II). Formula (I) represents a preferred embodiment of formula (II). In one embodiment of formula (II), x' is chosen such that Mn does not form <NUM>, <NUM>, <NUM> or <NUM> atom% of the entire composition, and y' is chosen such that Al does not form <NUM>, <NUM>, <NUM> or <NUM> atom% of the entire composition. In another embodiment (which can be combined witht the previous embodiment), z'is chosen such that carbon does not form <NUM> or <NUM> atom% of the entire composition. In yet another embodiment, the alloy of formula (II) does not have the composition Mn<NUM>Al<NUM>C<NUM>.

In the following, a more detailed description of the method steps for processing an alloy of formula (II) or (I) of the present invention is given. Notably, the alloy of formula (I) can also be prepared and processed differently, but is preferably also prepared and/or processed by the method of the present invention. For instance, the alloy of formula (I) may be subjected to a gas atomization process from the melt (as in step a. ), yet it is possible to produce a starting material for step b. and/step d differently, e.g. by strip casting.

It goes without saying that in addition to the method of the present invention as described below, additional process steps may or may not be conducted on the alloy of formula (II) or (I).

The method of the present invention comprises one or more of the following steps a.

In this step, the raw materials of the alloy (typically powders of manganese and aluminum as well as graphite) are provided, the raw materials are melted, and particles of the alloy are formed by gas atomization of the molten alloy.

The temperature of the alloy reached in the melting step is not particularly limited, but is typically in the range of <NUM> or higher, preferably <NUM> or higher, more preferably <NUM> or higher, such as <NUM> or higher. It is preferable to heat the material to the melt for a sufficient time in order to obtain a homogeneous melt by thermal diffusion of the components, e.g. for <NUM> - <NUM> minutes or longer at <NUM> or higher, such as <NUM>.

The step a. is preferably conducted under inert gas atmosphere (such as argon or nitrogen), in order to avoid the formation of oxides. For the same reason, the gas used for the gas atomization is preferably oxygen-free, and is preferably selected from argon, nitrogen, helium and neon. The atomization gas is normally argon or nitrogen (at a pressure between <NUM> to <NUM> bar, more preferably between <NUM> to <NUM> bar, such as <NUM> bar).

The formation of the particles by gas atomization has the advantage over other techniques for preparing alloy particles, such as grinding of a cast ingot, that the stress on the alloy is low, so that disturbances of the crystal structure can be minimized. This in turn allows obtaining a more homogeneous material.

Further, the obtained particles have an almost spherical shape, which is preferably for many later process steps and/or end applications. Specifically, the round shape of the obtained particles allows easy orientation in a magnetic field, which can be beneficial even after a milling step as it allows an increase in powder density and improvement of powder flow and facilitates alignment.

Finally, the gas atomization process allows obtaining a rapid cooling of the molten alloy, which avoids the formation of the τ phase and leads to particles having a high content (purity) of the ε phase. This, in turn, allows obtaining particles having a high τ phase content in a subsequent treatment step transforming the ε phase into the τ phase.

While step a. can be applied to all MnAI alloys having a composition within the scope of formula (II), it is preferably employed for a MnAI alloy having a composition of formula (I). This is due to the fact that in such cases a very pure ε phase can be obtained, as demonstrated by <FIG>.

<FIG> shows at the top the diffractogram (XRPD) of a sample having the composition Mn<NUM>Al<NUM>C<NUM>, and at the bottom of a sample having the composition Mn<NUM>Al<NUM>C<NUM>, both directly after gas atomization under identical conditions. As can be seen from the XRPD data, the sample at the top falling within the scope of formula (II) shows a reasonably good purity, yet some impurities (γ2, β) are also present as derivable from the minor peaks around <NUM> - <NUM> °in 2θ. Such peaks are also present in the alloy having a composition of formula (I) shown in the lower part, yet to a lesser extent. Interesting to note is also the similarity of the grain-matrix structure, as represented by the SEM pictures shown in <FIG> and b. Notably, the difference in composition between the grains and the matrix is smaller for the alloy in accordance with formula (I). This is reflected in the microstructure after high temperature treatment, as illustrated in <FIG> and d.

While thus step a. leads to a good product in terms of homogeneity and purity of the ε phase when an alloy having a low Mn content is used, an even better product is obtained when step a. is conducted on an alloy of formula (I) having a high Mn content of <NUM> to <NUM> (relative to Al and Mn being <NUM>).

However, even a material that does not conform to formula (I) can subsequently be transformed to a material having good magnetic properties and a high content of τ phase, by a suitable heat treatment at <NUM> to <NUM>, preferably <NUM> - <NUM> for a suitable time, such as <NUM> - <NUM> hours, preferably <NUM> to <NUM> hours, such as <NUM>-<NUM>, e.g. <NUM> hours. This is illustrated in <FIG>, showing the respective XRPD diagrams after <NUM> hours at <NUM>. This step is again preferably conducted under vacuum or under inert gas, such as argon. The microstructure after such treatment is very similar to the microstructure achieved directly after atomization (<FIG>).

As shown in <FIG>, the material at the top having the composition Mn<NUM>Al<NUM>C<NUM> has a high purity of τ phase of about <NUM>%. The magnetization at 9T was <NUM> emu/g at 9T, which is a very good value.

Conversely, the material of formula (I) shown at the bottom of <FIG> having the composition Mn<NUM>Al<NUM>C<NUM> shows actually a lower purity of the τ phase, even though the starting material had a higher purity of the ε phase. This is also reflected in the magnetic properties, as the sample had still a good, but slightly lower magnetization of <NUM> emu/g at 9T.

A heat treatment at <NUM> - <NUM> for a suitable time, such as <NUM> - <NUM> hours, is thus able to convert a material having a lower purity of ε phase after gas atomization into a material having actually a higher purity (and/or quality) of the τ phase. In consequence, the present invention also includes an embodiment of the method for producing an alloy wherein gas atomization step a. is conducted for an alloy of formula (II) wherein x' is less than <NUM>, e.g. <NUM> or lower but <NUM> or higher, preferably <NUM> or lower and <NUM> or higher, y' is higher than <NUM> but <NUM> or lower, preferably <NUM> or higher and <NUM> or lower (with x'+y' = <NUM>), and z' is <NUM> or less and <NUM> or more, preferably <NUM> to <NUM>, followed by a heat treatment <NUM> to <NUM>, preferably <NUM> - <NUM> for a suitable time, such as <NUM> - <NUM> hours. This allows obtaining a material rich in τ phase and having a magnetization of <NUM> emu/g (at <NUM> Tesla) or higher.

The particle size of the powders obtained by the gas atomization process is not particularly limited and can be adjusted by selected the appropriate conditions. Typically, the powders obtained from the gas atomization process have a particle size (or particle diameter) D50, as determined by a laser light scattering method, of <NUM> or less, preferably <NUM> or less, and more preferably in the range of <NUM> - <NUM>.

If desired, the material obtained from step a. may optionally be subjected to a further micronization treatment. This could for instance include a ball milling, e.g. the cryo milling of step c.

, the alloy of formula (II) (and preferably the alloy of formula (I)) is subjected to a heat treatment at <NUM> -<NUM>. This step is again preferably conducted under vacuum (i.e. at a pressure of less than <NUM> Pa, more preferably <NUM> Pa or less, such a <NUM> Pa or less), or under inert gas atmosphere, such as under argon.

It has been found that such a step b. is able to transform the ε phase into the τ phase. In consequence, the starting material for step b. is preferably an MnAI alloy of formula (I) that contains ε phase. More preferably, the content of the ε phase as determined by XRPD is <NUM>% or more, more preferably <NUM>% or more. For reference, the purity according to XRD/Rietveld is for <FIG> (bottom) about <NUM>% εand for <FIG> (bottom) ca <NUM>,<NUM>% τ, showing that essentially all of the ε phase is transformed into the τ phase.

may also minimize structure defects, such as APB (antiphase boundary) defects and the presence of twinned crystals, which are expected to impair the magnetic properties. In consequence, the τ phase resulting from step b. can show high crystal structure quality and can exhibit excellent magnetic properties.

The duration of the heat treatment in step b. depends on the size and shape of the alloy sample that is subjected to the treatment. It is easily recognizable that large lumps or ingots may require more time for the temperature to reach the interior thereof as compared to powders. In addition, larger segregates (i.e. large volume regions of compositional differences) require more time to minimize compositional gradients and homogenize microstructure because of the long range atomic diffusion that is necessary. A too long heat treatment time at <NUM> to <NUM> might however possibly be disadvantageous, as further reactions might occur. In consequence, it is preferred to conduct step b. on alloy of formula (II) or (I) that are present in the form of particles. The particles have preferably a diameter D50 of <NUM> or less, more preferably <NUM> or less, further preferably <NUM> or less, still further preferably <NUM> or less , such as in the range of <NUM> to <NUM> or <NUM> to <NUM>. Herein, the value of D50 defines that <NUM>% by weight of the particles have the indicated size or smaller, as determined by a laser light scattering method.

For particles having a such a particle size, a heat treatment time at <NUM> to <NUM> of <NUM> hours or less is generally sufficient, and the heat treatment time can also be <NUM> hours or less or <NUM> hours or less, such as <NUM> minutes or less, or even <NUM> minutes or less or <NUM> minutes or less. The minimum time is <NUM> minutes, but can also be <NUM> minutes or more or <NUM> minutes or more, such as <NUM> minutes or more, e.g. <NUM> minutes or more. In the present invention, the heat treatment time of step b. is defined as the time span from the heat-up when the alloy temperature reaches <NUM> to the time when the temperature drops below <NUM>. This time thus represents the dwell time within this temperature range from <NUM> to <NUM>.

The heat and cool rates are not particularly limited, but may be chosen appropriately in order to avoid formation of γ2 and β phases. While high (or low) heating rates at lower temperatures (i.e. below <NUM>, <NUM> or even below <NUM>) are not considered to be particularly disadvantageous as no major effects are expected at these temperatures, the temperature increase and decrease rates within the temperature range of from <NUM> (or <NUM>) to the lower limit of <NUM> of step b. is preferably from <NUM>/min or higher, more preferably <NUM>/min or higher, and preferably <NUM>/min or less, more preferably <NUM>/min or less, such as within the range of <NUM>/min to <NUM>/min.

The temperature within the range of <NUM> to <NUM> can be kept constant for a certain time, such as for <NUM> minutes or more. It is however also possible to provide for a temperature profile containing heating and cooling segments within the range of <NUM> to <NUM>. An exemplary temperature profile is shown in <FIG>.

can be performed with any alloy of formula (II) comprising ε phase, but is preferably performed with an alloy of formula (I). Further, step b. is preferably performed for an alloy comprising ε phase in particle form as defined above, and further preferably is performed with an alloy comprising ε phase in particle form obtained by gas atomization in accordance with the above step a. Yet, as outlined above, it is also possible to produce an alloy containing ε phase by both gas quenching and water quenching. Gas-quenched (as example <NUM>) gives lower oxygen amount (<<NUM>. 03wt% O) which is beneficial, however, water-quenched cooling (using e.g. degassed distilled water) could improve cooling rate and thus give higher ε-phase purity (<<NUM>. However, the lower quenching rate for gas-quenched powders may be compensated with a higher melt temperature, which decrease melt viscosity and atomized particle size and, thus, increases the quenching rate.

Most preferably, steb b. is conducted on an alloy of formula (I) in particle form obtained by gas atomization in accordance with the above step a. In all these embodiments, the content of the ε phase is preferably <NUM>% or more, more preferably <NUM>% or more (as determined by Rietveld methods on X-ray diffractograms).

The reason why step b. is preferably conducted for an ε phase containing alloy of formula (I) having a relatively high carbon and manganese content is that in particular for such compositions the effect of forming the τ phase is prominent, and thereby an alloy having extremely beneficial properties can be obtained. This is illustrated in <FIG>.

<FIG> shows at the top the XRPD diagram of the sample having the composition (Mn<NUM>Al<NUM>)C<NUM> (Example <NUM>-B6) and at the bottom of the sample having the composition (Mn<NUM>Al<NUM>)C<NUM> (Example <NUM>-A2), respectively, after gas atomization (ε phase, XRPD shown in <FIG>) and subsequent heat treatment at <NUM> for <NUM> minutes (heat rate and cooling rate up to <NUM>: <NUM> /min, vacuum). SEM images of the two materials after this heat treatment are shown in <FIG> and d.

As derivable from a comparison of <FIG> with <FIG>, the sample having a low Mn content with the composition (Mn<NUM>Al<NUM>)C<NUM> shows somewhat higher purity of the τ phase and somewhat better magnetic properties after a heat treatment at <NUM> for <NUM> under argon (<FIG>, top (Mn<NUM>Al<NUM>)C<NUM>): purity ca. <NUM>%, magnetization <NUM> emu/g at 9T; bottom (Mn<NUM>Al<NUM>)C<NUM>) purity ca. <NUM>%, magnetization <NUM> emu/g at 9T). This is observed despite the fact that the purity of the ε phase of (Mn<NUM>Al<NUM>)C<NUM>) as obtained after gas atomization is lower (see <FIG>, top: (Mn<NUM>Al<NUM>)C<NUM>), bottom:
(Mn<NUM>Al<NUM>)C<NUM>)).

Yet, for a heat treatment b. at <NUM> - <NUM>, the situation is actually reversed, as shown in <FIG>. Here, the sample having a low Mn content ((Mn<NUM>Al<NUM>)C<NUM>), x' in formula (II) = <NUM>) lead to a lower purity of the τ phase of ca. <NUM>%, and the SEM picture (<FIG>) shows severe compositional segregation with high content of impurity phases (β, γ2) at mainly the grain boundaries. This is reflected in less favorable magnetic properties, giving a magnetization of <NUM> emu/g at 9T. Conversely, the sample of formula (I) having a high Mn content ((Mn<NUM>Al<NUM>)C<NUM>), x in formula (I) = <NUM>) exhibited a purity of ca <NUM>% and excellent magnetic properties (magnetization <NUM> emu/g at 9T), as well as a more homogeneous microstructure (see <FIG>).

In consequence, step b. is applicable to all MnAI alloys of formula (II), yet better results are obtained if step b. is performed for an MnAI alloy also satisfying formula (I).

In consequence, the present invention also includes as one embodiment a method for producing an alloy of formula (II) as defined above wherein.

The heat treatment step b. is also preferably applied to MnAI alloys satisfying formula (I) that contain ε phase. The starting material for step b. may be a material that is obtained directly after step a. , in which case the method of the present invention includes both the steps a. The method comprising step b. is however not limited to methods wherein the starting material for step b. is obtained from step a. , and any MnAI alloy of formula (II), and preferably formula (I), containing ε phase can be used as starting material for step b.

The material obtained from step b. containing τ phase may in one embodiment be subjected to a micronization treatment, which may be the following step c. or any other milling procedure. Milling on the relatively soft τ phase can induce relatively more stress and can result in an amorphous state. This may be beneficial for a later recrystallization (and homogenization), e.g. by step d. described below, and may result in better magnetic properties.

is a step of milling an alloy represented by formula (II) at a temperature of -<NUM> or below, preferably -<NUM> or lower, further preferably -<NUM> or lower, and still further preferably - <NUM> or lower. This step is herein also referred to as "cryo milling". From a practical standpoint, it is most preferred to conduct the cryo milling at the boiling temperature of nitrogen under atmospheric pressure (-<NUM>).

It has been found by the present inventor that conventional milling operations, such as ball milling at room temperature, induce the formation of defects in the crystal structure and lead to deteriorated magnetic properties, in particular saturation magnetization, while leading to an increase in coercivity. Without wishing to be bound by theory, it is believed that the ε to τ transformation is highly dependent on the nucleation of the τ-phase at the interphase with the ε-phase, and hence, a microstructure with a high surface to volume aspect ratio of both grains and particles may be important in order to promote the τ-phase formation. Results indicate that the rate of formation of τ-phase is much faster from the smaller grained powder than that from coarse-grained ε-phase in the bulk.

The shape and surface morphology of the particles can be controlled by the type of milling technique and processing parameters. An optimal size, spherical shape and smooth surface of the particles may help or facilitate achieving a preferred magnetic alignment of the powders particles, e.g. in a compression tool prior to compaction. Such particles are better designed to respond and rotate more easily to external magnetic fields. The particles will align along the c-axis (easy axis) in a magnetic field. The alignment is preferably performed inside the compaction tool just before and/or during the compression or molding movement. Aligned particles will result in an enhanced remanent magnetization (Mr) of the magnet body, i.e. the magnetization left behind in the body after the external magnetic field has been removed.

It is observed that gas atomized powders are more spherical and have smooth surfaces as compared to water atomized powders. Thus, gas atomized powders are preferred as they may expose a higher degree of smooth spherical surfaces even after a subsequent milling operation. Optionally, milled powders of any synthesis method may be subjected to spherodization methods, e.g. plasma spherodization by Techna® Group.

Hence, the starting material for step c. may be obtained by gas atomization in accordance with step a. , to which optionally (and preferably) step b. has been conducted prior to the cryo milling. In this embodiment, the method of the invention comprises the steps a. , in this order. Yet, step c. can be conducted also on materials of formula (I) or (II) as starting materials that have been prepared via processes not including step a. and/or step b. Preferably, however, the starting material for step c. is an alloy of formula (I) or (II) containing τ phase.

Moreover, milling of brittle materials results in smaller and sharper particles as compared to milling of soft material. On the other hand, milling of soft material will result in relatively larger particles, or even flakes for the same energy input. Milling on softer material often gives smaller grains, a more stressed crystal structure, or even result in an amorphous state. This is a further reason why in one embodiment the method of the present invention comprises a step c. of cryo milling, preferably on an alloy of formula (I) or (II) containing τ phase.

As the hardness of ε-phase is significantly higher than the τ-phase, milling on either phase will have a decisive impact on the final properties. Also, cryogenic milling will make the powder even more brittle and, thus, influence the outcome of the milling in a similar way.

The starting material for step c. can thus be any alloy of formula (II). The starting materials can contain ε or τ phase, preferably in a predominant amount (e.g. more than <NUM>% or <NUM>%, as determined by XRD, see section <NUM>), and is in one preferred embodiment a material containing τ phase. For instance, it can be a τ phase containing MnAI alloy of formula (II) or formula (I) that is obtained by a cooling process, with or without any subsequent treatment for increasing the content or purity of the τ phase. It however can also be an ε phase containing MnAI alloy of formula (II) or (I) that is obtained from step a. , with or without a further treatment for e.g. purification of the ε phase, such as the gas atomized material obtained from only step a. , or a gas atomized material obtained after steps a. have been performed on a MnAI alloy of formula (I) as outlined above. It can also be an ε phase containing material of formula (II) wherein x' is lower than <NUM>, such as <NUM> or lower, and a value of y of higher than <NUM>, such as <NUM> or higher, and a value of z' of <NUM> or lower, on which optionally further a heat treatment at <NUM> - <NUM> for <NUM> - <NUM> hours, as outlined above, has been performed.

The milling technique can be freely chosen from conventional milling techniques, such as ball milling, jet milling, pin milling etc, or other high shear processes such as hot extrusion. In one embodiment, the cryo milling step employs ball milling at a ball-to-powder ratio (by volume) of <NUM>:<NUM> to <NUM>:<NUM> for a suitable duration, such as <NUM> to <NUM> minutes, preferably <NUM> to <NUM> minutes.

The milling is preferably conducted under vacuum or under inert gas atmosphere, such as under argon, in order to avoid the formation of oxides. The material is preferably milled to a small particle diameter (expressed as weight D50 and determined by a laser light scattering method) of e.g. <NUM> or less, preferably <NUM> or less, more preferably <NUM> or less, further preferably <NUM>µl or less, still further preferably <NUM> or less, such as <NUM> or less. In order to minimize distortions of the structure and deterioration, the cryo milling is preferably conducted for as short time as possible, but as long as necessary in order to obtain the required particle size and degree of micronization, and/or degree of crystal amorphization, which depends on the hardness of the material and the milling equipment used. For instance, the (cryo) milling is preferably performed for <NUM> hours or less, preferably for <NUM> hours or less, and still further preferably for <NUM> hours or less, but for <NUM> seconds or more, and often for <NUM> minute or more.

The powder prior milling could be coated with minor amounts (<<NUM>% by weight) of an organic substance, such as surfactants or fatty acids such as oleic acid, commonly known as surfactant-assisted milling. Without being bound to any theory, the additive may protect the surfaces for excess oxidation and minimize agglomeration during milling caused by adhesive or electrostatic forces between particles.

is a step wherein a heat treatment on particles of the alloy represented by formula (II) is performed at a temperature of <NUM> to <NUM> for a time of <NUM> to <NUM> minutes, preferably <NUM> to <NUM> minutes, more preferably <NUM> to <NUM> minutes. This step is also referred to as "flash heating" or "flash heat treatment".

It has been found by the present inventors that the structural distortions that are caused by processing operations such as milling can, at least in part, be reversed and the degree of purity of the τ phase be greatly improved by performing a flash heat treatment step as outlined above on particles of an MnA alloy of formula (II) or formula (I). This so-called "flash heat treatment" allows obtaining an excellent combination of high magnetic saturation and coercivity after the micronization process used for particle formation.

The flash heat treatment d. is preferably performed by first effecting a heating at relatively moderate temperatures at e.g. <NUM> - <NUM> in vacuum or reduced pressure to ensure the desorption of water, oxygen or oxygenated species.

Thereafter, the particles can be heated up further, either in vacuum or inert gases. In this respect, it is preferred to use high heating rates of e.g. <NUM>/min or higher, such as in the range of <NUM> - <NUM>/min, between <NUM> and <NUM> in order to maximize the time above the temperature at which the ε to τ transition occurs (<NUM> - <NUM> ).

The temperature is then held between <NUM> and <NUM> for <NUM> to <NUM>, preferably <NUM> to <NUM> minutes. It is preferred to keep the time as short as possible, in order to restrict grain growth and sintering of the particles. The suitable time depends also on the surface area, the degree of crystal strain as induced by the shear forces during micronization and the shape of the material, and often <NUM>-<NUM> minutes are sufficient to form the τ phase.

Subsequently, the material is cooled down, typically to room temperature. Again, the dwell time in the temperature range between <NUM> and <NUM> is again held as short as possible also during cooling, so that the cooling rate is typically also <NUM>/min or higher, such as in the range of <NUM> - <NUM>/min. Minimizing the time in the temperature range between <NUM> and <NUM> on both heating and cooling allows avoiding or minimizing the formation of impurities, such as β and γ2.

Optionally, the atmosphere pressure at higher temperatures (<NUM> or higher) can be increased in order to limit the loss of manganese and thus limit the risk of forming nonmagnetic impurities. This can be achieved by increasing the pressure, e.g. nitrogen or argon, at above 1bar.

The material used for step d. are particles of an MnAI alloy of formula (II) or formula (I). These particles contain at least one of ε and τ phase. The starting material for step d. is thus not particularly limited, and does not need to be a material that is obtained after any of the steps a. , and/or c. The starting material can however also be a material as obtained after the above step a. only, after the above step b. only, after the above step c. only, after conducting steps a. , after conducting steps a. , or after conducting steps a. , in this order. Preferably, the starting material for step d. is an MnAI alloy of formula (II) or more preferably formula (I) that is obtained after steps a. , and optionally c.

The starting material of step d. is preferably in a size and shape that does not require further processing after step d. The material is thus preferably particles having a small particle diameter (expressed as weight D50 and determined by a laser light scattering method) of e.g. <NUM> or less, preferably <NUM> or less, more preferably <NUM> or less, further preferably <NUM>µl or less, still further preferably <NUM> or less.

By this flash heat treatment step, τ phase is formed. This is shown in <FIG>, showing the XRPD data for a sample having the composition (Mn<NUM>Al<NUM>)C<NUM>.

At the bottom of <FIG>, the XRPD data of the sample obtained after gas atomization (step a. , ε phase), cryo milling (step c. , <NUM> minutes) and flash heat treatment (step d. ) is shown. In other words, a cryo milled ε phase was the starting material for the flash heat treatment. As is clearly derivable form the data, the material after flash heat treatment is relatively pure τ phase, even though the starting material was mainly ε phase. This material showed a magnetization of <NUM> emu/g (at 9T).

At the top of <FIG>, the same material is shown, wherein additionally between the steps of gas atomization (step a. ) and cryo milling (step c. ) a heat treatment for transforming the ε into the τ phase has been conducted (step b. In other words, a cryo milled τ phase was the starting material for the flash heat treatment. The obtained material after flash heat treatment shows also high purity of the τ phase and a magnetization of <NUM> emu/g at 9T.

This shows that a flash heat treatment step in accordance with step d. is able to provide particles having the desired shape (as present in the particles as starting material) having a high content and high quality of τ phase. This represents a major advantage over prior art processes wherein the τ phase was formed prior to further processing steps (e.g. milling), as such processing steps lead to deteriorated properties. With a flash heat treatment step d. , however, it becomes possible to obtain a final material having simultaneously the desired shape and good magnetic properties, including magnetization.

The present invention will be described in more detail by way of the following Examples, to which the invention is however not limited.

An Mn<NUM>Al<NUM>C<NUM> alloy ingot of <NUM> was synthesized by a drop synthesis process starting from a melt at <NUM>, similar to a synthesis method described in <NPL>. The raw materials were all of high purity, using Mn (Institute of Physics, Polish Academy of Sciences, purity <NUM>%), C (Highways international, <NUM>%) and Al (Gränges SM, purity <NUM>%).

First, Aluminum was heated and melted with carbon black in a alumina crucible at <NUM> after the atmosphere was evacuated and high vacuum established (<<NUM> bar). Small pieces of Mn metals were subsequently dropped into the melt of Al and C, then the eddy current power was increased to enable the Mn pieces to react with Al-C liquid immediately. The melt was kept at <NUM> for <NUM> minutes to ensure that the Mn-Al-C liquid forms a homogeneous alloy solution.

The Mn-AI-C alloy was cooled down to room temperature by cooling the ingot in vacuum over a water-cooled Cu plate. The achieved cooling rate allows the alloy to form τ-MnAl of high purity directly. The resulting material is in the following also referred to as "drop synthesized" material.

Cryogenic milling was performed at liquid nitrogen temperatures (-<NUM>) using a SPEX Freezer/Mill® <NUM>. The starting material prepared above was placed in a stainless steel vial with a stainless steel cylindrical impactor. The mass ratio between the impactor and the powder was <NUM>:<NUM>. Before the milling was started, the vial was allowed to cool down for <NUM> in the liquid nitrogen bath of the Freezer/Mill®. The milling was then carried out at an impact frequency of <NUM> for a total of <NUM> (CM2) or <NUM> hours (CM4). Each milling run consisted of <NUM> minutes milling and <NUM> minutes pause cycles.

A smooth spherical particle shape is easiest to align along the easy magnetization axis in a magnetic field. Cryo milling allows obtaining particles that may undergo such an easy alignment, as it does not lead to the formation of sheet-like particles with random orientation as was observed when milling is performed at ambient temperatures (see e.g. <NPL>, or <NPL>).

The smooth and spherical particle shape obtained by cryo millling was confirmed by visual inspection of particles obtained after <NUM> hours and <NUM> hours cryo milling using SEM imaging. Here, it was confirmed that the average particle size of the particles remained about <NUM> after <NUM> and <NUM> hours cryo milling, respectively. The surface smoothness increased somewhat at longer milling times.

X-ray powder diffraction (XRD) was performed at a Bruker Twin-Twin diffractometer, with a Cu double Kα radiation (λ1 = <NUM>Å, λ2 = <NUM>Å). The neutron powder diffraction was carried at JEEP-II reactor of IFE (Institution for Energy) at Kjeller, Norway. The neutron diffraction patterns were detected by a high pressure neutron diffractometer. The crystal structure and phase analysis were treated by Fullproof™ software through the Rietveld method as described in <NPL>). The peak shape of the diffraction pattern was characterized by the Thompson-Cox-Hastings pseudo-Voigt function.

The phase transition behaviors of <NUM> (CM2) and <NUM> (CM4) hours cryo milled Mn<NUM>Al<NUM>C<NUM> samples when heated and cooled at different rate was investigated by in situ synchrotron X-ray diffraction at the P02. <NUM> beamline at PETRA III (λ = <NUM>Å). The powder cryo milled Mn<NUM>Al<NUM>C<NUM> samples were loaded in a single crystal sapphire tube, the tube was wounded by Kanthal wire and heated up to <NUM> in vacuum (<NUM>/min), dwelled at <NUM> for <NUM> minutes, then cooled (<NUM>/min) to room temperature. The temperature was monitored by a K Type thermocouple insert from one side of the sapphire tube with close contact to the sample. The sample to detector distance and X-ray beam wavelength was determined and calibrated by the NIST LaB6 standard sample. The X-ray diffraction patterns were recorded by a PerkinElmer XRD1621 fast area detector. The diffraction patterns of 2D pictures were transformed to 1D diffractograms by the Fit2D™ program.

For the flash heating process, samples were first put into Al<NUM>O<NUM> crucibles, the crucibles were sealed in evacuated quartz tubes. Then, the ampoules were transferred to a preheated resistance furnace and "flash heated" at <NUM> for a total time of <NUM> minutes, <NUM> minutes, and <NUM> minutes respectively, followed by cooling the ampules in ambient air.

Powder samples were placed in gelatin capsules with varnish. The capsule and varnish together contribute a paramagnetic moment at <NUM> and account for < <NUM> % of the saturation magnetic moment at <NUM> T. Samples were measured in a Physical Properties Measurement System (PPMS) from Quantum Design equipped with a <NUM> T superconducting magnet or a MPMS from Quantum Design. Magnetization in SI units and µB were calculated from the sample weight and using the lattice parameters obtained from the XRD/NPD refinements.

The refined powder diffraction data of as-synthesized (<FIG>), <NUM> cryo milled (<NUM>, <FIG>) and <NUM> cryo milled (<NUM>, <FIG>) samples are shown in <FIG>.

From the XRPD data (<FIG> a clear decrease in the peak intensities combined with a pronounced peak width broadening is observed with longer milling time. In addition, several of the weaker peaks (i.e., <NUM>, <NUM> and <NUM>Å-<NUM>) related to the τ-phase (i.e., the (<NUM>) and the (<NUM>) and the (<NUM>) planes), gradually disappears with longer milling time. However, the remaining strong reflections (i.e., between <NUM>Å-<NUM> ≤ Q ≤ <NUM>Å-<NUM>) from XRPD indicate that a crystalline phase is still preserved. On the contrary, the NPD data (<FIG>) show a strong decrease of the reflection intensities for the <NUM> sample, while no reflections are observed for the <NUM> sample (<FIG>), reminiscence of an amorphous phase.

The combined XRPD and NPD data (<FIG>) of the DS, <NUM> and <NUM> samples were used to refine the lattice parameters in the space group P4/mmm. The NPD data was, however, not refined for the <NUM> due to the lack of peaks. From Table <NUM> with data obtained from the Rietveld refinement of the combined XRPD and NPD data, the occupancy of the Mn and Al-sites are found to vary with increased milling time. It is found that the Mn content at the Mn 1a (<NUM>, <NUM>, <NUM>) site decreases from <NUM>% to <NUM>%, while the Al content increases from <NUM>% to <NUM>% after <NUM> hours of cryo milling. The opposite is observed at the Al 1d (½,½,,½) site, where the Al content decreases from <NUM>% to <NUM>% after <NUM> hours of cryo milling (Table <NUM>).

<FIG> shows refined XRPD patterns of the flash heated <NUM> and <NUM> powders. It is clearly observed that the flash heating process recrystallizes the powder significantly after only <NUM> (<FIG>) to produce peaks comparable to the original DS sample (<FIG>).

All the XRPD patterns for the flash heated samples contain detectable amounts of the y2-phase (<FIG>) that are mainly observed ~ <NUM>Å-<NUM>, but the amount of these phases is quite low (<<NUM>%). The structural model with the same Mn and Al occupancies as in the DS sample indicates that a reordering of the Mn and the Al on the two crystallographic sites takes place upon the heat treatment (Table <NUM>).

The above results show that the flash heating process of the present invention is able to re-form the τ phase, even after processing steps such as milling that lead to a loss of the previously present τ-phase. These results are also in agreement with the XRPD and NPD patterns shown in <FIG> and <FIG>, where <FIG> shows the refined XRPD of cryo milled and flash heated samples; <FIG>) <NUM> hours cryo milled and <NUM> +<NUM> minutes flash heated; <FIG>) <NUM> hours cryo milled and <NUM> +<NUM> minutes flash heated; <FIG>) <NUM> hours cryo milled and <NUM> +<NUM> minutes flash heated; <FIG>) <NUM> hours cryo milled and <NUM> +<NUM> minutes flash heated).

To further investigate the stability range of the τ-phase as a function of heating rate and temperature, the <NUM> sample was analyzed in situ by synchrotron radiation (λ = <NUM>Å). During the measurement, the powder was subjected to a heating rate of <NUM>/min from room temperature up to <NUM> and was kept for <NUM> minutes before being cooled down to room temperature again, with a rate of <NUM>/min. It was observed that the <NUM> powder decomposes into a mixture of β-phase and τ-phase at ~<NUM>. At T > <NUM> the powder transforms fully into pure ε-phase. However, during the cooling process at a cooling rate of <NUM>/min, pure τ-phase is reformed again at T < <NUM>.

The effect of the cryo milling process on the magnetic properties (i.e., Hc and Ms) was evaluated from magnetization versus magnetic field measurements. From the magnetic hysteresis loop of the DS sample (cf. <FIG>) the coercive field is obtained as µ<NUM>Hc≈<NUM> mT, while the value of the saturation magnetization Ms≈<NUM> kA/m is close to the theoretical maximum of the magnetization. Furthermore, the effect on the magnetic properties from cryo milling followed by flash heating (<NUM>, <NUM> and <NUM>) is illustrated by the M-H measurements of <NUM> and <NUM> in <FIG>,d.

Overall, the magnetization decreases significantly with increasing milling time. After <NUM> of flash heating at <NUM>, the Ms for <NUM> and <NUM> only recovers to ~<NUM>% and ~<NUM>% of the original DS sample. The Hc is, however, ~<NUM>% and ~ <NUM>% higher than the original DS sample, respectively (see Table <NUM>). Further increase of the flash heating time from <NUM> to <NUM> at <NUM> results in an Ms value ~<NUM>% and ~<NUM>% of the DS sample for the <NUM> and <NUM> samples, respectively. The Hc values for the <NUM> and <NUM> samples are however, only ~ <NUM>% and ~ <NUM>% higher than the original DS sample after <NUM> of flash heating at <NUM>. Moreover, heating for <NUM> only resulted in minor changes as seen in Table <NUM>.

From the magnetic properties of the DS sample, the Ms value ~<NUM> kA/m (~<NUM> Emu/g) reach nearly the theoretical limit <NUM> kA/m for, indicating that the ε to τ transformation is nearly complete. Yet, the Hc value is low at <NUM> mT, which may not be sufficient for many industrial application. Notably, Hc is greatly increased by cryo milling, and the cryo-milled and flash-heated material shows both good Ms and Hc. This shows that the combination of cryo milling and flash heating allows obtaining a material that has both sufficient Ms and Hc, while beneficial effects are also observed for cryo milling alone (increase of Hc) and flash heating (restoration of the τ phase, i.e. increase in Ms close to that of the untreated DS material). Without wishing to be bound by theory, it is believed that the flash heating (besides minimizing impurities such as β and γ2 phases) re-orders the Mn and Al in the crystal lattice, and is in particular effective to heal defects in this ordering that have been induced by the milling procedure.

First, aluminum metal (Stena Aluminium AB, Sweden, ><NUM> wt%) was melted together with graphite (Carbomax AB, Sweden, ><NUM> wt%) at a melt temperature of about <NUM> in an argon atmosphere. Thereafter, the manganese metal (Manganese metal company Ltd, SA, ><NUM> wt%) was added and the melt temperature was adjusted to <NUM> and held there for <NUM> minutes prior to gas atomization under an argon atmosphere at a pressure of about <NUM> bar. This caused a rapid cooling of the gas-atomized particles to form ε-MnAl alloy of high purity and low oxygen content (ca <NUM> wt% for gas-quenched). The raw materials were carefully adjusted to give the compositions of the samples presented in Table <NUM>.

The gas-atomized particles were then milled using an MM <NUM> mixer mill (Retsch GmbH) capable for dry, wet or cryogenic grinding, optionally under argon atmosphere. A <NUM> hardened steel vial, <NUM> balls and an impact frequency of <NUM> for <NUM> with <NUM>:<NUM> ball-to-powder ratio were used.

The powders were prior coated with <NUM> wt% oleic acid (as applied by acetone that was allowed to vaporize prior milling). Before the cryo-milling was started, the sealed vial was allowed to cool down for <NUM> in a bath of liquid nitrogen. Between each <NUM> of milling the vial was let to cool down again for <NUM>. The cry milling was performed under the conditions outlined in section <NUM> above.

XRD measurements were performed using a Panalytical X'Pert Pro PW3040 Multi Purpose Diffraction system equipped with X-Celerator solid state line detector. Experimental data were processed using Panalytical B. X'Pert HighScore Plus software, version <NUM>. The measurement conditions were Cu Kα <NUM>, accelerating voltage <NUM> kV, current <NUM> mA, Ni-filter, 2θ-scan range: <NUM>-<NUM>°, divergence slit <NUM>, time <NUM>. The quantitative analysis by refinement of the phases was performed manner as described above in section <NUM>.

The samples were characterized with respect to their magnetic characteristics with a LakeShore <NUM> VSM (maximum applied field <NUM>. 8T) and a Quantum Design PPMS (maximum applied field 9T). Samples were placed in gelatin capsules and fixed with a weakly paramagnetic varnish (< <NUM>% of the moment at <NUM> T). A density of <NUM>/m3 was assumed.

The melting and cooling behavior of the powders were measured with a Simultaneous Thermal Analysis instrument (TGA & DCS) from Netzsch (Jupiter STA <NUM> F3); method <NUM>/min to <NUM> in argon gas, sample size ca <NUM>.

Particle samples were mounted in fina met with added carbon CFU-<NUM> and bakelite PhenoCure™, grinded and polished down to <NUM> by standard process. It was polished with OPS-S Non-drying colloidal silica suspension for <NUM> minutes in order to improve planeness. The sample was subsequently etched in a diluted modified Keller's reagent: with HCl <NUM>, HNO<NUM> <NUM>, HF <NUM> and H<NUM>O <NUM>. A thin layer of Au was sputtered on the surface to improve surface conductivity. Sample was analyzed in a Field Emission Scanning Electron Microscopy (FE-SEM) Hitachi SU6600 equipped with an Electron Discharge Spectroscopy (EDS) system (Bruker EDX XFLASH <NUM>).

A high temperature treatment as applied on as-synthesized powders or pieces, either in vacuum or argon. The temperature profile is shown in <FIG>.

The flash heat treatment was performed in a sealed tube furnace (L <NUM>, Ø4cm; Entech AB) using the temperature profile shown in <FIG>. First, the sample was degassed by repeatedly filling with argon and vacuum, and then the temperature was allowed to increase slowly in vacuum (<NUM>/min). The vacuum was maintained at <NUM> until low and stable pressure was reached. In a second stage, the temperature was increased rapidly to <NUM> (><NUM>/min) to perform a "flash treatment". The time above <NUM> was adjusted to about <NUM> minutes, followed by rapid cooling to ambient temperature. In the second stage the atmosphere was either vacuum or argon.

Following the synthesis scheme above, the following samples were prepared and subjected to DSC analysis:.

Of the above samples, those named "<NUM>-B" do not comply with formula (I) as defined in claim <NUM>. while those named "<NUM>-A" comply with formula (I).

<FIG> shows exemplary DSC graphs of materials as synthesized above (<FIG>: Example <NUM>-B1, <FIG>: Example <NUM>-A3). The exothermic peaks at about <NUM>~<NUM> represent the martensitic transition from ε-phase to τ-phase. At about <NUM>~<NUM> the ε-phase reforms, and at ca <NUM> to <NUM> the melting of the alloy starts.

The total area of the peaks between <NUM>~<NUM> correlates with the amount of carbon-stabilized τ-phase, and the total area covered by this/these peak(s) is indicated in the table above in J/g. Here, a higher area value represents a higher amount of carbon-stabilized τ-phase. However, in alloys containing up to <NUM>% Mn, only up to <NUM>% of carbon can be dissolved. Increasing the amount of Mn allows increasing the amount of carbon, which in turn leads to a greater stability of the τ phase. Further, the carbon-stabilized τ-phase is only stable for further processing (milling and heat treatment) if the Mn/C-ratio is close to <NUM> (in the range <NUM>~<NUM>). i.e. a carbon-saturated or close to carbon-saturated τ-phase.

It is apparent that a higher stability of the τ phase. i.e. higher peak area in combination with a carbon-saturated τ-phase (Mn/C-ratio preferably between <NUM>~<NUM>), could be obtained when the requirements of formula (I) are met.

Further, it is apparent from the results above that the alloys in accordance with formula (I) show only one peak in the DSC in the temperature range of <NUM> - <NUM> (or <NUM> - <NUM>), denoted as Peak <NUM>. This indicates a higher homogeneity of the τ phase, as only one apparent transition back to the ε phase is observed.

The following additional observations were made, illustrating that adjusting the composition of the alloy in accordance with formula (I) and adjusting the Mn/C ratio is of importance:
Sample <NUM>-B3 ((Mn<NUM>Al<NUM>)C<NUM>) has a high amount of carbon and low amount of Mn. This sample contained a lot of undissolved carbon that forms carbides after heating to <NUM> or <NUM>. This was confirmed by significant peaks for Mn<NUM>AlC in the XRPD (see also <FIG>). Further, the total peak area of <NUM> J/g indicates an amount of carbon dissolved in the τ phase of less <NUM>%, and the material is not stable upon heat treatment. Interestingly, the material does not at all exhibit a DSC peak <NUM> in the temperature range of <NUM> - <NUM>, but only a DSC peak <NUM> at higher temperatures.

Sample <NUM>-B6 ((Mn<NUM>Al<NUM>)C<NUM>) contains carbon in an amount just sufficient to stabilize the τ phase, but the amount of Mn is too low to satisfy the requirements of formula (I). As explained previously and shown in <FIG> and <FIG>), this sample has somewhat inhomogeneous composition and consequently upon a heat treatment a heterogeneous microstructure, which is shown to impair the magnetic properties.

Samples <NUM>-B11 and <NUM>-B12 have a high Mn/Al ratio, but lack sufficient carbon. These samples are not stable upon high temperature treatment and showed decomposition into β and γ<NUM>.

Sample <NUM>-B5 also has an even higher Mn/Al ratio, and also sufficient carbon. Yet, this sample is at the limit where τ phase can at all be formed, potentially favoring the formation of β-phase and impairing the purity of the τ-phase (and thereby magnetic properties).

The results above show that the desired balance of properties, including inter alia high purity of the ε phase and high stability and purity of the τ phase, stability of the material against high temperature processing, good magnetic properties and improved homogeneity on a microscopic level, can be simultaneously obtained by adjusting the composition of the MnAl alloy within the boundaries of formula (I) and its preferred embodiments.

The results further show that the process steps of the method of the present invention, separately or in combination, allow improving the synthesis of MnAI alloys for magnetic applications, respectively the properties of the obtained material.

Claim 1:
An alloy represented by the formula (I)

        (MnxAly)Cz     (I)

the alloy consisting of aluminum (Al), manganese (Mn), and carbon (C), and optionally unavoidable impurities;
wherein
x = <NUM> to <NUM>
y = <NUM> to <NUM>
x + y = <NUM>, and
z = <NUM> to <NUM>, which has a saturation magnetization Ms of <NUM> emu/g or more, wherein the ratio of x to z (x/z) is in the range of <NUM> to <NUM> and the alloy is in the form of particles.