Patent Application
20200325564
The present invention relates to a soft magnetic alloy, in particular a high permeability soft magnetic alloy.
Non-grain-oriented electrical steel with approx. 3 wt % silicon (SiFe) is the most common crystalline soft magnetic material used in laminated cores in electric machines. As the electric-powered vehicle sector progresses, more efficient materials that performance better than SiFe are needed. In addition to sufficiently high electrical resistance, this signifies that a higher level of induction in particular is desirable to provide high torques and/or compact components.
Even more efficient materials are desirable for use in sectors such as the automotive industry and electric-powered vehicles. Soft magnetic cobalt-iron (CoFe) alloys are also used in electric 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. In compositions of this type both a saturation induction of approx. 2.35 T and a high electrical resistance of 0.4 μΩm are achieved. It is, however, also desirable to reduce the material and production costs of CoFe alloys resulting, for example, from the high Co content, additional manufacturing steps and scrap content.
The object of the present invention is therefore to provide an FeCo alloy that has lower material costs and is easy to work in order to reduce the production costs of the alloy, up to and including laminated cores, and at the same time to achieve high power density.
According to the invention, a soft magnetic alloy, in particular a high permeability soft magnetic FeCo alloy, is provided that consists essentially of:
residual iron, wherein Cr+Si+Al+Mn≤3.0 wt %, and up to 0.2 wt % of other impurities. The alloy has a maximum permeability μmax≥5,000, preferably μmax≥10,000, preferably μmax≥12.000, preferably μmax≥17,000. Other impurities include, for example, B, P, N, W, Hf, Y, Re, Sc, Be and other lanthanides other than Ce. (wt % denotes weight percent)
Owing to the lower Co content, the raw material costs of the alloy according to the invention are less than those of an alloy based on 49 wt % Fe, 49 wt % Co, 2% V. The invention provides for an FeCo alloy with a maximum cobalt content of 25 percent by weight that offers better soft magnetic properties, in particular appreciably higher permeability, than other FeCo alloys with a maximum cobalt content of 25 percent by weight such as the existing commercially available FeCo alloys e.g. VACOFLUX 17, AFK 18 and HIPERCO 15. These existing commercially available alloys have a maximum permeability of less than 5000.
The alloy according to the invention has no significant adjustment in order and can therefore, unlike alloys with over 30 wt % Co, be cold rolled without first undergoing a quenching process. Quenching is a difficult process to control, particularly where large quantities of materials are concerned, as it is hard to achieve sufficiently fast cooling rates and ordering with the resulting embrittlement of the alloy may therefore take place. The lack of an order-disorder transition in the alloy according to the invention simplifies industrial-scale production.
Marked order-disorder transitions in alloys like that observed in CoFe alloys with a Co content greater than 30 wt % can be determined by means of differential scanning calorimetry (DSC) because they cause a peak in the DSC measurement. No such peak is observed in a DSC measurement carried out under the same conditions for the alloy according to the invention.
At the same time, however, in addition to an appreciably higher permeability level never previously attained for this type of alloy, this new alloy offers both significantly lower hysteresis losses than previously known commercially available alloys with Co contents of between 10 and 30 wt % and higher saturation. The FeCo alloy according to the invention can also produced cost-effectively on an industrial scale.
Owing to its higher permeability, the alloy according to the invention can be used in applications such as rotors and stators in electric motors in order to reduce the size of the rotor or stator and thus of the electric motor, and/or to increase output. For example, it makes it possible to generate a higher torque at the same physical size and/or weight, a solution that would prove advantageous if used in electrically-powered or hybrid motor vehicles.
In addition to a maximum permeability μmax≥5,000, preferably μmax≥10,000, preferably μmax≥12.000, preferably μmax≥17,000, the alloy can also have an electrical resistance ρ≥0.25 μΩm, preferably ρ≥0.30 μΩm, and/or hysteresis losses PHys≤0.07 J/kg, preferably hysteresis losses PHys≤0.06 J/kg, preferably hysteresis losses PHys≤0.05 J/kg, at an amplitude of 1.5 T, and/or coercive field strength Hc of ≤0.7 A/cm, preferably a coercive field strength Hc of ≤0.6 A/cm, preferably a coercive field strength Hc of ≤0.5 A/cm, preferably Hc of ≤0.4 A/cm, preferably Hc of ≤0.3 A/cm, and/or an induction B≥1.90 T at 100 A/cm, preferably B≥1.95 T at 100 A/cm, preferably B≥2.00 T at 100 A/cm.
The hysteresis losses PHys are determined from the re-magnetisation losses P at an amplitude of induction of 1.5 T across the y-axis intercept in a plot P/f over the frequency f by linear regression. The linear regression is carried out using at least 8 measured values distributed approximately evenly over a frequency range of 50 Hz to 1 kHz (e.g. at 50, 100, 200, 300, 400, 500, 600, 700, 800, 900 and 1000 Hz).
In one embodiment, the alloy has a maximum permeability μmax≥μmax≥10,000, an electrical resistance ρ≥0.28 μΩm, hysteresis losses PHys≤0.055 J/kg at an amplitude of 1.5 T, a coercive field strength Hc of ≤0.5 A/cm and an induction B≥1.95 T at 100 A/cm. This combination of properties is particularly advantageous for use as or in a rotor or stator of an electric motor in order to reduce the size of the rotor or stator and thus of the electric motor, and/or to increase output, or to generate higher torque at the same weight.
The soft magnetic alloy can therefore be used in an electric machine, e.g. as or in a stator and/or rotor of an electric motor and/or generator, and/or in an transformer and/or in an electromagnetic actuator. It can be provided in the form of a sheet with a thickness of 0.5 mm to 0.05 mm, for example. A plurality of sheets made of the alloy can be stacked together to form a laminated core to be used as a stator or rotor.
The alloy according to the invention has an electrical resistance of at least 0.25 μΩm, preferably a minimum of 0.3 μΩm. Eddy current losses can be reduced to a lower level by selecting a slightly smaller strip thickness.
The composition of the soft magnetic alloy is set out in greater detail in further embodiments, with 10 wt %≤Co≤20 wt %, preferably 15 wt %≤Co≤20 wt % and 0.3 wt %≤V≤5.0 wt %, preferably 1.0 wt %≤V≤3.0 wt %, preferably 1.3 wt %≤V≤2.7 wt % and/or 0.1 wt %≤Cr+Si≤2.0 wt %, preferably 0.2 wt %≤Cr+Si≤1.0 wt %, preferably 0.25 wt %≤Cr+Si≤0.7 wt %.
In one embodiment, the sum is defined in greater detail, with 0.2 wt %≤Cr+Si+Al+Mn≤1.5 wt %, preferably 0.3 wt %≤Cr+Si+Al+Mn≤0.6 wt %.
The soft magnetic alloy may also contain silicon, with 0.1 wt %≤Si≤2.0 wt %, preferably 0.15 wt %≤Si≤1.0 wt %, preferably 0.2 wt %≤Si≤0.5 wt %.
Aluminium and silicon can be interchanged such that in one embodiment the total Si and Al (Si+Al) is 0 wt %≤(Si+Al)≤3.0 wt %.
The alloys according to the invention are almost carbon-free and contain at most 0.02 wt % carbon, preferably ≤0.01 wt % carbon. This maximum carbon content should be regarded as an unavoidable impurity.
In the alloys according to the invention calcium, beryllium and/or magnesium may be added in small amounts of up to 0.05 wt % only for deoxidation and desulphurisation. In order to achieve particularly good deoxidation, it is possible to add up to 0.05 wt % cerium or cerium Mischmetal.
According to the invention, the improved magnetic properties can be achieved by heat treatment geared to the composition as described below. It has been shown, in particular, that ascertaining the phase transition temperatures for the selected compositions and determining the heat treatment temperatures and cooling rate in relation to the phase transition temperatures thus ascertained leads to improved magnetic properties. The fact that the alloys according to the invention with a cobalt content of at most 25 percent by weight have no order-disorder transition so that the manufacturing process does not require quenching to avoid ordering and the resulting embrittlement, is also taken into account.
Conventionally, CoFe alloys are used in strip thicknesses ranging from 0.50 mm to a thin 0.050 mm. In processing the strip, the material is conventionally hot rolled and then cold rolled to its final thickness. During cooling after hot rolling an embrittling adjustment in order takes place at approx. 730° C. and to ensure sufficient cold rollability special intermediate annealing followed by quenching is therefore also required to suppress the adjustment in order. The alloy according to the invention does requires no quenching since it has no order-disorder transition. This simplifies production.
To achieve the magnetic properties, CoFe alloys are subjected to a final heat treatment also referred to as final magnetic annealing. The stock is heated to the annealing temperature, held at the annealing temperature for a certain length of time and then cooled at a defined speed. It is advantageous to carry out this final annealing at the highest possible temperatures and in a clean, dry hydrogen atmosphere since at high temperatures, firstly, the reduction of impurities by means of hydrogen is more efficient and, secondly, the grain structure becomes rougher and so soft magnetic properties such as coercive field strength and permeability improve. In practice, the annealing temperature in the CoFe system has an upper limit since a phase transition from the magnetic and ferritic BCC phase to the non-magnetic and austenitic FCC phase takes place at approx. 950° C. in the binary system. When elements are added to the alloy, a two-phase region in which both phases coexist occurs between the FCC phase and the BCC phase. The transition between the BCC phase and the mixed two-phase or BCC/FCC region occurs at a temperature TÜ1 and the transition between the two-phase region and the FCC phase occurs at a temperature TÜ2, where TÜ2>TÜ1. The position and size of the two-phase region also depends on the nature and scope of the alloy making process. If annealing takes place in the two-phase region or in the FCC region, remnants of the FCC phase may impair the magnetic properties after cooling and incomplete retransformation. Even if retransformation is complete, the additional grain boundaries created have an damaging effect since coercive field strength behaves inversely proportionately to grain diameter. Consequently, the known commercial available alloys with Co contents of approx. 20 wt % undergo final annealing at temperatures below the two-phase BCC+FCC region. As a result, the recommendation for AFK 18 is 3 h/850° C. and that for AFK 1 is 3 h/900° C., for example. The recommendation for VACOFLUX 17 is 10 h/850° C. At such low final annealing temperatures and owing to the relatively high magneto-crystalline anisotropy (K1 approx. 45,000 J/m3 at 17 wt % Co), the potential for particularly good soft magnetic properties in these FeCo alloys is limited.
With VACOFLUX 17 strip, for example, the maximum permeability that can be reached at a typical coercive field strength of 1 A/cm is approx. 4,000 and its application is therefore limited.
In contrast to these known final annealing processes, the composition according to the invention permits a heat treatment that produces better magnetic properties than the standard single-step annealing with furnace cooling used with FeCo alloys, irrespective of the temperature range in which the single-step annealing takes place. The additives are selected such that the lower limit of the two-phase region and the BCC/FCC phase transition are pushed upwards to allow annealing at high temperatures, e.g. above 925° C. in the BCC-only region. Annealing heat treatments at such high temperatures are not conceivable with the FeCo alloys known to date. Moreover, the width of the two-phase region, i.e. the difference between the lower transition temperature TÜ1 and the upper transition temperature TÜ2 is kept as narrow as possible owing to the composition according to the invention. As a result, the advantages of high final annealing, i.e. the removal of potential magnetically unfavourable textures, the cleaning effect in H2 and the growth of large grains, are maintained by final annealing above the two-phase region in conjunction with cooling through the two-phase region followed by a holding period or controlled cooling in the BCC-only region without the risk of magnetically damaging remnants of the FCC phase.
It has been found that compositions with a phase transition between the BCC-only region and the mixed BCC/FCC region exhibit appreciably improved magnetic properties at higher temperatures, e.g. above 925° C., and with a narrow two-phase region, e.g. of less than 45K. Compositions with this specific combination of phase diagram features are selected according to the invention and heat treated accordingly in order to guarantee a high permeability of greater than 5000 or greater than 10,000.
Vanadium was identified as one of the most effective elements in an Fe—Co alloy, increasing electrical resistance and at the same time pushing the two-phase region up to higher temperatures. With a lower Co content, vanadium is more efficient at increasing transition temperatures. With the Fe-17Co alloy, it is even possible to increase the transition temperatures above the value of the binary FeCo composition by adding approx. 2% vanadium.
In the Fe—Co system, from approx. 15% cobalt the BCC/FCC phase transformation takes place at temperatures lower than the Curie temperature. Since the FCC phase is paramagnetic, the magnetic phase transition is now determined by the BCC/FCC phase transformation rather than the Curie temperature. Sufficiently large amounts of vanadium push the BCC/FCC phase transformation over the Curie temperature Tc, making the paramagnetic BCC Ophase visible.
However, if the vanadium content is too high, the width of the mixed region is increased. These compositions have lower maximum permeability values even though the phase transition between the mixed BCC/FCC region and the BCC-only region takes place at higher temperatures. Consequently, it has been established that that the composition has an influence both on the temperatures at which the phase transitions take place and on the width of the mixed region, and should therefore be taken into account when selecting the composition. In order to achieve the highest permeability values, the heat treatment temperatures can be selected in relation to the temperatures at which the phase transitions for this composition take place.
It has thus been found that a more precise determination of the temperatures at which the phase transitions take place is advantageous for a certain composition wen optimising the production process. These temperatures can be determined by means of differential scanning calorimetry (DSC) measurements. The DSC measurement can be carried out with a sample mass of 50 mg and at a DSC heating rate of 10 Kelvin per minute, and the phase transition temperatures thus determined can be used when heating and cooling the sample to determine the temperatures for heat treatment.
Chromium and other elements can be added in order, for example, to improve electrical resistance or mechanical properties. Like most other elements, chromium reduces the two-phase region of the binary Fe-17Co alloy. As a result, the amount of element to be added in addition to vanadium is preferably selected such together with vanadium it produces an increase in the two-phase region as compared to the binary FeCo alloy. In addition, the impurities and elements that have a particularly strong stabilising affect on the austenite (e.g. nickel) must be kept as low as possible.
The following contents have proved preferable in achieving very good magnetic properties:
cobalt content of 5 wt %≤Co≤25 wt %, with contents of 10 wt %≤Co≤20 wt % being preferred and contents of 15 wt %≤Co≤20 wt % being very particularly preferred;
vanadium content of 0.3 wt %≤V≤5.0 wt %, with contents of 1.0 wt %≤V≤3.0 wt % being preferred, and the following sum: 0.2 wt %≤Cr+Si+Al+Mn≤3.0 wt %.
The alloys according to the invention are almost carbon-free and have at most 0.02 wt % carbon, preferably ≤0.01 wt % carbon. This maximum carbon content should be regarded as an unavoidable impurity.
Only small amounts of calcium, beryllium and/or magnesium up to 0.05 wt % can be added to the alloys according to the invention for deoxidation and desulphurisation.
To achieve particularly good deoxidation and desulphurisation up to 0.05 wt % Cer or misch metal can be added.
The composition according to the invention allows a further improvement. Cobalt has a higher diffusion coefficient in the paramagnetic BCC phase than in the ferromagnetic BCC phase. As a result, by separating the two-phase region and the Curie temperature Tc, vanadium allows a further temperature range with high self diffusion, thereby allowing a larger BCC grain structure and thus better soft magnetic properties due to heat treatment in this range or cooling through this range. In addition, the separation of two-phase region and Curie temperature Tc signifies that during cooling both the passage through the two-phase BCC/FCC region and the transition to the region of the BCC-only phase take place completely in the paramagnetic state. This also has a positive effect on the soft magnetic properties.
According to the invention, a method is provided for the production of a soft magnetic FeCo alloy, this method comprising the following. A preliminary product (precursor) is provided with a composition consisting substantially of:
residual iron, where Cr+Si+Al+Mn≤3.0 wt %, and up to 0.2 wt % of other impurities due to melting. The other impurities may, for example, be one or more of the elements B, P, N, W, Hf, Y, Re, Sc, Be or other lanthanides other than Ce. In some embodiments the preliminary product has a cold-rolled texture or a fibre texture.
The preliminary product or the parts manufactured from the preliminary product are heat treated. In one embodiment, the preliminary product is heat treated at a temperature T1 and then cooled down from T1 to room temperature.
In an alternative embodiment, the preliminary product is heat treated at a temperature T1, then cooled down to a temperature T2 that is above room temperature, and further heat treated at temperature T2, where T1>T2. The preliminary product is not cooled to room temperature until it has been heat treated at temperature T2.
The preliminary product has a phase transition from a BCC phase region to a mixed BCC/FCC region to a FCC phase region, as the temperature increases the phase transition between the BCC phase region and the mixed BCC/FCC region taking place at a first transition temperature TÜ1 and, as the temperature continues to increase, the transition between the mixed BCC/FCC region and the FCC phase region taking place at a second transition temperature TÜ2, where TÜ2>TÜ1. Temperature T1 is above TÜ2 and temperature T2 is below TÜ1.
The transition temperatures TÜ1 and TÜ2 are dependent on the composition of the preliminary product. The transition temperatures TÜ1 and TÜ2 can be determined by means of DSC measurements, the transition temperature TÜ1 being determined during heating and the transition temperature TÜ2 being determined during cooling. In one embodiment, at a sample mass of 50 mg and a DSC heating rate of 10 Kelvin per minute the transition temperature TÜ1 is above 900° C., preferably above 920° C., and preferably above 940° C.
In one embodiment, the solidus temperature of the preliminary product is taken into account when selecting temperatures T1 and T2. In one embodiment, 900° C.≤T1<Tm, preferably 930° C.≤T1<Tm, preferably 940° C.≤T1<Tm, preferably 960° C.≤T1<Tm, and 700° C.≤T2≤1050° C. and T2<T1, Tm being the solidus temperature.
In one embodiment, the difference TÜ2−TÜ1 is less than 45K, preferably less than 25K.
In one embodiment, the cooling rate over at least the temperature range from T1 to T2 is 10° C./h to 50,000° C./h, preferably 10° C./h to 900° C./h, preferably 20° C./h to 1000° C./h, preferably 20° C./h to 900° C./h, preferably 25° C./h to 500° C./h. This cooling rate can be used with both of the heat treatments described above.
In one embodiment, the difference TÜ2−TÜ1 is less than 45K, preferably less than 25K, T1 is above TÜ2 and T2 is below T, 940° C.≤T1<Tm, where 700° C.≤T2≤1050° C. and T2<T1, Tm being the solidus temperature, and the cooling rate is 10° C./h to 900° C./h at least over the temperature range T1 to T2. This combination of properties of the alloy, i.e. TÜ2 and TÜ1, can be used with the heat treatment temperatures T1 and T2 to achieve particularly high permeability rates.
In one embodiment, the preliminary product is heat treated at above TÜ2 for a period of over 30 minutes and then cooled to T2.
In one embodiment, the preliminary product is heat treated at T1 for a period t, where 15 minutes≤t1≤20 hours, and then cooled from T1 to T2. In one embodiment, the preliminary product is cooled from T1 to T2, heat treated at T2 for a period t2, where 30 minutes≤t2≤20 hours, and then cooled from T2 to room temperature.
In embodiments in which the preliminary product is cooled down from T1 to room temperature, the preliminary product may than be heated up from room temperature to T2 and heat treated at T2 according to one of the embodiments described here.
As the alloy has no order-disorder transition, no quenching is carried out over the temperature range from 800° C. to 600° C. The cooling rate from 800° C. to 600° C. may, for example, be between 100° C./h and 500° C./h. However, a slower cooling rate can, in principle, also be chosen. The aforementioned cooling rates can also quite easily be carried out until room temperature is reached.
The preliminary product can be cooled from T1 to room temperature at a rate of 10° C./h to 50,000° C./h, preferably from 10° C./h to 1000° C./h, preferably from 10° C./h to 900° C./h, preferably from 25° C./h to 900° C./h, preferably from 25° C./h to 500° C./h.
The cooling rate from T2 to room temperature has less influence on magnetic properties so the preliminary product can be cooled from T2 to room temperature at a rate of 10° C./h to 50,000° C./h, preferably 100° C./h to 1000° C./h.
In a further alternative embodiment, the preliminary product is cooled from T1 to room temperature at a cooling rate of 10° C./h to 900° C./h. In embodiments with slow cooling from T1 to room temperature, e.g. with a cooling rate of less than 500° C./h, preferably less than 200° C./h, a further heat treatment at temperature T2 can be dispensed with.
Following heat treatment, the soft magnetic alloy may have the following combinations of properties:
a maximum permeability μmax≥5,000, and/or an electrical resistance ρ≥0.25 μΩm, and/or hysteresis losses PHys≤0.07 J/kg at an amplitude of 1.5 T, a coercive field strength Hc of ≤0.7 A/cm and an induction B≥1.90 T at 100 A/cm, or
a maximum permeability μmax≥10,000, and/or an electrical resistance ρ≥0.25 μΩm, and/or hysteresis losses PHys≤0.06 J/kg at an amplitude of 1.5 T, and/or a coercive field strength Hc of ≤0.6 A/cm, preferably Hc≤0.5 A/cm and/or an induction B≥1.95 T at 100 A/cm, or
a maximum permeability μmax≥≥12.000, preferably μmax≥17,000 and/or an electrical resistance ρ≥0.30 μΩm, and/or hysteresis losses PHys≤0.05 J/kg at n amplitude of 1.5 T, and/or a coercive field strength Hc of ≤0.5 A/cm, preferably Hc≤0.4 A/cm, preferably Hc≤50.3 A/cm and/or an induction B≥2.00 T at 100 A/cm.
In certain embodiments the soft magnetic alloy has one of the following combinations of properties:
a maximum permeability μmax≥5,000, an electrical resistance ρ≥0.25 μΩm, hysteresis losses PHys≤0.07 J/kg at n amplitude of 1.5 T, a coercive field strength Hc of ≤0.7 A/cm and an induction B≥1.90 T at 100 A/cm, or
a maximum permeability μmax≥10,000, an electrical resistance ρ≥0.25 μΩm, hysteresis losses PHys≤0.06 J/kg at an amplitude of 1.5 T, a coercive field strength Hc of ≤0.6 A/cm and an induction B≥1.95 T at 100 A/cm, or
a maximum permeability μmax≥12.000, an electrical resistance ρ≥0.28 μΩm, hysteresis losses PHys≤0.05 J/kg at an amplitude of 1.5 T, a coercive field strength Hc of ≤0.5 A/cm and an induction B≥2.00 T at 100 A/cm,
a maximum permeability μmax≥17,000, an electrical resistance ρ≥0.30 μΩm, hysteresis losses PHys≤0.05 J/kg at an amplitude of 1.5 T, a coercive field strength Hc of ≤0.4 A/cm, preferably Hc of ≤0.3 A/cm and an induction B≥2.00 T at 100 A/cm.
In one embodiment, the maximum difference in coercive field strength Hc after heat treatment measured parallel to the direction of rolling, measured diagonally (45°) to the direction of rolling or measured perpendicular to the direction of rolling between two of these directions is at most 6%, preferably at most 3%. In other words, the maximum difference in coercive field strength Hc measured parallel to the direction of rolling and measured diagonally (45°) to the direction of rolling is at most 6%, preferably at most 3%, and/or the maximum difference in coercive field strength Hc measured parallel to the direction of rolling and measured perpendicular to the direction of rolling is at most 6%, preferably at most 3%, and/or the maximum difference in the coercive field strength Hc measured diagonally (45°) to the direction of rolling or measured perpendicular to the direction of rolling between these two directions is at most 6%, preferably at most 3%. In rotor and stator applications, this anisotropy, which is extremely low for soft magnetic FeCo alloys, leads to uniform properties along the periphery and there is therefore no need to rotate rotor and stator sheets by layer to provide sufficient isotropy of the magnetic properties in the laminated core.
The heat treatment may be carried out in a hydrogen-containing atmosphere or in an inert gas.
In one embodiment, heat treatment is carried out in a stationary furnace at T1 and in a stationary furnace or a continuous furnace at T2. In another embodiment, heat treatment is carried out in a continuous furnace at T1 and in a stationary furnace or a continuous furnace at T2.
Prior to heat treatment the preliminary product may have a cold-rolled texture or a fibre texture.
The preliminary product may be provided in the form of a strip. At least one strip may be manufactured from the strip by stamping, laser cutting or water jet cutting. In one embodiment, heat treatment is carried out on stamped, laser-cut, electrical discharge machined or water jet-cut laminations manufactured from the strip material.
In one embodiment, after heat treatment a plurality of sheets are stuck (adhered) together using insulating adhesive to form a laminated core, or surface oxidized to form an insulating layer and then stuck, or laser welded together to form a laminated core, or coated with an inorganic-organic hybrid coating and then processed further to form a laminated core.
In some embodiments, the preliminary product takes the form of a laminated core and the laminated core is heat treated according to one of the embodiments described here. The heat treatment can thus be carried out on stamp bundled cores (progressively stacked cores) or welded laminated cores manufactured from laminations.
The preliminary product can be produced as follows. A molten mass may, for example, be provided by vacuum induction melting, electroslag remelting or vacuum arc remelting, this molten mass consisting substantially of:
residual iron, where Cr+Si+Al+Mn≤3.0 wt %, and up to 0.2 wt % of other impurities.
Other impurities may be one or more of the other B, P, N, W, Hf, Y, Re, Sc, Be or other lanthanides other than Ce. The molten mass is solidified to form an ingot and the ingot is mechanically formed to form a preliminary product with final dimensions, this mechanical forming being carried out by means of hot rolling and/or forging and/or cold forming.
In one embodiment, the ingot is mechanically formed to form a slab with a thickness D1 by means of hot rolling at temperatures between 900° C. and 1300° C. and then mechanically formed to form a strip with a thickness D2 by means of cold rolling, where 1.0 mm≤D1≤5.0 mm and 0.05 mm≤D2≤1.0 mm, where D2<D1. The degree of cold working by cold rolling may be >40%, preferably >80%.
In one embodiment, the ingot is mechanically formed to form a billet by means of hot rolling at temperatures between 900° C. and 1300° C. and then mechanically formed to form a wire by means of cold drawing. The degree of cold working due to cold drawing may be >40%, preferably >80%.
Intermediate annealing may be carried out in a continuous furnace or a stationary furnace at an intermediate dimension in order to reduce work hardening and so to set the desired degree of cold working.
The Curie temperature of the alloy may be taken into account when selecting the temperatures T1 and/or T2. For example, TÜ1>Tc, where Tc is the Curie temperature and Tc≥900° C. In one embodiment, TÜ1>T2>Tc.
In compositions in which there is a separation of the two-phase region and the Curie temperature Tc, there is a further temperature range with higher self diffusion. This allows a larger BCC grain structure and thus better soft magnetic properties as a result of heat treatment in this region or cooling through this region. The separation of the two-phase region and the Curie temperature Tc also signifies that during cooling both the passage through the two-phase BCC/FCC region and the transition to the BCC-only phase region take place entirely in the paramagnetic state. The soft magnetic properties can be further improved by selecting temperature T2 so that TÜ1>T2>Tc.
In one embodiment, the average grain size after final annealing is at least 100 μm, preferably at least 200 μm, preferably at least 250 μm.
In one embodiment, the measured density of the annealed alloy is more than 0.10% lower than the density calculated using the rule of three from the average atomic weight of the metallic elements of the alloy, the average atomic weight of the metallic elements of the corresponding binary FeCo alloy and the measured density of this annealed binary FeCo-alloy.
Owing to the heat treatment, the sulphur content in the finished alloy may be lower than that in the molten mass. For example, the upper limit of the sulphur content in the molten mass may be 0.025 percent by weight, while in the finished soft magnetic alloy the upper limit is 0.015 percent by weight.
In one embodiment, the preliminary product is also coated with an oxide layer for electrical insulation. This embodiment may, for example, be used if the preliminary product is used in a laminated core. The laminations or the laminated core can be coated with an oxide layer. The preliminary product may be coated with a layer of magnesium methylate or preferably zirconium propylate that transforms into an insulating oxide layer during heat treatment. The preliminary product may be heated treated in an atmosphere containing oxygen or water vapour to form an electrically insulating layer.
In one embodiment, laminations stamped, laser-cut or electrical discharge machined from the preliminary product are also subjected to final annealing, after which the annealed single sheets are then stuck together by means of an insulating adhesive to form a laminated core, or the annealed single sheets are surface-oxidised to form an insulating layer and then stuck, welded or laser-welded together to form a laminated core, or the annealed single sheets are coated with an inorganic-organic hybrid coating such as Remisol C5, for example, and then further processed to form a laminated core.
The soft magnetic alloy according to any one of the preceding embodiments, which can be produced using any one of the methods described here, may be used in an electric machine, e.g. as 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.
Embodiments of the invention are described in greater detail below with reference to the drawings and the following examples.
According to the invention, a soft magnetic alloy is provided that consists essentially of:
and up to 0.2 wt % of other impurities due to melting. The impurities may, for example, be one or more of the elements B, P, N, W, Hf, Y, Re, Sc, Be or other lanthanides other than Ce.
In order to increase electrical resistance, it is also possible, in addition to the alloy element vanadium, to add one or more of the group of Cr, Si, Al and Mn in an amount that satisfies the following sum:
0.05 wt %≤Cr+Si+Al+Mn≤3.0 wt %.
The alloy according to the present invention is preferably melted in vacuum induction furnaces, though it can also be processed using vacuum arc remelting or electroslag remelting. The molten mass first solidifies into an ingot from which the oxide skin is removed and then forged or hot rolled at temperatures between 900° C. and 1300° C. Alternatively, the removal of the oxide skin can also take place on bars that have previously been forged or hot rolled. The desired dimensions can be achieved by hot working strips, billets or bars. Surface oxides can be removed from hot rolled stock by blasting, grinding or stripping. Alternatively, however, the desired final dimensions can also be achieved by cold working strips, bars or wires. In the case of cold rolled strips, a grinding process can be integrated to remove embedded oxides caused by the hot rolling process. If cold working leads to excessive solidification, one or more intermediate annealing processes may be carried out at temperatures between 400° C. and 1300° C. for recovery and re-crystallisation. The thickness or diameter for the intermediate annealing should be selected such that cold working of preferably >40%, particularly preferably >80%, is achieved by the final thickness.
The last processing step is heat treatment at temperatures between 700° C. and the solidus temperature Tm (typically at most 1200° C.), which is also referred to as final magnetic annealing. Final annealing is preferably carried out in a clean, dry hydrogen atmosphere. Annealing in an insert gas or vacuum is also possible.
In variant 1, which is illustrated by the continuous line in
According to the invention, annealing may therefore take place either in two steps or by controlled cooling from a temperature above the upper transition temperature. Controlled cooling signifies that there is a defined cooling rate for creating the optimum soft magnetic properties. In all cases, one of the annealing steps takes place in the FCC region. The annealing processes according to the invention may be carried out in either a continuous furnace or a stationary furnace.
During the annealing process according to the invention, the alloy is annealed at least once at a temperature above TÜ2 between 900° C. (if TÜ2>900° C., then above TÜ2) and Tm in the austenitic FCC region in order to produce a large grain, to exploit the cleaning effect of the hydrogen and to remove potential magnetically disadvantageous textures. This final annealing step above TÜ2 takes place either in a stationary annealing process or in a continuous furnace. Alternatively, this heat treatment step may also take place on the strip stock in a continuous furnace. The alloy is then cooled at a rate of 10 to 50,000° C. per hour, preferably at a rate of 20 to 1000° C. per hour, to room temperature or to a temperature between 700° C. and 1000° C. in the BCC region.
A second annealing step may comprise either heating up or maintaining the temperature at between 700° C. and 1000° C. (if TÜ1<1000° C., then below TÜ1) in the ferritic BCC region in order to remove any potential remnants of the FCC phase. Following completed final magnetic annealing, the alloy is then cooled from the annealing temperature at a rate of 10 to 50,000° C. per hour, preferably at a rate of 20 to 1000° C. per hour.
The alloys according to the invention exhibit a phase transition from a BCC phase region to a mixed BCC/FCC region and at a slightly higher temperature a further phase transition from the mixed BCC/FCC region to a FCC phase region, as the temperature increases the phase transition taking place at a first transition temperature TÜ1 between the BCC phase region and the mixed BCC/FCC region and, as the temperature continues to increase, the transition taking place at a second transition temperature TÜ2 between the mixed BCC/FCC region and the FCC phase region, as shown in
The temperature at which the phase transitions from a BCC phase region to a mixed BCC/FCC region and from the mixed BCC/FCC region to an FCC phase region occur can be determined by means of DSC measurements.
The figures that follow show the results of DSC measurements carried out using a dynamic heat-flow differential scanning calorimeter from the company Netzsch. Two identical corundum (Al2O3) crucibles are placed in a furnace, one containing a real measuring sample, the other containing a reference calibration sample. Both crucibles are subjected to the same temperature programme, which may consist of a combination of heating, cooling or isothermal sections. The thermal flow difference is determined quantitatively by measuring the temperature difference at a defined heat conduction path between sample and reference. The various maxima and minima (peaks) determined by DSC measurement can be allocated to certain types of phase transformations on the basis of their curve shapes. The result is typical curve shapes that are material-specific but also dependent on the measurement conditions, in particular on the sample mass and the heating and cooling rates. To guarantee the comparability of the measurements, identical instrument heating and cooling rates and identical sample masses were used. The heating and cooling rates used in these tests were 10 K/min; the sample mass was 50 mg.
The transition temperatures TÜ1 and TÜ2 are determined by means of DSC measurement by heating a sample of a defined mass at a defined heating rate. In this measurement the transition temperatures are represented by the first onset. This parameter is defined in DIN 51005 (“Thermal analysis”) and is also referred to as the extrapolated peak onset temperature. It represents the onset of the phase transformation and is defined as the intersection point of the extrapolated baseline with the tangent through the linear part of an increasing or decreasing peak flank.
The advantage of this parameter is that it is independent of sample mass and heating and cooling rates. The width of the two-phase region is defined as the difference between the first onset temperatures:
The influence of composition on the transition temperatures TÜ1 and TÜ2 is determined by means of DSC measurement.
The peak Curie temperatures Te of heating (DSC) and cooling (DSC) are indicated by diamonds. For the special molten masses with lower V contents, Tc is the temperature of the phase transition. The highest measured maximum permeability μmax (triangle) is plotted on the secondary axis. The highest maximum permeabilities are achieved for V contents of between 1 and 3 wt %.
Final annealing is carried out to set the soft magnetic properties. In this test it was always carried out in a H2 protective atmosphere. The H2 quality used was always stets hydrogen 3.0 (or technical hydrogen) with a H2 percentage >99.9%, where H2O≤40 ppm-mol, O2≤10 ppm-mol, N2≤100 ppm-v.
The magnetic properties of the alloys were tested using strip stock manufactured from 5 kg heavy ingots. The alloys were melted in a vacuum and then poured into a flat mould at approx. 1500° C. Once the oxide skin had been milled off the individual ingots, they were hot rolled into 3.5 mm thick strips at a temperature of approx. 1000° C. to 1300° C. The resulting hot-rolled strips were then pickled to remove the oxide skin and cold rolled to a thickness of 0.35 mm. Sample rings were stamped and resistor strips were cut out of the strip in order to characterise the magnetic properties. The electrical resistance ρ was determined on the resistor strips. Maximum permeability μmax, coercive field strength Hc, inductions B at field strengths of 20, 25, 50, 90, 100 and 160 A/cm, remanence Br and hysteresis losses PHys were measured on the sample rings in the annealed state at room temperature. Hysteresis losses were determined by measuring the losses at an induction amplitude of 1.5 T for various frequencies. The axis intercept determined by linear regression in the plot P/f over f gives the hysteresis losses.
A disc was sawn off the ingots to analyse the elements. The results of the analysis appear in Tables 1 to 4. Table 1 shows the wet-chemical analysis of the metallic elements in order to determine the basic composition. Residual iron and other elements <0.01% are not indicated, the data being given in wt %. Table 2 shows the analysis by hot gas extraction of non-metal impurities in the batches from Table 1, the data being given in wt %. Table 3 shows the wet-chemical analysis of the metallic elements in order to fine-tune the basic composition and to limit the composition ranges and impurities. Residual iron and other elements <0.01% are not specified. Data is given in wt %. In batches 930502 and 930503 the feed material used was iron with a high level of impurities. Table 4 shows the analysis by hot gas extraction of non-metallic impurities in the batches from Table 3, the data being given in wt %.
Table 3 also shows the analysis of the metallic elements in two large melts. Residual iron and the P content of large melt 76/4988 is 0.003 wt %, the P content of large melt 76/5180 is 0.002 wt %, other elements <0.01% are not specified. Table 4 also shows the analysis by hot gas extraction of non-metallic impurities in the two large melts from Table 3, the data being given in wt %.
These figures indicate that vanadium reduces low inductions less strongly than chromium and that chromium increases electrical resistance more strongly than vanadium at the same decrease in saturation (B160). Co increases saturation (B160) but has less influence on low induction values and on electrical resistance.
Table 7 shows annealing variants according to the invention of batch 93/0330 with a strip thickness of 0.35 mm in comparison to annealing variants not according to the invention (see
Table 6 shows DSC transition temperatures and Curie temperatures Tc. Temperatures are given in ° C. #NV signifies that no signal is discernible in the DSC measurement.
One of the reasons for the very good soft magnetic properties is the grain structure achieved in the FCC region after annealing, which is unusually large for Fe—Co alloys.
After a short period of annealing of 4 h at 1050° C. in batch 93/0330 (Fe-17Co-2V), for example, grain sizes of 354 to 447 μm were determined. Similarly large grains could only be achieved by annealing in the BCC range after annealing lasting several days.
Batch 930330 was tested by way of example to compare the aforementioned annealing variants. Table 8 shows the results after step annealing in the first annealing step (batch 93/0330) (see
To find the optimum annealing temperature, samples are annealed at different annealing temperatures and then measured. If the number of annealing processes required is greater than the number of samples available, the same set of samples is generally annealed at different temperatures. This so-called “step annealing” starts at a low starting temperature and anneals at successively higher temperatures. Step annealing can be used to detect precipitation regions, recrystallization temperatures and phase transformations, for example, that have a direct influence on magnetic characteristics.
The magnetic properties were measured for alloys of various compositions after various annealing processes. The results are given in Tables 10 to 24, giving values B20, B25, B50, B90, B100, B160 (T) Hc (A/cm), μmax, Br (T) and PHys 1.5 T (Ws/kg).
Table 10 shows the results after annealing a selection of batches at 850° C. for 4 h at a cooling rate of 150° C./h. These embodiments are not in accordance with the invention.
Table 11 shows the results after annealing a selection of batches for 10 h at 910° C. at a cooling rate of 150° C./h. No demagnetisation was carried out prior to measuring the static values. These embodiments are not in accordance with the invention.
Table 12 shows the results after annealing a selection of batches for 10 h at 910° C. and cooling to room temperature, followed by annealing for 70 h at 930° C. The cooling rate is 150° C./h. No demagnetisation was carried out prior to measuring the static values. These embodiments are not in accordance with the invention.
Table 13 shows the results after annealing a selection of batches for 4 h at 1000° C.
Cooling rate 150° C./h. No demagnetisation was carried out prior to measuring the static values.
Table 14 shows the results after annealing a selection of batches in the first annealing step for 4 h at 1000° C. with cooling to room temperature, following by a second annealing step for 10 h at 910° C. The cooling rate is 150° C./h. No demagnetisation was carried out prior to measuring the static values.
Table 15 shows the results after annealing all the Fe—Co—V—Cr batches for 4 h at 1050° C. Cooling rate 150° C./h. No demagnetisation was carried out prior to measuring the static values. The resistances of batches 930322 to 930339 were measured after annealing for 4 h at 850° C. In V-rich batches 930422 and 930423 TÜ2 was only just below 1050° C. Adjusted annealing steps are indicated in Table 18.
Table 16 shows the results after annealing all the Fe—Co—V—Cr batches in a first annealing step for 4 h at 1050° C. with cooling to room temperature, followed by a second annealing step for 10 h at 910° C. Cooling rate 150° C./h. Demagnetisation was carried out prior to measuring. In the batches marked in grey, TÜ1 is either not far enough above or too far above 910° C. Adjusted annealing steps are indicated in Table 17.
Table 17 shows the results after adjustment of the annealing processes on the batches in which the transition temperatures of the DSC measurement (Table 6) do not or only just coincide with annealing for 4 h at 1050° C.+10 h at 910° C. (Tables 15 and 16). The cooling rate is 150° C./h. When annealing was carried out for 4 h at 1050° C. no demagnetisation was carried out prior to measuring. In all other cases demagnetisation was carried out prior to measuring.
Table 18 shows the results after annealing of batch 930423 in various phase regions to clarify the influences of the ferromagnetic and paramagnetic BCC region on magnetic properties (see also
Table 19 shows the results after annealing a selection of batches for 4 h at 1050° C. followed by slow cooling to room temperature at 50° C./h. No demagnetisation was carried out prior to measuring the static values.
Table 20 shows the results after annealing a selection of batches for 4 h at 1050° C. with slow cooling to room temperature at 50° C./h and a second annealing step for 10 h at 910° C. with furnace cooling at approx. 150° C./h. No demagnetisation was carried out prior to measuring the static values.
Table 21 shows the results after annealing a selection of batches for 4 h at 1100° C. The cooling rate is 150° C./h. No demagnetisation was carried out prior to measuring the static values except on batches 930422 and 930423.
Table 22 shows the results after annealing a selection of batches in a first annealing step for 4 h at 1100° C. and cooling to room temperature followed by a second annealing step for 10 h at 910° C. The cooling rate is 150° C./h. No demagnetisation was carried out prior to measuring the static values.
Table 23 shows the results after annealing a selection of batches for 4 h at 1150° C. The cooling rate is 150° C./h. No demagnetisation was carried out prior to measuring the static values except on batch 930442.
Table 24 shows the results after annealing a selection of batches in a first annealing step for 4 h at 1150° C. and cooling to room temperature followed by a second annealing step for 10 h at 910° C. The cooling rate is 150° C./h. No demagnetisation was carried out prior to measuring the static values.
Table 25 shows the data for maximum permeability and density for various Fe-17Co alloy compositions with various additives. Based on the binary alloy Fe-16.98Co, its measured density of 7.942 g/cm3 and its average atomic weight of 56.371 g/mol (calculated from the metallic alloy element contents analysed), the fictitious density of Fe-17Co alloys with added V, Cr, Mn, Si, Al and other metallic elements is calculated using their average atomic weights and compared with the measured density. For the alloy Fe-17.19Co-1.97V (batch 93/0330), for example, the average atomic weight is 56.281 g/mol. It is then possible, using the rule of three (7.942 g/cm3×56.281/56.371=7,929 g/cm3), to calculate the fictitious density that this alloy Fe-17.19Co-1.97V should have if its lattice constant were unchanged in relation to the binary Fe-16.98Co alloy. In reality, however, the density measured for this alloy, 7.909 g/cm3, is −0.26% lower than the fictitious density of 7.929 g/cm3. This signifies that the lattice constant of this alloy must be approx. 0.085% greater than that of the binar alloy.
Table 26 shows the data for selected batches and annealing processes that have both particularly high maximum permeabilities and low hysteresis losses at the same time as a very high level of induction B at 100 A/cm (B100).
Table 27 shows the data for the impurities C and S in ppm for selected batches and annealing processes. These impurities are effectively reduced by annealing at 1050° C. in hydrogen.
Table 28 shows magnetic values for the two large melts 76/4988 and 76/5180. The letters A and B refer to ingots A and B; the molten masses were poured into two moulds. The specific resistance of batch 76/4988 is 0.306μΩm; that of batch 76/5180 is 0.318μΩm.
Table 29 shows for various batches the transition temperatures TÜ1 and TÜ2 and the best coercive field strength Hc achieved for these Fe-17Co special melts with different V contents, including details of the annealing treatment. The alloys also contain up to a total of 0.6 wt % Cr and/or Si.
In the first annealing step (
In the second annealing step (
In summary, it can be said that the best magnetic properties are achieved if the first annealing step takes place at above TÜ2 and the second annealing step takes place at below TÜ1.
The influence of the degree of cold deformation on the magnetic properties is tested.
Cold deformation (KV) on strip stock with a final thickness D2 is defined as the percentage reduction in thickness in relation to a non-cold-deformed starting thickness D1 since expansion during rolling can be disregarded. The non-cold-deformed starting thickness D1 may, for example, be achieved by hot rolling or by intermediate annealing (ZGL or int. anneal).
KV[%]=[(D1−D2)/D1]×100
In
Assuming a constant D1 of 3.5 mm (hot rolling thickness), cold deformation achieved by a high degree of rolling to 0.20 mm and 0.10 mm once again results in an increase in Hc, as indicated by the broken line. This can be explained by the fact that too many nucleation sites for grains occur at the highest degrees of cold deformation and the grains obstruct one another's growth during annealing. As a result, the alloy in batch 930329 (Fe-17Co-0.5Cr-1.5V) (in wt %) produced without intermediate annealing after final annealing for 4 h at T1=1000° C. and for 10 h at T2=910° C. has an average grain size of 0.25 mm at a final thickness of 0.35 mm; an average grain size of 0.21 mm at a final thickness of 0.20 mm; and an average grain size of 0.15 mm at a final thickness of 0.10 mm. There is therefore an optimum degree of cold deformation of approx. 90%.
In order to test whether texture formation is a significant factor for magnetic properties, the texture was determined by means of X-ray diffraction on sheets measuring 50 mm×45 mm.
Here the sample was subject to angle-dependent Cu-Kα=0.154059295 nm radiation and the diffracted intensity was measured with a 2 mm pinhole aperture. A Lynxexe semi-conductor strip detector with 2° angular range and energy-dispersive operation was used as the detector. As shown by the (200) pole figures, for example, a rolling texture is present in the unannealed, full hard state that dissolves completely after annealing in the FCC region for 4 h at 1050° C. in H2.
The lack of texture also corresponds to the measurements of the directional Hc. Five Hc strips with dimensions of 50 mm×10 mm were taken from various directions relative to the direction of rolling (longitudinally=0°, diagonal=45°, transversely=90°) and measured in a Förster coercimeter.
Following annealing for 4 h at 910° C., the mean values exhibit anisotropic behaviour, though this anisotropy is not significant if statistical errors are taken into account. However, this slight anisotropy corresponds to residual texture from the corresponding pole figure (top centre image in
Below, the magnetic properties of the alloy according to the invention are compared with comparative alloys based on the example of batches 930329 (Fe-17Co-1.5V-0.5Cr) and 930330 (Fe-17Co-2.0V) according to the invention. The comparative alloys shown are TRAFOPERM N4 (Fe-2,5Si—Al—Mn), a typical electrical steel; three FeCo VACOFLUX 17 alloys (Fe-17Co-2Cr—Mo—V—Si); VACOFLUX 48 (Fe-49Co-1.9V) and a HYPOCORE special melt. The HYPOCORE special melt was melted according to the composition published by Carpenter Technologies (Fe-5Co-2.3Si-1Mn-0.3Cr—values in wt %).
In order to test the properties of the alloys according to the invention on a production scale, two large melts were carried out using the normal manufacturing process. 2.2 t of the desired composition were melted in a vacuum induction furnace and, once the exact composition had been set and analysed, poured into two round moulds with a diameter of 340 mm. After solidification and cooling, the round ingots were removed from the moulds and heated to a temperature of 1170° C. for hot rolling in a gas-fired rotary hearth furnace. The heated ingots were then hot rolled on a blooming roll to form slabs with a cross section of 231×96 mm2. These slabs were then ground on all sides to a dimension of 226×93 mm2 to remove the oxide skin.
Both slabs obtained from batch 76/4988 in this manner were rolled out on a hot rolling mill to form hot strip. To this end, the slabs were first heated at a temperature of 1130° C. and then, once sufficiently warmed through, rolled to form hot strip. The final thickness chosen for one of the strips was 2.6 mm. The final rolling temperature of this band was 900° C., the reeling temperature 828° C. The final thickness chosen for the other strip was 1.9 mm. The final rolling temperature of this strip was 871° C., the reeling temperature 718° C. Both hot strips were then blasted to remove the oxide skin. One part of the hot-rolled strip was intermediate annealed for 1 h at 750° C. in an H2 inert gas atmosphere. Another part of the hot-rolled strip was intermediately annealed for 1 h at 1050° C. in a H2 inert gas atmosphere. A remaining part of the hot-rolled strip did not undergo intermediate annealing. The strips were then rolled to their final thicknesses, oxides being removes from both sides of the strips at an intermediate thickness. Before the strip was hot rolled, sections with a thickness of 15 mm were also sawn off the slabs and made into a strip by hot rolling (to a thickness of 3.5 mm), pickling the hot strip thus obtained and then cold rolling in the pilot plant. The results obtained are also presented for the purposes of comparison.
In the case of batch 76/5180, a disc with a thickness of 15 mm was sawn off either end of the two slabs. These discs were preheated at 1200° C. and then hot rolled to form a strip with a thickness of 3.5 mm. The hot strips obtained in this manner were picked to remove oxides, then cold rolled to a thickness of 0.35 mm.
Stamped rings were produced from all the strips obtained in this way and then subjected to an annealing process. Table 28 shows the results obtained for the magnetic values. The specific resistance of batch 76/4988 is 0.306 μΩm; that of batch 76/5180 is 0.318 μΩm.
As is apparent from Table 28, better magnetic properties are measured for samples from the large melt than for the commercially available alloys with a Co content of below 30 percent by weight such as VACOFLUX 17. For a sample from the large melt 76/5180B, a maximum permeability of above 20,000 was measured. The alloy according to the invention is therefore suitable for the industrial-scale production of strip stock with improved magnetic properties.
The alloy according to the invention exhibits higher inductions than VACOFLUX 17 for all field strengths. At inductions above the inflection point, the new alloy lies between TRAFOPERM N4 and VACOFLUX 48. For both batches, the air flow-corrected induction B at a field strength of 400 A/cm close to magnetic saturation is 2.264 T (corresponding to a polarisation J of 2.214 T). In the operating range of typical electric motors and generators torque for the new alloy will therefore be higher than for VACOFLUX 17 and TRAFOPERM N4.
A comparison of 930329 and 930330 indicates that vanadium in conjunction with the heat treatment described above increases the rectangularity of the hysteresis loop to such an extent that, depending on the additive, maximum permeability is almost as high as that of VACOFLUX 48. This is surprising, not to say astounding, since the anisotropy constant K1 shows a zero crossing at approx. 50% Co that is not present at 17% Co. By contrast, at 17% Co the anisotropy constant K1 in the Fe—Co system is very high.
Very good soft magnetic properties are also apparent in the hysteresis losses, which are on a level comparable with those of TRAFOPERM N4. As frequency rises, TRAFOPERM N4 losses at identical strip thickness increase due to the higher electrical resistance, though less strongly than with the new alloy. It is, however, possible to compensate for this effect by selecting a somewhat smaller strip thickness with correspondingly lower eddy current losses.
In summary, a high permeability soft magnetic alloy is provided that offers both better soft magnetic properties, e.g. appreciably higher permeability and lower hysteresis losses, and higher saturation than existing, commercially available FeCo alloys. At the same time, however, this new alloy also offers significantly lower hysteresis losses than previously known commercially available alloys with Co contents between 10 and 30 wt % and, above all, an appreciably higher level of permeability never previously achieved for this type of alloy. The alloy according to the invention can also be produced cost effectively on an industrial scale.
2)
1)
| Number | Date | Country | Kind |
|---|---|---|---|
| 10 2017 009 999.5 | Oct 2017 | DE | national |
| 10 2018 112 493.7 | May 2018 | DE | national |
This U.S. national phase patent application claims the benefit of PCT/EP2018/079339, filed Oct. 25, 2018, which claims the benefit of DE application no. 10 2017 009 999.5, filed Oct. 27, 2017, and DE application no. 10 2018 112 493.7, filed May 24, 2018, the entire contents of which are incorporated herein by reference for all purposes.
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/EP2018/079339 | 10/25/2018 | WO | 00 |
