EP4027358B1 - Soft magnetic alloy and method for producing a soft magnetic alloy - Google Patents

Description

The invention relates to a soft magnetic alloy and a method for producing a soft magnetic alloy.

Soft magnetic cobalt-iron alloys (CoFe) are used, among other things, in electrical machines due to their extremely high saturation induction. Commercially available CoFe alloys typically have a composition of 49 wt% Fe, 49 wt% Co and 2% V. With such a composition, a saturation induction of around 2.35 T is achieved with a high electrical resistance of 0.4 µΩm. Due to their higher permeability, these alloys can be used in applications such as a rotor or a stator of an electric motor in order to reduce the size of the rotor or stator and thus of the electric motor and/or increase the performance compared to an FeSi alloy. For example, with the same size and/or the same weight, a higher torque can be generated, which would be advantageous when used in electric or hybrid-powered motor vehicles.

The DE 10 2018 112 491 A1 discloses a highly permeable soft magnetic FeCo alloy with 5 to 25 weight percent Co, which can have a maximum permeability of more than 17,000 and lower hysteresis losses. Due to the lower Co content, the raw material costs of this alloy are reduced compared to an alloy based on 49% Fe, 49% Co, 2% V. At the same time, this alloy does not have any significant order setting, so that, in contrast to alloys with over 30% by weight of Co, it can be cold-rolled without an upstream quenching process, which simplifies large-scale production.

EP 3 730 286 A1 discloses a method for producing a soft magnetic FeCo alloy, wherein a precursor for electrical insulation is completely coated with an oxide layer, for example with a layer of magnesium methylate or zirconium propylate, which is converted into an insulating oxide layer during a heat treatment and wherein the precursor in a sheet metal package is used.

However, it is desirable to be able to achieve these good magnetic properties more reliably, particularly in large-scale production.

The task is therefore to provide a soft magnetic CoFe alloy with a lower Co content and good magnetic properties, which can be produced more reliably using large-scale manufacturing processes.

According to the invention, there is provided a method for producing a soft magnetic alloy, comprising the following. A precursor is provided which has a composition which essentially consists of 2% by weight ≤ Co ≤ 30% by weight 0.3% by weight ≤ V ≤ 5.0% by weight 0% by weight ≤ Cr ≤ 3.0% by weight 0% by weight ≤ Si ≤ 5.0% by weight 0% by weight ≤ Mn ≤ 5.0% by weight 0% by weight ≤ Al ≤ 3.0% by weight 0% by weight ≤ Ta ≤ 0.5% by weight 0% by weight ≤ Ni ≤ 1.0% by weight 0% by weight ≤ Mon ≤ 0.5% by weight 0% by weight ≤ Cu ≤ 0.2% by weight 0% by weight ≤ Nb ≤ 0.25% by weight 0% by weight ≤ Ti ≤ 0.05% by weight 0% by weight ≤ Ce ≤ 0.05% by weight 0% by weight ≤ Approx ≤ 0.05% by weight 0% by weight ≤ Mg ≤ 0.05% by weight 0% by weight ≤ C ≤ 0.02% by weight 0% by weight ≤ Zr ≤ 0.1% by weight 0% by weight ≤ O ≤ 0.025% by weight 0% by weight ≤ S ≤ 0.015% by weight

The rest is iron, and up to 0.2% by weight of other melt-related impurities. The preliminary product has a phase transition from a BCC phase region into a BCC/FCC mixed region to an FCC phase region, with the phase transition between the BCC phase region and the BCC/FCC mixed region occurring at a first transition temperature T _α/α as the temperature increases _+γ and as the temperature continues to rise, the transition between the BCC/FCC mixed region and the FCC phase region takes place at a second transition temperature T _α+γ/γ , where T _α+γ/γ > T _α/α+γ .

In some embodiments, the difference T _α+γ/γ - T _α/α+γ is less than 45K, preferably less than 25K.

Other impurities include B, P, N, W, Hf, Y, Re, Sc, Be, other lanthanides except Ce.

This precursor is partially coated with a ceramic-forming layer, with 20% to 80% of the total surface of the precursor remaining free of the ceramic-forming layer. The partially coated precursor is then heat treated. This heat treatment is also described as final annealing, since the mechanical deformation steps, for example hot rolling and/or cold rolling, with which the preliminary product is produced have been completed, so that the preliminary product is no longer mechanically deformed after the heat treatment.

In some embodiments, the precursor is planar and has a first surface and a second opposing surface, with at least between 20% and 80%, preferably between 30% to 70%, particularly preferably between 50% to 70% of the first surface and between 20% and 80%, preferably between 30% to 70%, particularly preferably between 50% and 70% of the second surface remains free of the ceramic-forming layer. The preliminary product can be a strip, a sheet or a sheet with its final contour.

• Aufheizen des Vorprodukts und danach
• Wärmebehandeln des Vorprodukts in einer ersten Stufe mit einer Gesamtzeitdauer t₁, wobei in der ersten Stufe das Vorprodukt bei einer Temperatur in einem Temperaturbereich zwischen T_α+γ/γ und T₁ wärmebehandelt wird, und danach Abkühlen des Vorprodukts bis Raumtemperatur, wobei T₁ oberhalb T_α+γ/γ liegt und t₁ die Gesamtzeitdauer bei Temperaturen oberhalb T_α+γ/γ bedeutet.

In one embodiment, heat treating includes:

• Heating the preliminary product and afterwards
• Heat treating the preliminary product in a first stage with a total time period t ₁ , wherein in the first stage the preliminary product is heat treated at a temperature in a temperature range between T _α+γ/γ and T ₁ , and then cooling the preliminary product to room temperature, where T ₁ is above T _α+γ/γ and t ₁ means the total time period at temperatures above T _α+γ/γ .

• Aufheizen des Vorprodukts und danach
• Wärmebehandeln des Vorprodukts in einer ersten Stufe mit einer Gesamtzeitdauer t₁, wobei in der ersten Stufe das Vorprodukt bei einer Temperatur im Temperaturbereich zwischen T_α+γ/γ und T₁ wärmebehandelt wird und danach
• Abkühlen des Vorprodukts auf eine Temperatur T₂, und direkt anschließend
• Wärmebehandeln des Vorprodukts in einer zweiten Stufe bei der Temperatur T₂ für eine Zeitdauer t₂, und danach
• Abkühlen des Vorprodukts auf Raumtemperatur, wobei T₁ > T₂ ist, T₁ oberhalb T_α+γ/γ liegt und T₂ unterhalb T_α/α+γ liegt, und t₁ die Gesamtzeitdauer bei Temperaturen oberhalb T_α+γ/γ bedeutet.

In an alternative embodiment, heat treating includes:

• Heating the preliminary product and afterwards
• Heat treating the preliminary product in a first stage with a total time period t ₁ , wherein in the first stage the preliminary product is heat treated at a temperature in the temperature range between T _α+γ/γ and T ₁ and thereafter
• Cooling the preliminary product to a temperature T ₂ , and immediately afterwards
• Heat treating the preliminary product in a second stage at the temperature T ₂ for a period of time t ₂ , and thereafter
• Cooling of the preliminary product to room temperature, where T ₁ > T ₂ , T ₁ is above T _α+γ/γ and T ₂ is below T _α/α+γ , and t ₁ is the total time period at temperatures above T _α+γ/γ means.

In both exemplary embodiments, the heat treatment is carried out at least temporarily in a hydrogen-containing atmosphere, during which the exposed parts of the surface of the precursor are in direct contact with the hydrogen-containing atmosphere.

The temperature T ₁ denotes a temperature at which the soft magnetic alloy is in the FCC phase region, and T ₂ denotes a temperature at which the soft magnetic alloy is in the BCC phase region. The actual temperature of the oven can deviate from the value T ₁ or T ₂ , so that T ₁ includes temperatures T ₁ with a maximum deviation of +/-20 ° C, which are above T _α+γ/γ and T ₂ temperatures T ₂ with a maximum deviation of +/-20°C, which is below T _α/α+γ .

During large-scale production, several primary products are conventionally heat-treated at the same time. These precursors are conventionally coated with a ceramic layer so that the precursors do not weld together during heat treatment. For example, preliminary products in the form of a strip or sheet are stacked on top of each other and thus heat treated. A ceramic layer is typically placed between the strips or sheets to prevent the strips or sheets from welding together.

It was surprisingly found that for these alloys when heat treated at temperatures above the transition temperature T _{α + γ / γ} and thus in the FCC or γ phase region, the magnetic properties depend on the proportion of exposed surface of the precursor. If a portion of the precursor is in direct contact with the hydrogen-containing atmosphere at least temporarily during the heat treatment, the good magnetic properties can be achieved more reliably even with larger batches. However, when the surface is completely covered with the ceramic layer, which is actually advantageous in order to avoid the risk of welding of the precursors, it was found that the measured magnetic properties are not so good. It was found that the good magnetic properties are related to the formation of a texture in the soft magnetic alloy and that this formation of the texture depends on the proportion of the exposed surface that is in direct contact with a hydrogen-containing atmosphere during the final annealing.

The preliminary product is therefore only partially coated with the ceramic-forming layer, which is converted into a ceramic layer during the subsequent heat treatment. The applied coating can, for example, have a sol with ceramic nanoparticles distributed in an organic matrix, or metal ions from a metal oxide or metal hydroxide, so that a ceramic layer is not yet present in the applied form. During the heat treatment, a ceramic layer is formed, which can also contain a metal oxide and/or a metal hydroxide and only partially covers the surface.

In some exemplary embodiments, the preliminary product has the shape of a sheet with a first surface and a second opposite surface, with at least between 20% and 80%, preferably between 30% to 70%, particularly preferably between 50% to 70% of the first surface and between 20% and 80%, preferably between 30% to 70%, particularly preferably between 50% and 70% of the second surface is free of the ceramic-forming layer.

In one embodiment, the composition is defined in more detail and essentially consists of 5% by weight ≤ Co ≤ 25% by weight 0.3% by weight ≤ V ≤ 5.0% by weight 0% by weight ≤ Cr ≤ 3.0% by weight 0% by weight ≤ Si ≤ 3.0% by weight 0% by weight ≤ Mn ≤ 3.0% by weight 0% by weight ≤ Al ≤ 3.0% by weight 0% by weight ≤ Ta ≤ 0.5% by weight 0% by weight ≤ Ni ≤ 0.5% by weight 0% by weight ≤ Mon ≤ 0.5% by weight 0% by weight ≤ Cu ≤ 0.2% by weight 0% by weight ≤ Nb ≤ 0.25% by weight 0% by weight ≤ Ti ≤ 0.05% by weight 0% by weight ≤ Ce ≤ 0.05% by weight 0% by weight ≤ Approx ≤ 0.05% by weight 0% by weight ≤ Mg ≤ 0.05% by weight 0% by weight ≤ C ≤ 0.02% by weight 0% by weight ≤ Zr ≤ 0.1% by weight 0% by weight ≤ O ≤ 0.025% by weight 0% by weight ≤ S ≤ 0.015% by weight

The remainder is iron, with Cr+Si+Al+Mn ≤ 3.0% by weight, and up to 0.2% by weight of other melt-related impurities. Other impurities include B, P, N, W, Hf, Y, Re, Sc, Be, other lanthanides except Ce.

In some exemplary embodiments, the ceramic-forming layer is applied to the preliminary product in the form of a structure. The structure can be formed by a pattern of stripes or dots or a network. The ceramic-forming layer can be applied selectively to the surface of the preliminary product, for example by means of printing or profile printing or profile rolling. Alternatively, a mask with openings can first be applied to the surfaces, the ceramic-forming layer can be applied over a large area and the mask can be removed so that only the parts of the surface that are exposed in the openings in the mask are coated.

In some embodiments, the maximum width of the coated areas is less than 2mm, preferably less than 1.2mm, particularly preferably less than 0.8mm. It has been found that the magnetic properties can be achieved more reliably if the area of the coated parts is limited.

In some exemplary embodiments, the preliminary product has the shape of a ribbon. The structured coating can be applied to one or both sides of the belt.

In some exemplary embodiments, between 30% to 70%, preferably between 50 to 70%, of the total surface of the preliminary product remains free of the ceramic-forming layer.

In some embodiments, after the heat treatment, the ceramic layer formed has a metal oxyhydrate and/or metal oxide and/or metal hydroxide.

In addition to the partial coating or instead of the partial coating, in some exemplary embodiments the preliminary product is covered with a ceramic powder, a ceramic plate, or a metal plate during the heat treatment. The ceramic powder may have Al ₂ O ₃ or MgO or ZrO ₂ or a mixture of ZrO ₂ , SiO ₂ and Al ₂ O ₃ . The ceramic plate may have Al ₂ O ₃ or MgO or ZrO ₂ or a mixture of ZrO ₂ , SiO ₂ and Al ₂ O ₃ .

A ceramic plate or a metal plate can be used to ensure the planarity of precursors in the form of sheets or strips or stacks of strips or sheets.

It was also found that a covering with ceramic powder alone, i.e. without a partial coating of the preliminary product, also enables the surface to be in direct contact with the hydrogen-containing atmosphere and thus a texture and good magnetic properties can be achieved. However, covering with ceramic powder alone is not so practical in large-scale commercial manufacturing.

It was surprisingly found that contact with the hydrogen-containing atmosphere during the cooling phase of the final annealing has a greater effect on texture formation and magnetic properties than during the heating phase and in the first annealing stage above T _α+γ/γ . In particular, it was found that the use of a hydrogen-containing atmosphere with more than 5 vol.% hydrogen during cooling over the temperature range from T _{α + γ / γ} to T _{α / α + γ} is advantageous for adjusting the soft magnetic properties.

In some exemplary embodiments, during the heating of the preliminary product at least in the temperature range from T _{α / α + γ} to T ₁ , the heat treatment takes place in a protective gas atmosphere with less than 5 vol.% hydrogen, preferably less than 1 vol.% hydrogen and the cooling of T ₁ at least in the temperature range from T _α+γ/γ to T _α/α+γ in a hydrogen-containing atmosphere with more than 5 vol.% hydrogen. This embodiment can also be used with a completely uncoated preliminary product.

In some exemplary embodiments, after cooling the preliminary product to a temperature T ₂ , where T ₂ is below T _α/α+γ , the preliminary product is kept at the temperature T ₂ for a period of time t ₂ , and only then cooled further.

In some exemplary embodiments, the heat treatment of the preliminary product in the first stage is carried out for the total period of time t ₁ in a protective gas atmosphere with less than 5 vol.% hydrogen, preferably less than 1 vol.% hydrogen.

In some exemplary embodiments, the preliminary product is cooled from T ₁ to T ₂ in a hydrogen-containing atmosphere with more than 5 vol.% hydrogen.

In some exemplary embodiments, the preliminary product is cooled from T ₁ to room temperature in a hydrogen-containing atmosphere with more than 5 vol.% hydrogen.

Argon or nitrogen alone, or a mixture of argon or nitrogen with less than 5 vol.% hydrogen can be used as the protective gas.

In some embodiments, the hydrogen-containing atmosphere with more than 5 vol.% hydrogen has an initial dew point of less than -40 ° C.

In some embodiments, the hydrogen-containing atmosphere with more than 5 vol.% hydrogen also includes argon.

In some exemplary embodiments, the following values apply: T _α+γ/γ ≤ T ₁ ≤ T _α+γ/γ + 50 ° C and 15 minutes ≤ t ₁ ≤ 10 hours, preferably 15 minutes ≤ t ₁ ≤ 4 hours and 700 ° C ≤ T ₂ ≤ 1050°C and 30 minutes ≤ t ₂ ≤ 20 hours.

In some exemplary embodiments, the heat treatment further comprises a subsequent second final annealing in a hydrogen-containing protective gas atmosphere at a maximum temperature that is below the first transition temperature T _{α / α + γ} .

An ideal soft magnetic material has no preferred magnetic direction. As soon as a preferred direction exists, the magnetization processes are more difficult because additional energy is necessary to create the magnetic moments out of the preferred direction. This direction-dependent energy is referred to as magnetocrystalline anisotropy energy in crystalline soft magnetic materials. Due to the existing symmetry relationships, the magnetocrystalline anisotropy energy in a cubic crystal is expressed by a series expansion, the first-order coefficient of which is referred to as the anisotropy constant K ₁ .

FeCo alloys have a preferred magnetic direction that influences the magnetic properties. In FeCo alloys with Co contents below 42% by weight, the cube edges <100> are the magnetically light axes, while the space diagonals <111> represent the magnetically hard axes. The surface diagonals <110> have medium soft magnetic properties.

For use in transformers, stators and rotors, soft magnetic materials in the form of thin sheets are used to reduce the formation of eddy currents during magnetic reversal. For such applications, it is therefore advantageous if the magnetically favorable cube edges <100> lie, if possible, in the plane of the sheet metal, while the magnetically unfavorable orientations <111> lie, if possible, not in the plane of the sheet metal.

If there is a preferred orientation of the crystal orientations in a polycrystalline material, it is called a texture. In the case of the cube surface texture, the cube surfaces {001} also lie in the sheet metal plane, but the orientation of the cube edges is not fixed, but varies evenly between the rolling direction and the direction transverse to the rolling direction. The cube face texture is therefore described by {100}<uvw>. This results in significantly more homogeneous magnetic properties. In principle, this texture represents the best variant for use in stators or rotors, at least for the alloys considered here, where K ₁ is always significantly greater than zero.

It was surprisingly found that in practice the formation of the {111}<uvw> texture has a large influence on the magnetic properties, so that the proportion of this texture should be reduced in order to reliably achieve good magnetic properties. A high proportion of the {100}<uvw> texture alone does not lead to the best magnetic properties if at the same time the proportion of the {111}<uvw> texture is too high. It was surprisingly found that the proportion of the {111}<uvw> texture is reduced by the method according to the invention, ie the use of a final annealing which is temporarily carried out in the FCC phase region, the partial covering of the surface and the temporary annealing under hydrogen leads to suppression of the formation of the {111}<uvw> texture.

In some exemplary embodiments, the alloy after heat treatment has an area proportion of a {111}<uvw> texture that is a maximum of 13%, preferably a maximum of 6%, with grains with a tilt of up to +/- 10 ° or even better up to +/- 15° relative to the nominal crystal orientation of the {111} planes.

In some embodiments, the alloy after heat treatment has an area fraction of a {100}<uvw> cube surface texture that is at least 30%, preferably at least 50%, with grains having a tilt of up to +/- 15° or even better up to +/- 10° relative to the nominal crystal orientation of the {100} planes.

In some exemplary embodiments, the heating rate is 10 K/h to 1000 K/h, preferably 20 K/h to 100 K, over at least the temperature range from T _α+γ/γ to T _{α /α+γ} , preferably 900° C. to T 1 /H.

In some exemplary embodiments, the cooling rate over at least the temperature range from T _{α / α + γ} to T _{α + γ / γ} is preferably T ₁ to 900 ° C 10K / h to 200 K / h and preferably 20K / h to 100 K / h.

In some embodiments, the preliminary product is weighted with an additional weight and the preliminary product is subjected to heat treatment with the weight.

In some exemplary embodiments, the weight of the weighting is at least 10%, preferably at least 30%, of the weight of the preliminary product.

In some exemplary embodiments, after the heat treatment, the preliminary product is subjected to a further heat treatment in an atmosphere containing oxygen or water vapor to form the electrically insulating layer.

In some exemplary embodiments, the preliminary product has the shape of several stacked metal sheets, metal sheets in final contour or one or more sheet metal packages. In the case of the shape of several stacked metal sheets, after the heat treatment, in some exemplary embodiments, at least one laminated core is manufactured from the stacked metal sheets by eroding, laser cutting or water jet cutting.

In the case of the shape of sheets in the final contour, the sheets are glued to a laminated core after heat treatment using an insulating adhesive, or oxidized on the surface to create an insulating layer and then glued to the laminated core or laser welded or coated with an inorganic-organic hybrid coating and then to form a sheet metal package, for example by stacking and gluing or laser welding.

In some exemplary embodiments, after heat treatment, the soft magnetic alloy has a maximum permeability µ _max ≥ 6,000 and/or an electrical resistance ρ ≥ 0.25 µΩm, hysteresis losses P _Hys ≤ 0.07 J/kg at an amplitude of 1.5 T and/ or a coercivity H _c of ≤ 0.8 A/cm and/or induction B ₂₀ ≥ 1.70 T at 20 A/cm.

In some exemplary embodiments, after heat treatment, the soft magnetic alloy has a maximum permeability μ _max ≥ 10,000 and/or an electrical Resistance ρ ≥ 0.25 µΩm and/or hysteresis losses P _Hys ≤ 0.06 J/kg at an amplitude of 1.5 T and/or a coercivity H _c of ≤ 0.5 A/cm and induction B ₂₀ ≥ 1, 74 T at 20 A/cm.

In some exemplary embodiments, after heat treatment, the average grain size of the soft magnetic alloy is at least 100 μm, preferably at least 200 μm, most preferably at least 250 μm.

In some embodiments, after heat treatment, the average grain size of the soft magnetic alloy is 1.0 to 2.0 times, preferably 1.0 to 1.5 times, the strip thickness.

According to the invention, a soft magnetic alloy is also made, which essentially consists of 2% by weight ≤ Co ≤ 30% by weight 0.3% by weight ≤ V ≤ 5.0% by weight 0% by weight ≤ Cr ≤ 3.0% by weight 0% by weight ≤ Si ≤ 5.0% by weight 0% by weight ≤ Mn ≤ 5.0% by weight 0% by weight ≤ Al ≤ 3.0% by weight 0% by weight ≤ Ta ≤ 0.5% by weight 0% by weight ≤ Ni ≤ 1.0% by weight 0% by weight ≤ Mon ≤ 0.5% by weight 0% by weight ≤ Cu ≤ 0.2% by weight 0% by weight ≤ Nb ≤ 0.25% by weight 0% by weight ≤ Ti ≤ 0.05% by weight 0% by weight ≤ Ce ≤ 0.05% by weight 0% by weight ≤ Approx ≤ 0.05% by weight 0% by weight ≤ Mg ≤ 0.05% by weight 0% by weight ≤ C ≤ 0.02% by weight 0% by weight ≤ Zr ≤ 0.1% by weight 0% by weight ≤ O ≤ 0.025% by weight 0% by weight ≤ S ≤ 0.015% by weight

Rest iron, and up to 0.2% by weight of other melt-related impurities, provided. The soft magnetic alloy also has a surface area of a {111}<uvw> texture which is a maximum of 13%, preferably a maximum of 6%, with grains with a tilt of up to +/- 10° or preferably up to +/ - 15° from the nominal crystal orientation can be included.

The combination of this composition and texture makes it possible to ensure the magnetic properties of the soft magnetic alloy more reliably.

In one exemplary embodiment, the soft magnetic alloy has an area proportion of a {100} <uvw> cube surface texture that is at least 30%, preferably at least 50%, with grains with a tilt of up to +/- 15 ° or preferably up to + /- 10° from the nominal crystal orientation can be included. The magnetic properties are achieved with the combination of an area proportion of a {111}<uvw> texture that is a maximum of 13%, preferably a maximum of 6% and an area proportion of a {100}<uvw> cube surface texture that is at least 30%, preferably at least 50% is further improved.

In some embodiments, the average grain size of the {100}<uvw>-oriented grains is at least 1.5 times as large, preferably at least 2.0 times as large, as the average grain size of the {111}<uvw>-oriented grains.

In some exemplary embodiments, the area fraction of the {100}<uvw>-oriented grains is at least 3 times as large, preferably at least 7 times as large, as the area fraction of the {111}<uvw>-oriented grains.

In some exemplary embodiments, the soft magnetic alloy has a maximum permeability μ _max ≥ 6,000 and/or an electrical resistance ρ ≥ 0.25 μΩm, hysteresis losses P _Hys ≤ 0.07 J/kg at an amplitude of 1.5 T and/or a coercive field strength H _c of ≤ 0.8 A/cm and/or induction B ₂₀ ≥ 1.70 T at 20 A/cm, or a maximum permeability µ _max ≥ 10,000 and/or an electrical resistance ρ ≥ 0.25 µΩm and/or hysteresis losses P _Hys ≤ 0.06 J/kg at an amplitude of 1.5 T and/or a coercive field strength H _c of ≤ 0.5 A/cm and induction B ₂₀ ≥ 1.74 T at 20 A/cm.

• 10 Gew.-% ≤ Co ≤ 20 Gew.-%, vorzugsweise 15 Gew.-% ≤ Co ≤ 20 Gew.-%, gilt, und
• 0,5 Gew.-% ≤ V ≤ 4,0 Gew.-%, vorzugsweise 1,0 Gew.-% ≤ V ≤ 3,0 Gew.-%, vorzugsweise 1,3 Gew.-% ≤ V ≤ 2,7 Gew.-%, gilt, und/oder
  0,1 Gew.-% ≤ Cr ≤ 2,0 Gew.-%, vorzugsweise 0,2 Gew.-% ≤ Cr ≤ 1,0 Gew.-%, vorzugsweise 0,3 Gew.-% ≤ Cr ≤ 0,7 Gew.-%, gilt, und/oder
• 0,1 Gew.-% ≤ Si ≤ 2,0 Gew.-%, vorzugsweise 0,15 Gew.-% ≤ Si ≤ 1,0 Gew.-%, vorzugsweise 0,2 Gew.-% ≤ Si ≤ 0,5 Gew.-%, gilt und/oder
• die Summenformel 0,1 Gew.-% ≤ Cr+Si+Al+Mn ≤ 1,5 Gew.-%, vorzugsweise 0,2 Gew.-% ≤ Cr+Si+Al+Mn ≤ 0,6 Gew.-%, gilt.

In some embodiments, the composition of the soft magnetic alloy is defined in more detail, where

• 10% by weight ≤ Co ≤ 20% by weight, preferably 15% by weight ≤ Co ≤ 20% by weight, applies, and
• 0.5% by weight ≤ V ≤ 4.0% by weight, preferably 1.0% by weight ≤ V ≤ 3.0% by weight, preferably 1.3% by weight ≤ V ≤ 2, 7% by weight, applies, and/or
  0.1% by weight ≤ Cr ≤ 2.0% by weight, preferably 0.2% by weight ≤ Cr ≤ 1.0% by weight, preferably 0.3% by weight ≤ Cr ≤ 0, 7% by weight, applies, and/or
• 0.1% by weight ≤ Si ≤ 2.0% by weight, preferably 0.15% by weight ≤ Si ≤ 1.0% by weight, preferably 0.2% by weight ≤ Si ≤ 0, 5% by weight, applies and/or
• the molecular formula 0.1% by weight ≤ Cr+Si+Al+Mn ≤ 1.5% by weight, preferably 0.2% by weight ≤ Cr+Si+Al+Mn ≤ 0.6% by weight , applies.

In one embodiment, the composition consists essentially of 5% by weight ≤ Co ≤ 25% by weight 0.3% by weight ≤ V ≤ 5.0% by weight 0% by weight ≤ Cr ≤ 3.0% by weight 0% by weight ≤ Si ≤ 3.0% by weight 0% by weight ≤ Mn ≤ 3.0% by weight 0% by weight ≤ Al ≤ 3.0% by weight 0% by weight ≤ Ta ≤ 0.5% by weight 0% by weight ≤ Ni ≤ 0.5% by weight 0% by weight ≤ Mon ≤ 0.5% by weight 0% by weight ≤ Cu ≤ 0.2% by weight 0% by weight ≤ Nb ≤ 0.25% by weight 0% by weight ≤ Ti ≤ 0.05% by weight 0% by weight ≤ Ce ≤ 0.05% by weight 0% by weight ≤ Approx ≤ 0.05% by weight 0% by weight ≤ Mg ≤ 0.05% by weight 0% by weight ≤ C ≤ 0.02% by weight 0% by weight ≤ Zr ≤ 0.1% by weight 0% by weight ≤ O ≤ 0.025% by weight 0% by weight ≤ S ≤ 0.015% by weight

The rest is iron, whereby Cr+Si+Al+Mn ≤ 3.0% by weight, and up to 0.2% by weight of other melt-related impurities may be present. Other impurities include B, P, N, W, Hf, Y, Re, Sc, Be, other lanthanides except Ce.

A laminated core consisting of several stacked electrically insulated sheets of a soft magnetic alloy according to one of the above exemplary embodiments is also provided. In some exemplary embodiments, the laminated core has a filling factor F ≥ 90%, preferably > 94%, whereby a greater power density can be guaranteed.

In some exemplary embodiments, the laminated core has at least two sheets, each of which has a thickness of 0.05 mm to 0.50 mm, and the electrical insulation layer, e.g. a ceramic layer or an oxide layer, between adjacent sheets has a thickness of 0.1 μm to 5, 0µm, preferably 0.5µm to 3.0µm.

The invention also relates to the use of the laminated core according to one of the preceding exemplary embodiments in an electrical machine, for example as a stator tooth, stator segment, or in a stator and/or rotor of an electric motor and/or a generator, and/or in a transformer and/or in an electromagnetic actuator.

Fig. 1

zeigt beispielhaft ein Phasendiagramm für eine der erfindungsgemäßen Legierungsvarianten mit unterschiedlichen Co-Gehalten und mit 2% V.

Fig. 2

zeigt eine schematische Darstellung eines Ablaufs einer erfindungsgemäßen Wärmebehandlung.

Fig. 3

zeigt die magnetische Induktion B20, wobei B20 die Induktion bei einer Feldstärke H von 20 A/cm bezeichnet, für frei liegende und mit Pulver abgedeckte Proben.

Fig. 4

zeigt die Koerzitivfeldstärke H_c für frei liegende und mit Pulver abgedeckte Proben.

Fig. 5

zeigt eine Neukurve B(H) für frei liegende und mit Pulver abgedeckte Proben.

Fig. 6

zeigt die Kurve der Permabilität µ(H) für frei liegende und mit Pulver abgedeckte Proben.

Fig. 7

zeigt (001) Polfiguren für frei liegende und mit Pulver abgedeckte Proben.

Fig. 8

zeigt Orientierungsdichteverteilungsfunktionen (ODF) für frei liegende und mit Pulver abgedeckte Proben.

Fig. 9

zeigt den Zusammenhang zwischen der magnetischen Induktion B(20 A/cm) und dem Flächenanteil A(001).

Fig.10

zeigt den Zusammenhang zwischen der magnetischen Induktion B(20 A/cm) und dem Flächenanteil A(111).

Fig. 11

zeigt den Zusammenhang zwischen der magnetischen Induktion B20 und der mittleren Korngröße GS aller Körner.

Fig. 12

zeigt den Zusammenhang zwischen Induktion B20 = B(20 A/cm) und der mittleren Korngröße GS(001) aller Körner mit (001)-Orientierung.

Fig. 13

zeigt den Zusammenhang zwischen Induktion B20 = B(20 A/cm) und der mittleren Korngröße GS(111) aller Körner mit (111)-Orientierung.

Fig. 14

zeigt die jeweiligen Flächenanteile A(001) und A(111).

Fig. 15

zeigt beispielhafte Abbildungen für Oberflächenmuster einer strukturierten Beschichtung.

Fig. 16

zeigt EDX-Analyse für eine streifige Beschichtung der Oberfläche vor und nach einer Glühung.

Fig. 17

zeigt anhand eines Beispiels das Verhältnis der Streifenbreite zur Korngröße im geglühten Zustand.

Embodiments and examples will now be explained using the drawings.

Fig. 1

shows an example of a phase diagram for one of the alloy variants according to the invention with different Co contents and with 2% V.

Fig. 2

shows a schematic representation of a heat treatment process according to the invention.

Fig. 3

shows the magnetic induction B20, where B20 denotes the induction at a field strength H of 20 A/cm, for exposed samples and those covered with powder.

Fig. 4

shows the coercivity H _c for exposed samples and those covered with powder.

Fig. 5

shows a new curve B(H) for exposed and powder-covered samples.

Fig. 6

shows the permability curve µ(H) for exposed and powder-covered samples.

Fig. 7

shows (001) pole figures for exposed and powder-covered samples.

Fig. 8

shows orientation density distribution functions (ODF) for exposed and powder-capped samples.

Fig. 9

shows the relationship between the magnetic induction B(20 A/cm) and the surface area A(001).

Fig.10

shows the relationship between the magnetic induction B(20 A/cm) and the surface area A(111).

Fig. 11

shows the connection between the magnetic induction B20 and the average grain size GS of all grains.

Fig. 12

shows the connection between induction B20 = B(20 A/cm) and the average grain size GS(001) of all grains with (001) orientation.

Fig. 13

shows the connection between induction B20 = B(20 A/cm) and the average grain size GS(111) of all grains with (111) orientation.

Fig. 14

shows the respective area shares A(001) and A(111).

Fig. 15

shows exemplary images for surface patterns of a structured coating.

Fig. 16

shows EDX analysis for a streaky coating on the surface before and after annealing.

Fig. 17

uses an example to show the relationship between strip width and grain size in the annealed state.

Figure 1 shows a schematic representation of a phase diagram of an FeCo alloy with the composition of the precursor, in which the BCC phase region, which is also referred to as a ferritic α region, the BCC / FCC mixed region, which is also referred to as a two-phase α + γ region , and the FCC phase region, which is also referred to as the austenitic γ region, as well as the transition temperatures T _α/α+γ and T _α+γ/γ are shown. The transition temperatures T _α/α+γ and T _α+γ/γ as well as the difference T _α+γ/γ - T _α/α+γ depend on the composition of the alloy.

According to the invention, the FeCo alloy has a composition which essentially consists of: 2% by weight ≤ Co ≤ 30% by weight 0.3% by weight ≤ V ≤ 5.0% by weight 0% by weight ≤ Cr ≤ 3.0% by weight 0% by weight ≤ Si ≤ 5.0% by weight 0% by weight ≤ Mn ≤ 5.0% by weight 0% by weight ≤ Al ≤ 3.0% by weight 0% by weight ≤ Ta ≤ 0.5% by weight 0% by weight ≤ Ni ≤ 1.0% by weight 0% by weight ≤ Mon ≤ 0.5% by weight 0% by weight ≤ Cu ≤ 0.2% by weight 0% by weight ≤ Nb ≤ 0.25% by weight 0% by weight ≤ Ti ≤ 0.05% by weight 0% by weight ≤ Ce ≤ 0.05% by weight 0% by weight ≤ Approx ≤ 0.05% by weight 0% by weight ≤ Mg ≤ 0.05% by weight 0% by weight ≤ C ≤ 0.02% by weight 0% by weight ≤ Zr ≤ 0.1% by weight 0% by weight ≤ O ≤ 0.025% by weight 0% by weight ≤ S ≤ 0.015% by weight

The rest is iron, and up to 0.2% by weight of other melt-related impurities.

The alloy has a phase transition from a BCC phase region into a BCC/FCC mixed region to an FCC phase region, with the phase transition between the BCC phase region and the BCC/FCC mixed region at a first transition temperature T _α/α as the temperature increases _+γ and as the temperature continues to rise, the transition between the BCC/FCC mixed region and the FCC phase region takes place at a second transition temperature T _α+γ/γ , where T _α+γ/γ > T _α/α+γ , as in Figure 1 is shown. The composition is selected so that the difference T _α+γ/γ - T _α/α+γ is less than 45K, preferably less than 25K.

In order to improve the magnetic properties of the alloy, a heat treatment is carried out on a precursor of the above composition, which is known as shot annealing since it takes place after all the deformation steps such as hot rolling and cold rolling. The preliminary product can have the shape of a strip or a sheet. After heat treatment, the alloy is not subjected to further cold working. According to the invention, this final annealing includes a heat treatment at a temperature at which the alloy is found in the FCC phase region.

Figure 2 shows a schematic representation of a heat treatment process according to the invention according to an exemplary embodiment for an alloy with the above-mentioned composition. The alloy can be heat treated in the form of a strip or sheet. The process is divided into the three processes “heating” (E), “holding” (H) and “cooling down” (A). Within these processes, a distinction is also made as to which crystallographic phase the precursor product is in. α denotes the ferritic region (BCC), γ denotes the austenitic region (FCC) and α+γ denotes the two-phase (BCC/FCC) mixed region.

In the first heating phase E(α) the material is completely in the ferritic phase. After passing through the phase transition T(α→α+γ), the material passes through the two-phase mixing region α+γ in phase E(α+γ). This relatively narrow area is limited at the top by the phase transition T(α+γ→γ), so that the heating phase E(γ) is reached. The holding phase H(γ) is located entirely in the austenitic γ region. After the cooling A(γ) from the first annealing stage in the austenitic γ region and reaching the phase transition T(γ→α+γ), the cooling A(α+γ) takes place temporarily in the two-phase region α+γ, which passes downwards the phase transition T(α+γ→α) is limited. During the remaining cooling time A(α), the material is in the ferritic α region. According to the invention, the maximum temperature of the heat treatment is in the FCC area.

In the Figure 2 The temperature ramps shown during heating or cooling serve to achieve a defined heating or cooling of the entire annealing material in a technical process. The temperatures at which the phase transitions T(α→α+γ) and T(α+γ→γ) take place during heating and at which the phase transitions T(α+γ→γ) and T(γ→α+γ) occur during cooling, are basically identical. In practice, however, due to finite heating and cooling rates, shifts in the range of a few degrees Celsius can occur, ie during heating Certain temperatures may be slightly higher than the temperatures determined during cooling.

It was surprisingly found that the build-up of the precursors during the heat treatment, which includes a step above the transition temperature T _α+γ/γ , has an effect on the measured magnetic properties. To further investigate this observation, samples with uncovered surfaces in a free-hanging setup or samples covered with a powder were placed in the three phase regions identified in the Figure 1 are shown schematically, heat treated. The results are listed in Table 1 which illustrates this.

For this purpose, stamping rings measuring 28.5 mm x 20.0 mm were made from 0.35 mm thick strip of the above-mentioned alloy VACOFLUX X1 with a nominal composition: 16.8% Co, 2.3% V, no Si additive, balance Fe manufactured. The samples were annealed in a tube furnace with dry hydrogen flushing. Some samples (labeled H) were heat treated hanging so that all surfaces are in contact with the hydrogen. Some samples (labeled P) are covered with powder and heat treated. Table 1 sample Annealing (H ₂ ) T _max in °C Phase (T _max ) Construction H1 4h 910°C 910 α hanging H2 H1 + 4h 930°C 930 α hanging H3 H2 + 4h 950°C 950 α hanging H4 H3 + 4h 970°C 970 α+γ hanging H5 H4 + 4h 990°C 990 α+γ/γ hanging H6 H5 + 4h 1010°C 1010 γ hanging H7 H5 + 4h 1030°C 1030 γ hanging H8 H6 + 4h 1050°C 1050 γ hanging P1 4h 910°C 910 α powder P2 P1 + 4h 930°C 930 α powder P3 P2 + 4h 950°C 950 α powder P4 P3 + 4h 970°C 970 α+γ powder P5 P4 + 4h 990°C 990 α+γ/γ powder P6 P5 + 4h 1010°C 1010 γ powder P7 P6 + 4h 1030°C 1030 γ powder P8 P7 + 4h 1050°C 1050 γ powder

The Figure 3 shows the magnetic induction B20, where B20 denotes the induction at 20 A/cm, and Figure 4 the coercivity Hc measured for these samples. After annealing in the γ region, for example at 1010°C, the freely hanging rings annealed without powder and with very good hydrogen flushing from series H show induction values B20 similar to the samples that were annealed in the α region. The coercive field strength is even in the range of the best glow in the α region. Thanks to the narrow two-phase region, it was possible to completely avoid a deterioration in the soft magnetic characteristics, despite phase transformation during heating and cooling for the samples annealed in the γ region.

The rings of the P series, which are annealed close together in the powder, however, show a completely different behavior. After annealing in the γ region (1010°C) there are significantly increased induction values B20, which are more than 100 mT above the best value after annealing in the α region. The coercivity H _c continues to decrease as the temperature increases. This means that after annealing at 1050°C, even the best Hc value obtained after annealing in the α region is undercut.

Texture studies are carried out using EBSD (Electron Backscatter Diffraction). Four states with the composition 17.25% Co 1.49% V 0.23% Si, balance Fe were examined, see Table 2. Sample A as a reference was not annealed. Sample B was annealed at 930°C in the α region as a reference. Sample C was According to the invention, annealed in the γ region at 1000 ° C, with the sample rings lying in ceramic annealing powder. As a reference, sample D was annealed like sample C, but without powder, but hanging freely. Table 2 sample T in °C T in h the atmosphere Construction A - - - - b 930 4 H2 powder C 1000 4 H2 powder D 1000 4 H2 without powder, hanging freely

In addition to the magnetic samples (punching rings 28.5 mm x 20.0 mm), two strips measuring 50 mm x 32 mm were made from each sample, which were also annealed in the powder or exposed and on which EBSD tests were carried out to determine the texture. In the case of exposed annealed sheets, the texture determination was carried out on the upper surface that was exposed during annealing.

The measurement of the magnetic properties included the new curve B(H) or µ(H), the coercivity H _c and the remanence B _r . The area proportions A(001) and A(111) of the grains with (001) and (111) orientation at a maximum tilt of +/- 10° were determined from the ESBD measurements. In addition, the surface area A'(001) was determined for samples B, C and D, for which the tilt was max. +/- 15°. The average grain size GS was also determined from the EBSD measurements. Table 3 shows the magnetic characteristics, the area proportions of the texture and the average grain size GS. For sample A, no texture components were initially determined and only the parameter A' (001) was redetermined. Additionally are in Figure 5 or. Figure 6 the new curve B(H) or the permability curve µ(H) is shown. Table 3 sample B20 in T Br in T A(001) in % A'(001) in % A(111) in % GS in µm A 0.528 0.70 - 22.9 - - b 1,682 1.16 21.7 30.8 29.5 161 C 1,783 1.41 51.9 69.2 4.3 917 D 1,643 1.31 27.9 38.1 13.6 459

The unannealed sample A has, as expected for a heavily cold-formed material, very low inductions and a very large H _c > 1500 A/m.

Sample B, annealed in the α region, shows significantly improved properties, but with B(20 A/cm) of 1.682 T, it is still below the target value of at least 1.70 T. It can be assumed that there is little or no texture here , since the area proportions A(001) and A(111) of the grains of different orientations are similar in size.

Sample C, annealed in powder in the γ region, shows the best magnetic properties. The induction B (20 A/cm) is extremely high at 1.783 T. The area share A(001) is larger and the area share A(111) is smaller. The very low coercivity of 31 A/m and the high permeability also indicate large grain diameters.

The sample D, annealed in the γ region without powder, on the other hand, has a significantly lower induction B(20 A/cm) < 1.65 T, similar to sample B. In contrast to this, the H _c is lower and µ _max higher, which is probably due to the increased grain growth due to the higher annealing temperature and the associated higher diffusion rate in the γ region. The area portion A(001) is smaller and the area portion A(111) is larger than in sample C. It was therefore found that the differences in the magnetic properties between the two test series can be attributed to the different annealing textures.

In Figure 7 The pole figures 001 are shown as an example, as they show the essential symmetries. The other pole figures 110 and 111 can be found in the appendix. A1 denotes the rolling direction (RD, “ Rolling Direction ”) and A2 the transverse direction (TD, “ Transverse Direction ”). (Axis A1 || Rolling direction)

The rolling texture (sample A) changes due to the annealing in the α region (sample B) into an annealing texture that contains both (001) components (intensities in the middle and at the outermost edge) as well as (111) components ( intensities along half the diameter). The pole figure looks completely different after annealing in the γ region (sample C): This clearly corresponds to a cube surface texture (001). After annealing in the γ region without powder (sample D), not only parts with (001) and (111) orientations are found, but also a variety of other diffuse orientations that are in the range of the typical annealing and rolling textures of bcc materials .

In Figure 8 the ODF (orientation density distribution function) at ϕ2 = 45° for samples A to D are shown.

Sample A (reference, rolled-hard) has all α-fiber components, i.e. from {001}<110> to {112}<110> to {111}<110>. This is typical of a cold rolled texture of bcc metals. In addition, parts of the γ-fiber can be seen, i.e. {111}<121> and {111}< 112>.

In sample B (reference, annealing in the α region), this α fiber is shifted by 20° at an angle ϕ1, which corresponds to a rotation of the cube layer in the sheet plane. The parts of the γ-fiber that can be seen in the unannealed state are now intensified. In these orientations, the magnetically hard axis, the spatial diagonal <111>, lies in the plane of the sheet metal.

Sample C according to the invention shows all proportions of the θ fiber, ie from {001}<110> to {001}<010> to {001}<110>. This corresponds to a magnetically advantageous cube surface texture in which the magnetically easy axis, the cube edge <001>, lies in the sheet plane and the orientation of these cube edges varies between the rolling direction and the direction transverse to the rolling direction. Furthermore, there are hardly any shares of the magnetically unfavorable γ fiber. These magnetically unfavorable components could be significantly reduced through annealing.

The reference sample D, which was also annealed in the γ region, although with a different annealing structure, also has cube layers. In contrast to sample C, these are significantly less pronounced and are also more oriented. In addition, compared to sample C, there are significantly higher proportions of the magnetically unfavorable γ-fiber, i.e. proportions of {111}<121> and {111}<112>.

From the texture investigations it becomes apparent that the cause of the high induction values B20 of state C according to the invention lies, on the one hand, in the formation of a cube surface texture, i.e. the formation of orientations of the α-fiber, and on the other hand in the avoidance of portions of the γ-fiber.

To further investigate the dependence of the magnetic characteristics on the texture, a series of samples of the VACOFLUX X1 alloy from the same batch 7410163B were annealed at different temperatures, holding times, atmospheres and setups. The parameters are given in Table 4. Table 4 Condition T in °C Phase (T) t in h the atmosphere Construction annotation A1 750 α 2 H2 powder - A2 930 α 4 H2 powder - A3 955 α + γ 4 H2 powder - B1 1000 γ 4 H2 hanging - B2 1000 γ 4 vacuum stack - B3 1000 γ 4 H2 stack coated with HITCOAT, on both sides C1 970 γ 2 H2 powder - C2 970 γ 2 H2 powder with pre-heating 2h 750°C, H ₂ C3 970 γ 4 H2 powder - C4 1000 γ 0.5 H2 powder - C5 1000 γ 2 H2 powder - C6 1000 γ 4 H2 powder - C7 1000 γ 20 H2 powder - C8 1100 γ 4 H2 powder - C9 1000 γ 4 H2 stack Annealing with coated intermediate layers C10 1000 γ 4 H2 stack coated with HITCOAT, one side

Table 4 shows a list of the annealing carried out with general conditions and information about the sample. All samples were made from the same batch 74101563B, tape thickness 0.20 mm.

• A: Glühung im α-Gebiet oder im α+γ-Mischgebiet (Referenz)
• B: Glühung im γ-Gebiet mit unvorteilhaften Parametern (Referenz)
• C: Glühung im γ-Gebiet mit vorteilhaften Parametern (erfindungsgemäß)

The samples are divided into three rows depending on the annealing temperature T:

• A: Annealing in the α region or in the α+γ mixed region (reference)
• B: Annealing in the γ region with unfavorable parameters (reference)
• C: Annealing in the γ region with advantageous parameters (according to the invention)

The assignment of the phase is based on the phase transitions α→α+γ at 944°C and α+γ→γ at 965°C determined for this batch, which were determined from a DSC measurement based on the 1st onset during cooling or heating .

The annealing time, ie the holding time t at the annealing temperature, was varied between 0.5 h and 20 h, with most annealing being carried out with a holding time of 4 h. The atmosphere used was primarily dry hydrogen H ₂ with an initial dew point of -40°C or lower. In one case, example B2, annealing was carried out in a vacuum under a pressure of 10 ^-1 mbar.

The annealing structure was also varied: The annealing “in powder” was carried out in ceramic powder made from Al ₂ O ₃ . In the "hanging" setup, the annealed samples were threaded onto a ceramic tube at a short distance from each other. For annealing in a “stack”, the samples (rings or sheets) were placed on top of each other. This stack of sheets was placed on a base plate and weighted down with a cover plate. In order to avoid welding of the touching rings in uncoated samples, as in example C9, rings coated on both sides with HITCOAT were used as intermediate layers. Before the magnetic measurement, the intermediate layers, which only served as an annealing aid, were sorted out so that the magnetic values refer to the uncoated rings.

The tests were evaluated by measuring the magnetic properties and texture, see Table 6 where B20 = B(20 A/cm), B100 = B(100 A/cm), A(001) = area fraction (001), A( 111) = area fraction (111), GS = average grain size of all orientations, GS(001) = average grain size of the (001) orientation, GS(111) = average grain size of the (111) orientation. Table 5 Condition B20 in T Br in T H _c in A/m A(001) in % A(111) in % GS in µm GS(001) in µm GS(111) in µm A1 1,715 1.22 166 14.8 31.2 20 19.1 19.6 A2 1,682 1.16 58.5 21.7 29.5 161 139 176 A3 1,680 1.31 61.6 20.0 27.9 186 161 195 B1 1,643 1.31 39.3 27.9 13.6 459 373 505 B2 1,679 1.25 43.0 21.7 13.1 417 452 - B3 1,687 1.36 46.4 15.7 14.4 455 346 475 C1 1,794 1.38 34.6 65.6 1.7 502 511 135 C2 1,786 1.35 36.2 61.9 2.9 391 401 238 C3 1,802 1.41 30.5 60.8 1.8 728 305 762 C4 1,786 1.43 34.3 59.7 3.7 526 327 554 C5 1,784 1.41 30.9 61.3 3.2 789 459 824 C6 1,752 1.34 30.2 38.0 5.9 1238 764 1390 C7 1,710 1.32 31.1 16.9 10 1619 1021 1284 C8 1,711 1.31 30.3 10.0 12 2030 929 1226 C9 1,745 1.39 33.4 48.2 5.2 409 300 432 C10 1,720 1.36 40.4 25.2 9.6 451 361 472

For the magnetic measurements, punched rings measuring 28.5 mm x 20.0 mm were annealed. After annealing, the rings were wound with secondary and primary windings in a plastic trough and statically measured based on IEC 60404-4. The parameters determined are the maximum permeability µ _max , the magnetic inductions B20 = B(20 A/cm) and B100 = B(100 A/cm), the remanence B _r and the coercive field strength H _c .

To determine the texture, sheets measuring 20 mm x 30 mm were annealed in the same way. The sheets were measured using EBSD after annealing as described above. The determined parameters are the area proportions A(001) and A(111) of the orientations (001) and (111), respectively.

Figure 9 shows the connection between the magnetic induction B(20 A/cm) and the surface area A(001) of the magnetically advantageous (001) orientation. The line drawn illustrates the course of row C according to the invention. In the reference samples of rows A and B, an area proportion of max. 27.9% is formed. In the samples of series C according to the invention, however, area proportions A(001) of up to 65.6% are achieved. It can also be seen that a high proportion of the cube surface texture also leads to a high magnetic induction B (20 A/cm).

The area proportions were determined at a maximum tilt of +/- 10°. If one were to tolerate a larger tilt of +/- 15°, this would result in an area share A(001) of up to 82%.

Figure 10 shows the connection between the magnetic induction B(20 A/cm) and the surface area A(111) of the magnetically unfavorable (111) orientation. The reference samples from rows A and B show area proportions A(111) of at least 13%, which is associated with a low induction B20 of less than 1.70 T. Only state A1, which was annealed at a very low temperature of 750 ° C, shows an induction B20 of 1.715 T. However, since the grain size of this sample is extremely small, this also results in a very high coercivity Hc of 166 A/m.

The states of series C according to the invention, on the other hand, consistently show inductions greater than 1.70 T. The reason for this lies in the small proportion of the magnetically unfavorable (111) orientation. The samples with the lowest area proportions A(111) show the highest induction values B20. The line shown illustrates the course of series C according to the invention.

Figure 11 shows the connection between the magnetic induction B20 and the average grain size GS of all grains, regardless of their orientation. The samples from series A, which were annealed in the ferritic α region, only reach grain sizes of up to 200 µm. The samples from series B annealed in the austenitic γ region, on the other hand, show significantly larger grains in the range of 400 to 500 µm. Nevertheless, the desired high induction B20 does not occur here. The large grain itself is not the cause of the high B20 induction.

Only the samples from series C, annealed under advantageous conditions in the γ region, show induction values B20 > 1.70 T. For example, sample C1 has a grain size of 502 µm, i.e. about 0.5 mm, which is similar to the values of Row B lies, and in contrast shows a very high induction B20 of 1.794 T.

Furthermore, it is found that an increase in grain size within series C has a negative influence on the induction B20. The direct factors influencing the grain size are primarily the annealing temperature and the annealing time. For example, sample C8 has an average grain size of 2030 µm, ie approximately 2 mm, due to an increased annealing temperature of 1100°C. The induction B20 is therefore reduced to 1.711 T. The sample C7, which was only heat treated at 1000 ° C but with an extended annealing time of 20 h, also has coarse grains with an average grain size of 1619 µm, ie about 1.6 mm, and only shows a low induction B20 of 1.71 T

The states C9 and C10 according to the invention are listed separately in the figures, since the changed glow structure leads to a deviation here. The samples were annealed “in a stack”, i.e. the lamellas were placed on top of each other and weighted down with a cover plate. The inductions B20 achieved are still very good, i.e. greater than 1.70 T. However, compared to states C1 to C8, which were annealed in ceramic annealing powder, the inductions are lower for the same grain size. With such a structure, the annealing parameters must be changed in order to achieve optimal conditions.

The results on grain size can also be processed separately according to the orientation of the grains, cf. Figure 12 for the magnetically advantageous (001) orientation or Figure 13 for the magnetically unfavorable (111) orientation. The interpretation is analogous to the interpretation of the total grain size. Figure 12 shows the connection between induction B20 = B(20 A/cm) and mean grain size GS(001) of all grains with (001) orientation. The line drawn illustrates the course of the states C1 to C8 of the series according to the invention. Figure 13 shows the connection between induction B20 = B(20 A/cm) and mean grain size GS(111) of all grains with (111) orientation. The line drawn illustrates the course of the states C1..C8 of the series C according to the invention.

In summary, in order to achieve the highest possible induction, it is desirable to carry out annealing in the γ region while keeping the grain size to a minimum. Possibilities for this include, for example, lowering the annealing temperature as close as possible to the phase transition α+γ→γ or reducing the annealing time. When it comes to the annealing time, it should be noted that the heating and cooling ramps have an indirect influence on the annealing time, i.e. a slow heating or cooling ramp also leads to a longer residence time in the γ region. The glow structure also has a decisive influence.

Avoiding excessive grain growth also makes sense in order to ensure that the magnetization is as uniform as possible in all directions when a cube surface texture is present.

Finally, there is also a direct connection between the formation of the preferred cube position (001) compared to the magnetically unfavorable spatial diagonal (111). In Figure 14 the respective area proportions A(001) and A(111) are plotted against each other. The states A annealed in the α region show a very high proportion of (111) of over 25% and at the same time a low proportion of (001) of less than 25%. Both are disadvantageous from a magnetic point of view. The states B annealed under unfavorable boundary conditions in the γ region already show a lower proportion (111), which is still relatively high at over 13%. This is also the cause of the low magnetic induction of these states, i.e. B20 < 1.7 T.

Figure 14 shows the connection between the area proportions A(001) and A(111). The states C according to the invention, which were annealed in the γ region under advantageous boundary conditions, show an area proportion (111) that is consistently below 13%. A further distinction was made in the figure: The points marked C* (open symbols) correspond to the examples C7, C8 and C10. These are borderline in terms of parameters, ie C7 was annealed with an extremely long holding time of 20 hours and C8 was annealed at a very high annealing temperature of 1100°C. In both cases this results in strong growth of the unfavorable (111) orientation. Sample C10 was provided on one side with an annealing-resistant, continuous ceramic coating, so that when it cooled and passed through the phase transition α+γ→α, the advantageous (001) texture could only form on the uncoated side. Accordingly, the (111) content in all three samples is still 9.6 to 12%. Furthermore, the (001) proportion is not yet pronounced and is relatively low at 10.0 to 25.2%. Examples C* show a B20 value between 1.70 T and 1.72 T, which is not optimal, but is still significantly better than states A and B not according to the invention.

The points marked C (closed symbols) correspond to the remaining examples of group C. They are highly optimized with regard to annealing, i.e. the annealing temperature was above the phase transition α+γ→γ, but not more than 100 ° C above this temperature, and the holding time was a maximum of 4 hours. In addition, it was in In all examples, the surface was without a continuous coating, so that the (001) texture was able to form on both sides when cooling and passing through the phase transition α+γ→α. Samples C accordingly show the lowest area proportions (111) of less than 6%. At the same time you can see a clear proportion of the cube position (001), from at least 25% up to 66%. This favorable preferred orientation is also the reason that B20 inductions of over 1.72 T and up to 1.80 T were consistently achieved here.

The influence of the structure and the atmosphere during the heat treatment as well as in the heating, holding and cooling phases of the heat treatment was examined in more detail.

Argon was used in a first series of tests. Table 6 shows the magnetic values B20, B100, µ _max , Hc and Br of the alloy according to the invention after heat treatment in a tube furnace in ceramic annealing powder (exemplary embodiments R1, E1, E2, E3, E3, E4) and hanging (exemplary embodiments E5, E6, E7). With the freely hanging annealing setup, the samples are threaded onto a ceramic or thin metal rod and are rinsed very well from all sides.

All samples were first heated to 1000°C as quickly as possible and then kept at a temperature of 1000°C for 4 hours. The cooling took place up to an intermediate temperature of 900°C at 30 K/h and then via oven cooling, i.e. faster than 30 K/h.

The annealing atmosphere was varied in the individual phases, ie in addition to dry hydrogen (H ₂ ) with a dew point of -40°C or lower, argon (Ar) of quality 5.0 was sometimes used. For the purpose of better illustration, the same gas atmosphere was used within the individual stages, ie "E" denotes the atmosphere in the heating phases E(α), E(α+γ) and E(γ), "H" the holding phase H( γ) and "A" the cooling phases A(γ), A(α+γ) and A(α). Table 6 RE Construction E H A B20 in T B100 in T µ _max H _c in A/m Br in T A1 R powder Ar Ar Ar 1,672 1,955 8,887 46.6 1.28 A2 E powder H2 H2 H2 1,791 2,031 15,075 31.8 1.40 A3 E powder Ar H2 H2 1,804 2,042 16,306 34.2 1.45 A4 E powder Ar Ar H2 1,798 2,038 15,095 38.9 1.47 A5 E powder H2 H2 Ar 1,740 2,006 13,654 35.1 1.37 A6 E hanging H2 H2 H2 1,691 1,973 12,570 35.1 1.32 A7 E hanging Ar H2 H2 1,661 1,944 11,174 37.7 1.31 A8 E hanging Ar Ar H2 1,768 2,013 13,263 36.5 1.42

Sample A1 corresponds to an annealing process not according to the invention, which only took place under argon. Sample 1 has a very low induction B20 of 1.672 T, a very low remanence Br of 1.28 T and a high coercivity Hc of 46.6 A/m.

In the case of sample A2 according to the invention, the annealing was carried out completely while being flushed with dry hydrogen. The induction B20 of sample A2 is very high at 1.791 T, as is the remanence Br at 1.40 T. At the same time, there is a very low coercivity Hc of only 31.8 A/m.

In state A3 according to the invention, the inert gas argon was used in the heating phase and dry hydrogen was used for the holding and cooling phase. The characteristic values B20, B100, µ _max and especially Br are higher than the corresponding values of sample A2, which was only annealed under hydrogen. By using Ar alone in the heating phase, an improvement in the squareness of the hysteresis loop could be achieved. It is assumed that the phase transformations upon heating, ie α→α+γ→γ, also depend on the surface energy and thus on the gas atmosphere. In contrast to cooling down, it seems It is an advantage when heating that the surface is not completely reduced.

For sample A4, not only the heating phases but also the holding phase H(γ) were carried out under argon. Here too there are very good magnetic values, similar to sample A3.

Examples A3 and A4 show that the partial use of Ar leads to a slight increase in the coercivity compared to annealing only under hydrogen (sample A2). This is probably due to less pronounced grain growth. If lower coercive field strengths are necessary in the application, downstream grain growth can be initiated by an additional holding stage during cooling in the α-region or a subsequent heat treatment in the α-region.

For sample A5, the reduction of the surface area by H ₂ was carried out in the heating and holding stage. The cooling took place under argon. Sample A5 has an induction B20 of 1.74 T and a high remanence Br of 1.37 T.

The other exemplary embodiments A6 to A8 were annealed hanging. For sample A6, the annealing took place entirely under pure hydrogen. Due to the suspended structure, the sample was washed very well, which led to a strong reduction in contamination throughout the entire annealing period. Compared to the structure in the powder, the induction B20 is low at 1.691 T. The high remanence Br of 1.32 T and the high maximum permeability of 12,570 show that the magnetically unfavorable (111) orientation could be partially suppressed.

For sample A7, the protective gas Ar was used during heating and dry hydrogen was used for the rest of the annealing, ie for the holding stage and cooling. The resulting magnetic values are even slightly worse than those of sample E5, which was annealed purely under hydrogen. However, compared to the non-inventive reference sample R1, the remanence Br and the maximum permeability μ _max are still slightly higher.

In sample A8 according to the invention, H(γ) argon was used not only during heating but also during the four-hour holding phase. Only when it cooled down did the switch to hydrogen occur. This process made it possible to significantly improve the magnetic properties: the induction B20 reaches 1.768 T, the remanence Br is 1.42 T.

The exemplary embodiments show that the use of H ₂ during cooling during annealing leads to advantageous crystal orientations. H ₂ should therefore preferably be available during the cooling phase.

In principle, it is possible to use hydrogen only in the cooling phase A(α+γ). In practice, when using Ar, it is possible to switch to pure hydrogen at the end of the holding time, so that the complete cooling takes place under hydrogen. This ensures that, even in an industrial process, there is enough time to sufficiently flush the entire material to be annealed with hydrogen and thus ensure the reduction of the oxides near the surface.

The use of an inert protective gas during the heating or holding phase can be advantageous in order to avoid the formation of an unfavorable intermediate structure. Structures with very good flushing require a longer flushing with argon than structures with poorer flushing in powder or in the stack.

The results can be transferred to mixed gases, ie instead of using pure argon, a mixture of argon and hydrogen can also be used. For example, an H ₂ /Ar gas mixture with 20 vol.% argon and 80 vol.% hydrogen or an Ar/H ₂ gas mixture with 80 vol.% hydrogen and 20 vol.% argon can be used. The exact mixing ratio can be adjusted depending on the annealing setup, annealing duration and flushing conditions.

In a second series of experiments, nitrogen (N ₂ ) is used in the heating or holding phase. In contrast to Ar, N ₂ is not inert and annealing under nitrogen can lead to the formation of vanadium nitrides due to the vanadium content of the claimed alloys.

Vanadium nitrides preferentially precipitate at grain boundaries and thereby prevent grain growth. Therefore, the formation of nitrides in soft magnetic alloys is in principle undesirable, since a fine-grain structure leads to a high coercivity Hc. In addition, the presence of precipitates generally leads to an increase in Hc, since the nonmagnetic precipitates act as interference points for the domain wall movements.

In the context of the present invention, the suppression of grain growth during the heating and holding phase is to be assessed as a positive effect, i.e. there is therefore the possibility that an intermediate structure that is advantageous for the further cooling process is formed.

However, these precipitates are thermodynamically relatively stable at temperatures of 1000 ° C or lower, i.e. it is not possible to dissolve the nitrides again during the cooling phase under dry hydrogen. Therefore, no further grain growth can occur, which is disadvantageous with regard to the coercivity.

A compromise can be to limit the amount of precipitates, for example by limiting the time during which nitrogen is used for annealing, or by only mixing nitrogen with hydrogen in small quantities.

The variation of the annealing atmosphere was carried out analogously to the experiments on annealing with argon and is shown in Table 7. Table 7 HE Construction E H A B20 in T B100 in T µ _max H _c in A/m Br in T B1 R powder N2 N2 N2 1,617 1,933 1,624 274 1.44 B2 E powder H2 H2 H2 1,791 2,031 15,075 31.8 1.40 B3 E powder N2 H2 H2 1,808 2,038 13,901 44.1 1.50 B4 R powder N2 N2 H2 1,637 1,939 2,427 202 1.37 B5 R powder H2 H2 N2 1,600 1,928 1,670 225 1.37 B6 E hanging H2 H2 H2 1,691 1,973 12,570 35.1 1.32 B7 E hanging N2 H2 H2 1,818 2,045 13,797 40.1 1.52 B8 R hanging N2 N2 H2 1,637 1,939 2,427 202 1.37

Sample B1 is an annealing process not according to the invention, which took place continuously under N ₂ purging. The induction B20 is extremely low at 1.617 T, as is the maximum permeability µ _max at only 1.623. The coercive field strength is very high at 274 A/m.

Sample B2 according to the invention represents the reference annealing under pure hydrogen and corresponds to sample A2 from the examples with argon.

For sample B3, nitrogen was used in the heating phase, but dry hydrogen was used for the rest of the annealing phase. The nitrogen has a positive effect on the magnetic characteristics B20 and Br, which are 1.808 T and 1.50 T, respectively, higher than with pure hydrogen annealing of sample B2. At the same time, the short exposure time of the nitrogen minimized the negative effect on the coercivity Hc, i.e. the Hc at 44.1 A/m is still in an acceptable range for most applications.

In the case of sample B4 not according to the invention, both the heating and the holding phase were carried out under nitrogen. The long exposure time of the nitrogen results in very poor magnetic characteristics, in particular a very low maximum permeability of 2427 and a very high coercivity Hc of 202 A/m.

In the case of sample B5, which is also not according to the invention, the surface was initially reduced by H ₂ in the heating and holding stage, but the subsequent cooling took place under nitrogen. The resulting magnetic values are similarly poor as in condition B4. This exemplary embodiment shows that during the cooling process the surface of the material must not only be free of oxides, but also of nitrides or other coatings.

The other exemplary embodiments B6 to B8 were annealed while hanging.

The sample B6 was annealed in a freely suspended manner under pure hydrogen and corresponds to the sample A6 from the exemplary embodiments with argon. Due to the good flushing, there is a low induction B20 of only 1.691 T despite the high maximum permeability.

For sample B7, the sample was also suspended annealed, but nitrogen was used during the heating phase. Surprisingly, there is a very significant increase in the induction B20, which is 1.818 T, as well as a very high remanence Br of 1.52 T. It is assumed that the nitrides formed in the heating phase inhibit grain growth, so that despite very good By flushing with hydrogen during the holding phase, a favorable intermediate structure is created that allows the preferred formation of the (001) orientation or the suppression of the (111) orientation during the cooling phase.

Sample B8, which is not according to the invention, was annealed analogously to sample B7, with not only the heating phase but also the holding phase being carried out under nitrogen. Similar to example B4, the negative influence of nitrogen predominates and the desired magnetic properties can no longer be adjusted. In particular, the coercive field strength of 202 A/m is significantly too high.

The exemplary embodiments show that nitrogen in small amounts can be advantageous for increasing the induction B20 and the remanence Br. If nitrogen is offered in too large an amount or over a too long period of time during annealing, the deterioration in Hc is too great.

Optionally, after the cooling phase A(α+γ), a holding stage can still take place in the ferritic α region. The purpose of this is to further promote grain growth, which leads to a further reduction in the coercivity Hc and a further increase in the maximum permeability.

This optional holding stage can alternatively also take place in a second heat treatment, which should, however, take place completely in the α region. This subsequent heat treatment makes it possible to improve only part of the originally annealed material, e.g. the part that does not yet meet all magnetic requirements in the γ region after the first annealing.

Table 8 shows some exemplary embodiments that show the influence of a second heat treatment in the α region. Lines C1 to C5 each correspond to the states after the heat treatment in the γ region with a four-hour holding stage at 1000 ° C. For lines C1' to C5', the same samples were subjected to a second heat treatment in the α region with a four hour hold at 930°C. Table 8 HE Glow E H A B20 in T B100 in T µ _max H _c in A/m Br in T C1 E 4h 1000°C H2 H2 H2 1,691 1,973 12,570 35.1 1.32 C1' E 4h 930°C H2 H2 H2 1,695 1,967 14,805 33.0 1.38 C2 E 4h 1000°C Ar H2 H2 1,804 2,042 16,306 34.2 1.45 C2` E 4h 930°C H2 H2 H2 1,802 2,039 20,238 29.2 1.51 C3 E 4h 1000°C N2 H2 H2 1,776 2,017 6,267 84.2 1.47 C3' E 4h 930°C H2 H2 H2 1,780 2,019 7,747 77.1 1.54 C4 R 4h 1000°C N2 N2 N2 1,617 1,933 1,624 274 1.44 C4' R 4h 930°C H2 H2 H2 1,458 1,731 1,761 250 1.32 C5 R 4h 1000°C Ar Ar Ar 1,672 1,955 8,887 46.6 1.28 C5' R 4h 930°C H2 H2 H2 1,674 1,953 11,344 44.1 1.37

Example C1 corresponds to sample A6 already presented. Due to the suspended annealing under pure hydrogen, there is a high permeability µ _max and a low Hc, but the induction B20 is relatively low. Even after-annealing in the α region as in sample C1' does not lead to any significant improvement in B20.

Example C2 corresponds to the sample A3 already explained, which was annealed under Ar in the heating phase. The partial use of Ar results in very good magnetic values. The subsequent heat treatment in the α region, as shown in example C2', leads to a further significant reduction in Hc to 29.2 A/m and a significant increase in the maximum permeability to 20238.

Examples C3 and C3` show the influence of a second heat treatment on a sample that was annealed under nitrogen in the heating phase of the first heat treatment. All of the magnetic characteristics listed improve slightly, although even after this heat treatment the maximum permeability is still relatively low at 7747 and the coercive field strength is still relatively high at 77.1 A/m.

The non-inventive reference example C4 was annealed only under nitrogen in the first heat treatment. By omitting hydrogen, the magnetic characteristics could not be achieved and the long presence of nitrogen during annealing results in a very high Hc of 274 A/m. This is probably due to the formation of nitrides on the surface and in the material. These can no longer be dissolved by the second heat treatment at 930 ° C carried out in example C4', so that the magnetic values can no longer be improved.

The reference example C5 corresponds to the example A1, ie annealing not according to the invention only under Ar. Even the afterglow does not result in any significant improvement in the magnetics in example C5`.

The examples illustrate that the second heat treatment is particularly effective for the samples in which the first heat treatment in the γ region was partially carried out under Ar.

In large-scale production, the preliminary products are conventionally coated with a ceramic layer so that the preliminary products do not stick together and there is electrical insulation between the layers to minimize eddy current losses. In the case of a preliminary product in the form of a strip or sheet, the preliminary products are stacked so that a ceramic layer is arranged between the strips or sheets.

It was surprisingly found that during heat treatment at temperatures above the transition temperature T _α+γ/γ and thus in the FCC or γ phase region, the magnetic properties depend on the proportion of exposed surface of the precursor. However, if a portion of the precursor is in direct contact with the hydrogen-containing atmosphere at least temporarily during the heat treatment, the good magnetic properties can be achieved more reliably. It was found that these good magnetic properties are related to the formation of a texture in the soft magnetic alloy.

The preliminary product is therefore only partially coated with the ceramic-forming layer, which is converted into a ceramic layer during the subsequent heat treatment. In the case of a sheet or strip shape, one or both of the opposing major surfaces is only partially coated so that portions of one or both of the opposing major surfaces are free of the coating during heat treatment.

The preliminary product is partially coated with a ceramic-forming layer, with 20% to 80% of the total surface of the preliminary product remaining free of the ceramic-forming layer. The partially coated precursor is then heat treated. The applied coating can, for example, have a sol with metal ions, see above that ceramic is not yet present in the applied form. It is also possible for the applied layer to have ceramic nanoparticles in the form of a sol.

In some exemplary embodiments, the preliminary product has the shape of a sheet with a first surface and a second opposite surface, with at least between 20 and 80%, preferably between 30% to 70%, particularly preferably between 50% to 70% of the first surface and between 20% and 80%, preferably between 30% to 70%, particularly preferably between 50% to 70% of the second surface is free of the ceramic layer containing metal oxide or metal hydroxide.

Ceramic strip coatings are used for strips made of Fe-Co in order to avoid welding of touching metal surfaces during the necessary magnetic final annealing of sheets or sheet metal cuts. Examples include the DL1 coating based on Mg methylate, which converts into magnesium oxide during annealing, and the HITCOAT coating based on Zr propylate, which converts into zirconium oxide during annealing. After annealing, both coatings are present as a thin film on both sides with a thickness of typically 0.5 µm or thinner on each side. Since the coatings are applied in a thin liquid using a solvent, they are distributed evenly over the belt surface and form a continuous coating.

An influence of the coating on the soft magnetic properties was surprising for the Fe-Co material class. On the contrary, despite the relatively small thickness, the coatings mentioned result in a significant improvement in magnetization losses, since the electrical insulation leads to a reduction in eddy currents.

However, if you now coat a strip made of the VACOFLUX

This can be detailed in the examples given in Table 9, which shows the magnetic properties after final annealing for 4 hours at 1000°C depending on the coating. The table shows the magnetic characteristics µ _max , B3, B20, B100, Hc and Br for different coating variants. All tests were carried out on batch 7410163B, which has already been described in several places, in a strip thickness of 0.20 mm. Rings measuring 28.5 mm x 20.0 mm were punched from the strips and annealed in a chamber furnace at 1000 ° C and a holding time of 4 hours while being flushed with dry hydrogen. The structure was "stacked", meaning that around 20 rings were stacked on top of each other, placed on a ceramic base plate and covered with a ceramic cover plate so that they were still flat after annealing. Table 9 # variant µ _max B3 in M B20 in T B100 in T Hc in A/m Br in T E uncoated 12,814 1,537 1,770 2,020 37.3 1.43 F HITCOAT on both sides 10,240 1,437 1,687 1,963 46.4 1.36 G HITCOAT on both sides, thin on one side 10,793 1,468 1,719 1,986 43.3 1.34 H HITCOAT one-sided 12,467 1,493 1,737 2,001 39.8 1.39

Example E represents the uncoated reference sample. After annealing, very good soft magnetic characteristics are obtained. In particular, the very high induction B20 of 1.77 T is an indicator of a very high proportion of cube surface texture (001)[uvw], which was successfully formed by annealing in the γ region. Due to the advantageous orientation, there is also a very high maximum permeability of almost 13,000 and a low coercivity Hc of only 37.3 A/m.

In Example F, the strip was provided with a zirconium propylate coating on both sides before punching. During annealing, this transforms into flat zirconium oxide. This dense coverage suppressed the formation of the cube surface texture. As a result, the induction B20 is only 1.687 T. Also At 46.4 A/m, the coercivity Hc is significantly higher than the value of the reference sample E.

In order to enable the formation of the cube surface texture and still achieve sufficient separation of the sheets during annealing, one of the above-mentioned coatings (DL1 or HITCOAT) can still be used, as long as they are sufficiently thin or as long as the coating only takes place on one side of the strip .

In Example G, the tape was first coated regularly, i.e. on both sides, and then the coating on one side was chemically removed using a solvent. The free side was then recoated with a highly diluted coating solution so that one side of the band was regularly coated, with a coating thickness between 100 nm and 500 nm, while the second side only had a very thin coating, which is in the range below 100 nm. After annealing, the soft magnetic properties of this sample are better than those of the regular, thicker coating in Example B. At 1.719 T, the induction B(20) is still well below the reference value of the uncoated sample. In addition, the very thin coating leads to initial adhesions due to annealing, i.e. the main function of the coating, layer separation, is severely impaired.

In example H, a tape was first coated regularly and the coating on one side was chemically removed using a solvent. After annealing, magnetic properties emerge which indicate that a significant proportion of the cube surface texture could be formed here. In addition to a high induction B20 of 1.737 T, a maximum permeability of almost 12,500 is also achieved, which almost corresponds to reference state A.

The examples show that the gradation applies when setting a high induction B20: $$\mathrm{<figure-callout\; id="20"\; label="b"\; filenames="imgf0001.png,imgf0009.png"\; state="\{\{state\}\}">b</figure-callout>}20\left(\mathrm{uncoated}\right)\ge \mathrm{<figure-callout\; id="20"\; label="b"\; filenames="imgf0001.png,imgf0009.png"\; state="\{\{state\}\}">b</figure-callout>}20\left(\mathrm{one-sided}\right)\ge \mathrm{<figure-callout\; id="20"\; label="b"\; filenames="imgf0001.png,imgf0009.png"\; state="\{\{state\}\}">b</figure-callout>}20\left(\mathrm{bilaterally}\right)$$

The one-sided coating enables sheets to be annealed in stacks. The prerequisite for this is that the sheets are positioned in a position-oriented manner, meaning that, for example, the coated side of the sheets always faces upwards and there is therefore no contact between the uncoated and coated sheet sides. If necessary, it is necessary to use a sheet metal coated on both sides as a separating layer from the base plate and the cover plate, if used. Other methods can also be used to achieve one-sided coil coating:
In one example, when applying the coating using rollers, the coating is squeezed off on one side under pressure by a ground roller. This process can be used on a system that is also used for double-sided coating.

In another example, the strip is coated regularly, i.e. the coating is evenly present on both sides. The coating is then removed on one side by mechanical brushing.

The results of annealing in annealing powder show that the presence of a surface coating with loose ceramic particles is entirely compatible with the formation of a cube surface texture by annealing in the γ region in order to reliably achieve good magnetic properties. However, loose powders are unfavorable for industrial production because the powder particles can come loose during handling. On the one hand, this causes quality problems and can also pose a health risk when inhaled. Furthermore, ceramic oxides are very hard and therefore lead to very rapid wear of the tools in stamping processes. If the particles are too large, the cutting gaps in punching tools can become clogged and damage them. Even with powders where the median size distribution is in the small µm range, individual grains with grain sizes > 10 µm lead to the filling factor within a laminated core increasing disproportionately.

A well-adhering coating was therefore developed that consists of plastic-bound, very fine particles and in which the particles only convert into a thermally stable ceramic oxide during the final annealing.

Specifically, particles made of aluminum oxide hydroxide (boehmite) with a size of 10 to 300 nm, preferably 20 to 150 nm, are used in a binder based on aqueous acrylate dispersions, which also contains wetting agents and ammonia as a further component for adjusting and controlling the pH value and dispersed by topping up with demineralized water. The resulting dispersion is applied to both sides of a VACOFLUX X1 strip with a thickness of 0.20 mm using profiled rollers in a continuous process. This initially creates a stripy structure in the roll gap. Through the ceramic content and optionally through an additional rheologically effective additive, the viscosity of the coating dispersion is adjusted so that these stripes do not run or only run very little when they exit the roll gap. In the subsequent drying section, the coated strip is dried with warm air (280°C) so that the binder forms a film. The tape then has a well-adhering coating with a stripy structure.

• Die Oberfläche ist nur teilbedeckt.
• Die Partikelgröße liegt im Bereich unter einem µm.
• Die feinen Partikel sind in der Vorstufe gebunden.
• Das Aluminium liegt in der Vorstufe als weiches Böhmit vor (Moshärte 3,5) und wandelt erst bei der Schlussglühung in hartes Al₂O₃ um (Moshärte 9).

This coating achieves the following:

• The surface is only partially covered.
• The particle size is in the range below one µm.
• The fine particles are bound in the precursor.
• The aluminum is present in the preliminary stage as soft boehmite (Mos hardness 3.5) and only converts into hard Al ₂ O ₃ (Mos hardness 9) during the final annealing.

By varying the process, e.g. by changing the viscosity by adjusting the boehmite content and/or by changing the roller profile, other structures with similar properties can in principle also be set on the belt.

In Figure 15 exemplary images of surface patterns are shown, with examples a, b representing the reference states and the other examples being surfaces according to the invention.

Example a shows a sheet without coating. While the formation of a cube surface texture is possible here, uncoated sheets cannot be easily annealed because they weld together at high temperatures of typically over 900 ° C, and increased eddy current losses result when laminated cores made of such sheets are annealed by layer welding.

Example b shows a flat coating, without interruptions. This corresponds to coatings such as HITCOAT or DL1. While very good layer separation can be achieved with such a coating during final annealing in the γ region, it is not possible to form a significant proportion of cube surface texture.

Example c shows a stripy structure. The dark stripes correspond to areas with a very dense coating of Al-containing particles. The light areas between the stripes contain no or very few particles, so these layers typically appear transparent. Of course, there may also be binders in the intermediate areas in the unannealed state.

Example d also shows a stripy structure. In contrast to example c, the dark stripes enriched with Al particles are narrower than in example c. This allows a special design of the cube surface structure, but increases the risk of sheet metal adhesion due to annealing

Example e shows a grid structure in which the lines run diagonally to the direction of the strip.

Example f shows a coating in which local accumulations of particles have formed surrounded by free areas. The particles are bound before annealing so that the coating adheres.

Example g shows schematically the appearance of an actually applied coating. When coating, the dispersion used has a lower viscosity than in examples c and d, so that the coating had a longer opportunity to run laterally after application. The result is a streaky course with branches. In this representation you can also see that the free areas, which appear completely white in the idealized representations, in practice always contain a proportion of fine Al particles. However, the concentration in these areas is very low compared to the thick stripes.

Some of these abstract surface structures were made concrete in the form of the exemplary embodiments listed in Table 10. Magnetic characteristics of coated samples after a final annealing of 4 hours at 1000°C, H ₂ , heating rate of 900°C..1000°C with 20 K/h, cooling rate of 1000°C..900°C with 20 K/h are shown. Table 10 # Picture Tmax in °C Identifier B3 in M B20 in T B100 in T µ _max Hc in A/m Br in T 1 a 1000 1901543 1,537 1,770 2,020 12,814 37.3 1.43 2 b 1000 2001969 1,437 1,687 1,963 10,240 46.4 1.36 3 b 1000 2000230 1,406 1,663 1,940 11,417 48.1 1.38 4 b 1000 2000231 1,437 1,696 1,956 12,148 46.7 1.38 5 f 1000 2000200 1,507 1,734 1,998 14,523 39.9 1.42 6 9 1000 2000252 1,502 1,732 1,995 12,959 42.2 1.41 7 d 1000 2001968 1,493 1,727 1,995 12,694 40.8 1.40 8th d 1100 2002008 1,537 1,755 2,016 14046.6 35.8 1.39

VACOFLUX X1, batch 7610163B, strip thickness 0.20 mm was used as the starting material in all cases. Punching rings measuring 28.5 mm x 20.0 mm were produced for each coating condition.

For this purpose, 20 rings were placed on top of each other in a stack, placed on a ceramic plate, weighted with a ceramic plate and annealed in a chamber furnace with a flush of dry hydrogen. All annealing took place at a temperature T _max of at least 1000°C in the γ region, so that a throw surface texture can form when cooling through the two-phase region α+γ in the presence of an oxide-free surface. The indicator for the presence of a significant proportion of cube surface texture is an induction value B20 = B(20 A/cm) of at least 1.70 T, preferably at least 1.74 T.

The measurement of the magnetic induction B3 = B(3 A/cm), B20, B100, the maximum permeability µ _max , the coercivity Hc and the remanence Br was carried out in accordance with the measurement standard IEC 60404-4.

Sample #1 represents the uncoated reference sample. Due to the completely free surface, similar to figure a, the highest proportion of cube surface texture (001) [uvw] can be formed here through annealing. The induction B20 is 1.770 T. However, due to the lack of coating, it is not possible to anneal these sheets in contact with each other.

Sample #2 represents the reference sample with a flat, continuous coating, i.e. as in figure b. The strip was previously coated on both sides with HITCOAT, a zirconium propylate-based coating which, after final annealing, is present as ceramic zirconium oxide. Since the entire surface is covered here, there is no preferred formation of a cube surface texture, which is reflected in a very low induction B20 of 1.687 T.

Sample #3 and Sample #4 were coated with the new coating TX1, but are not according to the invention due to the dense surface.

The batch used for sample #3 had a relatively high ceramic content of 11%. The coating was applied in a laboratory test by manually pressing on a profiled roller. This composition and method of application resulted in thick stripes, similar to surface image c, but the high concentration in combination with the low contact pressure led to a dense base coating even between the stripes, so that this condition corresponds overall to surface image b. Here too, the flat, continuous coating as in sample #2 results in a very low induction after final annealing, which is a B20 of 1.663 T.

For sample #4, a lower ceramic content of 7% was chosen. At this concentration, the coating runs after exiting the roll gap, resulting in a continuous coating according to surface image b. The induction B20 after the final annealing was therefore only 1.696 T.

Samples #5, #6, #7 and #8 represent conditions according to the invention. In all of these conditions, the sheets could be separated without damage after annealing.

Sample #5 was coated with the same approach as sample #3, i.e. with a ceramic content of 11%. The strip ran through two profiled rollers with a contact pressure of 2 bar. This resulted in the streaky surface appearance d. Due to the partially exposed surface, a significant amount of cube surface texture was able to form during the final annealing; the induction B20 was 1.734 T.

For sample #6, a lower ceramic content of 7% was chosen, similar to sample #4. In order to obtain a non-flat coating despite the lower ceramic content, an additional additive was added to the coating, which increases the basic viscosity. This change made it possible to achieve a non-surface coating as shown in Figure g. Here too there was a high induction B20 of 1.732 T.

In sample #7 the ceramic content was reduced even further, i.e. to 4% and, as in sample #6, an additive was added. The surface image corresponded to image d, i.e. fine, well-separated lines. The induction B20 of 1.727 T is due to the formation of the cube surface texture.

Finally, sample #8 illustrates that the very good layer insulation of the coating also makes it possible to further increase the annealing temperature. Increasing the annealing temperature can be advantageous when annealing in the stack. In this specific case, a strip with the same coating as sample #7 (4% ceramic content + additive) was annealed at an increased holding temperature of 1100 ° C in the γ region, with the same holding time of 4 hours. This results in an increased induction B20 of 1.755 T compared to the other examples according to the invention.

The exemplary embodiments illustrate that it is not the chemistry of the coating that is decisive for the formation of the cube surface texture, but rather that it is an aid to setting the correct surface coverage. It is therefore important that the coating after annealing is present as a non-continuous layer with free surface areas in between.

This is shown in the example according to the invention Figure 16 illustrated. Figure 16 shows an example of a streaky surface after coating with TX1, 11% ceramic content (sample #5 according to the invention). Before annealing (top row), the analysis of the coating strips reveals proportions of the ceramic (Al and O) and the binder (C). After annealing (bottom row), the EDX analysis also shows a surface free of deposits between the strips, i.e. only elements of the base material VACOFLUX X1.

Figure 17 shows a photograph of a coating strip made of TX1 (11% ceramic content) after final annealing (sample #5 according to the invention). The coated areas are narrower than the resulting grains, so that sufficient free surface remains within each grain and enables the formation of the cube surface texture.

Claims (16)

1. Method for producing a soft-magnetic alloy, having:

   the provision of a precursor having a composition substantially consisting of 2 % w/w ≤ Co ≤ 30 % w/w 0.3 % w/w ≤ V ≤ 5.0 % w/w 0 % w/w ≤ Cr ≤ 3.0 % w/w 0 % w/w ≤ Si ≤ 5.0 % w/w 0 % w/w ≤ Mn ≤ 5.0 % w/w 0 % w/w ≤ Al ≤ 3.0 % w/w 0 % w/w ≤ Ta ≤ 0.5 % w/w 0 % w/w ≤ Ni ≤ 1.0 % w/w 0 % w/w ≤ Mo ≤ 0.5 % w/w 0% w/w ≤ Cu ≤ 0.2 % w/w 0 % w/w ≤ Nb ≤ 0.25 % w/w 0 % w/w ≤ Ti ≤ 0.05 % w/w 0 % w/w ≤ Ce ≤ 0.05 % w/w 0 % w/w ≤ Ca ≤ 0.05 % w/w 0 % w/w ≤ Mg ≤ 0.05 % w/w 0 % w/w ≤ C ≤ 0.02 % w/w 0 % w/w ≤ Zr ≤ 0.1 % w/w 0 % w/w ≤ O ≤ 0.025 % w/w 0 % w/w ≤ S ≤ 0.015 % w/w

   rest iron and up to 0.2 % w/w melting-related impurities,

   wherein the precursor has a phase transition from a BCC phase area into a BCC/FCC mixed area to an FCC phase area, wherein at increasing temperature the phase transition between the BCC phase area and the BCC/FCC mixed area occurs at a first transition temperature T_α/α+γ and at further increasing temperature the transition between the BCC/FCC mixed area and the FCC phase area occurs at a second transition temperature T_α+γ/γ and the difference T_α+γ/γ - T_α/α+γ is less than 45 K, preferably less than 25 K,

   the partial coating of the precursor with a ceramic-forming layer, wherein the precursor has the form of a sheet with a first surface and a second, opposite surface, wherein at least between 20% and 80%, preferably between 30% and 70%, particularly preferred between 50% and 70% of the first surface and between 20% and 80%, preferably between 30% and 70%, particularly preferred between 50% and 70% of the second surface remain free of the ceramic-forming layer,

   the heat treatment of the partially coated precursor, wherein the heat treatment comprises

   the heating of the precursor, and then

   the heat treatment of the precursor in a first stage with a total time t₁, wherein in the first stage the precursor is heat-treated at a temperature in a temperature range between T_α+γ/γ and T₁, and then

   the cooling of the precursor to room temperature,
   or

   the heating of the precursor, and then

   the heat treatment of the precursor in a first stage with a total time t₁, wherein in the first stage the precursor is heat-treated at a temperature in a temperature range between T_α+γ/γ and T₁, and then

   the cooling of the precursor to a temperature T₂, and then

   the heat treatment of the precursor in a second stage at the temperature T₂ for a time t₂, and then

   the cooling of the precursor to room temperature,

   wherein the heat treatment is at least intermittently performed in a hydrogenous atmosphere while the exposed parts of the surface of the precursor are in direct contact with the hydrogenous atmosphere, wherein the temperature T₁ > T₂, T₁ is above T_α+γ/γ and T₂ is below T_α/α+γ.
2. Method for producing a soft-magnetic alloy according to claim 1, having:
   the provision of a precursor having a composition substantially consisting of 5 % w/w ≤ Co ≤ 25 % w/w 0.3 % w/w ≤ V ≤ 5.0 % w/w 0 % w/w ≤ Cr ≤ 3.0 % w/w 0 % w/w ≤ Si ≤ 3.0 % w/w 0 % w/w ≤ Mn ≤ 3.0 % w/w 0 % w/w ≤ Al ≤ 3.0 % w/w 0 % w/w ≤ Ta ≤ 0.5 % w/w 0 % w/w ≤ Ni ≤ 0.5 % w/w 0 % w/w ≤ Mo ≤ 0.5 % w/w 0% w/w ≤ Cu ≤ 0.2 % w/w 0 % w/w ≤ Nb ≤ 0.25 % w/w 0 % w/w ≤ Ti ≤ 0.05 % w/w 0 % w/w ≤ Ce ≤ 0.05 % w/w 0 % w/w ≤ Ca ≤ 0.05 % w/w 0 % w/w ≤ Mg ≤ 0.05 % w/w 0 % w/w ≤ C ≤ 0.02 % w/w 0 % w/w ≤ Zr ≤ 0.1 % w/w 0 % w/w ≤ O ≤ 0.025 % w/w 0 % w/w ≤ S ≤ 0.015 % w/w rest iron, wherein Cr+Si+AI+Mn ≤ 3.0 % w/w, and up to 0.2 % w/w melting-related impurities.
3. Method according to claim 1 or claim 2, wherein the maximum width of the coated area is less than 2 mm, preferably less than 1.2 mm, particularly preferred less than 0.8 mm.
4. Method according to any of claims 1 to 3, wherein the ceramic-forming layer has a metal oxyhydrate and/or metal oxide and/or metal hydroxide.
5. Method according to any of claims 1 to 4, wherein, during the heating of the precursor, at least in the temperature range of T_α/α+γ to T₁, the heat treatment is carried out in a protective gas atmosphere with less than 5 % by volume of hydrogen, preferably less than 1 % by volume of hydrogen, and the cooling of T₁, at least in the temperature range of T_α+γ/γ to T_α/α+γ, is carried out in a hydrogenous atmosphere with more than 5 % by volume of hydrogen.
6. Method according to claim 5, wherein, following the cooling of the precursor to a temperature T₂, with T₂ being below T_α/α+γ, the precursor is held at the temperature T₂ for a time t₂ and only then cooled further.
7. Method according to any of claims 1 to 6, wherein the heat treatment of the precursor in the first stage is carried out for the total time t₁ in a protective gas atmosphere with less than 5 % by volume of hydrogen, preferably less than 1 % by volume of hydrogen.
8. Method according to any of claims 1 to 7, wherein the cooling of the precursor from T₁ to T₂ is carried out in a hydrogenous atmosphere.
9. Method according to any of claims 1 to 8, wherein T_{α+ γ/γ} ≤ T₁ ≤ T_α+γ/γ + 50°C and 5 minutes ≤ t₁ < 10 hours, preferably 15 minutes ≤ t₁ < 4 hours, and 700 °C ≤ T₂ < 1050 °C and 30 minutes ≤ t₂ < 20 hours.
10. Method according to any of claims 1 to 9, wherein the heat treatment furthermore comprises a subsequent final annealing in a hydrogenous protective gas atmosphere, which is carried out at a maximum temperature below the first transition temperature T_α/α+γ,
11. Method according to any of claims 1 to 10, wherein the heating rate, at least over the temperature range of 900°C to T₁, is 10 K/h to 1000 K/h, preferably 20 K/h to 100 K/h, and/or the cooling rate, at least over the temperature range of T₁ to 900°C, is 10 K/h to 200 K/h and preferably 20 K/h to 100 K/h.
12. Method according to any of claims 1 to 11, wherein the sheets, following the heat treatment,

    are bonded by means of an insulating adhesive to form a laminated core,

    or oxidised on the surface to provide a layer with insulating action, and then bonded to form the laminated core, or laser-welded,

    or coated with an inorganic/organic hybrid coating and then processed to form the laminated core.
13. Soft-magnetic alloy substantially consisting of 5 % w/w ≤ Co ≤ 25 % w/w 0.3 % w/w ≤ V ≤ 5.0 % w/w 0 % w/w ≤ Cr ≤ 3.0 % w/w 0 % w/w ≤ Si ≤ 3.0 % w/w 0 % w/w ≤ Mn ≤ 3.0 % w/w 0 % w/w ≤ Al ≤ 3.0 % w/w 0 % w/w ≤ Ta ≤ 0.5 % w/w 0 % w/w ≤Ni ≤ 0.5 % w/w 0 % w/w ≤ Mo ≤ 0.5 % w/w 0 % w/w ≤ Cu ≤ 0.2 % w/w 0 % w/w ≤ Nb ≤ 0.25 % w/w 0 % w/w ≤ Ti ≤ 0.05 % w/w 0 % w/w ≤ Ce ≤ 0.05 % w/w 0 % w/w ≤ Ca ≤ 0.05 % w/w 0 % w/w ≤ Mg ≤ 0.05 % w/w 0 % w/w ≤ C ≤ 0.02 % w/w 0 % w/w ≤ Zr ≤ 0.1 % w/w 0 % w/w ≤ O ≤ 0.025 % w/w 0 % w/w ≤ S ≤ 0.015 % w/w

    rest iron, wherein Cr+Si+AI+Mn ≤ 3.0 % w/w, and up to 0.2 % w/w melting-related impurities,

    characterised in that the soft-magnetic alloy has a surface proportion of a {111}<uvw>texture of maximally 13%, preferably maximally 6%, wherein grains with a tilt of up to +/- 10° or preferably up to +/- 15° against the nominal crystal orientation are enclosed, and a surface proportion of a { 100 } <uvw>cube face texture of at least 30%, preferably at least 50%, wherein grains with a tilt of up to +/- 15° or preferably up to +/- 10° against the nominal crystal orientation are enclosed.
14. Soft-magnetic alloy according to claim 13,

    in which the average grain size of the {100}<uvw>-oriented grains is at least 1.5 times, preferably twice, the average grain size of the {111}<uvw>-oriented grains, and/or

    in which the surface proportion of the {100}<uvw>-oriented grains is at least 3 times, preferably at least 7 times, the surface proportion of the {111}<uvw>-oriented grains.
15. Soft-magnetic alloy according to claim 13 or claim 14, wherein

    10 % w/w ≤ Co ≤ 20 % w/w, preferably 15 % w/w ≤ Co ≤ 20 % w/w, applies, and

    0.5 % w/w ≤ V ≤ 4.0 % w/w, preferably 1.0 % w/w < V ≤ 3.0 % w/w, preferably 1.3 % w/w ≤ V ≤ 2.7 % w/w, applies, and/or

    0.1 % w/w ≤ Cr ≤ 2.0 % w/w, preferably 0.2 % w/w ≤ Cr ≤ 1.0 % w/w, preferably 0.3 % w/w ≤ Cr ≤ 0.7 % w/w, applies, and/or

    0.1 % w/w ≤ Si ≤ 2.0 % w/w, preferably 0.15 % w/w ≤ Si ≤ 1.0 % w/w, preferably 0.2 % w/w ≤ Si ≤ 0.5 % w/w, applies, and/or

    the molecular formula 0.1 % w/w ≤ Cr+Si+AI+Mn ≤ 1.5 % w/w, preferably 0.2 % w/w ≤ Cr+Si+AI+Mn ≤ 0.6 % w/w, applies.
16. Laminated core of a plurality of stacked electrically insulated sheets of a soft-magnetic alloy according to any of claims 13 to 15.

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