Note : Les descriptions sont présentées dans la langue officielle dans laquelle elles ont été soumises.
CA 02792615 2012-09-10
Descriptive Report of Patented Invention: "TOOL STEEL FOR
EXTRUSION"
The present invention relates to a steel intended for use in various hot
form tools and dies, particularly for extrusion of aluminum alloys or other
non-ferrous
metals. Although initially designed for extrusion processes, the material can
also be
employed in other hot forming processes, in which the metal to be formed
withstands
temperatures above 600 C, although the said steel can be employed in
processes at
lower temperatures or even at ambient temperature. The composition of the
steel in
question allows it to be classified as hot work tool steel, whose primary
characteristic
the lower content of high-cost alloying elements, such as molybdenum and
vanadium, but with tempering resistance (or resistance to loss of hardness)
greater
than that of conventional steels of prior art concept. An additional
alternative to the
steel of the present invention is provided to increase hardness after
nitriding, and
may result in performance levels even greater than those of conventional
steels, at
the same time that the cost is kept low due to a simpler chemical composition.
Such
an effect is possible by carefully designing the alloy, and setting the
optimum ranges
of the elements: carbon, chromium, molybdenum and aluminum.
The term hot work tools is applied to a large number of hot-forming
operations, employed in industries and focused on the production of parts for
mechanical applications, especially automotive parts. The most popular hot-
forming
processes are forging of steel, and the extrusion or casting of non-ferrous
alloys.
Other applications performed at high temperature, typically above 500/600 C,
can
also be classified as hot work. In these applications, molds, dies, punches,
inserts
and other forming devices are classified by the generic term: hot work tools.
These
tools are usually made of steels, which require special properties to
withstand high
temperatures and the mechanical efforts of the processes in which those tools
are
employed.
Among their key properties of hot work steels, the following stand out:
resistance after high temperature tempering, the resistance to the loss of
hardness
called tempering resistance, the toughness, the hardenability and physical
properties
such as thermal conductivity and specific heat.
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The extrusion dies used for non-ferrous alloys, especially aluminum
alloys are main hot work target for applying the steel of the present
invention. These
typical dies comprise an important segment of the tool steel market both in
Brazil and
abroad. In this application, the steels are very standardized, based on steels
such as
ABNT H13 (see Table 1), with quality requirements not as strict as those of
other
applications, e.g., pressure die casting, but with emphasis on lower
production costs.
The increased cost of metal alloys, especially Mo and V, significantly
impaired this segment, making it eager for low-cost alternatives. Low-alloy
steels
have been employed, such as DIN 1.2714 (chemical composition given in Table
1).
However, their low wear resistance due to reduced hot strength and lower post-
nitriding hardness prevents them from being applied.
Recent developments, such as US 2009/0191086, were focused on
the reduction of alloying elements, by means of reduced Cr, Mo and V content.
However, negative effects are produced by reducing the Cr content. First, the
alloys'
composition is not sufficient to achieve high hardness after tempering (at
least 45
HRC after tempering at 600 C). Second, a reduced Cr content can also generate
lower hardness after nitriding, which is not suitable for extrusion
applications,
considering the apparent gain produced by nitridation in these applications
(virtually
all extrusion dies are currently nitrified).
Table 1: Typical chemical composition of steels of prior art concept. The sum
Mo + V
+ Co is shown because these elements have the highest cost, and are closely
related to the final cost of the tool steel. Content in percentage by mass and
Fe
balance. For all extrusion applications the W content is low, usually < 0.1 %.
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Designation C Si Mn Ni P Cr Mo V Al
H13 0.38 1.0 0.3 0.3 0.025 5.0 1.2 1.0 <0.05
DIN 1.2714 0.56 0.3 0.7 1.7 0.025 1.1 0.5 0.1 <0.05
US 2009/0191086 0.38 0.5 1.3 0.3 0.009 2.4 0.6 0.5 <0.05 or
0.53
A third problem of invention US 2009/0191086 relates to the hardness
of the die core, which may be lower due to decreased hardenability as a result
of
reduced Cr and Mo contents. To avoid this, the alloys of invention US
2009/0191086
have higher Mn content, which lead to higher hardenability, potential
segregation
problems (banding) and excessive austenite retention. Both effects may impair
the
final hardness and toughness and, thus, the tool life. A final aspect can also
be
mentioned, with regard to the high Mn content: Scrap from this steel can
hardly be
incorporated into the production of conventional, low-Mn-content hot work
steels.
Given all these drawbacks, invention US 2009/0191086 is considered
by the authors as a cost-reducing solution, but with inferior properties. In
the text of
the patent, the authors quantify the expected efficiency loss, of about 20 to
30%
lower than that of steel H13. Considering the machining and heat treatment
costs
associated with dies, this efficiency loss can be considered quite
significant, thus
requiring a reduction of the cost of the material by more than 30% to
compensate the
substitution. For example, considering that only 60% of the final die cost is
associated with the tool steel used, a 30% lower life can only be viable if
the cost of
the new material is half the cost of the conventional material. From 2005 to
2008,
when the cost of raw materials peaked, this could be true (though still
difficult to
occur, because the cost difference required is too high). However, for the
current
scenario, such cost reduction can hardly be achieved for steel H13,
considering only
the reduction of the Mo and Cr contents. Thus, reduction in cost associated
with
efficiency loss of the alloy of patent US 2009/0191086 can be currently
considered
impractical for such application.
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Given this scenario, it is evident the need for a tool steel which
effectively has a positive effect on tool life by means of an equivalent
performance,
but at a cost lower than that of steel H13. This is only possible if the steel
in question
has tempering resistance and hardness after tempering at 600 C (typical heat
treatment condition) equivalent to those of steel H13, but with lower content
of
alloying elements and suitable hardness after nitriding. In addition, the
material used
must have high hardenability, but free of problems associated with high Mn
content,
thus allowing it to be applied to tools larger than extrusion dies.
Therefore, the steel of the present invention will fulfill all these needs.
To achieve the cost reduction/zero quality loss goal, the effect of the
key elements related to hot strength, Cr and Mo, was studied separately. Apart
from
significant findings, this study also showed that the variation of the content
of these
elements is not sufficient to promote the hot strength required. Thus, the C
content
could be increased up to levels that did not impact toughness, especially whit
the
accompaniment of low P and Si contents. Finally, the Al effect was used to
compensate for the reduction of Cr, hence, a potential lower hardness after
nitriding.
This work also focused on this issue because the nitrified layer is critical
to providing
wear resistance to various hot forming tools, especially extrusion and hot
forging
tools.
Therefore, in order to satisfy the above conditions, the steel of the
present invention has a composition of alloying elements, which, in percentage
by
mass consists of:
= 0.40 to 0.60 C, preferably 0.45 to 0.55 C, typically 0.50 C
= 2.5 to 4.5 Cr, preferably 3.0 to 4.2 Cr, typically 3.8 Cr
= 0.30 to 0.90 Mo, preferably 0.50 to 0.70 Mo, typically 0.60 Mo.
Given its chemical similarity to W, Mo can be replaced with W, 2W:1 Mo ratio
by
mass.
= 0.1 to 1.0 V, preferably 0.3 to 0.8 V, typically 0.4 V; V can be
partially or fully replaced with Nb, following a 1 Nb:0.5 V ratio by mass.
= up to 1.0 Si, preferably up to 0.50 Si, typically 0.30 Si
Max 1.0 Mn, preferably max 0.80 Mn, typically max 0.50 Mn.
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As described below, Al can be added simultaneously to the alloys of
the present invention to provide gains in terms of hardness after nitriding,
but also
negative effects in terms of toughness and complexity of the steel-making
process.
Thus, the Al content must be dosed as follows, in percentage by mass:
- Max 1.0 Al, preferably max 0.80 Al, typically max 0.60 Al. For
compositions in which the effects of Al are not targeted, this element should
be
treated as residual impurity, limited to 0.10, typically < 0.05.
The compositions should be characterized by balance by Fe (iron) and
metallic or non-metallic deleterious substances inevitable to the steelmaking
process,
in which said non-metallic deleterious substances include but are not limited
to the
following elements, in percentage by mass:
= Max 0.030 P, preferably max 0.015 P, typically max 0.010 P.
= Max 0.10 S, preferably max 0.030 S, typically max 0.008 S.
= Max 1.5 Ni or Co, preferably up to 1.0 Ni or Co, typically below 0.5 Ni
and Co.
Next, we describe the ratios of the specification of the composition of
the new material. The percentages listed refer to percent by mass.
C: Carbon is primarily responsible for martensite hardening under low
temperature conditions. However, together with the alloying elements, carbon
also
plays a role in the secondary hardening, important for the hardening at high
temperature. In these cases, the C content is more important for hardness at
temperatures below 600 C, when hardness still depends on the martensite
hardness
or formation of cementite or Cr carbides. Furthermore, carbon is an important
hardenability-promoting element, and causes no increase in cost. It is also
considered important to increase hardness to 45 HRC and up, carbon contents of
at
least 0.40% are recommended, preferably above 0.45%. On the other hand, very
high C contents, cause excessive precipitation of grain-shaped carbides at the
time
of quenching (especially when Mo and V contents are high), as well as lead to
increased hardness and volume of secondary carbides. Thus, toughness is
generally
CA 02792615 2012-09-10
impaired., the C content should be limited to a maximum value of 0.60%,
preferably
below 0.55%. This limitation also plays a role in the reduction of the amount
of
retained austenite, preventing problems associated with dimensional
instability and
embrittlement.
Cr: The chromium content should be higher than 2.5%, preferably
greater than 3.0%, because this element favors hardenability, which is
important for
application in large tools. However, the Cr content should be limited. The
present
invention has incorporated the concept of reducing the Cr content to improve
tempering resistance. The mechanisms of this effect are not fully understood
but they
may be related to the formation of secondary Cr carbides, M7C3-type, which
dissolve
Mo and V are the first carbides to be formed. Therefore, the lower the Cr
content, the
lower the amount of M7C3 carbides and, thus, the greater the amount of Mo and
V
available for the formation of fine carbides M2C and MC, which are also
important for
secondary hardening. The end result is a significantly higher tempering
resistance in
steels with lower Cr content, thus enabling the reduction of the Mo content
when
compared to steels of prior art concept.
Mo and W: low concentrations of Mo have been employed in the
present invention not only for the purpose of cost reduction, but also to
promote the
highest secondary hardness and tempering resistance equivalent or even greater
than that of steel H13 in association with Cr and C contents. To do so, the
alloy of the
present invention must contain at least 0.30%, preferably above 0.50%. On the
other
hand, an extremely high Mo content might harm toughness due to deposition of
pro-
eutectic carbides during the quenching phase and can increase significantly
the alloy
cost, in an opposite direction to the cost-reduction goal of the present
invention.
Hence, the Mo content should be limited to 0.90%, preferably below 0.70%.
Tungsten and molybdenum produce similar effects in the tool steel of the
present
invention, forming M2C or M6C secondary carbides. Thus, they can be jointly
specified through the tungsten equivalent relationship (Weq) given by the sum
W +
2Mo, which normalizes the differences in atomic weight between the two
elements.
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V: Vanadium is primarily important for the formation of MC
secondary carbides. Because they are very thin, these carbides block the
movement
of dislocation lines, increasing mechanical strength. V also improves grain
growth,
allowing high austenitizing temperatures (above 1000 C). For such effects, V
must
be above 0.1%, preferably above 0.3%. However, excessively high V grades may
generate primary, difficult-to-solubilize carbides, thus reducing toughness,
and also
promote significant increase of costs. Hence, the V content should be lower
than
1.0%, preferably below 0.6%.
Si: silicon produces a strong effect on secondary hardening and
toughness. When a low Si concentration is used, toughness improves due to a
better
distribution of secondary carbides. Therefore, the Si content of the material
of the
present invention must be lower than 1.0%, typically below 0.5%.
Mn: high Mn contents may be considered undesirable for promoting
intense micro-segregation generating banding at different degrees of hardness,
and
for increasing the retained austenite content; therefore Mn is considered a
deleterious element in the present invention. Thus, the Mn content should be
limited
to 1.0%, preferably below 0.8%, typically below 0.50%.
Al: to promote greater hardness of the nitrified layer, the alloys' Al
content can be high. However, the Al content, under these conditions, should
be
limited to 1.0% because they lead to decreased toughness. Thus, Al contents
between 0.40% and 0.60% may be of interest for this purpose. However, for
applications in which the hardness of the nitrified layer is slightly lower
than that of
steel H13, but high toughness is required, the Al content of the alloy of the
present
invention can be < 0.1 %, typically below 0.05%.
Residual Elements: Other elements such as Ni and Co should be
considered as deleterious substances associated with the steelmaking
deoxidation
processes or inherent to the manufacturing processes. Hence, the Ni and Co
content
should be limited to 1.5%, preferably below 1.0%. In terms of formation of
inclusions,
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the sulfur content should be controlled, because such inclusions may lead to
cracking
during operation; therefore the S content should remain below 0.050%,
preferably
below 0.020%. Also, for high toughness purposes, embrittling elements such as
P
should be avoided, being desirable P < 0.030%, preferably P < 0.015%,
typically P <
0.010%. Indeed, a low Cr content also helps to reduce the P content in
electric arc
furnace steelmaking processes, thus leading to conclusions that are not
contradictory
to the cost reduction philosophy desired.
The alloy, as described above, can be produced as rolled or forged
products through conventional or special processes such as powder metallurgy,
spray forming or continuous casting, such as wire rods, bars, wires, sheets
and
strips.
The experiments carried out are described below, and reference is
made to the following attached figures:
- Figure 1A shows the effect of the Mo content on hardness after tempering at
600
C, while Figures 1 B and 1 C show the effect of the Cr content at 0.60% Mo on
usual
C contents (Figure 1B) and higher C contents (Figure 1C); the horizontal
dashed line
of Figures 1A, 1B and 1C indicates the Minimum Hardness desirable for the
application.
- Likewise Figure 1, Figures 2A, 2B and 2C show the effect of molybdenum (Fig.
2A)
and chromium (Fig. 2B and Fig. 2C) on tempering resistance. The higher the
hardness at high temperatures the greater the alloy's tempering resistance. In
all
cases, the alloys were first annealed at 600 C.
- Figures 3A and 3B show the CCT curve of the compositions of the present
invention, considering two Cr contents. Quantitative hardenability results can
be
obtained from the number of formed phases (pearlite and bainite) and, most
importantly, from the final hardness obtained per rate. The compositions are
summarized in Table 1, base 3, considering Cr contents of 3% and 4% for
comparison purposes. Figure 3A illustrates the CCT curve for a 0.50% C, 3.00%
Or
composition, and Figure 3B shows the CCT curve for a 0.50% C, 4.00% Cr
composition.
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- Figure 4 shows the CCT curve of H13 steel of the prior art concept, whose
data can
be compared to the results of the steel of the present invention. The same
data
concerning number of phases and hardness shown in Figure 3 can be assessed for
different cooling rates.
- In Figures 5A and 5B, the alloys with the final composition of the present
invention,
PI 1 to PI 3, are compared in terms of hardness after tempering (Fig. 5A) and
loss in
hardness vs. time (Fig. 5B) at 600 C (referred to in the tempering resistance
text).
- Figure 6 compares the results of impact toughness tests conducted for two
types of
transverse test specimens: unnotched (7 mm x 10 mm section, as per NADCA) or
Charpy V, with 10 mm x 10 mm section and V notch. All materials treated to
hardness 45 HRC according to the parameters of Figure 5a.
- Figure 7 shows the hardness profile of the nitrified layer of alloys PI 1,
PI 2 and PI 3
vs. steel H13. A plasma nitriding process was conducted for steel H13. Prior
to
nitriding, all sample alloys were quenched and tempered such to reach 45 HRC.
EXAMPLE 1: Effect of Molybdenum, Chromium and Carbon
For this work, samples of approximately 200 g were collected in an
experimental VIM furnace with varied composition for the same heat. Therefore,
three heats were produced by varying the Cr, Mo and C contents, as shown in
Table
1 below (details: Annex 1). Steel H11 served as a base for these alloys since
it
already has half of V content. The materials were always characterized after
special
annealing (austenitizing at 1010 C, oil solubilization and over-annealing at
810 C).
In this process we used annealing at 1020 C and tempering between 400 and
650 C. Steel H13, of typical industrial composition, was used as a base.
Hardness after tempering at 600 C is shown in Figure 1, highlighting the
effects of reduced Mo and Cr contents, and also the effect of higher C
content. With
regard to the Mo content, a lower Mo concentration results in lower hardness
after
tempering. However, if the Cr content drops, post-tempering hardness rises. A
possibility is that a lower Cr content reduces the amount of M7C3 which, in
turn,
dissolves Mo. Thus, a higher content of free Mo should be present in alloys of
lower
Cr content, which explains a more intense response to tempering.
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Despite this important Cr effect, just reducing its content is not sufficient
to
promote the required hardness (about 45 HRC). Possibly, the required hardness
can
be obtained by tempering at lower temperatures. However, this practice is
sometimes not feasible for hot work because the ideal tempering temperature
should
be 50 to 80 C above the working temperature to provide proper tempering
resistance. Thus, for hot work involving extruded and cast aluminum, the
typical
tempering temperature should be 600 C.
Table 1: Chemical compositions adopted for samples from the same heat with
variation of a single element. The asterisks used in the Cr and Mo fields of
the table
below indicate that several compositions using this base were produced for the
same
heat, increasing the content of this element, but keeping the base composition
of the
heat.
Base 1 Base 2 Base 3 H13
Variation of... Mo Cr C -
C 0.36 0.36 0.48 0.37
Si 0.32 0.32 0.32 0.92
Mn 0.26 0.28 0.27 0.31
P 0.007 0.006 0.006 0.022
S 0.001 0.002 0.001 0.001
Co 0.02 0.02 0.02 0.02
Cr 5.00 ** *"" 4.82
Mo * 0.65 0.6 1.17
Ni 0.15 0.06 0.06 0.11
V 0.4 0.41 0.41 0.79
W 0.01 0.01 0.01 0.09
Cu 0.02 0.03 0.03 0.03
Al 0.013 <0.005 <0.005 0.02
* Mo variation: 0.05; 0.30; 0.60; 0.90; 1.22; 1.51
** Cr variation, considering 0.36%C: 2.0; 3.0; 4.0; 5.1; 6.2; 7.1,
*** Cr variation, considering 0.48%C: 2.0; 3.0; 4.0; 5.1; 6.1; 7.0;
Therefore, to increase hardness after tempering at 600 C, we increased the
C content. As shown in Figure 1, the result was effective and hardness even
higher
than those from H13 were obtained. In this case, the C effect is related to
increased
formation of secondary carbides and, when associated with a lower Cr content,
it
provides the hardness required to start the work, even in alloys of lower Mo
content
(half of steel H13). In alloys of higher C content, a similar Cr effect can be
observed.
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Besides hardness after tempering, loss of hardness is also a key factor to
promote adequate response by the alloys in question to the high temperatures
they
are subjected to. The results shown in Figure 2 demonstrate the important Mo
effect
in this regard (Fig. 2a), and also that the reduction of the Cr content is
also an
interesting option to reduce the loss of hardness, which means re-plotting the
curves
to higher hardness levels (see Fig. 2b). In alloys with higher C content (Fig.
2c), this
effect is even stronger. Thus, the low Cr/ high C combination seems
interesting.
On the other hand, the Cr content cannot be too low, such that hardenability
is not reduced. This effect was studied in the curves of Figure 3 and compared
to
steel H13 in Figure 4. Quantitatively, hardness reached after 0.3 and 0.1 C/s
corresponds to steel H13 with 635 HV and 521 HV (Figure 4), whereas the 3% Cr
alloy corresponds to 595 HV and 464 HV under the same conditions (Figure 3a).
The
scenario changes for the 4% Cr alloy, which reaches hardness >_ H13, i.e., 696
HV
and 523 HV for rates of 0.3 and 0.1 C/s (Figure 3b). Therefore, Cr contents
close to
4% Cr seem to be more interesting. Extremely below this value, i.e., 3%Cr or
less,
the volume of bainite and the hardness after tempering may prevent the
application.
Thus, a 3.8% Cr content was selected for all other tests, production of pilot-
scale
billets and evaluation of mechanical properties.
EXAMPLE 2: Effect of Al content
After defining an alloy target, four heats (50 kg cast billets, 140 mm
average section) were produced and forged as plates (Table 2) with dimensions
of
65mm x 165mm. The materials were then annealed following the same process
described in Example 1 and their properties were evaluated as discussed below.
The results confirmed the initial results shown in Figures 1 and 2, as
shown in Figure 5. Thus, the new alloys can reach similar results in terms of
hardness at 600 C (Figure 5a), or even better, in terms of tempering
resistance, if
compared to steel H13 (Figure 5b).
Table 2: experimental 50 kg billets produced for the alloys of the present
invention (PI)
and steel H 13.
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P11 P12 P13 H13
C 0.50 0.49 0.51 0.38
Si 0.32 0.31 0.32 0.99
Mn 0.35 0.35 0.35 0.35
P 0.011 0.011 0.011 0.023
S 0.003 0.003 0.004 0.004
Co 0.01 0.01 0.01 0.02
Cr 3.76 3.78 3.81 5.25
Mo 0.62 0.64 0.61 1.32
Ni 0.14 0.13 0.13 0.13
V 0.40 0.39 0.40 0.85
W 0.01 0.01 0.01 0.02
Cu 0.05 0.05 0.05 0.05
Al 0.037 0.51 1.02 0.031
Another important point can be compared in Figure 6, in terms of
toughness. The toughness of the alloy of the present invention, when bearing
low Al
contents, is equivalent to that of steel H13. This demonstrates that the low
Si and P
contents of alloy P11 compensate for the loss of toughness likely to occur as
the C
content increases in relation to steel H13. Figure 6 also shows that toughness
is
inversely proportional to the Al content.
Al contents are responsible for a significant increase of hardness after
nitriding, as shown in Figure 7. Thus, for applications in which high hardness
of the
nitrified layer is considered more relevant than toughness (e.g., extrusion of
solid
shapes), alloy PI 2 becomes interesting for having toughness > 200J and
extremely
high hardness of the nitrified layer (almost 1400 HV). Alloy PI 3 does not
show gains
in terms of the nitrified layer, but toughness is far lower.
On the other hand, in applications highly susceptible to cracking, such
as pipe extrusion dies, toughness can be considered a key property. For these
cases, alloy P11 seems more appropriate, also showing hardness after nitriding
similar to that of steel H13, reaching more than 1000 HV on surface, which is
the
typical specification for extrusion tools. Furthermore, as previously shown in
Figure
5, alloy PI 1 also presents improved hot strength properties.
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Therefore, considering the properties required for hot work applications, the
alloys of
the present invention show results equivalent to or better than those of steel
H13.
Such results are quite relevant for non-ferrous alloy extrusion dies, e.g., Al
alloys, or
hot forging dies. Alloy PI 1 has improved tempering resistance, but hardness
after
nitriding and toughness equivalent to steel H13, while alloy PI 2 has lower
toughness,
but tempering resistance and hardness after nitriding significantly higher
than steel
H13. The alloy should be selected on the basis of the most critical properties
required
for the application. However, in all cases, significant cost reductions can be
obtained
due to the low Mo and V content of the alloys of the present invention.
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