Patent Application
20190217363
This disclosure is related to alloys and methods of developing yield strength distributions during the formation of metal parts. Formation of metal parts through procedures such as stamping, especially for complex geometries, involves cold formability which requires ductility. The alloys herein provide improved yield strength distributions after formation which reduce cracking and other associated problems in metal part formation.
Metal stamping involves a number of steps including successful forming of the stamping and achieving a targeted set of properties in the stamping. Successful forming of the stamping depends on the material properties including the global and local formability under a wide variety of stress states and strain rates. Sufficient cold formability is needed to produce the targeted geometry during the stamping operation after which a very limited material ductility remains in the stamping. This makes the stamping potentially susceptible to subsequent failure through various modes since the internal plasticity is not sufficient to develop an effective plastic zone in front of the crack tip to prevent crack propagation. Additionally, due to lack of remaining ductility, the metal stamping would also have a lack of toughness.
In metal stamping, the properties of the stamping are generally not specified as long as crack free stampings are produced. Instead, the properties of the sheet material utilized for stamping are stated. For conventional steels, properties in the stamped part are similar to that in the sheet material utilized since they undergo limited strain hardening during stamping operation and limited property changes.
As the development of steels has progressed, especially for autobody applications, it has been found that the increase in strength needed for lightweighting/gauge reduction results in the reduction in ductility/formability as shown by the “Banana plot” in
Accordingly, a need remains for the development of alloys and methods that would provide the ability to develop improved yield strength distributions during formation of metal parts, such that failure mechanisms such as cracking are eliminated or reduced, with an overall improvement in the number of successfully formed parts produced.
A method to develop yield strength distributions in a formed metal part comprising:
(a) supplying a metal alloy comprising at least 70 atomic % iron and at least four or more elements selected from Cr, Ni, Mn, Si, Cu, Al, or C, melting said alloy, cooling at a rate of <250 K/s, and solidifying to a thickness of 25.0 mm up to 500 mm;
(b) processing said alloy into sheet form with thickness from 0.5 to 10 mm wherein said sheet exhibits a yield strength of A1 (MPa), an ultimate tensile strength of B1 (MPa), a true ultimate tensile strength C1 (MPa), a total elongation D1(%);
(c) straining said sheet one or a plurality of times above said yield strength A1 at a strain rate of 100/s to 102/sec at an ambient temperature of 1° C. to 50° C. and forming a metal part having a distribution of yield strengths A2, A3, and A4, wherein:
A2=A1±100; (i)
A3>A1+100 and A3<A1+600; and (ii)
A4≥A1+600. (iii)
The detailed description below may be better understood with reference to the accompanying FIGS. which are provided for illustrative purposes and are not to be considered as limiting any aspect of this invention.
Alloys herein can be initially produced in a sheet form by different methods of continuous casting including but not limited to belt casting, thin slab casting, and thick slab casting with achievement of advanced property combinations by subsequent post-processing. After processing into a sheet form as a hot band or cold rolled sheet, which may or may not be annealed, a preferred thickness of 0.5 mm to 10.0 mm is produced.
In
The casting step can preferably be done in a wide variety of processes including ingot casting, bloom casting, continuous casting, thin slab casting, thick slab casting, belt casting etc. Preferred methods would be continuous casting in sheet form by thin slab casting or thick slab casting. To produce alloys herein in a sheet form, the casting processes can vary widely depending on specific manufacturing routes and specific targeted goals. As an example, consider thick slab casting as one process route to produce sheet product. The alloy would be cast going through a water cooled mold typically in a thickness range of 150 to 350 mm in thickness and typically processed through a roughing mill hot roller into a transfer bar slab of 25 to 150 mm in thickness and through the finishing mill into a hot band with thickness of 1.5 to 10.0 mm. Another example would be to preferably process the cast material through a thin slab casting process. In this case, after casting typically forms 25 to 150 mm in thickness by going through a water cooled mold, the newly formed slab goes directly to hot rolling without cooling down and the strip is rolled into hot band coils with typical thickness from 1.5 to 5.0 mm in thickness. Note that bloom casting would be similar to the examples above but higher thickness might be cast typically from 200 to 500 mm thick and initial breaker steps would be needed to reduce initial cast thickness to allow it to go through a hot rolling roughing mill.
Step 2 in
Preferably, sheet material from the alloys herein have a yield strength of A1 (250 MPa to 750 MPa), a tensile strength of B1 (700 MPa to 1750 MPa), a true ultimate tensile strength of C1 (1100 MPa to 2300 MPa), and exhibits a total elongation D1 (10% to 80%). While engineering stress is determined as the applied load divided by the original cross-sectional area of the specimen gauge, true stress corresponds to the applied load divided by the actual cross-sectional area (the changing area with respect to time) of the specimen at that load. True stress is the stress determined by the instantaneous load acting on the instantaneous cross-sectional area. True ultimate tensile strength (C1) is related to ultimate tensile strength (B1) and can be calculated from the test data for each alloy herein using Eq.1. Engineering strain is determined as the change in length divided by the original length. Calculated true ultimate tensile strength values vary from 1165 to 2237 MPa:
True Ultimate Tensile Strength(C1)=Ultimate Tensile Strength*(1+Engineering Strain) (Eq.1)
True strain at fracture corresponding to total elongation of each specimen was calculated by Eq.2. True strain at fracture was found to vary from 15.7 to 58.1%.
True Strain at Fracture=In(1+Engineering Strain) (Eq.2)
Depending on alloy chemistry, the magnetic phase volume percent generally varies from 0.2 to 45.0 Fe % for hot band or cold rolled and annealed sheet. Such magnetic phase volume is then increased as discussed more fully below.
Straining of the alloy sheet above its yield strength, which may preferably occur via stamping of the sheet from said alloy with the indicated influence on yield strength occurring during the stamping operation, is shown by Step 3 in
The alloys herein undergoing what is illustrated in
Areas of the microstructure in the initial sheet from the alloys herein with relatively stable austenite retain the austenitic nature but deform through primarily dislocation mechanisms supporting material ductility and formability during stamping and forming Microconstituent 2 in the final microstructure after deformation. Microconstituent 2 itself contains two components which are micron sized stable austenite particles, typically 1.0 to 10.0 microns in size (longest linear dimension) and nanoprecipitates typically 2 to 100 nm in size (longest linear dimension). Nanoprecipitates in either Microconstituent 1 or 2 can be directly observed through TEM microscopy and are observed to exhibit a spherical, elliptical, or rectangular shape in the size range indicated. To further identify, selected area diffraction in the TEM on the precipitates can be done to show that they have different structures (i.e. not FCC austenite or BCC Ferrite) than the matrix phases (i.e. austenite which is FCC or alpha ferrite which is BCC). Accumulation of dislocations within micron-sized austenite grains results in dislocation cell block boundaries, and dislocation cell formation leading to material strengthening. Additionally, as noted, nanoprecipitates with a size from 2 to 100 nm are present in both Microconstituents 1 and 2 also contributing to material strengthening.
The resulting volume fraction of Microconstituent 1 and Microconstituent 2 in the localized areas of the stamping, i.e., the final formed part, depends on alloy chemistry, the level of straining at particular location, and the level of strain hardening which occurs during the single or multistage stamping operation. Note that the microstructure and resulting properties will change in the stamped part from the starting sheet/blank depending on the local level of straining. Typically, as low as 1 volume percent and as high as 85 volume percent of the alloy structure after stamping will exist as the ferrite containing Microconstituent 1 with the remaining regions representing Microconstituent 2. Thus, Microconstituent 1 can be in all individual volume percent values from 0.5 to 85.0 in 0.1% increments (i.e. 0.5%, 0.6%, 0.7%, . . . up to 85.0%) while Microconstituent 2 can be in volume percent values from 99.5 to 15 in 0.1% increments (i.e. 99.5%, 99.4%, 99.3% . . . down to 15.0%). The volume percent of nanoprecipitates which occur in both microconstituents is anticipated to be 0.1 to 10%. While the magnetic properties of these nanoprecipitates are difficult to individually measure, it is contemplated that they are non-magnetic.
As ferrite is magnetic (i.e. ferromagnetic), and austenite is non-magnetic (i.e. paramagnetic), the volume fraction of the magnetic phases present provides a convenient method to evaluate the relative presence of Microconstituent 1. The magnetic phases volume percent is abbreviated herein as Fe %, which should be understood as a reference to the presence of ferrite and any other components in the alloy that identifies a magnetic response such as alpha-martensite. Note that the alpha-ferrite and alpha-martensite have similar magnetic responses and cannot be distinguished separately by the Feritscope so both will be identified as ferrite. Magnetic phase volume percent herein is conveniently measured by a Feritscope. The Feritscope uses the magnetic induction method with a probe placed directly on the sheet sample and provides a direct reading of the total magnetic phases volume percent (Fe %). After cold deformation, the volume fraction of Microconstituent 1 is estimated using the measured Fe % value which can include alpha-ferrite and/or alpha-martensite. Microconstituent 2 which is nonmagnetic and cannot be measured by the Feritscope, would then be considered the remaining constituent.
While the multiple mechanistic components of the NR&S mechanism described above support deformation of the sheet during its forming into targeted shape, sheet material from alloys herein undergoes a substantial strain hardening/strengthening which results in the presence of distributions (i), (ii), and (iii) in the formed parts provided in
Forming of the alloys herein can be done by various methods including but not limited to forming in single and/or progressive dies and with one stage or multiple stages up to 25 towards targeted final form using a combination of techniques, without external heating, including but not limited to stamping, roll forming, metal drawing, and hydroforming. In connection with such procedures the deformation that exceeds the yield strength may include hole expansion, hole extrusion drawing, bending and/or stretching. Common to all of these processing techniques is the introduction of a one or a plurality of deformations (introduction of strain) such that yield strength is exceeded with the result that all of the above referenced distribution of yield strengths are achieved in the formed part. The final formed part applications include but are not limited to automotive industry (a vehicular frame, vehicular chassis, or vehicular panel), and/or railroad industry (a storage tank, freight car, or railway tank car).
The chemical composition of the alloys herein is shown in Table 1 which provides the preferred atomic ratios utilized.
With regards to the above, and as can be further seen from Table 1, preferably, when Fe is present at a level of greater than 70 at. %, and one then selects the four or more elements from the indicated seven (7) elements, or selects five or more elements, or selects six or more elements or selects all seven elements to provide a formulation of elements that totals 100 atomic percent. The preferred levels of the elements, if selected, may fall in the following ranges (at. %): Cr (0.2 to 8.7), Ni (0.3 to 12.5), Mn (0.6 to 16.9), A1 (0.4 to 5.2), Si (0.7 to 6.3), Cu (0.2 to 2.7), and C (0.3 to 3.7). Accordingly, it can be appreciated that if four (4) elements are selected, two of the six elements are not selected and may be excluded. If five (5) elements are selected, one of the elements of the six can be excluded. Moreover, a particularly preferred level of Fe is in the range of 70.0 to 85.0 at. %. The level of impurities of other elements is in the range of 0 to 5000 ppm. Accordingly, if there is 5000 ppm of an element other than the selected elements identified, the level of such selected elements may then in combination be present at a lower level to account for the 5000 ppm impurity, such that the total of all elements present (selected elements and impurities) is 100 atomic percent.
The alloys herein were processed into a laboratory sheet by processing of laboratory slabs. Laboratory alloy processing is developed to mimic closely the commercial sheet production by continuous casting and include hot rolling and cold rolling. Annealing might be applied depending on targeted properties. Produced sheet can be used in hot rolled (hot band), cold rolled, annealed, or partially annealed states.
Alloys were weighed out into 3,000 to 3,400 gram charges according to the atomic ratios in Table 1 using commercially available ferroadditive powders and a base steel feedstock with known chemistry. Impurities can be present at various levels depending on the feedstock used. Impurity elements would commonly include the following elements; Co, N, P, Ti, Mo, W, Ga, Ge, Sb, Nb, Zr, O, Sn, Ca, B and S which if present would be in the range from 0 to 5000 ppm (parts per million) (0 to 0.5 wt %) at the expense of the desired elements noted above. Preferably, the level of impurities is controlled to fall in the range of 0 to 3000 ppm (0.3 wt %).
Charges were loaded into a zirconia coated silica crucible which was placed into an Indutherm VTC800V vacuum tilt casting machine. The machine then evacuated the casting and melting chambers and flushed with argon to atmospheric pressure twice prior to casting to prevent oxidation of the melt. The melt was heated with a 14 kHz RF induction coil until fully molten, approximately from 5 to 7 minutes depending on the alloy composition and charge mass. After the last solids were observed to melt it was allowed to heat for an additional 30 to 45 seconds to provide superheat and ensure melt homogeneity. The casting machine then evacuated the chamber and tilted the crucible and poured the melt into a water cooled copper die. The melt was allowed to cool under vacuum for 200 seconds before the chamber was filled with argon to atmospheric pressure.
A sample of between 50 and 150 mg from each alloy herein was taken in the as-cast condition. This sample was heated to an initial ramp temperature between 900° C. and 1300° C. depending on alloy chemistry, at a rate of 40° C./min. Temperature was then increased at 10° C./min to a max temperature between 1425° C. and 1510° C. depending on alloy chemistry. Once this maximum temperature was achieved, the sample was cooled at a rate of 10° C./min back to the initial ramp temperature before being reheated at 10° C./min to the maximum temperature. Differential Scanning Calorimetry (DSC) measurements were taken using a Netzsch Pegasus 404 DSC through all four stages of the experiment, and this data was used to determine the solidus and liquidus temperatures of each alloy, which are in a range from 1294 to 1498° C. (Table 2). Depending on the alloys chemistry, liquidus-solidus gap varies from 26 to 138° C. Thermal analysis provides information on maximum temperature for the following hot rolling processes that varies depending on alloy chemistry.
The density of the alloys herein was measured on samples from hot rolled material using the Archimedes method in a specially constructed balance allowing weighing in both air and distilled water. The density of each alloy is tabulated in Table 3 and was found to be in the range from 7.48 to 8.01 g/cm3. The accuracy of this technique is ±0.01 g/cm3.
The alloys herein were preferably processed into a laboratory hot band by hot rolling of laboratory slabs at high temperatures. Laboratory alloy processing is developed to simulate the hot band production from slabs produced by continuous casting. Industrial hot rolling is performed by heating a slab in a tunnel furnace to a target temperature, then passing it through either a reversing mill or a multi-stand mill or a combination of both to reach the target gauge. During rolling on either mill type, the temperature of the slab is steadily decreasing due to heat loss to the air and to the work rolls so the final hot band is formed at a reduced temperature. This is simulated in the laboratory by heating in a tunnel furnace to between 1100° C. and 1250° C., then hot rolling. The laboratory mill is slower than industrial mills causing greater loss of heat during each hot rolling pass so the slab is reheated for 4 minutes between passes to reduce the drop in temperature, the final temperature at target gauge when exiting the laboratory mill commonly is in the range from 800° C. to 1000° C., depending on furnace temperature and final thickness.
Prior to hot rolling, laboratory slabs were preheated in a Lucifer EHS3GT-B18 furnace. The furnace set point varies between 1100° C. to 1250° C., depending on alloy melting point and point in the hot rolling process, with the initial temperatures set higher to facilitate higher reductions, and later temperatures set lower to minimize surface oxidation on the hot band. The slabs were allowed to soak for 40 minutes prior to hot rolling to ensure they reach the target temperature and then pushed out of the tunnel furnace into a Fenn Model 061 2 high rolling mill. The 50 mm casts were hot rolled for 5 to 10 passes though the mill before being allowed to air cool. Final thickness ranges after hot rolling are preferably from 1.8 mm to 4.0 mm with variable reduction per pass ranging from 20% to 50%.
Hot band material was media blasted prior to cold rolling to remove surface oxides which could become embedded during the rolling process. The resultant cleaned sheet material was rolled using a Fenn Model 061 2 high rolling mill down to 1.2 mm thickness. Reductions before annealing ranged from 10% to 40%.
Once the final gauge thickness of 1.2 mm was reached, tensile samples were cut from the laboratory sheet by wire-EDM. The samples were annealed under conditions intended to simulate the thermal exposure expected during an industrial continuous annealing process representing final treatment of sheet material in Step 2 in
Tensile properties were measured on an Instron mechanical testing frame (Model 3369), utilizing Instron's Bluehill control and analysis software. All tests were run at ambient temperature in displacement control at a constant displacement rate of 0.036 mm/s. Tensile properties of 1.2 mm thick sheet from alloys herein after annealing at 850° C. for 10 minutes are listed in Table 4. The ultimate tensile strength values of the annealed sheet from alloys herein is in a range from 717 to 1683 MPa with total elongation recorded in the range from 17.1 to 78.9%. The 0.2% proof stress varies from 273 to 652 MPa, 0.5% proof stress varies from 295 to 704 MPa, and 1.0% proof stress varies from 310 to 831 MPa. True ultimate tensile strength calculated from the data for each alloy herein, which varies from 1188 to 2237 MPa with true strain at fracture from 15.7 to 58.1%.
As the exact point of yielding is difficult to determine, a range of proof tests were employed at 0.2%, 0.5% and 1.0% proof stresses. That is, the exact point where the deformation changes from elastic to plastic is complicated by the unique deformation mechanisms of the alloys herein, resulting in a curvature of the initial portion of the stress strain curve. The 0.2%, 0.5%, and 1.0% represent offset strains whereby at these strain levels, a parallel line is drawn to the stress strain curve and the resulting points of intersection is defined at the proof stress at the identified offsets respectively. At the 0.5% proof stress, more consistent and representative values are obtained so that the yield strength herein (A1, A2, A3, and A4) will be defined at the 0.5% proof stress. In
Incremental tensile testing was done on an Instron mechanical testing frame (Model 5984), utilizing Instron's Bluehill control and analysis software. All tests were run at ambient temperature in displacement control. Samples were tested at a displacement rate of 0.025 mm/s during initial loading to 2% strain and 0.125 mm/s for the remaining duration of the test. Due to the variation in sample length during testing effective strain rates generally ranged from ˜104/s to 10−3/s for the initial loading and after initial loading strain rates ranged from ˜10−3/s to ˜10−2/s. It should be noted that while the incremental tensile testing was done at these indicated strain rates, such incremental tensile testing is considered to support the yield strength distributions (i.e. values of A2, A3 and A4) and increase in magnetic phase volume for the alloys herein at the recited at strain rates (100/sec to 102/sec). See, e.g., Case Example #3 (stamping) and Table 13 (incremental tensile testing).
A control specimen from the same area of the sheet was tested up to failure from each alloy to evaluate initial sheet properties of the specific sample set used for incremental testing and the results are listed in Table 5 for each alloy herein. The ultimate tensile strength values are in a range from 745 to 1573 MPa with total elongation recorded in the range from 13.3 to 77.1%. The 0.5% proof stress or yield strength (A1) varies from 287 to 668 MPa and true ultimate tensile strength is in a range from 1175 to 2059 MPa. After each control specimen was tested, a new duplicate sample of each alloy was then strained approximately 5%, and then unloaded. The specimen dimensions were measured as well as the magnetic phases volume percent (Fe %) prior to the next increment of testing. Magnetic phases volume percent (Fe %) was measured by Fisher Feritscope.
Incremental test data for each alloy herein is listed in Table 6 through Table 39 and illustrated in
As can be seen from the above, the magnetic phases volume of the sheet is increased when exposed to one or a plurality of strains above the yield strength of the sheet. That is, for a given sheet material, having a magnetic phases volume that falls in the range of 0.2 Fe % to 45.0 Fe %, such value is observed to increase and the metal part that is formed indicates a magnetic phases volume that falls in the range of 0.5 Fe % to 85.0 Fe %. For example, for Alloy 1 that indicates in the sheet an initial magnetic phase volume of 0.7 Fe %, after nine (9) strains above the yield strength of the sheet indicates a magnetic phases volume of 67.5 Fe %. Alloy 2 sheet is initially 22.0 Fe % and after six (6) strains above the yield strength of the sheet indicates a magnetic phases volume of 67.1 Fe %. For each alloy provided herein, the properties including yield change as a function of applied strain in sheet form. In stamping operations, a wide range of strains rather than a singular strain is applied over the stamped part. This results in a wide range of localized strain and resulting properties in the stamped part which may include the entire range of properties found for example by the separately applied strains in the sequential cycles for each alloy.
These results show the key structural changes which lead to strengthening during cold deformation with commensurate increases in both yield and tensile strength during the deformation process.
Laboratory slabs with thickness of 50 mm were cast from Alloy 7 and Alloy 8 according to the atomic ratios in Table 1 that were then laboratory processed by hot rolling, cold rolling and annealing at 850° C. for 10 min as described in the Main Body section of the current application. Microstructure of the alloys in a form of processed sheet with 1.2 mm thickness after annealing corresponding to a condition of the sheet in annealed coils at commercial production was examined by SEM and TEM.
To prepare TEM specimens for a structural analysis of the annealed sheet from the alloys before deformation, the samples were first cut with EDM, and then thinned by grinding with pads of reduced grit size every time. Further thinning to make foils of 60 to 70 μm thickness was done by polishing with 9 μm, 3 μm, and 1 μm diamond suspension solution, respectively. Discs of 3 mm in diameter were punched from the foils and the final polishing was fulfilled with electropolishing using a twin-jet polisher. The chemical solution used was a 30% nitric acid mixed in methanol base. In case of insufficient thin area for TEM observation, the TEM specimens may be ion-milled using a Gatan Precision Ion Polishing System (PIPS). The ion-milling usually is done at 4.5 keV, and the inclination angle is reduced from 4° to 2° to open up the thin area. To analyze structure in the alloys after deformation, TEM samples were cut from the gauge section of the tensile specimens close to the fracture and prepared in the similar manner. The TEM studies were done using a JEOL 2100 high-resolution microscope operated at 200 kV. The TEM specimens were studied by SEM. Microstructures were examined by SEM using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc.
The microstructure in the Alloy 7 sheet before deformation is shown by SEM and TEM micrographs in
During tensile testing to failure, the initial structure undergoes NR&S leading to formation of the final structure, which is demonstrated for Alloy 7 and Alloy 8 by SEM and TEM micrographs in
Further details of the microstructure after deformation highlighting microstructural features of each microconstituent were obtained from structural analysis of the gauge section of the tensile specimen from Alloy 8 sheet after testing to failure. A TEM bright-field micrograph corresponding to Microconstituent 1 in the sheet material is shown in
A TEM bright-field micrograph corresponding to Microconstituent 2 in the sheet material is shown in
This Case Example demonstrates that the microstructure of the alloys herein undergo transformation during cold deformation through the NR&S mechanism leading to formation of the microstructure with distinct microconstituents resulting in material strengthening.
Sheet blanks from Alloy 8 with a thickness of 1.4 mm were used for stamping trial of a B-pillar at a commercial stamping facility with stamping speed estimated at 290 mm/s. Using an existing die, Alloy 8 sheet blanks were stamped into B-pillars. Non-destructive analysis of the B-pillar was done by Feritscope measurements of the local magnetic phases volume percent in different areas.
Feritscope measurements provide an indication of the structural changes occurring during deformation from stamping. As shown previously, in the Alloy 8 sheet, the initial sheet microstructure changes from non-magnetic (i.e. paramagnetic) to magnetic (i.e. ferromagnetic) microstructure during cold deformation through the NR&S mechanism. The baseline for the sheet in Feritscope measurements before stamping was <1 Fe %. Increase in the volume fraction of Microconstituent 1 results in higher Fe % measured. Feritscope measurements with ˜20 mm grid pattern were taken from two stamped B-pillars including one which underwent 4 out of 5 stamping hits and one which underwent 5 out of 5 stamping hits. The 5th hit is mainly a flanging operation so little structural or property change was expected in the B-pillar. The examples of the grid pattern on the different areas of the B-pillars are shown in
The summary of Fe % measurements of the B-pillar which underwent a total of 4 stamping hits is shown in
This Case Example demonstrates significant changes in magnetic phases volume percent in the stamping as compared to initial sheet. These changes correspond to microstructural transformation the unique NR&S mechanisms leading to sheet material strengthening as it deforms.
A sheet blank from Alloy 8 with a thickness of 1.4 mm were used for a stamping trial of a B-pillar at a commercial stamping facility with stamping speed estimated at 290 mm/s. Alloy sheet properties before stamping are shown in Table 40. Using an existing die, Alloy 8 sheet blanks were stamped into B-pillars.
For destructive analysis, tensile specimens were cut along the entire length of the B-pillar. The view of the B-pillar before and after specimen cutting is shown in
In total, 213 tensile specimens cut from the B-pillar were tested. Rockwell C hardness and Feritscope measurements were taken from each tensile specimen. Tensile property data for selected specimens are listed in Table 41. Examples of the stress—strain curves for specimens cut from the B-pillar with various levels of magnetic phases volume percent (Fe %) are presented in
The measured tensile properties were correlated to structural changes during stamping evaluated from direct Feritscope measurements on the grip sections of the tensile specimens after cutting from the B-pillar prior to testing. Correlation between the measured Fe % and tensile properties is shown in
Non-destructive analysis showed the maximum value of 31 Fe % in highly bent areas of the B-pillar that cannot be used for tensile specimen cutting. However, the current correlations based on 213 data points and shown in
This Case Example demonstrates a dramatic increase in both yield and tensile strength in the stamped part as a result of material cold deformation during stamping operation. Cold deformation activates NR&S mechanism in the alloys herein leading to material strengthening. The 213 tensile specimens measured over the surface of the stamped part illustrate the resulting change in properties resulting from the localized changes found in the stamped part. While the stamped part was not deformed until failure, the range of properties found in the stamped part, are similar to the range of tensile properties (prior to failure) found for the same alloy from incremental tensile testing as previously provided in Table 13.
A sheet blank from Alloy 8 with a thickness of 1.4 mm was used for stamping trial of a B-pillar at a commercial stamping facility. Detailed TEM analysis was done on the samples cut from different locations of the stamped part to demonstrate the structural response to the deformation during stamping.
To prepare TEM specimens for a structural analysis, the samples were first cut with EDM from the areas of interest, and then thinned by grinding with pads of reduced grit size every time. Further thinning to make foils of 60 to 70 am thickness was done by polishing with 9 am, 3 am, and 1 am diamond suspension solution, respectively. Discs of 3 mm in diameter were punched from the foils and the final polishing was fulfilled with electropolishing using a twin-jet polisher. The chemical solution used was a 30% nitric acid mixed in methanol base. In case of insufficient thin area for TEM observation, the TEM specimens may be ion-milled using a Gatan Precision Ion Polishing System (PIPS). The ion-milling usually is done at 4.5 keV, and the inclination angle is reduced from 4° to 2° to open up the thin area. To analyze structure in the alloys after deformation, TEM samples were cut from the gauge section of the tensile specimens close to the fracture and prepared in the similar manner. The TEM studies were done using a JEOL 2100 high-resolution microscope operated at 200 kV. The TEM specimens were studied by SEM. Microstructures were examined by SEM using an EVO-MA10 scanning electron microscope manufactured by Carl Zeiss SMT Inc.
In
TEM analysis of the microstructure was also done for the gauge section of the corresponding samples tested in tension from the same three locations. Bright-field TEM images of the microstructure after tensile testing are provided in
This Case Example demonstrates microstructural changes of the alloy herein during stamping operations corresponding to localized increases in magnetic phases volume percent consistent with the localized Feritscope measurements. These specific microstructural changes are consistent with the activation of the identified NR&S mechanism and conclusively show the material strengthening occurring in the stamping.
Nine specimens with reduced size were cut from the same Alloy 8 sheet that used for stamping trial of the B-pillar and used for incremental testing. Alloy sheet properties are shown in Table 40. Incremental tensile testing was done on an Instron mechanical testing frame (Model 3369), utilizing Instron's Bluehill control and analysis software. All tests were run at ambient temperature in displacement control. The specimen dimensions were measured as well as the magnetic phases volume percent (Fe %) prior to next increment of testing. Magnetic phases volume percent (Fe %) was measured by Fisher Feritscope.
Yield strength data collected from incremental testing of Alloy 8 sheet as well as that from tensile testing of specimens cut from the B-pillar during destructive analysis were correlated with magnetic phases volume percent (Fe %). 0.2 and 0.5% proof stress as a function of the Fe % is presented in
This Case Example shows good correlation between the changes in yield strength in incremental tensile specimens and that in specimens tested during destructive analysis of the B-pillar as a function of magnetic phases volume percent. Cold deformation results in structural transformation detected by an increase in Fe % leading to strengthening of alloys herein and to an increase in strength characteristic values.
Laboratory slabs with thickness of 50 mm were cast from Alloy 8 according to the atomic ratios in Table 1. The slabs were then processed by a mixture of hot and cold rolling to achieve the targeted sheet thickness of 0.5, 1.3, 3.0 and 7.1 mm. The thickest material was hot rolled only, while all other conditions were cold rolled to achieve the targeted thickness. After cold rolling the samples were wrapped in stainless steel foil to minimize oxidation and placed into an 850° C. furnace for 10 minutes then removed and allowed to cool in air. The details of each sheet processing are listed in Table 42.
Incremental tensile testing was done on an Instron mechanical testing frame (Model 5984), utilizing Instron's Bluehill control and analysis software. All tests were run at ambient temperature in displacement control. Samples were tested at a displacement rate of 0.025 mm/s during initial loading to 2% strain and 0.125 mm/s for the remaining duration of the test.
Each specimen was strained approximately 5%, and then unloaded. The specimen dimensions were measured as well as the magnetic phases volume percent (Fe %) prior to the next increment of testing. Magnetic phases volume percent (Fe %) was measured by Fisher Feritscope. Control specimen from the same sheet from each alloy was tested up to failure to evaluate initial sheet properties that are listed in Table 43 for sheet samples at each thickness.
Incremental test data for samples with each thickness herein is listed in Table 44 through Table 47. Incremental stress-strain curves along with engineering stress-strain curves and true stress-true strain curves are shown for Alloy 8 sheet with each thickness in
The incremental testing results also show an extensive increase in yield strength with increasing accumulated strain. The difference in yield strength values between first and last cycle of testing varies from 1112 to 1332 MPa confirming a significant material strengthening. Note that while this example highlights individual strains applied to the sheet in specific steps, the range of properties demonstrated are deemed simultaneously possible in a stamped part made from the alloys herein.
This Case Example demonstrates that the strengthening and strain hardening mechanisms occur in the sheet material with a range of thicknesses from 0.5 to 7.1 mm.
Sheet material from commercial steel grades of TRIP 780 and DP980 was used for incremental testing. TRIP 780 has the following chemistry (at %); 97.93 Fe, 1.71 Mn, 0.15 Cr, 0.12 Si, 0.05 C, and 0.04 Cu. DP980 has the following chemistry (at %); 96.86 Fe, 2.34 Mn, 0.42 C, and 0.38 Si. Incremental tensile testing was done on an Instron mechanical testing frame (Model 5984), utilizing Instron's Bluehill control and analysis software. All tests were run at ambient temperature in displacement control. Samples were tested at a displacement rate of 0.025 mm/s during initial loading to 2% strain and 0.125 mm/s for the remaining duration of the test.
Each specimen was strained approximately 5%, and then unloaded. The specimen dimensions were measured prior to the next increment of testing. Control specimen from the same sheet from each steel grade was tested up to failure to evaluate initial sheet properties that are listed in Table 48 for each grade. Magnetic phases volume percent (Fe %) in initial sheet and in the specimen gauge after testing was measured by Fisher Feritscope that is listed in Table 49. The measurement showed no changes in Fe % before and after testing the specimens from TRIP 780 and DP980.
Incremental test data for each steel grade is listed in Table 50 and Table 51 and illustrated in
This Case Example demonstrates less degree of strain hardening in commercial steel grades during deformation with no changes in magnetic phases volume percent (0 to 0.1 Fe % difference before and after deformation).
This application claims the benefit of U.S. Provisional Application 62/618,356 filed Jan. 17, 2018 which is fully incorporated herein by reference.
| Number | Date | Country | |
|---|---|---|---|
| 62618356 | Jan 2018 | US |
