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Method for producing a steel part and steel part
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.
W02016/158470 is Age 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
W02016/158470 lacks hydrogen embrittlement.
W02011/124851 is a mechanical steel part in steel with high characteristics,
characterized in that its composition, comprising in weight percentages, is
0.05`)/000.25`)/0; 1.2c)/oMn2c)/0; 1c)/oCr2.5(:)/0; wherein the contents of C,
Mn and Cr
are such that (830-270C %-90 Mn `)/0-70Cr %)560; O<Si1.55; O<Ni1`)/0;
O<Mo0.5`)/0;
O<Cu1`)/0; 0<V0.3%; O<A10.1`)/0; 0<B0.005%; O<Ti0.03`)/0; O<Nb0.06`)/0;
0<S0.1%; O<Ca0.006`)/0; O<Te0.03`)/0; O<Se0.05`)/0; O<Bi0.05`)/0; O<Pb0.1`)/0;
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
W02011/124851 is not demonstrate hydrogen embrittlement as well as reduction
area of
58% or more.
However, it is desirable to further improve the resistance to hydrogen
embrittlement
of the parts.
Therefore, an aim of the invention is to provide 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 :
- an ultimate tensile strength greater than or equal to 1100 MPa and
preferably
above 1150 MPa, or even above 1180 MPa,
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- a yield strength greater than or above 880MPa and preferably above 900MPa
- a total elongation of 12% or more and preferably 13% or more.
- Hydrogen Embrittlement index of less than 0.09 and preferably less than
or equal
to 0.08
- a reduction in area of more than 58% and preferably above 60% or more and
more preferably 62% or more.
In a preferred embodiment, the steel part shows a hardness from 360Hv to
405Hv.
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 `)/0
(wt%).
The steel part according to the invention has a composition comprising, by
weight:
0.05% C 0.15%
0.01% Si 1%
1.2% Mn 2`)/0
0.1% Cr 2%
0.001 Al 0.1%
0.003% N 0.01%
0 S 0.015%
0 P 0.015%
optionally
0% Ni 1%
0% Mo 1.0%
0% Nb 0.1%
0% Ti 0.04 c)/0
0 V 0.5%
0% B 0.01%
the remainder consisting of iron and unavoidable impurities.
Carbon is present in the steel of 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
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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 present invention from 0.01% to 1%. Silicon
imparts the
steel of 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 present invention. An
amount of more
than 1 `)/0 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 `)/0, manganese has a negative
effect on the
steel of 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 `)/0 and more
preferably
from 1.4% to 1.9%.
Chromium is present from 0.1% to 2% in the steel of 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 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
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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 (AIN), 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.
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The steel may optionally contains from 0.01 to 1wtc)/0 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.
5
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 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
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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
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the form of laths and including, between these laths, carbides such that the
number N of
inter-lath carbides larger than OA 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
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 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 present invention. However, when martensite and Residual
austenite
cumulatively presence is above 25%, it imparts excess strength but diminishes
the
elongation beyond acceptable limit for the steel of 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 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, 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:
- providing a semi-finished product made of steel;
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- annealing this semi-finished product at an annealing temperature strictly
lower than
the Ad 1 temperature of the steel;
- cold forming the semi-finished product into a cold formed product;
- subjecting the cold formed product to a heat treatment so as to obtain a
cold
formed steel part, the heat treatment comprising:
- heating the cold formed product to a heat treatment temperature greater
than or
equal to the full austenitization temperature (Ac3) of the steel
- quenching to room temperature
- and then optionally,
- holding the product at a holding temperature comprised from 100 C to 400 C
for
a time from 15 minutes to 2 hours.
The semi-finished product provided during the provision step has the following
composition, by weight:
0.05% C 0.15%
0.01% Si 1%
1.2% Mn 2%
0.1% Cr 2%
0.001 Al 0.1%
0.003% N 0.01%
0 S 0.015%
0 P 0.015%
optionally
0% Cu 1%
0% Ni 1%
0% Mo 1.0%
0% Nb 0.1%
0% Ti 0.04 A,
0 V 0.5%
0% B 0.01%
the remainder consisting of iron and unavoidable impurities.
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.
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As mentioned above, the annealing step is performed at an annealing
temperature
strictly lower than the Ad 1 temperature of the steel. As is conventional, the
Ad1
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 Ac 1-20 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 600MPa 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.
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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,
5 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.
10 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:
- heating the cold formed product to a heat treatment temperature greater
than or
equal to the full austenitization temperature Ac3 of the steel;
- quenching to room temperature
- and then, optionally, holding the product at a holding temperature
comprised from
100 C to 400 C for a time from 15 minutes to 2 hours.
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 pm, and in particular comprised from 8 to 15 pm. 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
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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.
Experiments
Laboratory tests were carried out on castings having the chemical compositions
II to
16 which are according to the present invention. RI to R4 are reference steels
composition
which are not according to the present invention.
Table 1: Chemical compositions of the castings
Steel C Mn Si Cr Al P S N Ni Cu Mo Nb Ti
11 0.107 1.750 0.344 1.020 0.016 0.005 0.008 0.0038 0.016 0.010 0.047 0.047
0.021 0.0021
12 0.124 1.760 0.093 0.789 0.019 0.003 0.003 0.0031 0.010 0.009 0.065 0.048
0.028 0.0013
13 0.121 1.750 0.334 1.010 0.026 0.008 0.007 0.0050 0.103 0.011 0.198 0.034
0.025 0.0034
14 0.130 1.800 0.370 1.250 0.026 0.009 0.008 0.0080 0.107 0.012 0.107 0.036
0.028 0.0035
15 0.127 1.800 0.348 1.250 0.022 0.007 0.006 0.0050 0.109 0.011 0.105 0.022
0.026 0.0037
16 0.118 1.590 0.350 1.380 0.016 0.007 0.007 0.0100 0.106 0.011 0.101 0.036
0.029 0.0034
IR1 0.299 0.898 0.072 0.270 0.031 0.005 0.003 0.0073 0.034 0.047 0.096 0.001
0.037 0.0020
M 0.443 0.940 0.169 1.480 0.025 0.005 0.005 0.0047 0.159 0.040 0.192 0.049
0.014 0.0008
M 0.364 0.875 0.035 1.040 0.030 0.010 0.005 0.0049 0.018 0.007 0.002 0.002
<0,002 <0,0003
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.
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Table 2 ¨ Process parameters
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 16mm. 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. Ad 1 and Ac3
for both
Inventive Steels (11 to 16) and reference steel (RI 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.
Table 2 ¨ Process parameters
Heating Heating
Heating rate Cooling Quenching Ad 1
Ac3
Trials temperature holding time
( C/s) C h) rate ( C/s) stop T ( C)
( C) ( C)
() (
11 1 900 0.5 20 25 736
868
12 1 900 0.5 20 25 723
860
13 1 900 0.5 20 25 750
855
14 1 900 0.5 20 25 751
856
15 1 900 0.5 20 25 747
847
16 1 900 0.5 20 25 741
851
RI 1 890 0.5 20 25 719
751
R2 1 890 0.5 20 25 729
807
R3 1 890 0.5 20 25 732
755
The underlined values are not according to the invention in Table 2.
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Table 3 : Mechanical properties
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:
TS (MPa) refers to the tensile strength measured by tensile test in the
longitudinal
direction relative to the rolling direction,
YS (MPa) refers to the yield strength measured by tensile test in the
longitudinal
direction relative to the rolling direction,
RA (`)/0) refers to the percent reduction of area measured by tensile test in
the
longitudinal direction relative to the rolling direction,
El (`)/0) refers to the total elongation measured by tensile test in the
longitudinal
direction relative to the rolling direction,
HV30 refers to the result of the hardness measurement,
Table 3: Mechanical properties of the samples after quenching
TS (MPa) YS (MPa) El (`)/0) RA (Y()) HV30
11 1225 901 15.0 64.3 383
12 1216 925 13.6 63.7 393
13 1231 1002 14.2 64.7 394
14 1280 1033 14.4 65.4 400
15 1280 1028 15.3 63.5 400
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16 1240 1013 14.5 64.7 393
R1 1013 819 16.1 66.2 364
R2 1253 1168 13.9 53.5 410
R3 1260 1172 7.9 53.5 448
the underlined values are not according to the invention.
Table 4 : Hydrogen Embrittlement Results
For each of the experiments 11 to 16 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 1= 0.8 mA/cm2 for 5 hours.
For each pair of samples (charged and uncharged), the embrittlement index IR,
relating to the percent reduction of area is calculated using the following
formula:
I Ra = 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 IR, lower than to 0.09 was
considered
satisfactory in view of the desired applications and embrittlement index IR,
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.
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Ra (`)/0) Ra (Y()) Hydrogen Embrittlement
without H2 with H2 index (I Ra )
Ii 67.5 66.2 0.02
12 67.3 65.0 0.03
13 65.2 64.0 0.02
14 65.1 64.0 0.07
15 67.4 65.0 0.04
16 65.8 64.0 0.03
R1 68.2 59.3 0.13
R2 54.3 49.6 0.09
R3 55.1 20.3 0.63
As can be seen from the above Table 4, the ductility for the inventive steel
is not
5 significantly affected by hydrogen.
The steels having compositions 11 to 16 exhibit a higher hydrogen resistance
than
the reference grade R1 toR4 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
10 is R1 to R4 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 R4).
15 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).
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Table 5: Microstructure
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.
Trials Bainite Martensite +RA
(%) (%)
11 85 15
12 82 18
13 85 15
14 95 5
95 5
16 90 10
R1 0 100
R2 0 100
R3 0 100
10 the underlined values are not according to the invention.
