MX2013008138A - High thermal diffusivity and high wear resistance tool steel. - Google Patents

Abstract

A tool steel family with outstanding thermal diffusivity, hardness and wear resistance has been developed, also exhibiting good hardenability. Also its mechanical strength, as well as its yield strength, at ambient and high temperature (superior to 600ºC) are high, due to a high alloying level in spite of the high thermal conductivity. Because of its high thermal conductivity and good toughness, steels of this invention have also good resistance to thermal fatigue and thermal shock. This steels are ideal for discontinuous processes where it is interesting to reduce cycle time and that require high hardness and/or wear resistance (plastic injection molding, other plastic forming processes and curing of thermosets, hot forming of sheet...). These tool steels are also appropriate for processes requiring high wear resistance and good resistance to thermal fatigue (forging, hot stamping, light-alloy injection...).

Description

STEEL OF TOOLS WITH A STRENGTH OF WEAR AND EXTRAORDINARY THERMAL DIFUSIVITY.

Field of the invention The present invention relates to a tool steel with a very high thermal diffusivity and a high resistance to wear, mainly abrasive. Tool steel also has good hardenability.

Description Tool steels often require the combination of various properties considered opposed. A typical case are for example the elastic limit and the tenacity. For many industrial forming applications in which heat extraction of the manufactured product takes place in a discontinuous system, thermal diffusivity is of vital importance. Traditionally this property has been considered opposed to the hardness and wear resistance of tool steels. In plastic injection, hot stamping, even forging, metal injection, compound curing and other forming processes, wear resistance and high or very high thermal diffusivity are often required simultaneously. In many of these applications, large-sized tools are required, so the hardenability of the tool material is also extremely important. The thermal diffusivity (a) is related to other fundamental properties of the material such as density (p), specific heat (Cp) and thermal conductivity (?) As follows: ? = p · cp · a Or if you prefer: a =? / (p - cp) The wear in the forming processes is mainly of abrasive and adhesive origin, but sometimes other types of erosion, cavitation, etc. are also present. To counteract abrasive wear generally hard particles are required in the tool steel, these are generally ceramic particles such as carbides, nitrides, borides or combinations of the above. Thus, the volumetric fraction, the hardness and the morphology of said hard particles will determine the wear resistance of the tool material for a given application. Also the hardness of use of the tool material is of great importance to determine the durability of the material under conditions of stress with abrasive wear. The morphology of the hard particles determines its adherence to the matrix and the size of the exogenous abrasive particle that can be counteracted without detaching from the matrix of the tool material. To counteract adhesive wear, it is best to use functionally graded materials (FGM), usually in the form of a ceramic coating on the tool material. In this case, the most important thing is to provide a good support to said coating, which is generally quite fragile. To provide good support to the coating, hardness and hard particles are required in the tool material. Thus, for various 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 (in the case of having to counteract large abrasive particles).

Thermal gradients are the cause of thermal shock and thermal fatigue. In many applications, steady transmission rates are not achieved due to low exposure times or limited amounts of power from the source causing a temperature gradient. The magnitude of the thermal gradient in the tool materials is also a function of its thermal conductivity (inverse proportionality applies for all cases with a sufficiently small Biot number).

Thus, in a given application with a given thermal flux density function, a material with higher thermal conductivity is subjected to a lower surface charge, since the resulting thermal gradient is lower. The same occurs when the coefficient of linear thermal expansion is smaller and when the Young's modulus is smaller.

Traditionally, in many applications in which thermal fatigue is the main failure mechanism, as in many cases of casting or extrusion of light alloys, it is desirable to maximize conductivity and toughness (generally fracture toughness and CVN). The steels of the present invention prioritize resistance to wear and diffusivity compared to CVN, which is considered very important also for some applications and therefore it is also intended to maximize but without a great rejection of the other two properties. Generally, when increasing the hardness of a tool steel, it is expected that thermal tenacity and diffusivity will decrease and wear resistance is expected to increase. In the steels of the present invention, a higher level of diffusivity has been achieved for a given level of hardness, generally accompanied by good hardenability and in some cases with an excellent tenacity compromise.

In many applications thick tools are used, especially when sufficient mechanical strength is required to involve a heat treatment. In this case it is also very convenient to have a good hardenability to reach the desired level of hardness, on the surface and preferably up to the core. The hardenability is also very interesting for steels for hot work, because it is much easier to achieve high tenacity with a tempered martensite microstructure than with a tempered bainite. Thus, the higher the hardenability, the lower the brusqueness required in cooling the temper. An abrupt cooling is more difficult to achieve, and also more expensive and since the shapes of the tools and components manufactured are often complex, it can lead to the breakage of the parts treated thermally or to severe deformations. The wear resistance and the mechanical resistance are usually proportionally inverse to the tenacity. Thus, it is not easy to achieve a simultaneous increase of both properties. The thermal conductivity helps in this case, since it allows a large increase in thermal fatigue resistance, even if the CVN has been reduced to increase wear resistance or mechanical strength.

There are many other desirable, if not necessary, properties for a hot working steel that do not necessarily influence the longevity of the tool, but in its production costs, such as: ease of machining, welding or repair in general, Support provided to the coating, costs, ...

The authors have discovered that the problem of simultaneously obtaining a high thermal diffusivity, resistance to wear and hardenability, with good levels of tenacity can be solved by applying certain standards 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 % A1 = 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 consists of iron and unavoidable impurities, in which o / oCeq =% C + 0.86 *% N + 1.2 *% B, (C *, > 0.36) In the present invention it is always the case that: % ?? + ½ ·% W > 3.0 Some rules of alloy selection within the range and thermo-mechanical treatments required to obtain the desired high thermal diffusivity at high level of hardness and wear resistance are presented in the detailed description of the invention. Obviously a detailed description of all possible combinations is out of place. The thermal diffusivity is regulated by the mobility of the transporters of the heat energy, which unfortunately can not be correlated to a compositional range and a thermo-mechanical treatment.

In a further aspect, the invention relates to a process for the manufacture of a steel for hot working tools, characterized in that the steel is subjected to a martensitic, bainitic or martensitic-bainitic temper with at least one cycle of tempering at a temperature above 590 ° C, so that a steel with a hardness above 47 HRc with a low dispersion structure characterized by a diffusivity of 9 mm2 / s or more. In another embodiment, a steel with a hardness of more than 53 HRc can be obtained with a low dispersion structure characterized by a diffusivity of 9 mm2 / s or more. In a further form of this process, the steel is subjected to at least one annealing cycle at a temperature above 600 ° C so that a steel having a hardness greater than 50 HRc or more can be obtained with a structure of Low dispersion characterized by a diffusivity of 5.8mm2 / s or more at 600 ° C. State of the art Until the development of high thermal conductivity tool steels (EP 1887096 Al), the only known way to increase the thermal conductivity of a tool steel was to keep low alloy and, therefore, with poor mechanical characteristics, especially at high temperatures . Tool steels capable of exceeding 42 HRC after tempering at 600 ° C or more were considered with an upper limit of thermal conductivity of 30 W / mK and with a thermal diffusivity of 2 2 8 mm / s and 6.5 mm / s for hardness above 52 HRc. The tool steels of the present invention have a thermal diffusivity greater than 8mm2 / s, and often higher than 12mm2 / s for hardness of more than 52 HRc, and even more than 6 mrn2 / s for hardness of more than 42 HRc, also presenting a very good resistance to wear and good hardenability. Thermal diffusivity is considered the most relevant thermal property because it is easier to measure accurately, and because most tools are used in cyclic processes, so thermal diffusivity is much more important to evaluate performance than conductivity thermal The tool steels of the present invention have a higher wear resistance and hardness level than the steels described in EP2236639A1. The latter have a higher hardenability in the pearlitic zone and higher CVN with respect to the steels for high thermal conductivity tools of the present invention. So for applications where the main failure mechanism is thermal fatigue and there is no wear, it is better to use the steels of EP2236639A1, but for applications where the resistance to wear is important, the tool steels of the present invention have great advantage. In addition, the steels of the present invention have greater thermal diffusivity for the same level of hardness. This is largely due to the fact that in EP2236639A1 M3Fe3C type carbides are practically exclusively used, where M corresponds to Mo and / or W, partly due to the presence of% Ni in the matrix that penalizes thermal diffusivity in favor of the hardenability, tenacity (CVN) and lower coefficient of linear thermal expansion. In the present invention,% Ni minor are present and the carbides are often partially replaced by harder carbides, even when the harder carbide forming elements tend to be solubilized in the carbides of Mo and / or W, as is the case with the % V.

The tool steels of the present invention can achieve much higher levels of thermal diffusivity than the tool steels of WO2004 / 046407 Al, where the high levels of% Cr impose very strict restrictions that are not observed in the compositions to be considered within the scope of the invention. proposed range and the next small process window during thermomechanical processing to achieve high levels of carrier mobility.

There are other inventions that may have partial coincidence of the compositional range but that have nothing to do with the present invention because the rules to select the composition within the range and / or the thermo-mechanical treatments required to achieve a structure with a matrix poor in elements in solid solution with great dispersion capacity of heat energy carriers, and with carbides with a high level of crystalline perfection and consequently very low dispersion of heat energy carriers (mainly electrons and phonons). This could be the case of JP04147706 where the inventors, looking for an optimized surface oxide layer, use% Cr levels lower than normal (around 0.5%) to allow said oxidation with specific treatments at high temperature. In the present invention% Cr has a tendency to dissolve in the W and / or Mo carbides causing dispersion of the heat energy carriers and therefore their presence is also undesirable. This is the only coincidental point of coincidence that also in the case of JP04147706 does not entail high thermal diffusivity in any of the described examples. To a lesser extent is the case of JP11222650 where the inventors seek the presence of a large quantity of primary carbides to withstand the massive wear as is the case of a fast steel but with an exceptionally low% C content to allow the cold locking .

Other cases can be deceived by not making special mention or have a generic mention on the levels of non-functional elements for the aforementioned application, this is often the case of% Cr,% Si and% Mn. In fact it is expensive to achieve a low level of certain elements in steels. For example, a steel supposedly devoid of Cr (0% Cr in nominal composition), especially if it is an alloyed steel, will have very probably% Cr > 0.3% if the steel requires for some reason from selected scrap; in case you can use normal scrap, significantly cheaper, you would expect% Cr > 0.5% If% Cr is not mentioned in a composition it means that its presence is not considered important, but neither is its absence. In this case, the% Cr content does not require the use of special scrap and if there are no other elements that require it then we can expect% Cr > 0.5 Even more important is where is placed this% Cr, which will be predominantly dissolved in the carbides if the necessary measures are not taken.

The case of% If it is a little different, since it is possible to reduce its content through a refining process such as, p. eg, the ESR, although due to the narrow window of the process in this case it is technologically very difficult (and expensive, and therefore only performed in the case of pretending a certain functionality) reduce the% If below 0.2% and at the same time reach a low level of inclusions (especially oxides).

There are many tool steels that have a composition with the potential to achieve high thermal conductivity and, in fact, they do not. This is mainly due to the following two reasons: - The thermo-mechanical treatments used do not aim to maximize the mobility of the heat energy carriers. The thermal conductivity is not properly chosen as a desirable main characteristic or, for previously developed materials, the knowledge was deficient in the way of reaching a desired level of thermal diffusivity before the publication of patent EP 1887096 Al, and therefore the phases present in the final microstructure is chosen depending on the optimization of some other desirable properties for the application, usually a certain commitment of the mechanical properties relevant to the application. Therefore, within a composition, hardening mechanisms are often chosen that are very detrimental to thermal diffusivity.

- In the fusion process, in secondary metallurgy or re-fusion, not enough attention is paid to what happens beyond the micrometric and nanometric scale, so that unfavorable dispositions take place at the atomic scale, not necessarily in all present phases, and that lead to a strong dispersion of carriers. Again, this is mainly due to the lack of knowledge prior to the publication of patent EP 1887096 Al.

There are several families of tool steels that, with their nominal range of composition, could have the potential to achieve high thermal diffusivity when the correct strategy is employed during the thermo-mechanical process according to the present application and the EP document. 1887096 Al, but do not end with compositions capable of developing high or very high thermal diffusivity. This is mainly due to the following reasons: - The ratio of% C to that of carbide formers is not well balanced to be able to minimize solid solutions in the metallic matrix, especially% C, and- Levels are provided that can not be handled properly afterwards by the treatments Thermo-mechanical devices used to pursue the maximization of the mobility of heat energy carriers.

- The nominal levels of certain critical elements are far from the values of the actual content in the embodiment. For example, 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 limit descriptions, such as% Cr < 1 (or even without mentioning the% Cr, a fact that can lead to misunderstanding that is 0%) and in the same way as is normally the case with% Si < 0.4 ends up being% Cr > 0.3 and% Si > 0.25. This is also valid for all trace elements with a strong impact on the conductivity of the matrix and even more those with a high solubility in carbides and great distortion potential of the carbide structure. In general, with the exception of% Ni, and for some applications% Mn, no element is desirable in solid solution with the matrix in a percentage higher than 0.5%. Preferably the percentage of these individually in solid solution should not exceed 0.2%. If the main objective for the application is to maximize the thermal conductivity, then any metallic element (obviously including the transition metals), except for the% Ni and in some cases the% C and% Mn, in solid solution with the matrix should not Exceed 0.1% or better even 0.05%.

Detailed description of the invention In order to obtain tool steels with high thermal diffusivity and wear resistance at high hardness levels with good hardenability, it has been observed that within the compositional range indicated above, several rules and general considerations in the composition selection must be taken into account within the range and the thermo-mechanical treatments to be used, some of which are described below. The thermal diffusivity is a consequence of the dispersion mechanisms of the phases present for all types of carriers present. The perfection of the network plays an important role, but also other mechanisms of dispersion are relevant. In this document the thermal diffusivity itself will be used as a measure of the structure reached. Within the same chemical composition different structures can be obtained and therefore also different levels of thermal diffusivity.

The steels of the present invention excel by their high thermal diffusivity and wear resistance, mainly. The wear resistance and the tenacity tend to be inversely proportional, but different microstructures reach different ratios, that is, a function of the microstructure is the one that allows to obtain different levels of tenacity for the same elastic limit and level of hardness at a certain temperature , and for a given type of material, hardness tends to correlate with wear resistance unless the volume fraction of wear resistant particles or their morphology is changed markedly. In this order of things, it is well known that for the majority of tool steels with medium carbon, the microstructure of pure, 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 such as stable ferrite-perlite or metastable bainite during the subsequent cooling to austenitization in the thermal treatment process. Therefore, rapid cooling rates will be required and, if higher hardenability is required, some alloying elements should be used to retard the kinetics of the formation of these more stable structures, and all possible alternatives should be used. with the minimum negative effect on thermal diffusivity.

A strategy to obtain resistance to wear and a higher elastic limit at high temperatures and, at the same time, obtain a high thermal conductivity is the use of carbides with a high electron density type M¾Fe3C secondary and sometimes even primary carbides (M- should be only Mo or W for greater conductivity thermal). There are other carbide species of (Mo, W, Fe) with high electron densities and tendency to solidify with great crystalline perfection. Some elements such as Zr and Hf, and to a lesser extent the Ta for example in comparison with the Cr, when they dissolve in this type of carbides, they do not contribute great distortion to the crystalline structure and the dispen of charge carriers is small with what is also the effect in the thermal conductivity. In addition, these carburizing elements have a tendency to form separate MC-type carbides, due to their high affinity for C.

In fact, in the present invention it has been observed that the effect can be quite positive if a moderate amount of% V is used and compensated for by the presence of some strongly carburigenic element (preferably Zr and / or Hf). It has been seen that you can have amounts of% V up to 0.9% without practically forming primary carbides (obviously depends on the Ceq and the presence of other carbides, so for higher contents of Ceq you have to reduce the percentage of% V maximum to 0.8 and even 0.5 or 0.4 to avoid the presence of primary carbides or the massive dissolution in these) and with little dissolution in the carbides of (Fe, Mo, W) especially if they are used simultaneously more strongly carburigenos elements, in addition you get the additional benefit to displace more carbon out of the matrix with the consequent benefit for the general thermal diffusivity (the benefit is remarkable in this case with% Hf +% Zr +% Ta superior to 0.1, and very remarkable if it is higher than 0.4 or 0.6 depending on the quantities presents of% Ceq and% V). In fact, this combination is very desirable since both% V and% Zr, as% Hf and% Ta tend to significantly improve the wear resistance compared to a steel that only has carbides of (Fe, Mo, W), is also the case for% Nb. The effect becomes remarkable with a% V = 0.1 and remarkable with% V = 0.3 or 0.5 depending on the level of % Ceq. If extreme wear resistance is sought with the presence of primary carbides, as is the case in applications with large abrasive particles as in the hot stamping of uncoated sheets, then higher amounts of% V up to 1.5% or even up to 1.5% can be used. 2% is possible maintaining a good level of thermal diffusivity especially if it is compensated with strongly carburizing elements. In this case it may be convenient to have levels of strongly carburgenic elements combined with high% V (% V +% Nb +% Hf +% Zr) greater than 1.2 or even 2 weight percent (for applications with high wear resistance even 3 , but then the cost of the alloy is high). In this case, rarely any strongly carburgenic element (% V,% Nb,% Ta,% Zr,% Hf) will exceed 3% individually, except for% V where the upper limit is usually 4% by weight (for applications where wear resistance is prioritized at the expense of thermal conductivity loss) or 1.8% for applications requiring very high thermal diffusivity and Nb which, due to its negative effect on thermal diffusivity, tends to be used only as a controller of the grain size and when it is used as a primary carbide former, it will rarely be above 1.5%. It is desirable to have most of the strong carbide formers bonded in the carbides and not dissolved in the matrix, therefore the level of% Ceq has to be adjusted precisely as explained below to minimize the amount of forts carbide formers and Ceq% in solid solution. As an example in most applications of this invention if% Ceq is less than 0.35% then% V should be kept below 1.7%. In general, it is desired to have a majority of carbides (Fe, Mo, W) (where obviously part of the C can be replaced by N or B), normally more than 60% and optimally more than 80% or even 90% of this type of carbides. The dissolution of other metallic elements in this type of carbides (obviously in the case of carbides the metallic elements are mainly transition elements) may exist, but it is desired small to guarantee a high phononic conductivity. Normally no other metallic element, apart from the main Fe, Mo and W, should exceed 20% of the weight of all the metallic elements of the carbide, in this type of mainly desired carbides. Preferably it should not be more than 15% and even better 5%. This is because they tend to form structures with extremely low solidification deficit densities even for fast solidification kinetics (hence less structural elements to cause scattering of charge carriers).

As previously mentioned, the only exception is the presence of a limited amount of strongly carburgenic elements, although it is preferred that these form independent carbides. In this case, the Mo and the W provide sufficient obstacles for the formation of stable structures (pearlite and ferrite), although bainite formation is very rapid. In some steels, it is possible to form superbainitic structures by applying a type of heat treatment, which consists in the complete solubilization of alloying elements followed by rapid cooling at a certain temperature (to avoid the formation of ferrite) in the order of a formation. Lower bainitic, and a long maintenance of the temperature to obtain a 100% bainitic structure. For most steels a pure martenistic structure is desirable, so in this system some elements have to be added to delay the bainitic transformation, since the Mo and the W are very inefficient in this aspect. As a general rule,% Cr is usually used for this purpose, but it has an extremely negative effect on conductivity thermal for this system since it dissolves in the M3Fe3C carbides and causes a great distortion, so that it is much better to use strongly carburious elements and elements not soluble in carbides. These last elements will reduce the conductivity of the matrix and, therefore, those that have the least negative effect should be used. Thus, a natural candidate is Ni, but others can also be used in parallel, (special mention must be made to% B for this effect with very low concentrations). Since carbide formers having a high affinity with carbon tend to be used in the present invention for their positive effect on wear resistance, the necessary and desirable amount of elements for retarding the transformation kinetics to stable structures during tempering is less. Normally, an amount of up to 1% and for large sections up to 3% by weight will suffice to obtain sufficient hardenability and contribute to the increase in toughness without excessively damaging the conductivity. Higher amounts of% Ni provide greater tenacity and a reduction of the coefficient of linear thermal expansion, but the priority in the present invention is the wear resistance combined with the thermal diffusivity, so only for some special applications can the strategy of using high contained in% Ni, with a maximum of 3.8%. There are applications in which lower amounts of% Ni already have the desired effect, especially if the contents of% Mn and / or% Si are somewhat higher (% Mn does not usually exceed 3%) or the sections of material used are smaller.

The use of% Mo as the only carbide-forming element (obviously together with Fe), is somewhat advantageous in terms of maximizing thermal conductivity, but has the disadvantage of providing a higher coefficient of thermal expansion, and therefore decreases the resistance to thermal fatigue. Therefore it is usually preferred to have a ratio of 1.2 to 3 times more Mo than W but not the absence of W. The exception is the applications where you only want to maximize the thermal conductivity together with the toughness but not particularly the resistance to thermal fatigue. The hardenability and cost of the alloy given the high volatility of the% o and% W prices can lead to change the preferences regarding that% W is the main element in% MOeq, being% M eq =% Mo + ½ ·% W.

High contents of Moeq with high levels of Ceq can be used, resulting in an alloy of high cost, low tenacity, very difficult solderability, complicated hardenability for large parts and limited machinability. But very high levels of wear resistance can be achieved with good diffusivity thermal For applications where the highlighted drawbacks are not determining, they can be alloys of interest. This may be the case, for example, of some cutting applications. In this case, C ^ levels are usually higher than 0.5% and often higher than 0.6%. % MOeq levels are often above 5% and often above 6% and even 9%. Also the limits of the Moeq / Ceq ratio are often displaced to levels higher than the rest of the alloys of the present invention. Values greater than ló are possible, and higher than 13 probable.

Throughout the document the term carbide refers to both primary and secondary carbides unless specifically indicated.

The more restrictive it can be with% Si and% Cr, the higher the thermal conductivity will be but the solution is more expensive (in addition, some properties, which could be important for certain applications, and therefore it would be desirable to maintain them, could worsen with the reduction of these elements below certain levels, such as, for example, the tenacity of trapped oxide inclusions if Al, Ti, Si, and any other deoxidizer are used, in insufficient quantities, or some cases of corrosion resistance if% Cr or% Si are too low). Therefore, it is often necessary to look for a compromise between increasing costs, reducing toughness, resistance to corrosion or other relevant characteristics for certain applications, and the benefit of greater thermal conductivity. The maximum thermal conductivity can be obtained only if the levels of% Si and% Cr are below 0.1% or better even if they are below 0.05%. To maximize the thermal diffusivity, also the levels of all other elements except for% C,% Mo,% W,% V,% Zr,% Hf,% Ta,% Nb and in some instances% Mn and% Ni have to be as low as possible (below 0.05 is technically possible with an affordable cost for most applications, although a maximum of 0.1 is naturally less expensive). For some applications where toughness is especially important, less restrictive levels of% Si must be adopted (it is the least harmful to the thermal conductivity of all iron deoxidizing elements), and therefore give up a certain thermal conductivity, for ensure that the level of inclusions is not too high. Depending on the levels of% C,% Mo, and% W used, hardenability may be sufficient, especially in the pearlitic zone. For cases of large pieces where it is not possible to avoid the formation of bainite during tempering it may be interesting to use elements in solid solution that prevent the formation of coarse cementite precipitates (Fe3C) that cause very low tenacities, such as % A1 and% Si. Generally below 0.4, exceptionally with levels of around 1% and very exceptionally above 2% in the case of% A1, reaching 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 the smallest possible amount of these elements remains in solid solution. The same is true for the other carbide formers that could be used to obtain a certain tribological response (such as% V,% Zr,% Hf,% Ta, ...).

In some applications, a certain environmental resistance can be interesting, and therefore it is desirable to have some% Cr or% Si (resistance to oxidation at high temperatures) in solid solution. The negative effect on thermal diffusivity can be mitigated by carbon fixation with more strongly carburizing elements. Without the latter the% Cr should not exceed 2%, and preferably 1.5%. But 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 reached, maintaining good thermal diffusivity, and up to 1.4% Si. In fact, for most applications% Cr < 2.8% is necessary if the thermal diffusivity has to be high. Many compositions require% Cr < 2.5% to be able to reach a high thermal diffusivity with the appropriate thermo-mechanical processing (which is dependent on the composition, as explained). At this level, the effect of protecting the environment is only a little noticeable if Cr% is left mainly in solid solution in the matrix. Finally, a much larger range of the compositions can reach a high thermal diffusivity when the appropriate thermo-mechanical treatment is applied, if% of Cr is limited to remaining below 1.9%.

The simplest compositional rule to describe the compositions within the range capable of achieving a high thermal diffusivity simultaneously with a high resistance to wear can be based on a relation R = Moeq / Ceq, where% Moeq =% Mo + ½% W and% Ceq =% C + 0.86 *% N + 1.2 *% B. This rule applies only for sufficiently large contents of % Ceq (normally 0.32 min, preferably 0.5 min and more accurately when at least 0.38% Ceq) and% Moeq (normally 3.2 min, preferably 3.4 min and more accurately when at least 3.6% Moeq). It is also a rule that can only be used for low Cr% content, typically% Cr < 2.5%, and desirably% Cr < 1.9%. The minimum value of the results of R when calculating the minimum% Moeq to apply the rule divided by 0.9, which is the maximum Ceq% for the present invention (for example for a minimum Moeq = 3.2 then the minimum value of R it turns out to be 3.56). The maximum value of R has been observed to be possibly 11.5, preferably 10.8 and optimally 10.5 for low values of Ceq%. Low values Ceq% are for this rule those under 0.35%, sometimes under 0.36% or even below 0.37%. For high values Ceq% the maximum value of R has been observed to be possibly 16.8, preferably 16.0 and optimally 15. High values Ceq% 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 of R has been observed to be 14, preferably 13, and optimally 12. In general, it is convenient to maximize thermal diffusivity only (ie there are no other properties of great importance) , observe the following alloy standard (to minimize the% C in solid solution), if you want to obtain a martensite or bainite microstructure that is resistant to mechanical demands. The formula has to be corrected if carbide formers with high affinity for% C (such as Hf, Zr or Ta, and even Nb) are used, the formulation must also be modified if% Cr > 0.2 o Moeq > 7: 0. 02 < xCeq - solC - AC · [(xMo-solMo) / (3 · AMo) + (xW-solW) / (3 · AW) + (xV-solV) / AV] > 0.265 where: xCeq - percentage in carbon weight; xMo - weight percentage molybdenum; xW - wolfram weight percentage; xV - percentage in vanadium weight; AC - carbon atom (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 - percentage of carbon in solid solution; solMo - percentage of molybdenum in solid solution; solW - percentage of tungsten in solid solution; solV - percentage of vanadium in solid solution.

For 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 better yet: 0. 09 < C ^ - solC - AC · [(xMo-solMo) / (3 · AMo) + (xW-solW) / (3 · AW) + (xV-solV) / AV] > 0.18 To rectify the presence of other carbide formers with great avidity for% C, an extra term must be added to the formula for each type of high avidity carbide former: -AC * xM / (R * AM) where: xM - percentage by weight carbide-forming; AC - carbon atom (12.0107 u); R - number of carbide former units per carbide unit (eg 1 if the carbide type is MC, 23/7 if the carbide type was M23C7 ...) AM - carbide-forming atomic mass.

This balance provides an extraordinary thermal conductivity if the formers of ceramic reinforcing particles, including the non-metallic part (% C,% B, and% N) are brought to the carbides (as an alternative nitrides, borides or intermediates). Then the appropriate heat treatment has to be applied. This heat treatment will have a phase in which most of the elements are dissolved (austenitization at a sufficiently high temperature, normally around 1080 ° C for moderate levels of ?? ^, 1 120 ° C for medium levels of Moeq and 1240 ° C for high levels of Moeq. Exceptionally, if the distortion in the heat treatment is of great importance for the application, lower austenitization temperatures can be used), a sudden cooling will follow, the intensity of which will be determined by the desired mechanical properties, although stable microstructures should be avoided because they imply phases with large amount of% C and carbide formers in solid solution. The metastable microstructures are even worse, since the distortion in the microstructure caused by the carbon is even greater, and therefore the lower thermal conductivity, although once these metastable structures have relaxed, is when the carbide formers are placed in the desired position. The martensite and bainite tempered according to this process will be the desired microstructures in this case. The highest possible replacement of Fe in the secondary carbides by Mo, W and all the carbide forming elements with more affinity for carbon than Cr is desired, which is why the tempering strategy chosen has a great influence on the final thermal diffusivity, with special importance for the final tempering temperature and the minimum tempering temperature. For hardness above 40 HRc, the highest possible temperature for the last temper is desirable if the thermal diffusivity is to be maximized, and with this criterion the intermediate temperings must be set. That is, the same level of final hardness can be achieved with different tempering sequences and the one that uses a higher final tempering temperature is chosen, if the sole objective is to maximize the thermal diffusivity at a certain level of hardness. Thus, abnormally high tempering end temperatures are usually used, often above 600 ° C even when hardness of more than 50 HRc is chosen. In the steels of the present invention it is customary to achieve hardness of 47 HRc, even more than 52 HRc, and often more than 53 HRc and with the embodiments considered as particularly advantageous due to their resistance to wear, hardnesses are possible above 54HRc and often more than 56 HRc with even an annealing cycle above 590 ° C presenting a low dispersion structure characterized by a thermal diffusivity of more than 8 mm2 / s and, generally, more than 9 mm / s or even more than 10 mm2 / s; when it runs particularly well then greater than 11 mm2 / s, even greater than 12 mm / s and occasionally above 12.5 mm2 / s. As well as achieving hardness above 42 HRc, even more than 50 HRc with a last tempering at more than 600 ° C, often at more than 640 ° C, and sometimes even above 660 ° C, presenting a low structure dispersion characterized by a thermal diffusivity greater than 10 mm2 / s, even at 12 mm / s. when they are especially well executed then greater than 14 mm / s, even greater than 15 mm2 / s, and occasionally above 16 mm2 / s. These alloys can have an even higher hardness by reducing the tempering temperatures, but for most of the intended applications a high resistance to tempering is highly desirable. As can be seen in the examples with some very particular embodiments with high carbon content and high alloy, which leads to a high hard particle volume fraction, in the present invention a hardness higher than 60 HRc is possible with structures of low dispersion characterized by a thermal diffusivity above 8 mm2 / s and generally more than 9 mm2 / s.

To achieve 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 4%, preferably above 5.5% and for some high wear applications, even above 9% The size of the primary hard particles is very important to have an effective wear resistance and yet is not excessively low tenacity. The inventors have observed that for a volume fraction of hard particles given the total resilience of the material decreases as the size of the hard particles increases, as would be expected. A little more surprisingly It has also been observed that when the size of the hard particles increases, the tenacity to the total fracture increases if the fracture resistance of the particles themselves is maintained. When it comes to abrasive wear resistance, the existence of a critical hard particle size has been observed, 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 in which the abrasive particles are small (usually below 20 microns), it may be desirable to have primary hard particles of less than 10 microns or even less than 6 microns, but in no case with an average size not less that 1 miera. For applications where large abrasive particles cause wear, large primary hard particles will be desirable. Therefore, for some applications it is desirable to have some primary hard particles larger than 12 microns, often larger than 20 microns and for some particular applications, even larger than 42 microns. For applications where mechanical strength rather than wear resistance is important, and it is desirable to achieve such mechanical strength without compromising too much hardness, the volume fraction of small secondary hard particles is of great importance. The small secondary hard particles, in this document, are those having a maximum equivalent diameter (diameter of a circle with an equivalent surface as the cross section of the maximum surface of the hard particle) below 7.5 nm. Then it is desirable to have a volume fraction of secondary small hard particles for such applications above 0.5%. It is believed that the saturation of the mechanical properties for hot working applications occurs around 0.6%, but the inventors have observed that for some applications that require high resistance to plastic deformation at somewhat lower temperatures, it is advantageous to have amounts greater than this 0.6%, often more than 0.8% and even more than 0.94%. Since the morphology (including the size) and the volume fraction of secondary carbides change with the heat treatment, the values presented here describe the values achievable with an adequate heat treatment.

Cobalt has often been used in hot working steels mainly due to the increase in mechanical strength, and in particular the increase and maintenance of creep resistance up to considerably high temperatures. This increase in creep resistance is achieved through the solid solution and therefore has a rather negative effect on toughness. The common amounts of Co used for this proposal is 3%. In addition to the negative effect on toughness, the negative effect on thermal conductivity is also well known. The inventors have seen that within the composition ranges of the present invention, it is possible to use Co, and achieve an improved strength / hardness ratio since the Co can promote the nucleation of secondary hard particles and thus maintain their small size. It has also been found that, for some compositions of the present invention, the addition of Co effectively decreases the thermal diffusivity at room temperature, but may actually increase at higher temperatures (typically above 400 ° C) if the treatment is applied. thermo-mechanical correct. The inventors have seen that the best results were found when% Co is above 1.3%, preferably above 1.5% and optimally above 2.4%. Also% C must exceed 3.2%, preferably 3.4% and optimally 3.6%. If the thermal conductivity at high temperatures is of paramount importance for the application, special care must be taken of not having excessive% V, which 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 up to 4.0%. The heat treatment has to be selected with a rather high austenitization temperature and abnormally high tempering temperatures. In fact, it is often achieved more than 55 HRc with at least one tempering cycle at 630 0 C or even above, it can be maintain 50 HRc even with an annealing cycle at 660 ° C or more. The adequate thermomechanical processing together with the compositional rules that have just been explained, have to be implemented to minimize dispersion at high temperatures, the optimized regime is characterized by providing diffusivity of more than 5.8 mm2 / s, often more than 6.1 mm2 / s, and even more than 6.5 mm2 / s in measuring temperatures up to 600 ° C When we mostly stay in the MoxW3-xFe3C carburigenic system, one of the preferred ways to balance the contents of% W,% Mo and% C in the present invention, is by adhering to the following alloy standard: % Ceq = 0.4 + (% MOeq (reair4) * 0.04173 where: ?? ^?) ^% Mo + (AMo / AW) *% W. with : Amo - molybdenum atomic mass (95.94 u); AW - tungsten atomic mass (183.84 u); so that in the end: Mo "(rai) =% Mo + 0.52 *% W.

If the expression is normalized in a parameter =% Ceq / (0.4 + (% Moeq (rcai) -4) * 0.04173), the desirable values for this parameter, in the present invention are as follows: It has been observed that when the carbon content is low (ie,% Ceq <0.39, preferably Ceq% <0.36% and optimally Ceq <0.35), the parameter should exceed 0, 75, preferably 0.76, more preferably 0.86 and 0.88 optimally. In fact, for some embodiments for applications requiring a very high wear resistance, it will normally be greater than 0.92. A very good performance is obtained as already described, at the expense of a higher cost, with the addition of elements that strongly bind the carbon to the carbides. In the case here, 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% . Since Nb can be quite detrimental to thermal diffusivity for some applications it will not be desirable (% Nb <0.09) and then the content 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 hard carbide former, then contents above 0.14%, preferably above 0.2% and even above 0.4% will be desirable. In these cases with the presence of strong carbon binders the restrictions on K can be relaxed around 3% to 5% for the alloys with low carbon content as described here. On the other hand, when the carbon content is not low (ie% Ceq> = 0.39, preferably% Ceq> = 0.36 and optimally% Ceq> = 0.35), the K parameter must exceed 0.6, preferably 0.75, more preferably 0.84 and optimally 0.87. In this case, if elements are used that strongly bind carbon (nitrogen or boron) to the carbides (nitrides, borides or mixtures) in the manner 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 a good combination of wear resistance and thermal diffusivity can be obtained for very high values of K if all the other rules of thermal processing and alkeation are met, normally in its most restricted version, but of course they are obtained the best results when K does not exceed 3, preferably 1.5 and 1.3 optimum.

An especially interesting realization, when the main objective for the chosen application is the maximization of thermal diffusivity for the highest possible level of hardness, arises when this alloy rule is applied together with very low levels of Cr%, especially in dissolution with the carbides, as described above. It has also been noted by the authors that it is possible to achieve a considerably high thermal diffusivity and wear resistance when much higher levels of% Mo and W% are used than those described in the last two paragraphs. The level of thermal diffusivity for a given level of hardness can not be optimized for such high values when applying the alloy rules described above. In addition this has a considerably higher cost, so, obviously, it is not the preferred way for most applications, but it can be an advantage for some very specific cases. For example, if a special oxidation color is convenient, or when the hardenability of the perlite ferrite is to be expanded and the use of other much more effective elements is not recommended. In such a case, the parameter K has to be selected to be quite low, in fact it should be less than 0.81, preferably less than 0.79, and optimally less than 0.75. This has to happen for sufficiently large values of Ceq%, usually greater than 0.33, even larger than 0.35 and sometimes greater than 0.41.

The instructions of this invention can be applied to the compositional range described 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 instructions work for values higher than 3.3%, and even more generalized in terms of applications for the values of% Moeq (real) > 3.6% and when Moeq% (real) > 3.8%, then the density of compositions that can reach a high thermal diffusivity and wear resistance within the range is much higher, and covers most applications (an exception is, for example, applications with exceptional hardness or wear resistance) . In the same way when it comes to% Ceq, while the instructions of the present invention already work for values higher than 0.31%, when% Ceq > 0.33%, and even more for Ceq% > 0.36% the density of compositions that can reach a high thermal diffusivity and resistance to wear within the range is much higher, and covers most applications (an exception is, for example, applications with exceptional hardness and wear resistance).

To increase the machinability can be used S, As, Te, Bi or even Pb, Ca, Cu, Se, Sb or others, with a maximum content of 1, with the exception of Cu, which can be up to 2. The substance More commonly, sulfur has, by comparison, a slight negative effect on the thermal conductivity of the matrix at the levels normally used to increase the machinability. However, its presence has to be balanced with the Mn to try to have it all in the form of spherical manganese disulphide, less harmful to the tenacity, as well as the smallest possible amount of the two remaining elements in solid solution in case thermal conductivity has to be maximized.

Another hardening mechanism can be used looking for some combination of specific mechanical properties or resistance to degradation caused by the work environment. Always try to maximize the desired property, but trying to have the least possible negative effect on thermal conductivity. The solid solution with Cu, Mn, Ni, Co, Si ... (including some carbide formers with less affinity to carbon such as Cr) and solid interstitial solution (mainly C, N and B). For this purpose, precipitation can also be used, with an intermetallic formation such as Ni3Mo, NiAl, Ni3Ti ... (and in addition to Ni and Mo, the Al, Ti elements can be added in small quantities, special care must be taken with Ti, which dissolves in M3Fe3C carbides and a maximum of 2% will be used). Finally, other types of carbides can also be used, but it is often quite difficult to maintain a high level of thermal conductivity with them, unless the carbide formers have a very high affinity with carbon, as discussed throughout this document. The Co can be used as a hardener in solid solution or as a catalyst for the precipitation of Ni intermetallic, rarely in contents higher than 6%. Some of these elements are also not so harmful when they dissolve in the carbides M3Fe3C, or other carbides of (Fe, Mo, W), this is especially the case of Zr and Hf, and with a smaller measure of Ta and these can also limit the solubility of V and Nb.

When the quantity is measured in weight percentages, the atomic mass and the type of carbide formed determine whether the quantity of an element used must be large or small. So, for example, 2% V is much more than 4% W. V tends to form types of MC carbides, a unless dissolved in other existing carbides. Thus, to form a carbide unit, only one unit of V is needed, and the atomic mass is 50.9415. W tends to form carbides of the M3Fe3C type in steels for hot work. Thus, three units of W are needed to form a carbide unit, and the atomic mass is 183.85. Therefore, 5.4 times more carbide units can be formed with 2% V than with 4% W.

The 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, electric furnace melting, vacuum induction casting. It is also possible to use powder metallurgy processes together with any type of atomization and, eventually, compaction such as HIP, CIP, cold or hot pressing, sintering (with or without liquid phase and despite the way in which the sintered, both simultaneously in all material, layer by layer or localized); laser cusing, spray forming, thermal projection or thermal coating, cold projection to mention some of them. The alloy can be obtained directly with the desired shape or improved by other metallurgical processes. Any refining metallurgical process can be applied, such as VD, ESR, AOD, VAR ... The slab or the laminate are frequently used to increase the toughness, even the three-dimensional forging of blocks. The tool steel of the present invention can be obtained in the form of a bar, wire or powder (among others to be used as welding or welding alloy). A matrix with a low cost casting alloy can even be made and the steel of the present invention applied to the critical parts of the matrix by bar or wire welding made from a steel of the present invention. It can also be welded by laser, plasma or electron beam using powder or wire made with the steel of the present invention. The steel of the present invention could also be used with a thermal spray technique to apply on 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 if it is embedded as an independent phase, or is achieved as one of the phases in a multiphase material. Also if it is used as a matrix in which other phases or particles are embedded whatever the method of realization of the mixture (for example, mechanical mixing, attrition, projection with two or more hoppers of different materials ...).

The tool steel of the present invention can also be used for the manufacture of parts subjected to large thermomechanical loads and wear, or basically for any part prone to failure due to wear and thermal fatigue, or with requirements of high wear resistance and that takes advantage of its high thermal conductivity. The advantage is a faster thermal transport or a lower working temperature. As an example: components for combustion machines (such as rings of the engine block), reactors (also in the chemical industry), heat exchange devices, generators or, in general, any machine for the transformation of energy. Matrices for the floor (in open or closed matrix), extrusion, lamination, casting and thixing of metals. Matrices for the plastic forming of thermoplastic and thermostable materials in all their forms. In general, any type of die, tool or piece can benefit from an increase in wear resistance and thermal fatigue. Also matrices, tools or pieces that benefit from a better thermal management, as it is the case of the matrices of conformed or of cut of materials with liberation of big quantities of energy (like the stainless steel or the TRIP steels) or working at high temperatures (hot cutting, sheet hot forming).

Further embodiments are described in the dependent claims.

EXAMPLES Some examples indicate how the steel composition of the invention can be specified more precisely for different types of hot working applications: Example 1 For sheet hot forming dies (Hot Stamping or Press Hardening). In this case the maximum possible thermal diffusivity is desired at high hardness levels. The desirable wear resistance depends on the coating of the formed sheet.

- Plates coated with Zn, AISi or other inorganic coatings (these same compositions are optimized for the manufacture of matrices for the injection of thermoplastics, especially if the steels described below are made by powder metallurgy): For this purpose, in the framework 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%) Yes: < 0.15% (preferably% Si <0.1, but with an acceptable level of oxide inclusions) Mn: < 0.5% ?? ^: 3.5 - 5.5 where ?? «, =% Mo + l / 2% W and Cec =% C + 0.86 *% N + 1.2 *% B The other elements should remain as low as possible and, in any case, always be below 0.15%, with the exception of strong carbide formers (% Ta,% Zr,% Hí). All values are indicated in weight percentages.

Three examples illustrate the properties that can be obtained: In all cases, the heat treatment that maximizes the dirusivity at the indicated hardness has been applied, minimizing the presence of elements in solution with the matrix except for% Cr, and especially minimizing the presence of% C and to a lesser extent% V in matrix. In all cases this implies very high austenitizing temperatures, from 3 to 5 tempers with the last in the range 600-640 ° C.

Advanced optimization is obtained by using elements that react strongly with% C to form carbides (also with% N and% B). Several examples illustrate the properties that can be obtained.

* In all cases the heat treatment has been applied, which maximizes the diffusivity to the indicated hardness, minimizing the presence of elements in solution with the matrix except for% Cr, and especially minimizing the presence of% C and to a lesser extent% V in the matrix. In all cases this involves very high austenitization temperatures, from 3 to 5 tempers with the last in the range 610-680 ° C. Being% Hf: 0.10 - 0.22,% Zr: 0.05 - 0.18 and% Nb: around 0.07, unless specifically indicated.

- Non-coated sheets and therefore with iron oxides which can be large: For this purpose, in the framework 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%) Yes: < 0.5% Mn: < 1.0% Mo ^: 3.5 - 9 where ?? ^ -% Mo + l / 2% W and Ceq =% C + 0.86 *% N + 1.2 *% B The other elements should remain 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.

Two examples illustrate the properties that can be obtained: * In all cases, the thermal treatment that maximizes the diffusivity to the indicated hardness has been applied, minimizing the presence of elements in solution with the matrix, and especially minimizing the presence of% C and to a lesser extent% V in the matrix. In this case, the largest possible presence of primary carbides is also sought. In all cases this implies very high austenitization temperatures, from 2 to 4 tempering with the last in the range 550-620 ° C. Being% Hf: 0.10 - 0.22,% Zr: 0.05 - 0.18 and% Nb: around 0.07, unless specifically indicated.

Example 2 For the forge in closed matrix. In this case, a simultaneous optimization of the wear resistance and the thermal fatigue has to be achieved, thus a maximum thermal diffusivity and wear resistance (presence of primary carbides) are desirable, keeping the CVN maximized as well. For matrices or large pieces subjected to shock or thermal fatigue, it is convenient to maintain a good CVN, even when the treatment can not be totally martensitic, in which case Si or Al is used to hinder the precipitation of coarse cementite (Fe3C), or % Ni is used to improve the hardenability in the ferritic-pearlitic zone and decrease the coefficient of linear thermal expansion. In this case, tool steels can be used in the following range (powder metallurgical steels except for applications where the abrasive particles present are large). The steels of the present invention are especially interesting for applications where wear is the predominant failure mechanism: For this purpose within the framework of the present invention a compositional range of the following type can be used: Cec,: 0.3 - 0.6 Cr < 0.1% (preferably Cr < 0.05%) Yes: < 1.4% Al: 0 - 2% Mn: < 1.5% ?? «,: 3.0 - 7.0 where ?? «, =% Mo + l / 2% W The other elements should remain 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.

Five examples illustrate the properties that can be obtained: * In all cases, the thermal treatment that maximizes the diffusivity to the indicated hardness has been applied, minimizing the presence of elements in solution with the matrix, and especially minimizing the presence of% C and to a lesser extent% V in the matrix. In this case, the largest possible presence of primary carbides is also sought. In all cases this involves very high austenitization temperatures, from 3 to 5 tempers with the last in the range 590-660 ° C. Being% Hf: 0.10 - 0.22,% Zr: 0.05 - 0.18 and% Nb: around 0.07, unless specifically indicated.

Example 3 Some applications of closed forge, predominantly require resistance to high temperatures, good hardness, especially resistance to fracture and CVN, and the best possible wear resistance. When the contact times are long, or the temperature of the forged part is high, the thermal diffusivity at high temperatures and good resistance to tempering are of utmost importance. In this case, the correct use of% Co is very important. For this purpose, in the context of this invention, a composition range of the following type can be used: Ceq: 0.32 - 0.7 V: < 2.8% Yes: < 1.4% Mn: < 1.5% Co: 1.3 - 6% MOeq: 3.3 - 7.0 where Moeq =% Mo + l / 2% W The rest of the elements should be kept as low as possible and in no case, always below 0.15%, with the exception of strong carbide formers (% Ta,% Zr,% Hf). All values are given in percent by weight.

Five examples show the properties that can be obtained: In all cases, the heat treatment that maximizes the diffusivity at the indicated hardness has been applied, minimizing the presence of elements in solution with the matrix, and especially minimizing the presence of% C and to a lesser extent% V in the matrix. In all cases this implies very high austenitization temperatures, from 3 to 5 tempers with the last in the range 640-690 ° C. Being% Hf: 0.02 - 0.16,% Zr: 0.05 - 0.18 and% Nb: around 0.07, unless specifically indicated.

Example 4 For hot cutting of sheets. In this case, the wear resistance has to be maximized, with a good hardenability and tenacity (in this case resilience). The thermal conductivity is very important to keep the temperature at the cutting edge as low as possible. Weldability is less important in this case, and small size inserts are often used, so compositions with high alloy content can be used. For this purpose, in the framework of the present invention, the following compositional range can be used: Ceq: 0.5 - 0.9 Cr < 0.1% (preferably Cr < 0.05%) Yes: < 0.15% (preferably Si <0.1%) V: 0-2% for cases with Moeq > 5 and V: O - 4% for cases with MOeq < 5 MOeq! 5 - 10 where ?? ^ =% Mo + 1 II% W The other elements should remain as low as possible and, in any case, always be below 0.15%, except for carbide formers, strong (% Ta,% Zr,% Hf). All values are indicated in weight percentages.

Three examples illustrate the properties that can be obtained: * In all cases, the thermal treatment that maximizes the diffusivity to the indicated hardness has been applied, minimizing the presence of elements in solution with the matrix, and especially minimizing the presence of% C and to a lesser extent% V in the matrix. In this case, the largest possible presence of primary carbides is also sought. In all cases this implies very high austenitizing temperatures (1 120 ° C in the first two cases and 1240 ° C in the latter case), from 2 to 4 tempers with the last in the range 600-640 ° C.

Claims (26)

REIVI DICATIONS:

1. A steel, in particular a steel for working tools in hot following composition, all the percentages indicated being percentages of weight: % Ceq = 0.3 - 0.9% C = 0.3 - 0.9% N = 0- 0.6% B = 0- 0.6 % Cr < 2.8% Ni = 0-3.8% Si = 0- 1.4% Mn = 0-3 % A1 = 0-2.5% Mo = 0-10% W = 0-10% 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, where % Ceq =% C + 0.86 *% N + 1.2 *% B, characterized because % ?? + ½ ·% W > 3.0.

2. A steel according to claim 1, wherein: when% Ceq is < 0.35, then K > 0.75, or when% Ceq is > = 0.35, then K > 0.84, or when% Ceq is > = 0.35, then% Hf +% Zr +% Ta +% Nb > = 0.01, being: K =% C6q / (0.4 + (% Moeq (real) -4) * 0.04173), and roMOeq ^ eal) =% ?? + 0.52 *% W.

3. A steel according to claims 1 or 2, wherein: "/ oMOeq ^ l) > 3.3%

4. A steel according to any one of claims 1 to 3, wherein: % V +% Nb +% Hf +% Zr > 0.1

5. A steel according to any one of claims 1 to 3, wherein: % V +% Nb +% Hf +% Zr > 1.2

6. A steel according to any one of claims 1 to 5, wherein: % Ceq > 0.32 and% C > 0.32

7. A steel according to any one of claims 1 to 5, wherein: % Ceq > 0.36

8. A steel according to any of claims 1 to 5, wherein: C > 0.4

9. A steel according to any of claims 1 to 8, wherein: % Mo + ½|% W < 10.0

10. A steel according to any of claims 1 to 8, wherein: % Mo + ½ ·% W < 4.5 with% Mo = 0 - 4.5 and% W = 0 - 9

11. A steel according to any of claims 1 to 10 with the proviso that: When% Ceq < 0.35, then% V < 1.7

12. A steel according to any of claims 1 to 10, wherein: % V < 1.8

13. A steel according to any of claims 1 to 12, wherein: % Nb < 0.09

14. A steel according to any of claims 1 to 13, wherein: % Ni < 2.99

15. A steel according to any of claims 1 to 13, wherein: % Ni < 1.0

16. A steel according to any of claims 1 to 15 wherein: when% Cr > 2, then% Nb +% Ta +% Zr +% Hf > 0.2

17. A steel according to any of claims 1 to 16 wherein: % Ceq > 0.32 % MOeq > 3.2 and % Cr < 2.5, with the condition of: when Ceq < = 0.36, then: 3.56 < % MOeq /% Ceq < 1 1.5, or when 0.36 < Ceq < = 0.38, then: 3.56 < % Moeq /% Ceq < 14, or when 0.38 < Ceq, then: 3.56 < % Moeq / < 16.8, being: ^ oMOeq =% ?? + ½ ·% W

18. A steel according to any of claims 1 and 3 to 17, wherein: % Ceq > = 0.33 and K < 0.81, being: K =% Ceq / (0.4 + (% MOeq (real) -4) * 0.04173), and % ?? ß, (, »?) =% Mo + 0.52 *% W.

19- A steel according to any of claims 1 to 18 which, when subjected to a martensitic, bainitic or martensitic-bainitic temper with at least one temper at a temperature higher than 590 ° C, a hardness greater than 47 HRc can be obtained with a structure of low dispersion characterized by a diusivity of 9 mm2 / s or more.

20. A steel according to any one of claims 1 to 19, when subjected to at least one tempering cycle at a temperature of 590 ° C, a hardness of 53HRc or more can be obtained with a low dispersion structure characterized by a thermal diffusivity by above 9 mm / s.

21. A steel according to any of claims 1 to 20 wherein: % C > 0.32, % Co > 1.3 and % V < 2.8.

22. A steel according to claim 21 wherein, when subjected to at least one annealing cycle at a temperature above 660 ° C, a hardness of 50HR.C or more can be obtained with a low dispersion structure characterized by a diffusivity 5.8 mm2 / s or more at 600 ° C.

23. A die, tool or part comprising at least partially a tool steel according to any one of claims 1 to 22.

24. A process for manufacturing a tool steel, characterized in that a steel according to any of claims 1 to 18 is subjected to martensitic, bainitic or martensitic-bainitic tempering with at least one tempering cycle at a temperature above 590 ° C, so that a steel with a hardness above 47 HRc can be obtained with a low dispersion structure characterized by a diffusivity of 9 mm2 / s or more.

25. A process for manufacturing a tool steel according to claim 24, wherein a steel with hardness above 53 HRc can be obtained with a low dispersion structure characterized by a diffusivity of 9 mm2 / s or more.

26. A process for manufacturing a tool steel according to claim 24, wherein the steel is subjected to at least one tempering cycle at a temperature above 660 ° C, so that a steel with a hardness of 50 HRc or more can be obtained with a low dispersion structure characterized by a diffusivity of 5.8 mm2 / s or more at 600 ° C.