Patent Application
20250207229
The invention relates to making thin steel strip in a twin roll caster and in downstream processing steps that is suitable for use in electrical steel applications.
In a twin roll caster, molten metal, typically steel, is delivered from a delivery system to a casting pool supported on casting surfaces of a pair of counter-rotated horizontal casting rolls, which are internally water cooled so that solidified metal shells form on the moving casting roll surfaces. The metal shells are brought together at a nip between them to produce a solidified strip product delivered downwardly from the nip between the casting rolls. The term “nip” is used herein to refer to the general region at which the casting rolls are closest together.
The molten metal may be poured from a ladle into a smaller vessel or series of smaller vessels from which it flows through a metal delivery nozzle or nozzles located above the nip, to form the casting pool of molten metal supported on the casting surfaces of the casting rolls above the nip and extending the length of the nip.
The casting pool is usually confined between side plates or side dams held in sliding engagement with end portions of the casting rolls to restrict the casting pool against outflow. The upper surface of the casting pool (generally referred to as the “meniscus” level) is usually above the lower end of the delivery nozzle so that the lower end of the delivery nozzle is immersed within the casting pool.
When casting steel strip by a twin roll caster, the thin strip exits the nip, passes across a guide table, through a pinch roll stand and then through a hot rolling mill, where the thin strip is reduced to a desired thickness. The hot rolled strip is then cooled to form strip with a required microstructure for end use applications. The cooled strip is then coiled, with a shear cutting the strip periodically upstream of the coiler to form required lengths of strip in each coil.
Depending on the steel compositions and the casting conditions, twin roll cast, hot rolled and cooled strip can be used in a range of end-use applications.
Electrical steels are the focus of the invention.
Electrical steels are low carbon iron-silicon soft magnetic steels which are widely used in electric motors, sensors, power generators, and transformers. Electrical steels can be classified into non-grain-oriented (NGO) and grain-oriented (GO) electrical steels. Cold-rolled non-grain-oriented steel (CRNGO) is generally less expensive than cold-rolled grain-oriented steel (CRGO). Thus, when the cost is important, or when the direction of magnetic flux for an end use application is not constant, NGO electrical steel is used.
When used in an electrical motor, NGO electrical steels are cut to thin laminations that are isolated with insulating coating layers and stacked to form the motor core to decrease eddy current losses. Thus, if thin strip NGO electrical steels can be continuously cast directly to a thin strip, that will be helpful to save the energy and time. The applicant has realized that the twin roll strip casting method has excellent potential because it can directly and continuously cast thin, typically 1-2 mm thick, steel strip.
When NGO electrical steels are cast by a conventional thick slab continuous casting process, it is difficult to maintain the θ-fiber texture which is always formed in a hot rolling step in the process. During subsequent cold rolling and recrystallization annealing steps in the process, strong α-fiber and γ-fiber textures are commonly observed. The formation of deformed matrix shear bands directly influenced the nucleation and growth of Goss grains in the recrystallization annealing process. Furthermore, because of the low local dislocation density and the sharp lattice curvature, the retained Goss grains after cold rolling can also promote the formation of Goss grains in the recrystallization annealing step in the process.
Based on the end use applications, it is important to control the magnetic properties of NGO electrical steels. The magnetic properties of NGO electrical steels are highly influenced by grain size and texture, which are in turn influenced by rolling and recrystallization annealing steps. It has been reported that, after final annealing, 2.0 wt. % Si NGO electrical steel has strong {1 1 0} (0 0 1) (Goss) texture. It has also been reported that, for 4.5 wt. % Si NGO electrical steel, two-step cold rolling can help to form coarse grains with a strong Goss and near Goss recrystallization texture in the annealing process.
As noted in the preceding paragraph, the magnetic properties of electrical steels are highly influenced by texture. In electrical steels, the <001> axis direction is easily magnetized, while the <111> axis direction is more difficult to magnetize. It is known that Goss and Cube orientations optimize magnetic properties. Brass and Goss orientations are reported to be formed in the BCC metals through shear deformation texture. It has also been reported that many Goss grains and Cube grains are formed at the shear bands within the γ-fiber deformed regions. It has also been reported that some Cube components are retained after the heavy cold rolling process because the Cube deformation bands also serve as the nucleation sites of the new Cube grains.
For NGO electrical steels, it is difficult to control texture during recrystallization annealing. It has also been reported that, in some cases, a phase transformation can be used to obtain the ideal orientation for magnetic properties. During annealing, because of anisotropic strain energy, some {100} oriented grains are formed when austenite transforms to ferrite. However, this transformation is not available for fully ferritic steel compositions. Methods to achieve ideal crystallographic orientations with chemical compositions without a phase transformation have also been studied in the literature. The influence of initial annealing on texture evolution and magnetic properties for a 3.4 wt. % Si electrical steel with 0.003 wt. % C has been studied. The recrystallization kinetics of a 3 wt. % Si electrical steel have also been studied.
It has also been reported that the rolling process has a significant influence on texture evolution, grain growth, and magnetic properties. In the recrystallization annealing process, some ideal textures evolve from deformed shear bands which are formed by the rolling process. Furthermore, the phase transformation during hot deformation also affects the recrystallization rate and grain size in the subsequent annealing process. It has also been reported that reported that for a 6.2 wt. % Si electrical steel with less than 0.01 C, hot rolling was beneficial to the final magnetic properties. It has also been reported that the beneficial effects of annealing prior to cold rolling on the electrical steel microstructure and magnetic properties.
Although there have already been some studies conducted that examine the influence of rolling and recrystallization annealing in manufacturing NGO electrical steels, there are few studies about the influence of rolling and annealing on thin electrical steel strip produced by a twin-roll casting method.
The above comments are not an admission of the common general knowledge in China or elsewhere.
In broad terms, the invention includes a method and an apparatus for producing an electrical steel strip and a product electrical steel strip. The method includes twin roll strip casting a thin strip of less than 3 mm in thickness from an electrical steel melt having a superheat temperature of at least 30° C. above the liquidus temperature Tliquidus of the melt, hot rolling, cooling and coiling the strip. The melt comprises: by weight, up to 0.015% carbon, between 1.0% and 2.0% manganese, between 2.70% and 3.80% silicon, silicon killed (with no purposeful addition of aluminium) containing less than 0.01% aluminum, and optionally any one or more of Cu, Cr, Ni, Mo, Ti, Nb, V, Sb, Sn, and the remainder iron, impurities and inclusions.
The invention includes a twin roll casting method of producing an electrical steel strip that includes:
The Mn concentration range of between 1.0% and 2.0% manganese makes it possible for the melt to have a low Al concentration (i.e. less than 0.01% aluminum), without an impact on resistivity of the resultant electrical steel strip. The melt is a silicon killed, as opposed to an Al-killed, steel.
The term “impurities” means elements in the melt as an inevitable result of steelmaking practices or as a consequence of the feed materials for steelmaking. These tend to be non-metallic elements. Examples of impurities are N, P, S, and H.
The term “inclusions” means compounds that form during steelmaking. Examples of “inclusions” include AlN and MnS.
The Cu concentration may be less than 0.3%.
The Cr concentration may be less than 0.2%.
The Cr concentration may greater than 0.01%.
The Ni concentration may be less than 0.2%.
The Mo concentration may be less than 0.1%.
The Mo concentration may be less than 0.06%.
The Mo concentration may be greater than 0.001%.
The concentration of Sb may satisfy the formula Sn+2*Sb<0.4%.
The concentration of Sb may be less than 0.2%.
The concentration of Sb may be less than 0.1%.
The Sb concentration may be greater than 0.001%.
The carbon concentration may be up to 0.0060%.
The carbon concentration may be up to 0.0080%.
The carbon concentration may be up to 0.01%.
The manganese concentration may be at least 1.10%.
The manganese concentration may be at least 1.20%.
The manganese concentration may be less than 1.70%.
The manganese concentration may be less than 1.80%.
The manganese concentration may be between 1.10% and 1.55%.
The silicon concentration may be at least 2.80%.
The silicon concentration may be at least 3.00%.
The silicon concentration may be no more than 3.70%.
The silicon concentration may be no more than 3.60%.
The silicon concentration may be no more than 3.50%.
The silicon concentration may be between 2.70% and 3.50%.
The concentrations of C, Mn and Si in the electrical steel melt may be as follows:
The concentrations of residuals in the electrical steel melt may be as follows in one embodiment:
The concentrations of residual gases in the electrical steel melt may be as follows:
The electrical steel in the coiled strip may have an at least substantially ferritic microstructure.
The microstructure of the electrical steel in the coiled strip may be at least substantially equiaxed grains.
The method may include cold rolling the electrical steel strip to further reduce the thickness of the strip.
The method may include cold rolling the electrical steel strip to further reduce the thickness of the strip to be no more than 0.50 mm, typically no more than 0.35 mm.
The method may include pickling the electrical steel strip before the cold rolling step.
The method may include annealing the electrical steel strip to obtain non-grain oriented electrical steel strip with desired magnetic properties.
The desired magnetic properties may be core loss or permeability. These terms are well-understood by persons skilled in the art of electrical steels and are covered by industry standards. Typically, a customer places an order with a steelmaker for an electrical steel having a particular thickness and core loss and/or permeability.
The desired magnetic properties may be as described in Chinese Standard GB/T 2521.1 entitled “Cold-rolled electrical steel delivered in the fully-processed state-Part 1: Grain non-oriented steel strip (sheet)”.
For example, the desired magnetic properties may be as described in the standard as 35W250, 0.35 mm, core loss P1.5/50 of 2.50 Watts/kg of electrical steel.
By way of further example, the desired magnetic properties may be as described in the standard as 35W300, 0.35 mm, core loss P1.5/50 of 3.00 Watts/kg of electrical steel.
The annealing step may be conducted in a controlled atmosphere, such as a mixed gas of hydrogen and nitrogen.
The annealing step may be conducted in a controlled atmosphere that reduces the carbon concentration to be no more than 0.0030%.
The electrical steel melt may be made by any suitable steelmaking method.
By way of example, the electrical steel melt may be made by and then transferred to the twin roll strip caster by the following steps: electric arc steel making furnace (EAF)→ladle furnace→tank or RH degasser→tundish→twin roll strip caster.
Alternatively, the electrical steel melt may be made by and then transferred to a twin roll strip caster by the following steps: basic oxygen steel making furnace (BOF)→ladle furnace→RH degasser→tundish→twin roll strip caster.
The method may include superheating the electrical steel melt to the superheat temperature before transferring the melt to the twin roll caster.
By way of example, the method may include superheating the electrical steel melt to the superheat temperature in any one or more of a steelmaking furnace, a ladle or tundish or other vessel that transfers the electrical steel melt from the steelmaking furnace to the twin roll caster.
By way of particular example, the method may include superheating the electrical steel melt to the superheat temperature in a tundish by means of tundish heating technology.
The method may include preparing the molten electrical steel melt with the superheat temperature.
The method may include delivering the molten electrical steel melt with the superheat temperature to the twin roll caster.
The superheat temperature may be Tliquidus+up to 120° C., i.e. up to 120° C. above the liquidus temperature.
The superheat temperature may be Tliquidus+up to 90° C., i.e. up to 90° C. above the liquidus temperature.
The superheat temperature may be Tliquidus+up to at least 35° C., i.e. at least 35° C.° C. above the liquidus temperature.
The superheat temperature may be Tliquidus+up to at least 40° C., i.e. at least 40° C. above the liquidus temperature.
The method may include controlling the hot rolling step so that there is a hot rolling mill exit temperature of 800-900° C.
The method may include supplying nitrogen to a hot box that encloses the electrical steel strip between the twin roll caster and the hot rolling mill at a rate of greater than 2500 m3/hr so that there is a high N concentration in the hot box, noting that the value of 2500 m3/hr may vary depending on air leakage into the hot box.
The method may include controlling the hot rolling step so that there is a hot rolling mill exit temperature of 720-820° C.
The method may include supplying nitrogen to a hot box that encloses the electrical steel strip between the twin roll caster and the hot rolling mill at a rate of less than 2500 m3/hr so that there is a low N concentration in the hot box, noting that the value of 2500 m3/hr may vary depending on air leakage into the hot box.
The method may include varying the nitrogen flow rate to a hot box that encloses the electrical steel strip between the twin roll caster and the hot rolling mill to control strip temperature, noting that lower N leads to more air leaking into the hot box, with the air oxidizing strip and increasing its emissivity and thus increasing heat loss and reducing strip temperature at a hot rolling mill entry.
The method may include hot rolling the cast strip with the cast strip entering the hot rolling mill at a mill entry temperature of 140-160° C., typically 150° C.° C. higher than the mill exit temperature.
The method may include coiling the hot rolled and cooled electrical steel strip at a coiler entry temperature in a range of 550-720° C.
The invention also provides an apparatus for producing an electrical steel strip, including
The apparatus may include a cold rolling mill to reduce the thickness of the electrical steel strip.
The apparatus may include an annealing unit for annealing the electrical steel strip to obtain non-grain oriented electrical steel strip with desired magnetic properties.
The electrical steel strip may be produced from the molten electrical steel melt as described above.
The twin roll strip caster may be configured to cast the molten electrical steel melt with a superheat temperature Tliquidus of up to 120° C. above the liquidus temperature.
The twin roll strip caster may be configured to cast the molten electrical steel melt with a superheat temperature Tliquidus of up to 90° C. above the liquidus temperature.
The twin roll strip caster may be configured to cast the molten electrical steel melt with a superheat temperature Tliquidus of at least 35° C. above the liquidus temperature.
The twin roll strip caster may be configured to cast the molten electrical steel melt with a superheat temperature Tliquidus of at least 40° C. above the liquidus temperature.
The invention also provides a twin roll strip cast and hot rolled electrical steel strip of less than 3 mm in thickness comprising a composition of by weight, up to 0.015% carbon, between 1.0% and 2.0% manganese, between 2.70% and 3.80% silicon, silicon killed containing less than 0.01% aluminum, up to 0.4% Cu; up to 0.3% Cr; up to 0.3% Ni; up to 0.2% Mo; up to 0.01% Ti; up to 0.005% Nb; up to 0.005% V, up to 0.3% Sb; up to 0.3% Sn; and the remainder iron, impurities and inclusions.
The concentrations of the residuals may be as described in Table 2.
The electrical steel strip may have a microstructure that is at least substantially a ferritic microstructure.
The microstructure of the electrical steel strip may be at least substantially equiaxed grains.
The Cu concentration may be less than 0.3%.
The Cr concentration may be less than 0.2%.
The Cr concentration may greater than 0.01%.
The Ni concentration may be less than 0.2%.
The Mo concentration may be less than 0.1%.
The Mo concentration may be less than 0.06%.
The Mo concentration may greater than 0.001%.
The concentration of Sb may satisfy the formula Sn+2*Sb<0.4%.
The concentration of Sb may be less than 0.2%.
The concentration of Sb may be less than 0.1%.
The Sb concentration may greater than 0.001%.
The carbon concentration may be up to 0.0060%.
The carbon concentration may be up to 0.0080%.
The carbon concentration may be up to 0.01%.
The manganese concentration may be at least 1.10%.
The manganese concentration may be at least 1.20%.
The manganese concentration may be less than 1.70%.
The manganese concentration may be less than 1.80%.
The manganese concentration may be between 1.10% and 1.55%.
The silicon concentration may be at least 2.80%.
The silicon concentration may be at least 3.00%.
The silicon concentration may be no more than 3.70%.
The silicon concentration may be no more than 3.60%.
The silicon concentration may be no more than 3.50%.
The silicon concentration may be between 2.70% and 3.50%.
The invention also provides a twin roll strip cast, hot rolled, cold rolled and annealed electrical steel strip of less than 3 mm in thickness comprising a composition of by weight, up to 0.015% carbon, between 1.0% and 2.0% manganese, between 2.70% and 3.80% silicon, up to 0.4% Cu; up to 0.3% Cr; up to 0.3% Ni; up to 0.2% Mo; up to 0.01% Ti; up to 0.005% Nb; up to 0.005% V, up to 0.3% Sb; up to 0.3% Sn; and the remainder iron, impurities and inclusions.
The electrical steel strip may have a microstructure that is at least substantially a ferritic microstructure.
The microstructure of the electrical steel strip may be at least substantially equiaxed grains.
The concentrations of carbon, manganese, silicon, and residuals may be as described above, noting that the carbon concentration may be lower than the melt concentration due to de-carburization when the electrical steel strip is annealed.
The electrical steel strip may have desired magnetic properties.
The desired magnetic properties may be as described in Chinese Standard GB/T 2521.1-2016 entitled “Cold-rolled electrical steel delivered in the fully-processed state-Part 1: Grain non-oriented steel strip (sheet)”.
For example, the desired magnetic properties may be 35W250 0.35 mm Core loss P1.5/50-2.50 Wkg-1 or 35W300 0.35 mm Core loss P1.5/50-3.00 Wkg-1.
It is noted that, whilst there may be 100% ferrite in the microstructures of (a) twin roll strip cast and hot rolled electrical steel strip and (b) twin roll strip cast and hot/cold rolled and annealed electrical steel strip, the invention extends to situations whether this is not the case and, for example, there is some austenite in the microstructure.
It is also noted that the microstructure will be a function of a range of factors, including composition, superheat, and selections of casting, hot rolling, cold rolling and annealing conditions.
In order that the invention may be described in more detail, some illustrative examples will be given with reference to the accompanying drawings in which:
The following description of embodiments of electrical steel strip produced by a twin roll strip casting method are not the only embodiments of the invention.
In addition, the following description of an embodiment of a twin roll caster is not the only embodiment of a twin roll caster suitable to produce electrical steel strip in accordance with the invention.
All other embodiments obtained by the ordinary person skilled in this art based on the described embodiments of the invention without any creative endeavors fall into the protection scope of the invention.
Unless defined otherwise, the technical terms or scientific terminology as used in the present disclosure should take the meaning usually understood by the ordinary person skilled in this art of invention.
Referring now to
The twin roll caster includes the pair of counter-rotatable casting rolls 12 having casting surfaces 12A laterally positioned to form a nip 18 there between. Molten metal, more particularly molten electrical steel described further below, is supplied from a ladle 13 through a metal delivery system to a metal delivery nozzle 17 (core nozzle) positioned between the casting rolls 12 above the nip 18. Molten metal thus delivered forms a casting pool 19 of molten metal above the nip 18 supported on the casting surfaces 12A of the casting rolls 12. This casting pool 19 is confined in the casting area at the ends of the casting rolls 12 by a pair of side closure plates, or side dams 20. The upper surface of the casting pool 19 (generally referred to as the “meniscus” level) may rise above the lower end of the delivery nozzle 17 so that the lower end of the delivery nozzle 17 is immersed within the casting pool 19. The casting area includes the addition of a protective atmosphere above the casting pool 19 to inhibit oxidation of the molten metal in the casting area.
The ladle 13 typically is of a conventional construction supported on a rotating turret 40. For metal delivery, the ladle 13 is positioned over a movable tundish 14 in the casting position to fill the tundish 14 with molten metal. The movable tundish 14 may be positioned on a tundish car 66 capable of transferring the tundish 14 from a heating station (not shown), where the tundish 14 is heated to near a casting temperature, to the casting position.
The movable tundish 14 may be fitted with a slide gate 25, actuable by a servo mechanism, to allow molten metal to flow from the tundish 14 through the slide gate 25, and then through a refractory outlet shroud 15 to a transition piece or distributor 16 in the casting position. From the distributor 16, the molten metal flows to the delivery nozzle 17 positioned between the casting rolls 12 above the nip 18.
The side dams 20 may be made from a refractory material such as zirconia graphite, graphite alumina, boron nitride, boron nitride-zirconia, or other suitable composites. The side dams 20 have a face surface capable of physical contact with the casting rolls 12 and molten metal in the casting pool 19. The side dams 20 are mounted in side dam holders (not shown), which are movable by side dam actuators (not shown), such as a hydraulic or pneumatic cylinder, servo mechanism, or other actuator to bring the side dams 20 into engagement with the ends of the casting rolls 12. Additionally, the side dam actuators are capable of positioning the side dams 20 during casting. The side dams 20 form end closures for the molten pool of metal on the casting rolls 12 during the casting operation.
The casting rolls 12 are internally water cooled as described below so that as the casting rolls 12 are counter-rotated, shells solidify on the casting surfaces 12A, as the casting surfaces 12A move into contact with and through the casting pool 19 with each revolution of the casting rolls 12. The shells are brought close together at the nip 18 between the casting rolls 12 to produce a cast thin strip product 21 delivered downwardly from the nip 18. The cast thin strip product 21 is formed from the shells at the nip 18 between the casting rolls 12 and delivered downwardly and moved downstream as described above.
In operation, the strip leaves the nip at temperatures of the order of 1400° C. and greater. To prevent oxidation and scaling of the strip, the metal strip is cast downwardly into the enclosure 27 supporting a protective atmosphere immediately beneath the casting rolls in the casting position. The enclosure 27 may extend along the path of the cast thin strip until the first pinch roll stand 31 and may extend along the path of the cast thin strip until the hot rolling mill 32 to reduce oxidation and scaling.
After the hot rolling mill 32, the rolled thin strip then passes into a cooling station 97 where the strip is cooled by water that is delivered by spray nozzles 90 of a plurality of rows of water spray assemblies extending across the run-out table 33 as the strip moves over the run-out table 33 in the cooling station 97. While spray nozzles atomize coolant to generate a spray, any other coolant discharge port may be employed in any embodiment in lieu of spray nozzles. In addition to generating a spray, other types of coolant discharge ports may discharge a non-atomized flow of coolant.
In the exemplary embodiment shown in
Finally, the cooled, hot rolled strip is coiled.
Further details of the twin roll caster described in relation to
The above-described embodiments of a twin roll caster and method are suitable for producing electrical steel strip less than 3 mm in thickness from a molten electrical steel melt with a superheat temperature of at least 30° C. above the liquidus temperature Tliquidus of the melt and up to a superheat temperature of Tliquidus+up to 120° C., with the melt comprising: by weight, up to 0.015% carbon, typically up to 0.0060% carbon, between 1.0% and 2.0% manganese, typically between 1.1% and 1.55% manganese, between 2.70% and 3.80% silicon, silicon killed containing less than 0.01% aluminum, up to 0.4% Cu; up to 0.3% Cr; up to 0.3% Ni; up to 0.2% Mo; up to 0.01% Ti; up to 0.005% Nb; up to 0.005% V, up to 0.3% Sb; up to 0.3% Sn; and the remainder iron, impurities and inclusions.
In this context, whilst not described above, there is a range of options for preparing the molten electrical steel melt with the superheat temperature and for delivering the superheated molten electrical steel melt to the twin roll caster that are known to a steelmaker and need not be described further here in detail. By way of example, the options include superheating the electrical steel melt to the superheat temperature in any one or more of a steelmaking furnace, a ladle or tundish or other vessel that transfers the electrical steel melt from the steelmaking furnace to the twin roll caster. One particular option is to superheat the electrical steel melt to the superheat temperature in the tundish 14 using tundish heating technology.
Typically, the hot rolling conditions are selected so that the cast strip leaves the hot rolling mill at a mill exit temperature of 800-900° C. in situations where there is a high N concentration in the hot box and at a mill exit temperature of 720-820° C. in situations where there is a low N concentration in the hot box. Typically, the mill entry temperature is selected to be 140-160° C.° C. higher than the mill exit temperature.
Typically, the cooled hot rolled cast strip is coiled at a coiler entry temperature in a range of 550 to 720° C.
In accordance with embodiments of the invention, the electrical steel strip can be optionally cold rolled to further reduce the thickness of the strip.
In accordance with embodiments of the invention, the cold rolled electrical steel strip can be annealed to develop the desired magnetic properties of the resultant non-grain oriented electrical steel strip.
Typically, the annealing step is conducted in a controlled atmosphere, such as a mixed gas of hydrogen and nitrogen.
Typically, the desired magnetic properties are as described in Chinese Standard GB/T 2521.1-2016 entitled “Cold-rolled electrical steel delivered in the fully-processed state-Part 1: Grain non-oriented steel strip (sheet)”. For example, the desired magnetic properties may be 35W250 0.35 mm Core loss P1.5/50-2.50 Wkg-1 or 35W300 0.35 mm Core loss P1.5/50-3.00 Wkg-1.
It is noted that it is often the case that there is a target C concentration for the steel in end use products. For example, for motors the target C concentration is 0.003%. The levers for obtaining a target concentration in end use products are (1) the C concentration in a steel melt and (2) the annealing step—to decarburize steel.
The applicant, via an external research organization, has carried out the experimental work summarized below to investigate hot rolling and annealing conditions for the invention.
In the experimental work, Fe-3.4 wt. % Si non-oriented electrical steel was produced using a vacuum assisted fast cooling sampling method to simulate the solidification conditions of the thin strip twin-roll casting process. The influence of rolling deformation on the magnetic properties, grain growth, and texture were analyzed.
Two non-oriented electrical steels with different carbon and sulfur contents were melted in a coreless medium frequency induction furnace under an Argon protective atmosphere.
Table 4 shows the chemical composition of these steels.
Steel 1 was produced using a high carbon and sulfur (C and S) composition, while Steel 2 was produced using a low C and S chemistry to avoid the negative influences of precipitates and phase transformations on NGO Si steel magnetic properties. To control the influence of MnS precipitates and the α to γ phase transformation during annealing, a low C and S chemical composition is commonly employed in industry for industrial NGO electrical steel production.
Charges of Steels 1 and 2 were melted in a 90 kg. coreless medium frequency induction furnace with an Ar protective atmosphere.
Samples were directly taken from the induction furnace at 100° C. superheat using a vacuum sampler, described below.
A vacuum assisted fast cooling (VA) sampling method was used to take samples that simulate the solidification conditions of an industrial twin-roll casting process. Specifically, a vacuum assisted process was used to draw liquid steel into a thin internal cavity inside a copper mold to produce a strip sample with high solidification cooling rate. The as-cast samples were 2 mm thick.
As-cast 2 mm thick samples were thermo-mechanical treated, simulating an embodiment of an industrial thin strip twin-roll casting process. Three thermo-mechanical processing routes were designed. One processing route included hot rolling with different hot rolling (HR) deformations (25% and 47%) at hot rolling temperatures in a range of 950-1000° C.° C. After hot rolling to deformations of 25% and 47%, the samples were cooled to ambient temperature via initial water spray and then furnace cooling and cold rolled to a final thickness of 0.35 mm. The other process route included cold rolling only-no hot rolling step. Higher HR deformation resulted in lower CR deformation with a combined total HR+CR reduction from 2 mm to 0.35 mm thickness. After the rolling process, samples were batch annealed at 1050° C. for various times (1, 6, 18, 24 hrs). A schematic diagram of the above-described thermo-mechanical processing route is shown in
Grain size, precipitates size distribution, magnetic properties, and texture distribution of samples with different annealing time were analyzed. The linear intercept method from ASTM E112-13 was used to measure the average grain size, a single sheet tester was used to measure the magnetic properties, automated feature analysis (AFA) on ASPEX PICA 1020 SEM was used in precipitate analysis, and electron backscatter diffraction (EBSD) analysis on Helios SEM equipment was used in analyzing texture distribution. Test samples for magnetic properties were cut along the rolling direction by 100×30 mm size. EBSD analyses were scanned along the RD-ND cross section (Rolling direction-Normal direction planes). The harmonic series expansion method was used in the orientation distribution function (ODF) calculations.
Samples were prepared metallographically, etched using water-based picric acid to reveal the dendrite structure and Nital etchant to reveal the grain structure. As noted above, the linear intercept method according to ASTM E112-13 was used for performing the secondary dendrite arm spacing (SDAS) and grain size measurement.
The solidification cooling rate of the VA as-cast samples was calculated from the measured SDAS. Based on the expected cooling rate and the chemical composition, the Suzuki's equation was used for this calculation in each case.
where S2 is the SDAS in μm, and r is the solidification cooling rate in K/s.
The optimum grain sizes were used to minimize the core losses for different test conditions. This was because it has been reported that exceptionally coarse grain sizes can lead to higher permeability, lower coercivity, and large domain size, which in turn can increase the core loss. It has been reported that the optimum grain size (GsOp) can be described as follows:
where c is an experimentally determined constant, p is the resistivity, B is the magnetic induction, t is the sample thickness, and f is the operating frequency.
Magnetic properties were measured using a single sheet tester, which was based on the ASTM A1036. The test samples were prepared by cutting processed material into 100 mm long, 30 mm wide strips in the rolling direction. Core loss was measured at 50 and 60 Hz, 1.5 and 1.0 T conditions (P1.5/50, P1.5/60, P1.0/50, P1.0/60). Magnetic induction was measured at 2500 and 5000 A/m conditions (B25, B50). The final recrystallized crystal orientations were analyzed using electron backscatter diffraction (EBSD) in a Helios SEM (30 kV, 11 nA). Scans were conducted on the RD-ND (rolling direction-normal direction planes) cross section. The harmonic series expansion method was used in the orientation distribution function (ODFs) calculations.
The dendrite structure of a VA as-cast sample is shown in
The structure of a 0% HR (100% CR) processed sample is shown in
To minimize the core loss of 0.35 mm thick 3.4 wt. % Si NGO electrical steel samples (4.70×10-7 Ω·m) at 1.5 T and 50 Hz condition, Eq. (2) was used to calculate the optimum grain size, which was determined to be ˜ 250 μm for the testing conditions. Thus, steel with a grain size close to 250 μm was expected to exhibit a decreased core loss at the 1.5 T and 50 Hz conditions.
The average grain sizes after 1, 6, 18, 24 hrs batch annealing at 1050° C. are shown in
For samples with high C and S concentrations, the final grain size was influenced by the α to γ phase transformation during the annealing process which likely retarded grain boundary migration.
As reported in a previous study, MnS and different types of oxides are the main precipitates founded in these NGO electrical steels strip samples. The recrystallization and grain growth in annealing process can be inhibited by the pinning effect of these precipitates on the grain boundaries.
It has been reported that the pinning force by small precipitates is as follows:
where γGB is the grain boundary energy, Fv is the precipitates volume fraction, and r is the precipitate average radius.
Initially, the pinning force is higher than the driving force for recrystallization and grain growth, Then, as shown in
According to Equation (1), the larger precipitates have a lower the pinning force, which reduces the grain boundary pinning effect.
The magnetic properties of different HR deformations and 1050° C. for 24 hrs batch annealed samples are shown in Table 5.
For all chemistries, the 25% HR and 47% HR samples met the magnetic properties requirement for the 35W250 NGO electrical steel at GB/T 2521.1, which is 2.50 W/kg for P1.5/50.
In general, samples processed at higher HR deformations (with same chemical composition), and the samples with lower C and S level (with the same HR deformations) showed better magnetic properties. The results are directly related to the average grain size that was discussed previously. Comparing the magnetic property results from Table 5 and the average grain size results from
Based on the inclusion analysis reported in a previous study, VA samples all had a similar inclusion size distribution. Thus, other than the grain size differences that were discussed previously, the magnetic properties were also influenced by texture, which will be discussed in the following subsection.
For the high C and S samples, increasing the HR deformation showed a decrease in the intensities of Goss orientation. There were high intensities of Cube orientation on the sample with 47% HR, while there was less on the sample with 25% HR and 0% HR. The γ-fiber had disappeared in each of the samples after annealing, which is beneficial to the magnetic properties. Compared to γ-fiber texture, the Cube and Goss textures were more ideal for magnetic properties.
For the low C and S samples, a rotated Goss orientation was observed in each of the samples. On the sample with 0% HR, there were low intensity Cube orientations and some rotated Goss orientations. This 0% HR sample was highly influenced by the presence of γ-fiber formed during the cold rolling process, while it was absent in the samples with 25% HR and 47% HR.
The recrystallization textures after final annealing were highly influenced by the deformation structure and texture during the rolling process. The strain induced boundary migration (SIBM) and subgrain growth at grain boundaries are considered to be the principal mechanisms for grain nucleation. The subgrain growth is usually observed in <111>//ND (γ-fiber) deformed grain, while the nucleation by SIBM always happens in <100>//ND (θ-fiber) deformed grain. With HR prior CR processing, the proportion of γ-fiber shear bands was decreased while the retention of {100} deformation microstructure was enhanced. Widespread shear bands within the γ-fiber deformed regions provided large number of new Goss grains.
Among samples at 25% HR deformation, the sample with high C and S showed a higher intensity of Goss orientations, while the intensity of the Cube texture was lower on the sample with low C and S. This appears to explain why at 25% HR deformation, despite the fact that the grain size of low C and S sample was coarser, it had similar core loss results compared to the high C and S sample at 1.5 T condition.
In
For the high C and S 0% HR samples (
As a comparison, the low C and S 0% HR samples (
This phenomenon can be explained by fiber texture evolution. The 0% HR samples were highly influenced by the α*-fiber and γ-fiber textures which formed before the annealing process. The formation of α*-fiber and γ-fiber textures are related to the high CR deformation and strain for the 0% HR sample. In the CR process, with the increase of deformation and strain, rotated Goss orientation will gradually rotate to {111}<110> orientation, and then form the γ-fiber texture. Finally, with the further increase of deformation and strain, α-fiber and α*-fiber textures will appear.
The main texture volume fraction of the 0% HR samples (
Theories in the literature for the thermodynamic driving force for grain growth and critical radius are helpful in explaining the texture evolution. Critical radius can be described as follows:
where γ is the boundary energy, and ΔGv is the driving force.
the velocity of grain boundary can be described as follows:
where γ is the boundary energy, R is radius, and M is grain boundary mobility (which depends on orientation of adjacent grains).
The texture evolution in
In the first step, the decrease in the volume fraction of Goss grains can be explained by the boundary energy differences between grains. High energy boundaries are more likely to appear around Goss grain. Thus, to reduce the system's total energy, it is energetically favorable to consume Goss grains early in the process of grain growth. This phenomenon can also be explained by critical radius difference. In the initial step, a large number of grains smaller than critical radius are consumed. According to Equation (4), higher boundary energy also assigns a larger critical radius to Goss grain. As reported in the literature, the critical radius of Goss grains is about 9% higher than grains with other textures. Thus, more Goss grains are consumed in the initial step. Furthermore, this phenomenon can also be explained considering the pinning effect of precipitates. At the first step, according to Equations (3) and (5), the pinning force from precipitates is still high and only grains with high energy boundaries can move.
Then, in the second step, with the precipitates coarsening at longer annealing times, the pinning force is decreased according to Equation (3). In this condition, pinning provides a mobility advantage to the survived Goss grains which are larger than the critical size. As indicated in Equation (5), this mobility advantage results in faster Goss grain growth by consuming other grains which are smaller than the critical size.
Furthermore, this main texture evolution is also related to the evolution of the α*-fiber and γ-fiber texture. With the longtime annealing, fiber textures nucleated and grew into other texture orientations. Goss orientation on BCC metals have been reported to be more likely to form from the shear deformation orientation and deformed fiber textures.
For the 25% HR samples (
The volume fraction of Cube grains at 1 hr can be considered to be influenced by the presence of Cube grains formed before the annealing process. In some cases, Cube texture components were retained after heavy cold rolling because deformed Cube grains serve as nucleation sites for new Cube grains. For the 1 to 6 hrs high C and S sample (
For the 47% HR samples (
In this case, the reversal in the fractions of Goss and Cube grains is related high volume fraction change in the final step (6 to 24 hrs) is considered to be caused by the formation of Cube grain from rotated Goss grains. It has been reported that the crystal volumes or crystallites of Cube orientation formed from the shear band of rotated Goss orientation. It has also been reported that, with the increase of strain, the Cube orientation is the most stable orientation formed from the shear bands of rotated Goss grains.
Core loss is influenced both by the grain size and texture distribution. It is difficult to separate their individual contributions to core loss. A qualitative model and equation for predicting core loss would be helpful to separate the contribution of grain size and texture distribution contributions to core loss.
For grain size influence on core loss, it has been reported that the influence by grain size can be formulated as follows [21]:
where d is the average grain size, A to C are positive constant depend on the chemical composition, precipitate size distribution, and test conditions.
According to the model, A represents the energy loss caused by eddy current and domain magnetic direction rotation in each single grain. It is influenced by chemical composition. Bd(−1) is the energy loss when the domain wall migrates inside the grain. This energy loss is influenced by grain size and precipitate size distribution. Cd(−2) represents the energy loss when eddy current passes through the grain boundaries. In practice, this energy loss is very small, because it is difficult for a domain wall to pass through a grain boundary.
For the influence by the texture distribution, Goss and Cube textures provide the main benefit to the magnetic properties. We assume that this effect is linear. The equation for the texture distribution influence is:
where D to E are positive constants, F_C and F_G are the percent of grains with Cube and Goss texture.
Combining Equation (6) and Equation (7), with the measured results, all the constants were determined using a MATLAB calculation. The resultant equation that relates core loss to grain size, texture distribution at 1.5 T 50 HZ for the Fe-3.4 wt. % Si non-oriented electrical steel is as follows:
A comparison between calculated core loss value and the measured core loss (P1.5/50) are shown in
To test the model, two additional groups of samples (Table 6) with similar average grain size were selected for evaluation. In group 1, a “Low C and 25% HR (6 hrs)” sample with larger grain size and a lower percentage of Goss and Cube grains was used. As calculated using Equation (6), the “Low C and S 25% HR (6 hrs)” sample have a higher core loss, as predicted by Equation (6). In group 2, the average grain sizes of samples are similar to each other. The “High C and 25% HR (24 hrs)” sample has a higher percentage of Goss and Cube grains, and in turn has a lower calculated core loss, as predicted. The measured results are in reasonable agreement with the calculated core loss values for these samples.
Although Equation (8) separates the influence by grain size and texture, but it is a simple semi-empirical model. For example, anomalous losses have not been considered in the calculation, and the influence by texture fraction is assumed to be linear, which still need more theoretical study. Further research is needed to make this calculation more accurate and can be used for different silicon steel compositions and service conditions.
In the above experimental work, Fe-3.4 wt. % Si non-oriented electrical steel strip samples were produced in the laboratory to simulate the solidification conditions of the thin strip twin-roll casting process. Thermo-mechanical processing routes with 0% HR, 25% HR, and 47% HR were studied on samples with high and low C and S.
The measured magnetic properties of the fully processed 25% HR and 47% HR samples all met the requirements for 35W250 NGO electrical steel in GB/T 2521.1, which is 2.50 W/kg for P1.5/50.
For the samples with the same HR deformation, the low C and S samples were observed to have a coarser average grain size after final annealing. It is likely that the high C and S samples were influenced by an α to γ phase transformation occurring during recrystallization annealing and the presence of austenite during hot deformation process. With an increase of the HR deformation, the average grain size after final annealing was also increased. This coarser grain size also led to lower core loss (P1.5.50, P1.5/60, P1.0/50, P1.0/60) and higher magnetic induction (B25, B50).
In some cases, Goss orientation seems to have a more positive effect on decreasing core loss than grain size. For example, samples at 25% HR with low C and S and high C and S show similar core loss (1.5 T condition) results despite the fact that the grain size in the former is coarser than in the latter. For the final annealed sample with high C and S, the intensities of Goss orientation were decreased with an increase in HR deformation. This observation is considered to be influenced by the decreased proportion of shear bands.
With increasing annealing times at 1050° C. from 1 to 24 hrs, the average grain size increased, and the core loss of the fully processed samples decreased. Furthermore, the increasing annealing time also had a strong influence on the evolution of grain texture. The texture evolution in the 0% HR samples was influenced by the presence of high intensity α*-fiber and γ-fiber textures formed before the annealing process from the high percentage of cold reduction used in this processing path.
Texture evolution can be divided into several different stages. In the initial stage of annealing, the fraction of grains smaller than critical radius decreased, but Goss grains were consumed more rapidly. Then, in the second stage of annealing, the fraction of Goss grains increased by consuming other grains.
It is proposed that with the coarsening of precipitates, decreased pinning provides a mobility advantage to the surviving Goss grains which are larger than the critical size for growth. Finally, with a further decrease of pinning force and an increase of grain size, normal grain growth is established. In this step, Cube grains can form from rotated Goss grains.
A simple model of core loss was developed to explain the influence of grain size and texture distribution on core loss. By comparing two groups of results, this equation successfully separated the influence of grain size and texture distribution on core loss.
The applicant has carried out additional experimental work on the following electrical steel melt: 0.0034% carbon, 1.23% manganese, 2.82% silicon, 0.0029% sulphur, 0.067% phosphorus, and 0.03% chromium.
Samples were produced by the above-described vacuum assisted fast cooling method, hot rolled with 25% and 47% reductions, cold rolled to 0.35 mm and annealed for 60 seconds at 950° C., 1000° C., and 1050° C.
The magnetic properties of the samples are in line with the results reported above.
While the principle and the mode of operation of the invention have been explained and illustrated with regard to a particular embodiment, it must be understood, however, that the invention may be practiced otherwise than as specifically explained and illustrated without departing from its spirit or scope.
| Number | Date | Country | Kind |
|---|---|---|---|
| 202210271658.0 | Mar 2022 | CN | national |
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/CN2023/082172 | 3/17/2023 | WO |
