Note: Descriptions are shown in the official language in which they were submitted.
CA 02824238 2013-07-09
WO 2012/095532
PCT/EP2012/050531
HIGH THERMAL DIFFUSIVITY AND HIGH WEAR RESISTANCE TOOL STEEL
Field of the invention
The present invention relates to a tool steel with very high thermal
diffusivity and high
wear resistance, mainly abrasive. This tool steel also shows good
hardenability.
Summary
Tool steels often require a combination of different properties which are
considered
opposed. A typical example can be the yield strength and toughness. For many
metal
shaping industrial applications in which there is a heat extraction from the
manufactured
product which is discontinuous, thermal diffusivity is of extreme importance.
Traditionally, for tool steels, this property has been considered opposed to
hardness and
wear resistance. During plastic injection, hot stamping, even forging, metal
injection,
composite curing and other metal shaping processes, wear resistance and high
or very high
thermal diffusivity are often simultaneously required. For many of these
applications, big
cross-section tools are required, for which hardenability of the material is
also of extreme
importance. Thermal diffusivity (a) is related to other fundamental material
properties like
the bulk density (p), specific heat (cp) and thermal conductivity (X) in the
following way:
= p = cp = a
or if preferable:
a = X / = cp )
Wear in material shaping processes is, primarily, abrasive and adhesive,
although
sometimes other wear mechanisms, like erosive and cavitative, are also
present. To
counteract abrasive wear hard particles are generally required in tool steels,
these are
normally ceramic particles like carbides, nitrides, borides or some
combination of them. In
this way, the volumetric fraction, hardness and morphology of the named hard
particles
will determine the material wear resistance for a given application. Also, the
use hardness
of the tool material is of great importance to determine the material
durability under
abrasive wear conditions. The hard particles morphology determines their
adherence to the
CA 02824238 2013-07-09
WO 2012/095532 - 2 -
PCT/EP2012/050531
matrix and the size of the abrasive exogenous particle that can be
counteracted without
detaching itself from the tool material matrix. The best way to counteract the
adhesive
wear is to use FGM materials (functionally graded materials), normally in the
form of
ceramic coating on the tool material. In this case, it is very important to
provide a good
support for the coating which usually is quite brittle. To provide the coating
with a good
support, the tool material must be hard and have hard particles. In this way,
for some
industrial applications, it is desirable to have a tool material with high
thermal diffusivity
at a relatively high level of hardness and with hard particles in the form of
secondary
carbides, nitrides and/or borides and often also primary hard particles (in
the case to have
to counteract big abrasive particles).
Thermal gradients are the cause of thermal shock and thermal fatigue. In many
applications
steady transmission states are not achieve due to low exposure times or
limited amounts of
energy from the source that causes a temperature gradient. The magnitude of
thermal
gradient for tool materials is also a function of their thermal conductivity
(inverse
proportionality applies to all cases with a sufficiently small Biot number).
Hence, in a specific application with a specific thermal flux density
function, a material
with a superior thermal conductivity is subject to a lower surface loading,
since the
resultant thermal gradient is lower. The same applies when the thermal
expansion
coefficient is lower and the Young's modulus is lower.
Traditionally, in many applications where thermal fatigue is the main failure
mechanism,
as in many casting or light alloy extrusions cases, it is desirable to
maximize conductivity
and toughness (usually fracture toughness and CVN). Steels of the present
invention
prioritize wear resistance and diffusivity to CVN, although it is also
considered very
important for some applications and, therefore, the intention is to try to
also maximize it
but without renouncing to the other two properties. Usually, increasing the
hardness of the
tool steel will decrease both toughness and thermal diffusivity and will
increase wear
resistance. A greater level of diffusivity for a given hardness level has been
achieved for
steels of the present invention, usually together with a good hardenability
and, for some
cases, with an excellent toughness compromise.
For many applications thick tools are used, especially when sufficient
strength is required
CA 02824238 2013-07-09
WO 2012/095532 - 3 -
PCT/EP2012/050531
as for to require a thermal treatment. In this case, it is also very
convenient to have a good
hardenability to be able to achieve the desired hardness level on surface and,
preferably, all
the way to the nucleus. Hardenability is also very interesting for hot work
steels, since it is
much easier to achieve high toughness with a quenched martensite structure
than with a
quenched bainite. Thus, the higher the hardenability the less abruptly the
quenching
cooling will need to be. A sudden cooling is more difficult to achieve and
also more
expensive and, since the forms of tools and components manufactured are often
complex,
can lead to breaking of the parts being heat treated or severe deformation.
Wear resistance and mechanical strength are often inversely proportional to
the toughness.
Thus, it is not easy to get a simultaneous increase in both properties.
Thermal conductivity
is a help in this case, since it allows a great increase of the thermal
fatigue resistance, even
if the CVN has been reduced to increase wear resistance or mechanical
strength.
There are many other desirable properties, if not necessary, for hot work
steels that do not
necessarily influence the longevity of the tool, but their production costs,
like: ease of
machining, welding or repair in general, support provided to the coating,
costs...
The authors have discovered that the problem to simultaneously obtain high
thermal
diffusivity, wear resistance and hardenability, together with good levels of
toughness, can
be solved applying certain rules of composition and thermo-mechanical
treatments within
the following compositional range:
%Ceq = 0.31 - 0.9 % C = 0.31 - 0.9 %N = 0 -0.6 %B = 0 -0.6
%Cr <2.8 %Ni = 0 - 3.8 %Si = 0 - 1.4 %Mn= 0 - 3
%Al = 0 - 2.5 %Mo= 0 - 10 %W = 0 - 12 %Ti = 0 - 2
%Ta = 0 - 3 %Zr = 0 - 3 %Hf = 0 - 3 %V = 0 - 4
%Nb = 0 - 1.5 %Cu = 0 - 2 %Co = 0 - 6 %S = 0 - 1
%Se = 0 - 1 %Te = 0 - 1 %Bi = 0 - 1 %As = 0 - 1
%Sb = 0 - 1 %Ca = 0 - 1,
the rest consisting of iron and unavoidable impurities, wherein
%Ceq = %C + 0.86 * %N + 1.2 * %B
In the present invention it is always the case that:
%Mo +1/2 = %W >3.0
CA 02824238 2013-07-09
WO 2012/095532 - 4 -
PCT/EP2012/050531
Some of the selection rules of the alloy within the range and thermo-
mechanical treatments
required to obtain the desired high thermal diffusivity to a high hardness
level and wear
resistance, are presented in the detailed description of the invention
section. Obviously, a
detailed description of all possible combinations is out of reach. The thermal
diffusivity is
regulated by the mobility of the heat energy carriers, which unfortunately can
not be
correlated to a singular compositional range and a thermo-mechanical
treatment.
In an additional aspect, the invention is related to a process to manufacture
a hot work tool
steel, characterised in that the steel is subjected to a martensitic, bainitic
or martensitic-
bainitic quench with at least one tempering cycle at temperature above 590 C,
so that a
steel having a hardness above 47 HRc with a low scattering structure
characterized by a
diffusivity of 9 mm2/s or more is obtainable. In another embodiment, a steel
having a
hardness above 53 HRc with a low scattering structure characterized by a
diffusivity of 9
mm2/s or more is obtainable. In an additional embodiment of this process, the
steel is
subjected to at least one tempering cycle at temperature above 660 C, so that
a steel having
a hardness of 50 HRc or more with a low scattering structure characterized by
a diffusivity
of 5.8mm2/s or more at 600 C is obtainable.
State of the art
Until the development of high thermal conductivity tool steels (EP 1887096
Al), the only
known way to increase thermal conductivity of a tool steel was keeping its
alloying content
low and, consequently, showing poor mechanical properties, especially at high
temperatures. Tool steels capable of surpassing 42 HRc after a tempering cycle
at 600 C or
more, were considered to be limited to a thermal conductivity of 30W/mK and
thermal
diffusivity of 8 mm2/s and 6.5 mm2/s for hardness above 42HRc and 52 HRc
respectively.
Tool steels of the present invention have a thermal diffusivity above 8 mm2/s
and, often,
above 12 mm2/s for hardness over 52 HRc, and even more than 16 mm2/s for
hardness over
42 HRc, furthermore presenting a very good wear resistance and good
hardenability.
Thermal diffusivity is considered the most relevant thermal property since it
is easier to
measure accurately and because most of the tools are used in cyclic processes,
so that the
thermal diffusivity is much more important for evaluating performance of the
tool than can
be thermal conductivity.
CA 02824238 2013-07-09
WO 2012/095532 - 5 -
PCT/EP2012/050531
Tool steels of the present invention have a wear resistance and hardness
higher than steels
described in EP2236639A1. The latter, on the contrary, show a higher
hardenability in the
perlitic region and higher CVN compared to the tool steels with high thermal
conductivity
of the present invention. Hence, for applications where the main failure
mechanism is
thermal fatigue and no wear is present is better to use steels of EP2236639A1
but, for
applications where wear resistance is important, tool steels of this invention
have great
advantage. Furthermore, the steels of the present invention exhibit higher
thermal
diffusivity for the same level of hardness. This is largely due to the fact
that in
EP2236639A1 carbides of the type of M3Fe3C, where M corresponds to Mo and/or
W, are
almost exclusively used, partly due to the presence of %Ni in the matrix that
penalizes the
thermal diffusivity in favour of hardenability, toughness (CVN) and lower
linear thermal
expansion coefficient. In the present invention there is lower %Ni and
carbides are often
partially replaced by harder carbides, even when the elements forming harder
carbides tend
to be solubilized in the Mo and/or W carbides, as is the case of %V.
The tool steels of the present invention can attain much higher levels of
thermal diffusivity
than the tool steels of W02004/046407 Al, where the high levels of %Cr impose
very
tight restrictions which are not observed, on the compositions to be taken
within the
proposed range and the small process window thereafter during the thermo-
mechanical
processing to attain high levels of carrier mobility.
There are other inventions that may have compositional range overlap but do
not have
anything to do with the present invention since rules for selecting the
composition within
the range and/or thermo-mechanical treatments required to achieve a structure
with a
matrix poor in elements in solid solution with great capacity to disperse heat
energy
carriers and having carbides with a high level of crystalline net perfection,
and
consequently a very low dispersion of heat energy carriers (mainly electrons
and phonons),
are not observed. This could be the case of JP04147706 here the inventors,
seeking an
optimized superficial oxide coating, are using levels of %Cr lower than the
normal ones
(around 0.5%) to allow the mentioned oxidation with some specific treatments
at high
temperature. In the present invention %Cr has the tendency to dissolve in the
W and/or Mo
carbides causing the dispersion of the heat energy carriers and thus their
presence is also
undesirable. This is the only point of coincidence that also, in the case of
JP04147706,
CA 02824238 2013-07-09
WO 2012/095532 - 6 -
PCT/EP2012/050531
does not lead to high thermal diffusivity in any of the examples described. At
an even
lower extent is the case of JP11222650, where the inventors look for the
presence of large
amounts of primary carbides to resist massive wear as is the case for high
speed steel but
with an exceptionally low content of %C to allow cold coining.
Other cases may be misleading because of not making special mention or having
a generic
reference levels of non-functional elements for the application mentioned,
this is often the
case for %Cr, %Si and %Mn. In fact, it is difficult to achieve a low level for
some
elements in steels. For instance, a steel supposedly lacking Cr (0%Cr in
nominal
composition), especially if it is an alloyed steel, will probably have %Cr >
0.3 if the steel is
required, for some reason, to be made of selected scrap. In the case where
normal scrap can
be used, significantly cheaper, a %Cr > 0.5 would be expected. If, for a
composition, the
%Cr is not mentioned then it means that its presence is not considered
important, but
neither its absence. In this case, the content of %Cr does not compel the use
of especial
scraps and, if there are not other elements that require so, then a %Cr > 0.5
can be
expected. Even more important is the placement of this %Cr, which will be
predominantly
dissolved in the carbides if no special measures are taken.
The case of %Si is slightly different, since it is possible to reduce its
content through a
refining process, such as ESR, although, due to the narrow window of the
process in this
case, it is technologically very difficult (and expensive, and therefore it is
only carried out
in the case of seeking a specific functionality) to reduce the %Si below 0.2
and, at the same
time, to reach a low level of inclusions (especially oxides).
There are many tool steels having a composition with the potential of
achieving a high
thermal conductivity and actually do not. This is mainly due to the two
following reasons:
- The thermo-mechanical treatments used do not pursue the maximization of
mobility of
the heat energy carriers. Thermal conductivity is not properly chosen as one
of the main
desirable characteristic or, for materials previously developed, the knowledge
was lacking
on how to attain a desired level of thermal diffusivity before the publication
of EP 1887096
Al, and thus the phases present in the final microstructure are chosen
according to the
optimization of some other properties desirable for the application, generally
a certain
compromise of relevant, to the application, mechanical properties. Thus,
within a
CA 02824238 2013-07-09
WO 2012/095532 - 7 -
PCT/EP2012/050531
composition, often strengthening mechanisms are chosen which are very
detrimental for
thermal diffusivity.
- In the melting, secondary metallurgy or re-melting process, not enough
attention is placed
on what is happening beyond the micrometric and nano-metric scales, and thus
unfavorable atomic scale arrangements take place, not necessarily in all
phases present,
that lead to strong carrier scattering. Again this is mainly due to the lack
of knowledge
before the publication of EP 1887096 Al.
There are several tool steels families that, with their nominal range of
composition, could
have the potential to achieve high thermal diffusivity when the correct
strategy during the
thermo-mechanical process is employed according to the present application and
EP
1887096 Al, but do not end up with compositions capable of developing high or
very high
thermal diffusivity. This is mainly due to the following reasons:
- The ratio of %C and that of carbide formers is not well balanced to be
able to minimize
solid solutions in the metal matrix, especially that of %C, and levels are
provided that
cannot afterwards be properly managed by the thermo-mechanical treatments used
to
pursue the maximization of mobility of the heat energy carriers.
- The nominal levels of certain critical elements are far away from the
real content values
in the embodiment. For instance, this is often the case for %Si and %Cr. While
the nominal
composition can describe a certain level, especially in the case of only upper
bound
descriptions, like %Cr < 1 (or even without mentioning the %Cr, which can lead
to the
erroneous assumption that is 0%) and in the same fashion as often the case %Si
< 0.4, it
ends up by being %Cr > 0.3 and %Si > 0.25. This applies also for all trace
elements with a
strong influence on the conductivity of the matrix and even more those with a
high
solubility in carbides and great potential for distortion of the carbides
structure. Usually,
with the exception of %Ni and for some applications the %Mn, no element is
desirable in
solid solution with the matrix at a level higher than 0.5%. Preferably, the
percentage of
these, individually in solid solution, should not exceed 0.2%. If the main
purpose of the
application is to maximize the thermal conductivity, then any metallic element
in solid
solution with the matrix (obviously including transition metals), with the
exception of %Ni
and in some cases the %C and %Mn, should not exceed 0.1% or, even better,
0.05%.
CA 02824238 2013-07-09
WO 2012/095532 - 8 -
PCT/EP2012/050531
Detailed description of the invention
To obtain tool steels with high thermal diffusivity and wear resistance to
high hardness
levels with good hardenability, it has been observed that, within the
compositional range
specified above, a number of rules and general considerations in the selection
of the
composition within the range and the thermo-mechanical treatments to be used,
some of
which are described below, have to be taken into account. Thermal diffusivity
is a
consequence of the scattering mechanisms on the phases present for all carrier
types
present. The perfection of the lattice plays an important role, but also other
scattering
mechanisms are of relevance. In this document the thermal diffusivity itself
will be used as
a measurement of the structure attained. Within a same chemical composition
different
structures can be attained and thus also different levels of thermal
diffusivity.
Tool steels of the present invention excel mainly because of their high
thermal diffusivity
and wear resistance. Wear resistance and toughness tend to be inversely
proportional,
although different microstructures reach different relationships, i.e., as a
function of
microstructure different levels of toughness for the same elastic limit and
hardness at a
given temperature can be reached and, for a specific type of material,
hardness tends to
correlate with wear resistance unless the volume fracture or the morphology of
wear
resistant particles is significantly changed. In this vein, it is well known
that for most tool
steels with medium carbon content, pure microstructure of tempered martensite
is the only
one that offers the best compromise of mechanical properties. This means that
it is
important to avoid the formation of other microstructures like stable ferrite-
perlite or
metastable bainite during cooling after the process of austenitization of the
heat treatment.
Therefore, fast cooling rates will be needed and, if higher hardenability is
required, some
alloying elements to delay the kinetics of the formation of these more stable
structures
should be used. From all possible alternatives those with less negative
effects on thermal
diffusivity should be used.
A strategy to obtain wear resistance and higher elastic limit at high
temperatures and, at the
same time, obtain high thermal conductivity, is the use of carbides with high
electron
density, as secondary carbides of the M3Fe3C type and sometimes even primary
carbides
(M- should only be Mo or W for a greater thermal conductivity). There are
other carbide
CA 02824238 2013-07-09
WO 2012/095532 - 9 -
PCT/EP2012/050531
types (Mo, W, Fe) with high electron densities and with tendency to solidify
with a good
crystalline perfection. Some elements like Zr and Hf and, at a less extent, Ta
for instance
comparing to Cr, when dissolving with this type of carbides do not provide
much distortion
to the crystalline structure and dispersion of charge carriers is small and so
is the effect on
thermal conductivity. Moreover, these high carbide forming elements tend to
form separate
MC type carbides, due to its high affinity for C.
In fact, in the present invention it has been observed that the effect can be
quite positive if
a moderate quantity of %V is used and it is balanced with the presence of
strong carbide
former (preferably Zr and/or Hf). It has been seen that there can be amounts
of %V up to
0.9 with practically no formation of primary carbides (obviously depending on
the Ceq and
the presence of other carbides, and for higher contents of Ceq is necessary to
reduce the
percentage of V at a maximum of 0.8 and even 0.5 or 0.4 to avoid the presence
of primary
carbides or massive dissolution in them) and with little dissolution in the
carbides of (Fe,
Mo, W), especially if used simultaneously with strong carbide forming
elements, also there
is a displacement of more carbon out of the matrix with the consequent benefit
to the
overall thermal diffusivity (in this case, the benefit is remarkable with
%Hf+%Zr+%Ta
greater than 0.1, and very significant if it exceeds 0.4 or 0.6, depending on
the quantities of
% Ceq and %V present). In fact, this combination is highly desirable as the
percentage of
V as the percentage of Zr, Hf and Ta tend to significantly improve the wear
resistance
compared to a steel that has only carbides (Fe, Mo, W), the same applied for
%Nb. The
effect becomes noticeable with %V = 0.1 and remarkable with %V = 0.3 or 0.5,
depending
on the level of %Ceq. If extreme wear resistance with the presence of primary
carbides is
to be achieved, as is the case in applications with large abrasive particles
such as in hot
stamping of uncoated sheet, then larger amounts of %V can be used, up to 1.5%
or even
2% is possible while maintaining a good level of thermal diffusivity,
especially if
compensated with strong carbide forming elements. In this case, it can be
convenient to
have high levels of strong carbide forming elements combined with %V
(%V+%Nb+%Hf+%Zr), above 1.2 or even 2.0 in weight percentage (for applications
where a good wear resistance is needed, even 3.0, but then the cost of the
alloy is
increased). In this case, rarely any strong carbide forming element (%V, %Nb,
%Ta, %Zr,
%Hf) will individually exceed 3%, with the exception of %V where the upper
limit is
usually 4% in weight (for applications where wear resistance is priorized at
the expense of
losing thermal diffusivity), or 1.8% for applications requiring very high
thermal diffusivity
CA 02824238 2013-C
WO 2012/095532 - 10 -
PCT/EP2012/050531
and Nb that, due to its negative effect on thermal diffusivity, tends to be
used only to
control grain size and when used as primary carbide former will rarely be
above 1.5%. It is
desirable to have most of the strong carbide formers bound in the carbides and
not
dissolved in the matrix, thus the level of %Ceq has to be finely adjusted as
explained later
to minimize both the amount of strong carbide formers and %Ceq in solid
solution. As an
example in most applications of this invention if %Ceq is smaller than 0.35
then %V
should be kept below 1.7%. In general it is desired to mostly have Fe, Mo and
W carbides
(where obviously part of the C can be replaced by N or B), usually more that
60% and,
optimally, more than 80% or even more than 90% of these type of carbides. The
dissolution of other metallic elements of these types of carbides (obviously
in the case of
carbides metallic elements are mainly transition elements) can exist, but it
is desired to be
small to guarantee a high phononic conductivity. Normally no other metallic
element, apart
from the principal Fe, Mo and W, should exceed 20% of the weight of all
metallic
elements of the carbide, for this type of mainly desired carbides. Preferably
should not be
more than 15% and even better a 5%. This is because they tend to form
structures with
densities of solidification defects extremely low even for fast solidification
kinetics
(therefore less structural elements to cause dispersion of carriers).
As discussed before, the only exception is the presence of a limited quantity
of strong
carbide forming elements, although the formation of independent carbides is
preferable. In
this case, Mo and W provide sufficient obstacles for the formation of stable
structures
(perlite and ferrite), although the formation of bainite is very fast. In some
steels
superbainitic structures can be formed applying a martempering heat treatment,
which
consists in the complete solubilisation of alloying elements followed of a
rapid cooling to a
specific temperature (to avoid ferrite formation) in the range of lower
bainitic formation
and an extended temperature maintenance to obtain a 100% bainitic structure.
For the
majority of the steels a pure martensitic structure is desirable, so that in
this system some
bainitic transformation delaying elements must be added, since Mo and W are
very
inefficient in this respect. Generally, for this purpose %Cr is commonly used,
but has an
extremely negative effect on the thermal conductivity for this system because
it dissolves
in the M3Fe3C carbides and causes a great distortion, so it's much better to
use strong
carbide forming elements and non soluble elements in carbides. These last
elements will
reduce the conductivity of the matrix and, thus, the ones with the minimum
negative effect
should be used. Accordingly, the natural candidate is Ni, but at the same time
others can be
CA 02824238 2013-07
WO 2012/095532 - 11 -
PCT/EP2012/050531
used (a special mention should be made to %B, for its effect with very low
concentrations).
Since in the present invention carbide formers with great affinity to C are
tend to be used
for their positive effect on wear resistance, the necessary and desirable
quantity of delaying
elements of the transformation kinetics to stable structures during quenching
is lower.
Usually, a quantity up to 1% in weight, and for large sections up to 3.0%,
will be enough to
get sufficient hardenability and contribute to the increase of toughness
without an
excessive detriment of conductivity. Higher %Ni quantities provide more
toughness and a
reduction in the linear thermal expanded coefficient, but the priority of the
present
invention is a combination of wear resistance with thermal diffusivity, thus,
only for some
special applications the strategy of using high contents of %Ni, with a
maximum of 3.8%,
can be used. There are applications where lower amounts of %Ni already lead to
the
desired effect, especially if the contents of %Mn and/or %Si are a bit higher
(%Mn usually
does not exceed 3%) or sections of the material used are smaller.
The use of %Mo as a single carbide former (obviously together with Fe), is
advantageous
when maximising thermal conductivity, but it has the disadvantage to provide a
higher
thermal expansion coefficient and, thus, it decreases the thermal fatigue
resistance. Hence
it is preferable to have a relation of 1.2 to 3 times more Mo than W, but not
the absence of
W. The exception are the applications where only thermal conductivity is to be
maximised
together with toughness, but not particularly thermal fatigue resistance.
Hardenability and
the alloy cost, due to the high volatility of Mo and W prices, can lead to
changing
preferences regarding the %W being the main element in %Moeq, (where %Moeq =
%Mo +
High contents of Moeq can be used with high levels of Ceq, resulting in an
increased cost
alloy, low toughness, very difficult to weld, complicated hardenability for
large parts and
limited machinability. But very high levels of wear resistance with good
thermal
diffusivity can be achieved. For applications where the highlighted drawbacks
are not
determinant these can be alloys of interest. This can be the case for some
cutting
applications. Here, levels of Ceq usually superiors to 0.5% are used and,
often, even over
0.6%. Levels of % Moeq are often above 5% and frequently above 6% or even 9%.
Also
the limits of the Moeq/Ceq ratio are shifted to superior levels compared to
the rest of the
alloys of the present invention. Values higher than 16 are possible and,
higher than 13, are
probable.
CA 02824238 2013-07-09
WO 2012/095532 - 12 -
PCT/EP2012/050531
In all document the term carbide is referring to the primary carbides as well
as the
secondary, unless otherwise specified.
The more restrictive the %Si and %Cr the higher the thermal conductivity,
although the
solution is more expensive (also, some properties, that could be important for
some
applications, and thus would be desirable to be maintained, could get worse
with the
reduction of these elements below certain levels as, for instance, for
toughness due to oxide
inclusions in the case Al, Ti, Si and any other deoxidizing, are used in
insufficient
quantities or, in some cases of corrosion resistance, if %Cr or %Si are too
low). Thus, often
there must be a compromise between increase of costs, toughness reduction,
wear
resistance or other relevant properties for certain applications and the
benefit of higher
thermal conductivity. Maximum thermal conductivity can be obtained only if
levels of %Si
and %Cr are below 0.1% or, even better, if below 0.05%. To maximise thermal
diffusivity,
also levels of the rest of the elements, with the exception of %C, %Mo, %W,
%V, %Zr,
%Hf, %Ta, %Nb and in some instances %Mn and %Ni, must be as low as possible
(below
0.05 is technically possible with an acceptable cost for most of the
applications, although a
maximum of 0.1 is, of course, less expensive). For some applications in which
toughness is
especially important less restrictive levels of %Si must be employed (is the
least
detrimental to the thermal conductivity of all iron deoxidants elements) and
thus give up to
some thermal conductivity, to ensure the inclusion level is not too high.
Depending on the
levels of %C, %Mo and %W used, there can be sufficient hardenability,
especially in the
perlitic zone. For cases of large components, where it is not possible to
avoid the formation
of bainite during quenching, the use of elements in solid solution to prevent
the formation
of coarse cementite precipitates (Fe3C) that entail very low toughness, such
as %Al and
%Si, may be interesting. Generally below 0.4, exceptionally with levels of
around 1% and,
very exceptionally, above 2% and for the %Al, up to a maximum of 2.5%. The
levels of
%Mo, %W and %C used to obtain the desired mechanical properties must be
balanced to
achieve a high thermal conductivity, so that within the matrix remain the
least amount of
these elements in solid solution. The same applies for the rest of carbide
formers that could
be used to obtain a certain tribological response (like %V, %Zr, %Hf,
%Ta,...).
For some applications some environmental resistance can be of interest and,
thus, be
desirable to have some %Cr or %Si in solid solution (oxidation resistance to
high
temperature). The negative effect on thermal diffusivity can be moderated
through carbon
CA 02824238 2013-07-09
WO 2012/095532 - 13 -
PCT/EP2012/050531
fixing with stronger carbide formers elements. Without the later, %Cr should
not exceed
2% and, preferable, 1.5%. Although in the presence of V, %Nb, %Ta, %Zr and
%Hf, and
preferably the last two or three, levels close to 3% of Cr can be achieved
maintaining a
good thermal diffusivity, and even 1.4% for the case of Si. In fact for most
applications
%Cr < 2.8% is required if the thermal diffusivity needs to be high. Many
compositions
require %Cr<2.5 % to be able to attain high thermal diffusivity with the
proper thermo-
mechanical processing (which is composition dependent, as explained). At this
level the
environmental protection effect is only somewhat noticeable if the %Cr is
mainly left in
solid solution in the matrix. Finally, a much greater range of compositions
can attain high
thermal diffusivity when the proper thermo-mechanical treatment is applied, if
%Cr is
restricted to remain below 1.9%.
The simplest compositional rule to describe the compositions within the range
that are
capable of attaining a high thermal diffusivity simultaneously to a high wear
resistance can
be based on a ratio R = Moeq / Ceq, where %Moeq = %Mo + 1/2 = %W and %Ceq = %C
+
0.86 * %N + 1.2 * %B. This rule applies only for big enough contents of %Ceq
(normally
0.32 min, preferably 0.35 min and most accurately when 0.38 minimum %Ceq) and
%Moeq
(normally 3.2 min, preferably 3.4 min and most accurately when 3.6 minimum
%Moeq). It
is also a rule that can only be used for lower %Cr contents, normally %Cr <
2.5%, and
desirably %Cr < 1.9%. The minimum value for R results when computing the %Moeq
minimum for the rule to apply divided by 0.9 which is the maximum %Ceq for the
present
invention (for example for a minimum Moeq = 3.2 then the minimum R value
results to be
3.56). The maximum value for R has been observed to be possibly 11.5,
preferably 10.8
and optimally 10.5 for low %Ceq values. Low %Ceq values are for this rule
those under
0.35%, occasionally under 0.36% or even under 0.37%. For high %Ceq values the
maximum value for R has been observed to be possibly 16.8, preferably 16.0 and
optimally
15. High %Ceq values are for this rule those above 0.38%, occasionally above
0.40% or
even above 0.45%. For intermediate values of %Ceq, the maximum value for R has
been
observed to be 14, preferably 13, and optimally 12.
Generally, to solely maximize thermal diffusivity (i.e. there are not other
properties of
great importance), it is convenient to observe the following alloying rule (to
minimize the
%C in solid solution), if a tempered martensite or bainite microstructure
withstanding
mechanical requirements wants to be obtained. The formula must be corrected if
carbide
CA 02824238 2013-07-09
WO 2012/095532 - 14 -
PCT/EP2012/050531
formers with high affinity for the %C (like Hf, Zr or Ta, even Nb) are used.
It must be also
modified if %Cr>0.2 or Moeq > 7:
0.02 < xCeq - solC ¨ AC = [(xMo-solMo) / (3. AMo) + (xW-solW) / (3. AW) + (xV-
solV) /
AV ] > 0.265
where:
xCeq - carbon weight percentage;
xMo - molybdenum weight percentage;
xW - tungsten weight percentage;
xV - vanadium weight percentage;
AC - carbon atomic mass (12.0107 u);
AMo- molybdenum atomic mass (95.94 u);
AW - tungsten atomic mass (183.84 u);
AV - vanadium atomic mass (50.9415 u);
solC - carbon percentage in solid solution;
solMo - molybdenum percentage in solid solution;
solW - tungsten percentage in solid solution;
solV - vanadium percentage in solid solution.
For an even higher thermal conductivity it is even more desirable to have:
0.04 < xCeq - solC - AC = [(xMo-solMo) / (3. AMo) + (xW-solW) / (3. AW) + (xV-
solV) /
AV] > 0.22
And still better:
0.09 < xCeq - solC ¨ AC = [(xMo-solMo) / (3. AMo) + (xW-solW) / (3. AW) + (xV-
solV) /
AV] >0.18
To compensate for the presence of other %C avid carbide formers, an extra term
must be
added to the formula for each type of %C avid carbide former:
-AC*xM/(R*AM)
where:
xM ¨ carbide former weight percentage;
AC ¨ carbon atomic mass (12.0107 u);
R ¨ number of carbide former units per carbide unit (for example: 1 if the
carbide type is
MC, 23/7 if the carbide type would be M23C7...)
CA 02824238 2013-07-09
WO 2012/095532 - 15 -
PCT/EP2012/050531
AM - carbide former atomic mass.
This balance provides an extraordinary thermal conductivity if the reinforcing
ceramic
particles formers, including the non metallic part (%C, %B and %N), are taken
into the
carbides (as an alternative nitrides, borides and intermediate substances).
Then, the
appropriated thermal treatment must be applied. This thermal treatment will
have a phase
in which most of the elements will be dissolved (austenitization to
sufficiently high
temperature, usually around 1080 C for moderated Moeq levels, 1120 C for
medium levels
of Moeq and 1240 C for high levels of Moeq, exceptionaly, if distortion of the
heat
treatment is of great importance for the application, lower austenitization
temperatures can
be used). An abrupt cooling will follow, its intensity will be determined by
the desired
mechanical properties, although stable structures should be avoided since
phases with big
quantities of %C and carbide formers in solid solution are implied. Metastable
microstructures are even worse, since the microstructure distortion caused by
carbon is
even greater, hence thermal conductivity is lower, although once these
metastable
structures have relaxed the carbide formers place themselves in the desired
position.
Martensite and bainite tempered following this procedure will be the desired
microstructures for this case. The largest possible carbide substitution of Fe
by Mo, W and
all carbide forming elements with greater affinity for carbon other than Cr
are desired, so
the tempering strategy selected has a great influence in the final thermal
conductivity, with
particular relevance to the final tempering temperature and minimum tempering
temperature. For hardness over 40 HRc, the highest possible temperature is
desirable for
the last tempering if thermal diffusivity is to be maximized, and this
approach is used to set
the intermediate tempering strategy. That is, the same final hardness level
can be achieved
with different sequences of tempering and the one using a higher final
tempering
temperature is chosen, if the only objective is to maximize the thermal
diffusivity at a
certain level of hardness. So, usually, unusually high final tempering
temperatures end up
being used, often above 600 C, even when hardness over 50 HRc are chosen. In
steels of
the present invention it is usual to achieve hardness of 47 HRc, even more
than 52 HRc,
and often more than 53 HRc and with the embodiments regarded as particularly
advantageous due to their wear resistance, hardness above 54HRc, and often
more than 56
HRc are possible with even one tempering cycle above 590 C, giving a low
scattering
structure characterized by a thermal diffusivity greater than 8 mm2/s and,
generally, more
than 9 mm2/s, or even more than 10 mm2/s, when particularly well executed then
greater
CA 02824238 2013-0=
WO 2012/095532 - 16 -
PCT/EP2012/050531
than 11 mm2/s, even greater than 12 mm2/s an occasionally above 12,5 mm2/s. As
well as
achieving hardness greater than 42 HRc, even more than 50 HRc with the last
tempering
cycle above 600 C, often above 640 C, and sometimes even above 660 C,
presenting a
low scattering structure characterized by a thermal diffusivity higher than 10
mm2/s, or
even than 12 mm2/s, when particularly well executed then greater than 14
mm2/s, even
greater than 15 mm2/s and occasionally above 16 mm2/s. Those alloys can
present even
higher hardness with lowering tempering temperatures, but for most of the
intended
applications a high tempering resistance is very desirable. As can be seen in
the examples
with some very particular embodiments with high carbon and high alloying,
leading to a
high volume fraction of hard particles, hardness above 60 HRc with low
scattering
structures characterized by thermal diffusivity above 8mm2/s and generally
more than
9mm2/s are possible in the present invention.
To attain the high levels of hardness and wear resistance often desirable in
the present
invention, considerably high levels of the volume fraction of hard particles
have to be used.
The volume fraction of hard particles (carbides, nitrides, borides and
mixtures thereof) is
often above a 4% preferably above a 5.5% and for some high wear applications,
even
above a 9%. Size of primary hard particles is very important to have an
effective wear
resistance and yet not excessively small toughness. The inventors have
observed that for a
given volume fraction of hard particles overall resilience of the material
diminishes as the
size of the hard particles increases, as would be expected. A bit more
surprisingly it has
also been observed that when the size of hard particles is increased, the
overall fracture
toughness increases if the fracture toughness of the particles themselves is
maintained.
When it comes to abrasive wear resistance it has been observed the existence
of a critical
hard particle size, below which the hard particle is not effective against the
abrasive agent.
This critical size depends on the size of the abrasive agent and the normal
pressure. For
some applications where the abrasive particles are of small size (normally
below 20
microns), it can be desirable to have primary hard particles smaller than 10
microns or
even smaller than 6 microns, but in any case with an average size not smaller
than 1
micron. For applications where big abrasive particles cause the wear, big
primary hard
particles will be desirable. Therefore, for some applications it is desirable
to have some
primary hard particles bigger than 12 microns, often greater than 20 microns
and for some
particular applications even greater than 42 microns.
CA 02824238 2013-07-09
WO 2012/095532 - 17 -
PCT/EP2012/050531
For applications where mechanical strength more than wear resistance are
important, and it
is desirable to attain such mechanical strength without compromising all too
much
toughness, the volume fraction of small secondary hard particles is of great
importance.
Small secondary hard particles, in this document, are those with a maximum
equivalent
diameter (diameter of a circle with equivalent surface as the cross section
with maximum
surface on the hard particle) below 7.5 nm. It is then desirable to have a
volume fraction of
small secondary hard particles for such applications above 0.5%. It is
believed that a
saturation of mechanical properties for hot work applications occurs at around
0.6%, but
ithas been observed by the inventors that for some applications requiring high
plastic
deformation resistance at somewhat lower temperatures it is advantageous to
have higher
amounts than these 0.6%, often more than 0.8% and even more than 0.94%. Since
the
morphology (including size) and volume fraction of secondary carbides change
with heat
treatment, the values presented here describe attainable values with proper
heat treatment.
Cobalt has often been used in hot work tool steels principally due to the
increase in
mechanical strength, and in particular the increase of yield strength
maintained up to quite
high temperatures. This increase in yield strength is attained trough solid
solution and thus
it has a quite negative effect in the toughness. The common amounts of Co used
for this
propose is 3%. Besides the negative effect in toughness it is also well known
the negative
effect in the thermal conductivity. The inventors have seen that within the
compositional
ranges of the present invention it is possible to use Co, and attain an
improved yield
strength/ toughness relation since Co can promote the nucleation of secondary
hard
particles and thus keep their size small. It has also been seen that for some
compositions of
the present invention, when adding Co the Thermal diffusivity does indeed
decrease at
room temperature, but then can actually increase at higher temperatures
(normally above
400 C) if the correct thermo-mechanical treatment is applied. The inventors
have seen
that the best results are encountered when %Co is above 1.3%, preferably above
1.5% and
optimally above 2.4%. Also %C should exceed 3.2%, preferably 3.4% qnd
optimally 3.6%.
If thermal conductivity at high temperatures is of outmost importance for the
application a
special care has to be taken not to have excessive %V, it should be kept below
2.8%,
preferably below 2.3% and optimally below 1.7%. Finally %Moeq should normally
exceed
3.3% often 3.5% and even 4.0%. Heat treatment has to be selected with a rather
high
austenitization temperature and an abnormally high tempering temperaures,
actually more
than 55 HRc commonly achieved with at least one tempering cycle at 630 C or
even
CA 02824238 2013-C"
WO 2012/095532 - 18 -
PCT/EP2012/050531
above, 50 HRc can be maintained even with one tempering cycle at 660 C or
more.
Proper thermo-mechanical processing together with the compositional rules just
explained
have to be implemented to minimize scattering at high temperatures, the
optimized
arrangements is characterized by providing diffusivities of more than 5.8
mm2/s, often
more than 6.1mm2/s and even more than 6.5mm2/s at measuring temperatures as
high as
600 C.
When mainly remaining in the carbide system MoxW3,Fe3C, one of the preferred
ways to
balance the contents of %W, %Mo and %C in the present invention is through the
adhesion
to the following alloying rule:
%Ceq = 0.4+(%Moeq(rea0-4)*0.04173
where: MOeq(real) = %M0 (AM0/AW)* %W.
with:
AMo - molybdenum atomic mass (95.94 u);
AW - tungsten atomic mass (183.84 u);
so that, at the end:
Moeq(reao = %Mo + 0.52 *%W.
If the expression is normalized in a parameter K = %Ceq (0.4+(%Moeq(rea0-
4)*0.04173),
the desirable values for this parameter, for the present invention, are as
follows:
It has been observed that when carbon content is low (that is to say
%Ceq<0.39, preferably
%Ceq<0.36 and optimally %Ceq<0.35), the parameter K should exceed 0.75,
preferably
0.76, more preferably 0.86 and optimally 0.88. In fact for some embodiments
for
applications requiring very high wear resistance, K will normally be higher
than 0.92. A
very good performance will be obtained as already described, at the expense of
a higher
cost, when adding elements that strongly bond carbon to the carbides. In the
case here
dealing with low %Ceq it is especially desirable that the added amount of %Hf,
%Zr, %Ta
and %Nb exceed 0.07%, preferably 0.09% and optimally 0.1%. Given that Nb can
be quite
detrimental for the thermal diffusivity for some applications it will not be
desirable (%Nb
<0.09) and then the contents of Hf, Zr and Ta in the sum should exceed 0.01%,
preferably
0.07% and in applications requiring high wear resistance with very high
thermal diffusivity
and where Zr is chosen as the main former of hard carbides, then contents
above 0.14%,
preferably above 0.2% and even above 0.4%, will be desirable. In these cases
with the
CA 02824238 2013-01
WO 2012/095532 - 19 -
PCT/EP2012/050531
presence of strong carbon binders the restrictions on K can be relaxed around
a 3% to 5%
for alloys with low carbon content as hereby described.
On the other hand, when carbon content is not low (that is to say %Ceq>=0.39,
preferably
%Ceq>=0.36 and optimally %Ceq>=0.35), the parameter K should exceed 0.6,
preferably
0.75, more preferably 0.84 and optimally 0.87. In this case if elements are
used that
strongly bond carbon (nitrogen or boron) to the carbides (nitrides, borides or
mixtures) in
the fashion described in the last paragraph, then the restrictions on K can be
relaxed very
severely, for some applications even eliminated.
The authors have observed that good combinations of wear resistance and
thermal
diffusivity can be obtained for very high values of K if all other alloying
and thermal
processing rules are observed, normally in their most stringed version, but of
course the
best results are obtained when K does not exceed 3, preferably 1.5 and
optimally 1.3.
An especially interesting embodiment, when the main goal for the chosen
application is the
maximization of the thermal diffusivity to the highest possible level of
hardness, arises
when applying this alloying rule together with very low levels of %Cr,
especially in
dissolution with the carbides, as described above.
It has also been observed by the authors that it is possible to attain
considerably high
thermal diffusivity and wear resistance when using much higher levels of %Mo
and %W
than described in the last couple paragraphs. The level of thermal diffusivity
for a given
hardness level cannot be optimized to such high values as when applying the
previously
described alloying rules. On top this comes at a considerably higher cost, so
obviously is
not the preferred way for most applications, but in can be advantageous for
some very
concrete cases. For example if a special oxidation color is desirable, or when
ferrite/perlite
hardenability wants to be extended and the usage of other much more effective
elements is
nor recommendable. In such case the parameter K has to be selected to be quite
low,
indeed it should be lower than 0.81, preferably lower than 0.79 and optimally
lower than
0.75. This has to happen for large enough values of %Ceq, normally larger than
0.33, even
larger than 0.35 and occasionally larger than 0.41.
CA 02824238 2013-0
WO 2012/095532 - 20 -
PCT/EP2012/050531
The teachings of this inventions can be applied to the described compositional
range for
alloys with a %Moeq > 3.0%. To be more precise it can be described in terms of
%Moeq(real), in which case for most applications the teachings work for values
superior to
3.3%, and even more generalized in terms of applications for values of
%Moeq(real) >
3.6% and when %Moeq(real) > 3.8% then the density of compositions which can
attain a
high thermal diffusivity and wear resistance within the range is significantly
greater, and
covers most applications (one exemption is for example applications with
exceptionally
high hardness or wear resistance). In the same way when it comes to %Ceq,
while the
teachings of the present invention work already for values higher than 0.31%,
when %Ceq
> 0.33%,and even more for %Ceq >0.36% the density of compositions which can
attain a
high thermal diffusivity and wear resistance within the range is significantly
greater, and
covers most applications (one exemption is for example applications with
exceptionally
high hardness or wear resistance).
To increase machinability S, As, Te, Bi or even Pb, Ca, Cu, Se, Sb or others
can be used,
with a maximum content of 1%, with the exception of Cu, than can even be of
2%. The
most common substance, sulfur, has, in comparison, a light negative effect on
the matrix
thermal conductivity in the normally used levels to increase machinability.
However, its
presence must be balanced with Mn, in an attempt to have everything in the
form of
spherical manganese bisulphide, less detrimental for toughness, as well as the
least
possible amount of the remaining two elements in solid solution in case that
thermal
conductivity needs to be maximized.
Another hardening mechanism can be used in order to search for some specific
combination of mechanical properties or environmental degradation resistance.
It is always
the intention to maximize the desired property, but trying to have minimal
possible adverse
impact on thermal conductivity. Solid solution with Cu, Mn, Ni, Co, Si, etc...
(including
some carbide formers with less affinity to carbon, like Cr) and interstitial
solid solution
(mainly with C, N and B). For this purpose, precipitation can also be used,
with an
intermetallic formation like Ni3Mo, NiAl, Ni3Ti... (also of Ni and Mo, small
quantities of
Al and Ti can be added, but special care must be taken for Ti, since it
dissolves in M3Fe3C
carbides and a 2% should be used as a maximum). Finally, other carbide types
can also be
used, but it is usually difficult to maintain high levels of thermal
conductivity, unless
carbide formers present a very high affinity with carbon, as it has been
described
CA 02824238 2013-07-09
WO 2012/095532 - 21 -
PCT/EP2012/050531
throughout this document. Co can be used as a hardener by solid solution or as
a catalyst of
Ni intermetallic precipitation, rarely in contents higher than 6%. Some of
these elements
are also not as harmful when dissolved in M3Fe3C carbides, or other carbides
of (Fe, Mo,
W), this is specially the case for Zr and Hf and, to a lesser extent, for Ta,
these can also
limit V and Nb solubility.
When amounts are measured in weight percentage, atomic mass and the formed
type of
carbide determine if the quantity of a used element should be big or small.
So, for instance,
2%V is much more than 4%W. V tends to form MC carbides, unless it dissolves in
other
existing carbides. Thus, to form a carbide unit only a unit of V is needed,
and the atomic
mass is 50.9415. W tends to form M3Fe3C carbides in hot work steels. So three
units of W
are needed to form a carbide unit, and the atomic mass is 183.85. Therefore,
5.4 more
times carbide units can be formed with 2%V than with 4%W.
Tool steel of the present invention can be manufactured with any metallurgical
process,
among which the most common are sand casting, lost wax casting, continuous
casting,
melting in electric furnace, vacuum induction melting. Powder metallurgy
processes can
also be used along with any type of atomization and eventually subsequent
compacting as
the HIP, CIP, cold or hot pressing, sintering (with or without a liquid phase
and regardless
of the way the sintering process takes place, whether simultaneously in the
whole material,
layer by layer or localized), laser cusing, spray forming, thermal spray or
heat coating, cold
spray to name a few of them. The alloy can be directly obtained with the
desired shape or
can be improved by other metallurgical processes. Any refining metallurgical
process can
be applied, like VD, ESR, AOD, VAR... Forging or rolling are frequently used
to increase
toughness, even three-dimensional forging of blocks. Tool steel of the present
invention
can be obtained in the form of bar, wire or powder (amongst others to be used
as solder or
welding alloy). Even, a low-cost alloy steel matrix can be manufactured and
applying steel
of the present invention in critical parts of the matrix by welding rod or
wire made from
steel of the present invention. Also laser, plasma or electron beam welding
can be
conducted using powder or wire made of steel of the present invention. The
steel of the
present invention could also be used with a thermal spraying technique to
apply in parts of
the surface of another material. Obviously the steel of the present invention
can be used as
part of a composite material, for example when embedded as a separate phase,
or obtained
as one of the phases in a multiphase material. Also when used as a matrix in
which other
CA 02824238 2013-07-09
WO 2012/095532 - 22 -
PCT/EP2012/050531
phases or particles are embedded whatever the method of conducting the mixture
(for
instance, mechanical mixing, attrition, projection with two or more hoppers of
different
materials...).
Tool steel of the present invention can also be used for the manufacturing of
parts under
high thermo-mechanical loads and wear resistance or, basically, of any part
susceptible to
failure due to wear and thermal fatigue, or with requirements for high wear
resistance and
which takes advantage of its high thermal conductivity. The advantage is a
faster heat
transport or a reduced working temperature. As an example: components for
combustion
engines (such as rings of the engine block), reactors (also in the chemical
industry), heat
exchange devices, generators or, in general, any power processing machine.
Dies for
forging (open or closed die), extrusion, rolling, casting and metal
thixoforming. Dies for
plastic forming of thermoplastics and thermosets in all of its forms. In
general, any matrix,
tool or part can benefit from increased wear resistance and thermal fatigue.
Also dies, tools
or parts that benefit from better thermal management, as is the case of
material forming or
cutting dies with release of large amounts of energy (such as stainless steel
or TRIP steels)
or working at high temperatures (hot cutting, hot forming of sheet).
Additional embodiments are described in the dependent claims.
EXAMPLES
Some examples indicate the way in which the steel composition of the invention
can be
specified with higher precision for different hot working applications:
Example 1
Dies for the stamping or press hardening of sheet. In this case maximum
possible thermal
diffusivity is desired at high hardness. The desired wear resistance depends
on the sheet
coating.
- Sheets coated with Zn, AlSi or other inorganic coatings (the same
compositions are
optimized for the manufacture of injection molds for thermoplastics,
especially when steels
described below are made by powder metallurgy):
CA 02824238 2013-07-09
WO 2012/095532 - 23 -
PCT/EP2012/050531
For this purpose in the context of the present invention the following
compositional range
can be used:
Ceq: 0.3 - 0.6 Cr < 3.0% (preferably Cr < 0.1%)
V: 0 - 0.9% (preferably 0.3 - 0.8%)
Si: <0.15% (preferably %Si <0.1, but with an acceptable level of oxide
inclusions)
Mn: <0.5% Moeq: 3.5 - 5.5
where Moeq= %Mo+1/2 %W and
Ceq= %C + 0.86 * %N+ 1.2 * %B
The rest of the elements should be kept as low as possible and, in any case,
always be
below 0.15%, with the exception of strong carbide formers (%Ta, %Zr, %Hf).
All values are given in weight percentage.
The following three examples show properties that can be obtained:
Hardness Therm. Diff.
%C /01Mo %W %V %Cr %Si %Mn Other
HRc mm2/s
0.40 3.6 1.4 0.3 <0.01 <0.05 <0.01 52 11.47
0.45 1.6 4.5 0.4 <0.01 <0.05 <0.01 52-53 10.96
0.41 3.5 1.4 0.8 1.3 <0.05 <0.01 50 9.32
* In all cases heat treatment which maximizes diffusivity at the indicated
hardness has been applied,
minimizing the presence of elements in solution with the matrix, except for %
Cr, especially minimizing the
presence of %C and, to a lesser extent, %V in the matrix. In all cases this
means very high austenitization
temperatures, from 3 to 5 tempering cycles, with the latest in the range 600-
640 C.
An advanced optimization is obtained when elements strongly reacting with %C
to form
carbides (also %N and %B) are employed. Several examples show the properties
that can
be obtained:
Max. Hard.
Therm.
Hardness
%C %Mo %W %V %Cr %Si %Mn
Other /Tem. T diff.
HRe
HRe / C
mm2/s
0.50 3.6 1.4 0.5 <0.01 <0.05 <0.01 Hf, Zr, Nb
56-57 12.73
0.50 3.6 1.4 0.5 <0.01 <0.05 <0.01 Hf, Zr, Nb
54 13.93
NbZr, ,
0.33 3.36 1.91 <0.01 <0.01 <0.05 0.4 Hf, 50
13.04
B=0.16
NbZr, ,
0.33 3.36 1.91 <0.01 <0.01 <0.05 0.4 Hf, 43
16.62
B=0.16
CA 02824238 2013-07-
WO 2012/095532 - 24 -
PCT/EP2012/050531
3,
0.42 3.45 1.4 0.6 2.2 <0.05 <0.01 Hf=0. 52
10.42
Zr=0.2
Hf, Nb, 56 / 600
0.36 3.67 1.33 0.46 <0.01 <0.05 <0.01 54
12.83
Zr=0.25
Zr, Nb, 57 / 600
0.36 3.75 1.34 0.5 <0.01 <0.05 <0.01 Hf=0.28 54.5
13.01
Zr = 0.22 54 / 615
0.32 3.67 1.67 0.23 <0.01 <0.05 <0.01 Hf=0.42 53.5
12.13
55 / 610
0.33 3.8 1.22 0.40 <0.01 <0.05 <0.01
Hf, Zr, Nb 42 16.01
Hf, Nb, 54 / 605
0.38 3.74 1.36 0.02 <0.01 <0.05 <0.01 51.5
13.34
Zr=0.55
Hf, Nb, 54 / 605
0.38 3.74 1.36 0.02 <0.01 <0.05 <0.01 44
16.04
Zr=0.55
Hf,Nb, 53 / 605
0.36 3.66 1.26 0.01 <0.01 <0.05 <0.01 51.5
12.10
Zr=0.44
* In all cases heat treatment which maximizes diffusivity at the indicated
hardness has been applied,
minimizing the presence of elements in solution with the matrix, except for %
Cr, especially minimizing the
presence of %C and, to a lesser extent, %V in the matrix. In all cases this
means very high austenitization
temperatures, from 3 to 5 tempering cycles, with the latest in the range 610-
680 C. Being %Hf: 0.10 - 0.22,
%Zr: 0.05 - 0.18 y %Nb: about 0.07, unless specifically indicated.
- Uncoated sheets and, therefore, with iron oxides that can be large:
For this purpose, in the context of the present invention, the following
compositional range
can be used:
Ceq: 0.4 - 0.9 Cr < 3.0% (preferably Cr < 0.1%)
V: 0 - 2.0% (preferably 0.4 - 0.8%)
Si: < 0.5%
Mn: < 1.0% Moeq: 3.5 - 9
where Moeq= %Mo+1/2 %W y
Ceq= %C + 0.86 * %N + 1.2 * %B
The rest of the elements should be kept as low as possible and, in any case,
always be
below 0.15%, with the exception of strong carbide formers (%Ta, %Zr, %Hf). All
values
are indicated in weight percentages.
The following examples show the properties that can be obtained:
Hardness Therm. cliff.
%C %Mo %W %V %Cr %Si %Mn Other
HRc min2/s
CA 02824238 2013-07-09
WO 2012/095532 - 25 -
PCT/EP2012/050531
0.60 3.6 1.2 0.62 <0.01 0.14 0.54 58
11.17
0.60 3.6 1.2 0.62 <0.01 0.14 0.54 47.5
12.47
0.85 6.48 4.0 <0.01 0.02 0.2 0.22 52.5
10.96
0.85 6.48 4.0 <0.01 0.02 0.2 0.22 45
12.87
0.45 3.87 1.67 0.49 <0.01 0.45 <0.01 51
10.65
* In all cases heat treatment which maximizes diffusivity at the indicated
hardness has been applied,
minimizing the presence of elements in solution with the matrix, especially
minimizing the presence of %C
and, to a lesser extent, %V in the matrix. In this case, also seeking the
highest possible presence of primary
carbides. In all cases this means very high austenitization temperatures, from
2 to 4 tempering cycles, with
the latest in the range 550-620 C. Being %Hf: 0.10 - 0.22, %Zr: 0.05 - 0.18 y
%Nb: about 0.07, unless
specifically indicated.
Example 2
For closed-die forging. In this case, a simultaneous optimization of wear
resistance and
thermal fatigue has to be achieved, therefore, maximum thermal diffusivity and
wear
resistance are desirable (presence of primary carbides) maintaining also
maximized CVN.
For dies or large parts subject to thermal shock or thermal fatigue a good CVN
should be
maintained, even when the treatment cannot be fully martensitic, in which case
Si or Al are
used to hinder the precipitation of thick cementite (Fe3C), or %Ni is used to
improve the
hardenability in the ferritic-perlitic zone and decrease the linear thermal
expansion
coefficient. In this case, tool steels in the following range can be used
(powder metallurgy
steels except for applications where the present abrasive particles are very
large). Steels of
the present invention are particularly attractive for applications where wear
is the
predominant failure mechanism:
For this purpose, in the context of this invention, a compositional range of
the following
type can be used:
Ceq: 0.3 -0.6 Cr < 0.1% (preferably Cr < 0.05%)
Si: < 1.4%
Al: 0 - 2%
Mn: < 1.5% Moeq: 3.0 - 7.0
where Moeq= %Mo+1/2 %W
CA 02824238 2013-01
WO 2012/095532 - 26 -
PCT/EP2012/050531
The rest of the elements should be kept as low as possible and, in any case,
always be
below 0.15%, with the exception of strong carbide formers (%Ta, %Zr, %Hf). All
values
are given in weight percentages.
Five examples show the properties that can be obtained:
Therm. diff.
Hardness
%C %Mo %W %V %Cr %Si %Mn Other mm2/s at
HRc
400 C
Hf, Zr,
0.37 3.3 1.01 <0.01 <0.01 <0.05 <0.01 45.5
11.14
Ni=2,9
Hf, Zr,
0.31 3.08 0.86 <0.01 <0.01 <0.05 0.16 44
12.69
Ni=2,3
0.5 3.65 1.27 0.45 <0.01 0.1 <0.01 Al=0.7 53 10.12
Zr,
0.5 3.73 1.52 0.17 <0.01 0.8 <0.01 Hf, 51 9.74
0.53 3.61 1.35 0.44 <0.01 <0.05 0.6 Al=0.8 55 9.62
* In all cases heat treatment which maximizes diffusivity at the indicated
hardness has been applied,
minimizing the presence of elements in solution with the matrix, especially
minimizing the presence of %C
and, to a lesser extent, %V in the matrix. In this case, also seeking the
highest possible presence of primary
carbides. In all cases this means very high austenitization temperatures, from
3 to 5 tempering cycles, with
the latest in the range 590-660 C. Being %Hf: 0.10 - 0.22, %Zr: 0.05 - 0.18 y
%Nb: about 0.07, unless
specifically indicated.
Example 3
Some closed-die forging applications, require predominantly yield strength at
high
temperatures, good toughness, specially fracture toughness and CVN, and as
good as
possible wear resistance. When the contact times are long, or the temperature
of the forged
piece high, thermal diffusivity at high temperatures and good tempering
resistance are of
outmost importance. In this case the correct usage of %Co is very important.
For this
purpose, in the context of this invention, a compositional range of the
following type can
be used:
Ceq: 0.32 - 0.7
V: <2.8%
Si: < 1.4%
Mn: < 1.5%
CA 02824238 2013-01 1
WO 2012/095532 - 27 - PCT/EP2012/050531
Co: 1.3 - 6%
Moeq: 3.3 -7.0
where Moeq= %Mo+1/2 %W
The rest of the elements should be kept as low as possible and, in any case,
always be
below 0.15%, with the exception of strong carbide formers (%Ta, %Zr, %Hf). All
values
are given in weight percentages.
Five examples show the properties that can be obtained:
Max. Hard.
Therm. din'.
Hardness
%C %Mo %W %V %Co %Mn
Other /Tem. T mm2/s at
HRe
HRe / C 600 C
0.32 3.36 1.52 0.45 2.66 <0.01 Hf,Zr, 55 /
600 51,5 6.05
Nb
0.32 3.36 1.52 0.45 2.66 <0.01 Hf,Zr, 55 / 600
39 6.37
Nb
Hf, Zr,
0.36 3.75 1.91 0.44 2.44 0.47 57 / 600
53 6.03
Si=0.2
Zr,
0.34 4.04 1.23 0.73 2.16 0.6 Hf, 56 / 600 41
6.14
Nb
Hf, Zr,
0.37 3.64 1.21 0.49 1.6 <0.01 55 / 605
42 6.04
Ni=2.7
Hf, Zr,
0.51 3.75 1.51 <0.01 2.1 <0.01 51 /600 44
6.42
Ni=2.9
0.36 3.28 0.91 0.55 3.1 0.58 Hf,Zr, 56.5 / 610
38 6.83
0.61 3.6 1.19 0.56 2.6 0.54 Hf,Zr, 59 /
615 40.5 6.55
0.43 3.22 0.96 0.04 2.8 0.5 Hf,Zr, 56 / 600
47 6.26
0.32 3.25 0.96 0.43 2.45 0.41 Hf, Zr 56 / 610
48 6.34
0.33 3.48 0.86 <0.01 2.49 0.16 Hf, Zr 54
/ 605 43 6.52
* In all cases heat treatment which maximizes diffusivity at the indicated
hardness has been applied,
minimizing the presence of elements in solution with the matrix, especially
minimizing the presence of %C
and, to a lesser extent, %V in the matrix. In all cases this means high
austenitization temperatures, from 3 to
5 tempering cycles, with the latest in the range 640-690 C. Being %Hf: 0.02 -
0.16, %Zr: 0.05 -0.18 y %Nb:
about 0.07, unless specifically indicated.
Example 4
CA 02824238 2013-07--
WO 2012/095532 - 28 -
PCT/EP2012/050531
For hot cutting of sheet. In this case wear resistance must be maximized with
a good
hardenability and toughness (fracture toughness, in this case). Thermal
conductivity is very
important to maintain the temperature at the cutting edge as low as possible.
Weldability is
less important in this case, and small inserts are often used, so compositions
with high
content of alloying elements can be used. For this purpose, in the context of
the present
invention, the following compositional range can be used:
Ceq: 0.5 - 0.9 Cr < 0.1% (preferably Cr < 0.05%)
Si: <0.15% (preferably Si <0.1%)
V: 0 - 2% for cases with Moeq > 5 and V: 0 - 4% for cases with Moeq < 5
Moeq: 5 - 10
where Moeq= %Mo+1/2 %W
The rest of the elements should be kept as low as possible and, in any case,
always be
below 0.15%, with the exception of strong carbide formers (%Ta, %Zr, %Hf). All
values
are given in weight percentages.
Three examples show the properties that can be obtained:
Therm. cliff.
Hardness
%C %Mo %W %V %C r %Si %Mn Other mm2/s at
HRc
400 C
0.59 6.7 4.6 <0.01 <0.01 <0.05 <0.01 55 12.39
0.69 7.89 3.95 0.7 <0.01 <0.05 <0.01 55 10.76
0.62 8.01 3.75 0.1 <0.01 <0.05 <0.01 Ni=0.28 57 10.10
28
0.75 6.11 3.4 0.5 <0.01 <0.05 <0.01 Hf=0. 61 9.87
Zr=0.14
23
0.87 6.92 4.4 0.7 <0.01 <0.05 <0.01 Hf=0. 64 9.03
Zr=0.15
* In all cases heat treatment which maximizes diffusivity at the indicated
hardness has been applied,
minimizing the presence of elements in solution with the matrix, especially
minimizing the presence of %C
and, to a lesser extent, %V in the matrix. In this case, also seeking the
highest possible presence of primary
carbides. In all cases this means very high austenitization temperatures (1120
C in the first two cases and
1240 C in the last one), from 2 to 4 tempering cycles, with the latest in the
range 600-640 C.
