Patent Application
20240301528
The present invention relates to a method for manufacturing through cold forming, in particular via cold heading, assembly parts, such as screws, bolts, etc., that the automotive industry commonly uses for chassis or wheel hub components of vehicles.
As is known, the automotive industry continually aims to decrease the vehicle weight; which can be done by modifying their safety assembly. The weight reduction requires an increasingly reduced size of the parts. These parts, however, remain subject to the same mechanical stresses, and must therefore have increasingly high mechanical properties, in particular tensile strength.
WO2016/158470 discloses hardening steel excellent in machinability before aging treatment and excellent in fatigue characteristics, toughness, and low cycle fatigue characteristics after aging treatment, that is, age hardening steel containing predetermined amounts of C, Si, Mn, S, Cr, Al, V, Nb, Ca, and REM, limiting contents of P, Ti, and N to predetermined amounts or less, having a balance of Fe and impurities, having an area ratio of bainite structures of 70% or more. But the steel of WO2016/158470 lacks hydrogen embrittlement.
WO2011/124851 discloses a mechanical steel part in steel with high characteristics, characterized in that its composition, comprising in weight percentages, is 0.05%≤C≤0.25%; 1.2%≤Mn≤2%; 1%≤Cr≤2.5%; wherein the contents of C, Mn and Cr are such that (830-270C %-90 Mn %-70Cr %)≤560; 0<Si≤1.55; 0<Ni≤1%; 0<Mo≤0.5%; 0<Cu≤1%; 0<V≤0.3%; 0<Al≤0.1%; 0<B≤0.005%; 0<Ti≤0.03%; 0<Nb≤0.06%; 0<S≤0.1%; 0<Ca≤0.006%; 0<Te≤0.03%; 0<Se≤0.05%; 0<Bi≤0.05%; 0<Pb≤0.1%; the remainder of the steel part being iron and impurities resulting from processing, and wherein the in that its structure of the steel is bainitic and contains no more than a total of 20% of martensite and/or pro-eutectoid ferrite and/or pearlite. But the steel of WO2011/124851 does not demonstrate hydrogen embrittlement as well as a reduction area of 58% or more.
However, it is desirable to further improve the resistance to hydrogen embrittlement of the parts.
The present invention provides a steel part which may be used as an assembly part for a motor vehicle, and which has improved resistance to hydrogen embrittlement while simultaneously having:
In a preferred embodiment, the steel part shows a hardness from 360 Hv to 405 Hv.
The invention will be better understood upon reading the description that follows, given solely by way of example.
In the entire patent application, the contents are indicated in weight % (wt %).
The steel part according to the invention has a composition comprising, by weight:
Carbon is present in the steel of the present invention from 0.05% to 0.15%. Carbon imparts strength to the steel by solid solution strengthening and carbon is gammagenous hence delays the formation of Ferrite. Carbon is the element that impacts the formation of Cementite-free Lath-Like Bainite. A minimum of 0.05% of carbon is required to reach a tensile strength of 1100 MPa but if carbon is present above 0.15%, carbon deteriorates ductility as well as machinability of the final product due to the formation of cementite. The carbon content is advantageously in the range 0.08% to 0.14% to obtain simultaneously high strength and high ductility and more preferably from 0.09% to 0.14%.
Silicon is present in the steel of the present invention from 0.01% to 1%. Silicon imparts the steel of the present invention with strength through solid solution strengthening. In particular, at the above-mentioned contents, the silicon has the effect of hardening the bainite microstructure through solid solution hardening. Silicon reduces the formation of cementite nucleation as silicon hinders precipitation and diffusion-controlled growth of carbides by forming a Si-enriched layer around precipitate nuclei. Therefore, resulting the cementite-free lath-like bainite. Silicon also acts as a deoxidizer. A minimum of 0.01% of silicon is required to impart strength to the steel of the present invention. An amount of more than 1% raises the activity of carbon in austenite promoting its transformation into pro-eutectoid ferrite, which can deteriorate the strength, and also resulting retardation for formation of bainite under continuous cooling thereby too much retained austenite at the end of cooling. The preferred limit for Silicon from 0.01 to 0.9% and more preferably from 0.01% to 0.5%
Manganese is added in the present steel from 1.2% to 2%. Manganese provides hardenability to the steel. It allows to decrease the critical cooling rate for which a bainitic transformation can be obtained in continuous cooling without any prior transformation and the manganese lowers the bainite start temperature of the steel, and therefore results in a refinement of the bainitic structure to form lath bainite and thus increases the mechanical properties of the part. A minimum content of 1.2% by weight is necessary to obtain the desired bainitic microstructure. But above 2%, manganese has a negative effect on the steel of the present invention as retained austenite can transform into MA islands or fresh martensite and these phases are detrimental for the properties. In addition, manganese forms sulphides such as MnS. These sulphides can increase machinability if the shape and distribution are well controlled. If not, they might have a very detrimental effect on Elongation. The preferred limit of manganese is from 1.3% to 1.9% and more preferably from 1.4% to 1.9%.
Chromium is present from 0.1% to 2% in the steel of the present invention. Chromium is an indispensable element in order to produce bainitic structure, especially lath bainite and impart Elongation and ductility to the steel of the present invention. Addition of Chromium promotes homogeneous and finer bainitic microstructure during the temperature range between Bs and room temperature. A minimum content of 0.1% of Chromium is required to produce the targeted bainitic microstructure and chromium also slows down the softening during the tempering treatment, allowing higher holding temperatures which favors degassing but also the formation of carbides that trap hydrogen. But the presence of Chromium content of 2% or more excessively increases the hardness of the steel, it makes it difficult to form it by cold forming, and in particular cold heading. It is advantageous to have Chromium from 0.2% to 1.6% and more preferably from 0.3% to 1.4%.
The steel aluminum is at a content from 0.001% to 0.1 wt %. Aluminum is a deoxidizer of the steel in the liquid state. It then contributes, in the form of nitrides, to controlling austenitic grain coarsening during hot rolling. On the other hand, present in too large an amount, it may lead to a coarsening of aluminate type inclusions in the steel which may prove damaging to the properties of the steel, especially its toughness. In particular, the aluminum content may be comprised at a content from 0.001 to 0.09 wt %.
In the steel according to the invention, the nitrogen content is comprised from 0% to 0.01 wt %. Nitrogen traps boron via the formation of boron nitrides, which makes the role of this element in the hardenability of the steel ineffective. Therefore, in the steel according to the invention, the nitrogen content is limited to 0.01 wt %. Nevertheless, added in small amounts, it makes it possible, via the formation in particular of titanium nitrides (TiN) and aluminum nitrides (AlN), to avoid excessive austenitic grain coarsening during heat treatments undergone by the steel. Similarly, it also allows, in this case, the formation of carbonitride precipitates that will contribute toward the trapping of hydrogen. Therefore, in the steel according to the invention, the nitrogen content is greater than or equal to 0.003 wt %.
The steel according to the invention comprises at most 0.015 wt % of phosphorus and at most 0.015 wt % of sulfur. The effect of phosphorus and sulfur are particularly harmful in the steels according to the invention, for several reasons. Indeed, since these elements are poisons for hydrogen recombination, they contribute to a higher concentration of atomic hydrogen capable of penetrating into the material, therefore to an increased risk of delayed fracture of the part in use. Moreover, by segregating at the grain boundaries, the phosphorus and the sulfur reduce the cohesion thereof. Their content must therefore be kept very low. For this purpose, measures must be taken to ensure that the steel is dephosphorized and desulfurized during its melting in the liquid state.
The steel may optionally contains from 0.01 to 1 wt % of nickel. This element provides an increase in the strength of the steel and has beneficial effects on the resistance to brittle fracture. It also improves, in a known manner, the corrosion resistance of the steel.
Boron is an optional element and can be present in the alloy at contents from 0.0003 to 0.01 wt %. By segregating at the prior austenitic grain boundaries, boron, even at very low contents, strengthens the grains boundaries, and makes it possible to increase the resistance to hydrogen-induced delayed fracture. The boron increases the cohesion of the grain boundary via its intrinsic effect, but also by making phosphorus segregation more difficult at these grain boundaries. The boron further strongly increases the hardenability of the steel and thus makes it possible to limit the carbon content needed to obtain the desired bainitic microstructure. Finally, boron acts in synergy with molybdenum and niobium, thus increasing the effectiveness of these elements and their own influence that their respective contents permit. An excess of boron (above 0.01 wt %) would however lead to the formation of brittle iron boro-carbides.
The molybdenum is an optional element and is comprised from 0.003 to 1 wt %. Molybdenum interacts strongly with phosphorus and limits the damaging effect of the phosphorus by limiting its segregation at the prior austenite grain boundaries. Furthermore, it displays a marked carbide-forming behavior. For given mechanical properties, it allows higher holding tempering temperature, which, as a result, favor the development of carbides that will be hydrogen traps. It is therefore an element that increases the resistance to delayed fracture.
Titanium is an optional element and present in the alloy at contents comprised from 0.01 to 0.04 wt %. Titanium is added to the liquid steel in order to increase the hardness of the material. Here, within the ranges indicated, it also increases the delayed fracture resistance in several ways. It contributes to austenitic grain refinement and forms precipitates that trap hydrogen. Finally, the hardening effect of the titanium makes it possible to carry out tempering operations at higher holding temperatures. The maximum titanium content is set here in order to avoid obtaining precipitates of too large a size which would then degrade the resistance of the steel to delayed fracture.
The steel of the present invention can optionally contain niobium at contents comprised from 0.01 to 0.1 wt %. Niobium improves the hydrogen resistance, as it can on the one hand limit the formation of borocarbides Fe3(C,B); Fe23(C,B)26 which consume, and therefore, lower the “free” boron content available for segregation at the grain boundaries, and, on the other hand, limits the austenitic grain growth by forming carbonitrides. The refinement of grains results in a higher total length of grain boundaries, and therefore in a better distribution of harmful elements, such as phosphorous and sulfur, in lower concentration. Furthermore, a decrease in austenitic grain size results in an acceleration of the kinetics of the bainitic transformation. The maximum niobium content is set in order to avoid obtaining precipitates of too large size which would then degrade the resistance of the steel to delayed fracture. Furthermore, when it is added in too large amount, niobium leads to an increased risk of “crack” defects at the surface of the billets and blooms as continually cast. These defects, if they cannot be completely eliminated, may prove very damaging in respect of the integrity of the properties of the final part, especially as regards fatigue strength and hydrogen resistance. This is why the niobium content is kept below 0.1 wt %.
Further optionally, the steel may comprise vanadium at a content lower than or equal to 0.5 wt %. When it is present, thanks to its hardening effect, the vanadium makes it possible to carry out tempering operations at higher temperatures. The maximum vanadium content is set to avoid obtaining precipitates of too large size which might degrade the resistance of the steel to delayed hydrogen fracture. In particular, the vanadium content may be comprised at a content from 0.05 to 0.5 wt %.
The rest of the composition is iron and unavoidable impurities, in particular resulting from the elaboration.
More particularly, the composition of the steel part consists of the above-mentioned elements.
The steel part according to the invention is more particularly a cold formed steel part, and more particularly a cold headed steel part.
The steel part has a microstructure comprising, in surface fractions or area %, of at least 80% bainite, and a cumulative presence of Residual Austenite and martensite from 1% to 25%.
Bainite is present in the steel according to the invention as a matrix phase and imparts strength to such steel. Bainite is present in the steel at least 80% by area fraction and preferably from 80% to 95% by area fraction and more preferably from 85% to 95%. Bainite is formed during quenching. Such bainite may include Cementite-Free Lath-Like Bainite and Lower bainite. The cementite-free Lath-Like bainite is consisting of bainite in the form of laths and including, between these laths, carbides such that the number N of inter-lath carbides larger than 0.1 micrometers per unit of surface area is less than or equal to 50000/mm2. This cementite-free lath-like bainite structure confers to the steel of the present invention a good resistance against hydrogen. The lower bainite is consisting of bainite in the form of laths and including, fine iron carbides sticks which are precipitated inside the laths. The lower bainite structure provide the steel of the present invention with elongation and tensile strength. The lath structure of both Lower bainite and cementite free lath-like bainite allow for a better distribution of the hydrogen which tends to segregate such an improved distribution of the hydrogen that may be present in the bainite areas of the microstructure therefore increases the resistance to hydrogen.
Residual Austenite and Martensite are cumulatively present from 1% to 25% by area fraction in the steel according to the invention. Martensite is formed during cooling after the soaking from the unstable austenite formed during annealing. Martensite is composed of fine laths elongated in one direction inside each grain issued from a primary austenite grain, in which fine iron carbides sticks which are 50 to 200 nm long are precipitated between the laths following the <111> direction. Martensite imparts ductility and strength to the Steel of the present invention. However, when martensite and Residual austenite are cumulatively present above 25%, they impart excess strength but diminishes the elongation beyond acceptable limit for the steel of the present invention due to the reason that martensite has same amount of carbon content as of Residual Austenite hence the fresh martensite is brittle and hard. Preferred limit for the cumulative presence of Residual Austenite and martensite for the steel of the present invention is from 4% to 22% and more preferably from 4% to 20%.
The steel parts according to the invention may advantageously be used as parts for chassis and wheel hub applications. In particular, these steel parts may be used as bolts and screws for such applications, and for example chassis bolts, hub to bearing bolts, rim to hub bolts.
The diameter of the steel part is for example lower than or equal to 22 mm, and more particularly lower than or equal to 20 mm, and even more particularly lower than or equal to 16 mm. More particularly, the diameter of the steel part is for example greater than or equal to 5.5 mm.
The steel part described above may, for example, be obtained using a method comprising:
The semi-finished product provided during the provision step has the following composition, by weight:
This composition corresponds to the composition previously described for the steel part.
The semi-finished product is in particular a wire, having, for example, a diameter comprised from 5 mm to 25 mm.
As mentioned above, the annealing step is performed at an annealing temperature strictly lower than the Ac1 temperature of the steel. As is conventional, the Ac1 temperature is the temperature at which austenite begins to form during heating.
The annealing step is intended for temporarily decreasing the tensile strength of the steel so as to prepare it for cold forming. For example, at the end of the annealing step, the steel has a tensile strength lower than or equal to 600 MPa. Such an annealing is called globulization or spherodization annealing.
More particularly, during the annealing step, the semi-finished product is heated to an annealing temperature greater than or equal to Ac1-20° C.
During the annealing step, the semi-finished product is preferably held at the annealing temperature for a time which is chosen, as a function of the annealing temperature, such that the tensile strength of the steel after annealing is lower than or equal to 600 MPa. For example, the holding time at the annealing temperature is comprised from 5 to 9 hours.
According to a particular example, the annealing step is performed at an annealing temperature equal to 720° C., and the holding time at the annealing temperature is equal to 5 hours.
The annealing step is preferably carried out in a neutral atmosphere, for example in an atmosphere consisting of nitrogen gaz.
After holding at the annealing temperature, the semi-finished product is cooled down to room temperature.
The cooling is preferably performed at a speed chosen so as to avoid the precipitation of pearlite and the formation of bainite, and thus so as to maintain a tensile strength smaller than or equal to 600 MPa after cooling. This cooling speed can be determined without difficulty using the CCT diagram of the steel.
According to a particular example, the cooling from the annealing temperature is performed in three stages: a first cooling stage from the annealing temperature to about 670° C., where the steel is cooled at a cooling speed smaller than or equal to 25° C./h, a second cooling stage from about 670° C. to about 150° C. at a cooling speed smaller than or equal to 250° C./s and a third cooling stage, from about 150° C. down to ambient temperature at a cooling speed corresponding to cooling in ambient or natural air. This three-step cooling and the corresponding temperatures and speeds are given only by way of example, and different temperatures and speeds may be used depending in particular on the composition of the steel and on the final tensile strength desired.
The cold forming step is, for example, a cold heading step, such that a cold headed product is obtained at the end of the cold forming step, and a cold headed steel part is obtained at the end of the heat treatment.
The method optionally comprises, between the annealing and the cold heading step, a step of cold drawing the annealed semi-finished product so as to reduce a diameter thereof. This cold drawing step is in particular a wire drawing step.
Preferably, the cold drawing step is preceded by a surface preparation comprising cleaning the surface of the semi-finished part, followed by a step of forming a lubricating coating on the surface of the semi-finished part.
The cleaning step for example comprises a degreasing and/or a mechanical or chemical descaling or pickling, optionally followed by a neutralization. In this context, neutralization is a cleaning process used to clean all the alien particles or substances from the surface of the steel in order to reduce the risk of corrosion.
The step of forming a lubricating coating for example comprises a phosphate treatment and a soaping.
After cold forming, the cold-formed steel part is subjected to a heat treatment comprising:
This optional heat treatment is a tempering heat treatment.
According to an example, during the holding step, the product is held at a holding temperature in a furnace. According to another example, the product can be held at the holding temperature by dipping in a molten salt bath
After the end of the holding step, the products are allowed to cool down to the ambient temperature in ambient or natural air.
The heating step is carried out in such a manner that the steel part has an entirely austenitic microstructure at the end of the heating step.
The average size of the austenite grains formed during this heating step is lower than or equal to 20 μm, and in particular comprised from 8 to 15 μm. This size is, for example, measured with a magnification of 500:1.
This small grain size results from the presence of micro-alloying elements in the steel which form precipitates able to pin the grain boundaries in order to avoid grain growth during the austenitizing step. This austenite grain size is the prior austenite grain size of the cold formed and quenched and tempered steel part according to the invention.
The heat treatment temperature is for example higher by a least 50° C. than the full austenitisation temperature Ac3 of the steel.
More particularly, during the heating step, the steel part is held at the heat treatment temperature for a time comprised from 5 minutes to 120 minutes.
Preferably the holding temperature during the holding step is comprised from 200 to 380° C.
At the end of the holding step, a cold formed, and more particularly cold headed, and quenched steel part is obtained.
The thus obtained steel part has the microstructure described above for the steel part.
Laboratory tests were carried out on castings having the chemical compositions I1 to I6 which are according to the present invention. R1 to R3 are reference steels composition which are not according to the present invention.
R1
0.299
0.898
R2
0.443
0.940
R3
0.364
0.875
In the above Table 1, the compositions are indicated in wt % and the underlined values are not according to the invention.
In all of the above compositions, the remainder of the composition consists of iron and unavoidable impurities.
The inventive steels and the reference steels are reheated at 1150° C. and then are hot rolled with a finishing temperature above 800° C. in the form of wire having a diameter of 16 mm. Thereafter all the wire rods (semi-finished product) for both inventive as well as reference steels were subjected to annealing comprising holding the wire rods at a temperature of 720° C. with a holding time of 5 hours, followed by cooling. Cooling was performed in three stages comprising cooling at a cooling speed of 25° C./h down to 670° C., followed by cooling at 250° C./h until 150° C., and finally natural or ambient air cooling down to room temperature. These cooling speeds were obtained by adjusting the heating conditions in the annealing furnace accordingly, the heating being reduced or turned off depending on the needs, in a manner known to the skilled person. Ac1 and Ac3 for both Inventive Steels (I1 to I6) and reference steel (R1 to R3) are calculated by a dilatometry study.
Thereafter the cold formed steel part is subjected to a heating and quenching heat treatment according to the table 2.
R1
R2
R3
The underlined values are not according to the invention in Table 2.
Tensile tests were performed directly on wire rods. The tensile testing was performed according to standard NF EN ISO 6892-1, i.e. with a cross head speed of 8 mm/min. Each value is the average of three measurements.
A hardness profile along the cross section of the samples was performed. Vickers hardness tests were carried out under a load of 30 kg for 15 seconds durations. The hardness was measured according to standard NF EN ISO 6507-1. Each value is the average of three measurements.
The results of these tests are summarized in Table 3 below.
Furthermore, the microstructure of the obtained products was analyzed based on cross-sections of these products. More particularly, the structures present in the cross-sections were characterized by light optical microscopy (LOM) and by scanning electron microscopy (SEM). The LOM and SEM observations were performed after etching using a Nital containing solution.
The results of these analyses are summarized in the following Table 4.
In Tables 3, the following abbreviations are used:
R1
1013
819
R2
53.5
R3
53.5
For each of the experiments I1 to I6 as well as R1 to R3, the hydrogen resistance of the corresponding samples was determined by comparison of the results of a slow strain rate tensile test conducted on the smooth test samples subjected to strain rate of 10−5 s−1 on an uncharged sample and then on a sample charged with hydrogen in accordance with NF A-05-304 standards.
More particularly, the inventors determined the ductility (through the percent reduction of area Ra) on the charged and uncharged samples and compared the results through an embrittlement index.
The total H2 content inside samples before charging was equal to about 0.3 ppm.
Hydrogen charging was performed through cathodic charging using an electrolytic solution composed of H2SO4 1N with the addition of an hydrogen promoter Thiourea 2.5 mg/L, with a current density I=0.8 mA/cm2 for 5 hours.
For each pair of samples (charged and uncharged), the embrittlement index IRa relating to the percent reduction of area is calculated using the following formula:
IRa=1−[RA(H2)/RA(H2=0)], where RA(H2) corresponds to the value of the percent reduction of area measured on the sample charged with hydrogen, and RA(H2=0) corresponds to the value of the percent reduction of area measured on the uncharged sample.
An embrittlement index close to 1 means that the grade is very sensitive to Hydrogen Embrittlement. An embrittlement index IRa lower than to 0.09 was considered satisfactory in view of the desired applications and embrittlement index IRa lower than or equal to 0.08 is advantageous for the desired applications.
The inventors further observed the fracture surface mode in each case.
The results of these tests are summarized in Table 4.
R1
0.13
R2
0.09
R3
0.63
As can be seen from the above Table 4, the ductility for the inventive steel is not significantly affected by hydrogen.
The steels having compositions 11 to 16 exhibit a higher hydrogen resistance than the reference grade R1 to R3 after quenching.
The comparison of the samples 11 to 16 having a bainite content greater than or equal to 80% as shown in table 5 with the sample having a martensitic microstructure that is R1 to R3 as shown in table 5 shows that the bainitic structure is less sensitive to hydrogen embrittlement than the martensitic structure.
It can finally be observed that the samples according to the invention (11 to 16) absorb less hydrogen under the same charging conditions than the comparative samples (R1 to R3).
Therefore, these experiments show that the steel parts according to the invention are particularly well adapted for applications as mentioned above, they have very good mechanical properties, and in particular a good tensile strength, associated with an improved resistance to hydrogen embrittlement as compared to prior art steel parts.
The method according to the invention further has the advantage that it allows obtaining, after annealing, a sufficiently low tensile strength so as to enable the use of conventional cold forming tools, and reduce the wear thereof, while at the time resulting in final parts having a high tensile strength (greater than or equal to 1100 MPa).
The microstructures of the steels were characterized using Light Optical Microscopy (LOM) Scanning Electron Microscopy (SEM) after 2% Nital etching. Quantitative X-ray analysis has been done to determine the fraction of retained austenite.
R1
0
100
R2
0
100
R3
0
100
the underlined values are not according to the invention.
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/IB2021/055331 | 6/16/2021 | WO |
