OPTIMIZING MICROSTRUCTURE FOR HIGH TOUGHNESS

OPTIMIZING MICROSTRUCTURE FOR HIGHTOUGHNESS COLD-WORK TOOL STEELS

D. Viale, J. Béguinot, F. Chenou and G. BaronUSINOR INDUSTEEL

Abstract Increasing toughness and machinability at given high levelof wear resistanceare consistently growing requirements for cold-work tool steels. Improv-ing microstructure characteristics, especially coarse carbides distribution andtheir chemical composition, reveals an appropriate way to meet such require-ments. Referring to the archetypal cold work tool steel AISID2, an improve-ment of coarse carbides hardness by higher alloying with strong carbidesformers allows a moderate reduction of their volume fraction, resulting inincreased toughness and machinability performances. Also, the increase ofultimate resistance of surrounding matrix by improved secondary hardeningpreventing premature pulling off of carbides in service contributes to longerservice life, while reasonably increased silicon content leads to still bettermachinability. A further step towards increased toughnessand machinabilitymay result from slightly refining the coarse carbides sizes through moderateaddition of fine titanium nitrides acting as precipitation promoters for M7C3

type carbides.

INTRODUCTION

Cold-work tool steels have been developed and used for more than acentury and have been designed, mainly on an empirical basis[1] in orderto cope with a large variety of often contradictory properties, among which:

high strength level to resist against permanent deformations resultingfrom high levels of applied stress

wear resistance during use, including abrasive wear, adhesive wear,surface fatigue

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300 6TH INTERNATIONAL TOOLING CONFERENCE

toughness (fracture resistance and also fatigue resistance)

dimensional stability during application (thermal treatment and sub-sequent use)

uniformity and isotropy of microstructure

acceptable machinability, at least in the annealed state

acceptable corrosion resistance, especially against pitting corrosion insome demanding applications

acceptably limited susceptibility to excessive hardeningand associ-ated crack sensitivity in thermal affected zones, related to EDM andweld deposits.

Conventional cold-work tool steel such as 12 % chromium, 1,5% carbon(AISI D2 type) have long proved to be a satisfying solution especially re-garding an equilibrated answer between deformation, wear and corrosionresistances and dimensional sensibility. On the other hand, the high vol-ume fraction (∼ 10 to 15 %) of coarse (∼ 20 µm) eutectic M7C3carbidesin these steels is largely responsible for low levels of toughness since thesecarbides are intrinsically prone to easy cracking and contribute thus to ex-cessive sensitivity of steel to fracture initiation and propagation. Timesgoing on, with steadily increasing deformation stresses applied to workingmaterial, it appeared that insufficient toughness becomes far more a causeof failure of tools with AISI D2 type steels than was an insufficient wearresistance. Accordingly, increasing demands from tool manufacturers andend users became in favour of improved fracture resistant cold work toolsteels. In this respect, and considering the major detrimental role of coarseeutectic M7C3chromium- molybdenum carbides on the generation cracksduring use, it seems quite logical to try to modify the steel composition ofD2 type steel in order to decrease the volume fraction of these eutectic car-bides, and in this scope, to reduce significantly the carbon and chromiumcontents which govern, in major part, this volume fraction.On the otherhand, as these carbides contribute strongly to the wear resistance, it seemsnecessary to compensate their lower volume fraction by a still higher intrin-sic hardness. This is likely to be obtained by a significant enrichment ofstrong carbides-forming elements like molybdenum and even, if necessary,

Optimizing Microstructure for High Toughness Cold-Work Tool Steels 301

by a complementary contribution of very strong MC carbides forming ele-ments such as vanadium or niobium, which are in moderate content in D2type steels. In addition, the higher bulk content of strong carbides-formingelements may also contribute to improve wear resistance of the steel throughtheir residual contents in the matrix. The stronger secondary precipitationof Mo, V, Nb-enriched fine carbides should make the matrix itself moreresistant to wear solicitations and thus reinforce the ability of the matrixto resist to the pulling off of the coarse carbides during service, in severeconditions. Also, to limit in an other way the detrimental effect of coarsecarbides on steel toughness, it seems as possibly useful to retain a mini-mum quantity of austenite (even after moderate tempering) embedding thecoarse carbides. Ductile austenite acts as a mean to reduce stresses concen-trations around these carbides during use and, by the way, the solicitationsfor premature cracking. At this point of view, sufficient silicon addition,which increases carbon solubility in austenite and thus acts indirectly as anaustenite stabilisating factor, may be considered (silicon is also interestingfor its generally recognised beneficial role on machinability answer). Even-tually, the alloys equilibrium for such a newly designed tool steel, especiallyregarding chromium and carbon contents, should be preferably optimisedregarding dimensional stability of the steel and also referring to corrosionresistance. Here, it is not the bulk alloys contents but the contents in thematrix, as depleted in carbon and carbides formers by the previous precip-itation of eutectic carbides, which has to be considered. AsD2 type steelsmay be considered as fairly well optimised regarding dimensional stability,the trend should be that the matrix composition of the new steels lay closeto that of D2 matrix. As concern corrosion resistance, the parallel decreaseof the bulk contents of chromium and carbon may lead to a substantiallyunchanged level of passivating soluble chromium in the matrix. In addi-tion, the significant increase of both bulk and solute molybdenum contentis intended to promote a higher pitting corrosion resistance [2]. As regardother properties, lowering carbon content looks positive as a way to reduceexcessive hardening and crack sensitivity of thermal affected zones, alongEDM cuts or welding deposits. Considering machinability, things are a lotcomplex since reduced volume fraction of carbides looks favourable and, onthe other hand, higher enrichment of carbides with molybdenum and vana-dium/niobium looks unfavourable. At least for the specific steel conceptionpresented hereafter, the balance reveals positive with significant improve-


ment of machinability observed in practice, both in annealed and end-treatedconditions. Incidentally this positive balance between contradictory effectsof reducing volume fraction and increasing intrinsic hardness of carbides isin agreement with similar observations on quite different types of steels [3].These metallurgical trends had substantiated considerable work, especiallyamong Japanese researchers, for example MATSUDA and SUDOH,andwere at the origin of the so called "8 % Cr. – 1 % C" new generation of coldwork tool steels. The "core composition" of this new conceptis: 1 % carbon,8 % chromium, 2.5 % molybdenum, 0.2% to 0.6% V + Nb/2, 1 % silicon, tobe compared to the typical composition of AISI D2 type steel:1.5 % carbon,12 % chromium, 0.8 % molybdenum, 0.25 % vanadium, 0.3 % silicon. Thisevolution of alloy composition puts in a concrete form the application ofthe metallurgical trends described above. Several tool steels manufacturerssubsequently derived a lot of variants, around this centralconcept.

TOWARDS A STILL FINER AND REGULARDISTRIBUTION OF COARSE EUTECTIC CARBIDESMICROSTRUCTURE

Undoubtedly, the reduction of coarse eutectic carbides volume fractionacts decisively for improvement of the cold work tool steelstoughness, ac-cording to the new concept [4]. But the accompanying decrease of coarsecarbides average sizes (reduced to∼ 5 – 10 µm as compared to∼ 15 –20 µm for D2 type steels) is likely to play also a significant role in this re-spect [5]. Indeed, refining the microstructure and especially coarse carbidessizes proves very efficient, for example when comparing results obtained bypowder metallurgy and by conventional processes, applied to the same steelcomposition with similar hot rolling ratios. Unfortunately, specific steelmaking routes, such as powder metallurgy, remain quite expensive. How-ever, from laboratory and industrial experiments, it was recently observedthan it is possible to improve the microstructure and toughness through spe-cific micro-additions. Namely, the addition, in small contents, of elementsof the titanium family, in the melt prior to casting, under severely controlledconditions, proves to be efficient in this respect. This refining effect is ten-tatively attributed to an indirect consequence of the fine precipitation ofsmall titanium nitrides directly in the melt, according to the very high ther-mal stability of these compounds. These fine nitrides particles in spite of


Table 1. Chemical analysis

C Mn Cr Mo V Others

X160Cr Mo V12 / D2 1.55 0.35 11.75 0.75 0.95 —TENASTEEL 1.0 0.35 7.5 2.6 0.3 Ti

their buoyancy in the melt, may act as promoters for carbidesprecipitation,accordingly making carbides more numerous and consequently finer. Infact this hypothetical mechanism would, logically, more readily address theprecipitation of primary carbides stricto-sensu rather than the more delayedprecipitation of eutectic carbides considered here. However, whatever be theactual mechanism involved, the average size observed for eutectic carbideswas 3.3 µm to be compared to 6–7 µm without titanium addition.Tough-ness was correspondingly improved from 20 % to near 40 % depending onthermal treatments applied.

METALLURGICAL CONCEPT FOR THEDEVELOPMENT OF THE NEW COLD WORK TOOLSTEEL OF USINOR INDUSTEEL : TENASTEEL

According to the previous considerations, the TENASTEEL differs fromstandard grade X160 Cr Mo V12 / D2, Table 1, by 3 main points :

decreasing in the carbon and chromium contents⇒ increasing of thetoughness

increasing of the molybdenum content to keep quenchability, hardnessand wear resistance

addition of titanium to refine the structure through fine precipitationof titanium nitrides.

High chromium and carbon contents always induce formation of coarseeutectic chromium carbides (Fig. 1) effective in term of wear resistance,but principal causes of brittleness of steel X160 Cr Mo V12 / D2 type.The concentration of these large carbides will be even more intense at midthickness of the products (segregated lines). Conversely,a low carbon andchromium content guarantees:


a much more finer carbide distribution (Fig. 1)

a better homogeneity in the thickness.

These are very good things to improve the toughness, the machinability orthe polishability of steel.

Figure 1. Structures of D2 and TENASTEEL

Moreover, a characterisation of the precipitates in the D2 and TENAS-TEEL grades was carried out by electron microprobe analysis. In the D2grade steel, there is only one carbide type with the M7C3 stoechiometrie.Of course, the Table 2 shows that it is chromium carbide with alittle bit ofvanadium, molybdenum and iron.

In the TENASTEEL, we can find this chromium carbide so, but moreoverthere are molybdenum carbides (M23C6 type) and titanium carbo-nitrides(M4(C,N3)) type). The substitution of chromium carbides by molybdenum


Table 2. Chemical composition of the carbides precipitates in D2 andTENASTEEL grade

Atomic % C Si Mn Cr V Mo Ti Fe

D2 M7C3 31,33 0,01 0,32 34,64 5,26 0,85 / 27,55TENASTEEL M7C3 31,91 0,03 0,36 31,01 3,14 3,82 0,09 29,64TENASTEEL M23C6 20,73 4,55 0,17 7,22 1,45 30,21 0,25 35,39TENASTEEL M4(C, N3) 0,33 0,03 0,55 4,09 1,35 39,5 1,98

carbides gives a good wear resistance to the TENASTEEL because of theirhardness (Mo carbides: 1800 HV, Cr carbides: 1500 HV).

HEAT TREATMENTS AND MECHANICALPROPERTIES OF TENASTEEL

To obtain the mechanical characteristics on a given steel, with a goodcompromise between strength and toughness, it is essentialto optimize itsmetallurgical structure as well as the size, the distribution, the density, thehomogeneity of the carbides which it contains. This double aim can beachieved by an adapted composition, as it is the case for TENASTEEL, butthat is not enough. The mechanical characteristics of a steel with a givencomposition can be widely improved by the heat treatments which it willundergo. While influencing its structure, they will be able to increase ordecrease the strength of metal and decrease its brittleness. A heat treatmentdoes not change the chemical composition of metal, but it modify (Fig. 2) :

its structure by controlling of carbide precipitation (size, distribu-tion…) as well as the control of the nature and the proportionof thecomponents (ferrite, austenite, martensite…) ;

the mechanical equilibrium in the metal (internal stress, expansion…).

Technically, a heat treatment is defined by a variation in temperatureaccording to time. A thermal cycle practised on steel can be divided intothree distinct stages :

a reheating to the desired temperature ;

a stage at the temperature defined according to the practicedheat treat-ment, and depending on its final purpose (homogenization, hardening,softening, increase in ductility, internal stress relaxation…)


Figure 2. Incidence of the structure on the steel properties.

a cooling speed will fix the structure of metal in terms of componentsand precipitation where the several speeds of cooling can follow oneanother before reaching the temperature of end of processing.

The implementation of the heat treatments thus requires thecomprehensionof the principal phenomena involved i.e. especially for tool steels, precipita-tion and dissolution of carbides, as well as the evolution ofstructures, theirtransformations and conditions under which they occur. In order to facilitateits machining, TENASTEEL is delivered in annealed condition to give a lowhardness structure. Softened steel can be formed, but a heattreatment willbe then necessary to give the final mechanical characteristics to the pieces.It is a quenching to harden the metal followed by temperings to eliminateits brittleness and to increase its toughness. The processing of hardeningconsists in:


slow heating, to limit deformation and to avoid cracking (due tostresses) up to a temperature just below AC1, then holding (time de-pending on thickness) to homogenize the temperature through thewhole thickness. Then, re-heating up to austenitization temperature(>AC3) ;

holding at austenitization temperature to get an homogeneous temper-ature in the whole piece, to transform the steel into austenite and todissolve a maximum of carbides previously formed ;

cooling in an adapted cooling medium to get a martensitic structure.In order to get a martensitic transformation, it is necessary to have acoolingspeedhigher than critical quenching rate of the steel (minimumspeed allowing cooling without transformation into ferrite-pearlite).

The lower the critical quenching rate is, the more steel willbe able to hardendeeply. The hardenability of a steel depends primarily on its chemical com-position. All the alloy elements, except cobalt, tend to increase hardenability.The TENASTEEL exhibits a critical quenching rate relatively low. Its hard-enability is comparable with that of steel D2. After hardening, the structureof steel is not completely martensitic, there remains a partof austenite calledretained austenite, and carbides. The more the steel is alloyed and the largerthe temperature and time of reaustenitization are, the morethere is retainedaustenite. This complex structure shelters internal stresses which increasethe brittleness of the steel. To decrease the harmful effects of hardening, anew heat treatment will be perform on the pieces : the tempering which con-sists in carrying them at a temperature lower than AC1, to avoid modifyingthe crystalline iron (α) structure, then to cool them quickly. The reheatingof martensite tends to bring back it in a state of balance because the carbonis rejected out of the structure and precipitates to give iron ε carbide (Fe2C)and cementite (Fe3C). This precipitation is accompanied by a contraction ofmetal and by a reduction in hardness (internal stress relaxation).

This softening due to the transformation of martensite, is attenuated by ahardening caused by the transformation of retained austenite in secondarymartensite or bainite during cooling ; this reaction is accompanied by anexpansion of metal. A second tempering is generally practiced to transformthis new martensite. Lastly, during temperings carried outat high tempera-ture (starting from 500◦C), a secondary hardening is also produced by thespecial carbides precipitation : vanadium and molybdenum carbides in the


case of TENASTEEL. This new hardening is also accompanied byan ex-pansion of metal. In summary, softening with tempering results from severalsimultaneous phenomena (Fig. 3):

the softening of martensite (primary then secondary)

the transformation of residual austenite

the special carbide precipitation if the tempering is carried out at hightemperature.

The thermal cycle of hardening - tempering to TENASTEEL takes intoaccount of these metallurgical considerations as schematised Fig. 4. Thereheating of austenitization will be practiced under vacuum, or at least in acontrolled atmosphere to prevent the risks of decarburization of steel. Thetemperature of reaustenitization can be selected between 1000 and 1100◦C.The hardness evolution of TENASTEEL after complete heat treatment isshown on Fig. 5 according to austenitizing and tempering temperatures. Itshould be noted that this very great interval of austenitizing temperaturesmakes it possible to be compatible with the temperatures usually used formany other steels (D2 in particular). This allows an optimization of furnaceproductivity, and thus a reduction of the costs of heat treatment, as well as areduction of the risks of errors related to the non-observance of austenitizingtemperatures.

Whatever the austenitizing temperature, a hardness range between 58and 62 HRC (standard of use for this type of steel) can be obtained if theTENASTEEL undergone two temperings between 500/550◦C(930/1020°F)to575◦C(1065°F). While anaustenitization between1000 and 1100◦C(1830/2010°F)leads to a good hardness of our steel, the best properties will be obtainedafter reheating around 1030 - 1050◦C(1885- 1920°F). Indeed, Fig. 6 showsthat the toughness of metal is maximum in this range of temperature.

Indeed, if the austenitizing temperature is too low, a largepart of finechromium molybdenum carbides will not be dissolved. They will remaincoarse and will not increase hardness and wear resistance ofsteel. In ad-dition, for the highest reheating temperatures, the hardening obtained bythe refinement of subsequent secondary precipitation of carbides is counter-balanced by softening due to the increasing in retained austenite rate afterhardening. A third tempering should be then necessary to completely desta-bilise this retained austenite. Number and temperature of temperings used


Figure 3. Various metallurgical phenomena leading to a softening with the tempering.

Figure 4. Thermal cycles of hardening and tempering practices on TENASTEEL.


Figure 5. Evolution of hardness of TENASTEEL with austenitizing and tempering tem-peratures.

to soften martensitic structure after quenching will allowTENASTEEL toobtain final mechanical characteristics and final using properties. The soft-ening curves of the TENASTEEL are compared with those of the D2 Fig. 7for an austenitizing temperature of 1050◦C(1920°F).

These softening curves of TENASTEEL make it possible to drawsomeinteresting conclusions :

TENASTEEL and D2 grades are treated in the same ranges of tem-peratures,

for an identical temperature of tempering, the TENASTEEL isharderthan D2,

lastly, TENASTEEL makes it possible toobtainhighhardnesses (> 60 HRC)after tempering at high temperature (500 – 550◦C).

This last possibility is a very good advantage regarding theaptitude forthe surface coating which requires for nitriding (gas, bathof salts, ionic...)or PVD, for example, relatively long holding time at high temperatures.These curves of soften show that for processing performed between 550 and


Figure 6. Effect of austenitizing temperature on the toughness of TENASTEEL

Figure 7. Softening curve of TENASTEEL compared with that of the D2.


575◦C, the TENASTEEL is able to keep a hardness of the matrix higherthan 60 HRC, whereas the Z160CDV12 sees its hardness breaking down (50– 58 HRC) in this temperature range.

The evolution of mechanical characteristics obtained after heat treatmentis shown on Fig. 8.

Figure 8. Evolution of mechanical characteristics of TENASTEEL following hardnessobtained after heat treatment.

The toughness of TENASTEEL strongly grows up with the reduction inthe hardness of steel, whereas its abrasive wear resistanceincreases slightlywith hardness. The best compromise, for a standard application is obtainedafter a double tempering between 525 and 575◦C. Moreover, the Fig. 9shows than the toughness of TENASTEEL is twice better than ofD2 steel.And this is always true in the range of the hardness used for these coldwork tool steels : 58 – 62 HRC. In the other hand, the abrasive wear resis-tance of the two steels is comparable, then, TENASTEEL exhibits the bestcompromise between wear resistance and toughness.


Figure 9. The toughness of TENASTEEL grade is almost twice better thanof D2.

SURFACE TREATMENT

In order to increase their resistance to seizing up and to minimize thefriction in service, tool steels are more and more frequently surface treatedby nitriding or coated by metal deposits. This surface treatment also makes itpossible to increase the surface hardness of pieces and to increase the tool lifesubjected to abrasive and/or adhesive wear. Nitriding is a thermochemicalprocess of hard facing by atomic nitrogen diffusion on the surface of thepieces previously treated by hardening and tempering. The insertion ofnitrogen atoms and the nitride formation with steel alloying elements, inducea hardening of surface (750 to 1400 HV) bringing the requiredproperties :

improvement of resistance to wear and seizing up of materials ;

increasing of the stress limit of material because of the compressivestresses created by the processing ;

maintain of metallurgical structures of the material and thus of itsinternal mechanical characteristics if the tempering has been carriedout at a temperature higher than that of the surface treatment.


Comparative gaseous nitriding tests were performed at 525◦Con TENAS-TEEL and X160 Cr Mo V12 / D2, both heat-treated to 60 HRC. NitridedTENASTEEL layer appears homogeneous in depth and morphology, whilethat of D2 reveals a lot of carbides and exhibit an irregular depth (Fig. 10).For TENASTEEL, depth of nitrided layers measured by micro-hardness

Figure 10. After gaseous nitriding at 525◦C, TENASTEEL exhibits layers thicker andmore homogeneous in depth and morphology than X160CrMoV12 /D2.

readings or shown on micrographs are coherent, and the values obtained arerespectively 60, 80 and 120 µm after treatment times of 4, 8 and 16 hours.For X160 Cr Mo V12 / D2, maximum depth reach only 50, 70 and 80 µmfor same treatment times. But the Fig. 10 shows that in some places thenitrided layer can drop down to 20 µm after 4h of treatment. Moreover, dueto the quantity of coarse carbides in the nitrided layer of X160 Cr Mo V12/ D2 and its interface with the substrate, a poor adhesion canbe expectedwith possible shipping of the nitrided layer. The second very important


Figure 11. Comparative hardness of TENASTEEL matrix and X160 Cr Mo V12/D2 matrixafter gaseous nitriding at 525◦C.

point to note is the influence of the hardness of the core of thepiece duringnitriding. The Fig. 11 shows hardness records measured on the matrix ofTENASTEEL and X160 Cr Mo V12 / D2 after 4, 8 16 hours of nitridingat 525◦C(975°F). Both steels were heat-treated to 60 HRC before nitriding.TENASTEEL keeps its initial hardness after gaseous nitriding, a treatmenttime of 16 H induce only a drop of 1 HRC. Conversely, the hardness of X160Cr Mo V12 / D2 is strongly affected as it drops down from 60 to respectively56, 55, and 50 HRC after 4, 8 and 16 hours of nitriding. This softening isnot surprising, looking at Fig. 7 showing the evolution of hardness versusthe holding time at 525◦C. Some other tests were carried out with ionicnitriding process at 500◦C(930°F). Like after gaseous nitriding, the layer atTENASTEEL surface seems to be homogeneous as well in thickness as inmorphology. Moreover, as already mentioned, layers on X160Cr Mo V12/D2 include much carbides and present a very irregular thickness.

For TENASTEEL, measurement of nitrided layers thickness obtained onmicrographs of the Fig. 12 or by micro-hardness measurements, are coher-ent and gives values of about 100 and 140 µm for 6 and 24 hours treatmenttime respectively. For X160 Cr Mo V12 / D2, thickness have a maximum


Figure 12. After ionic nitriding at 500◦C, TENASTEEL exhibits layers thicker and morehomogeneous in depth and morphology than X160CrMoV12 / D2 do.

size of 50 and 100 µm for a same treatment duration. TENASTEELallowsto minimise nitriding time (6h to obtain 100 µm on TENASTEEL and 24h to the same thickness on D2) and to obtain a same thickness layer allover the surface of the sheet. It has to be mentioned that the presence ofcoarse carbides in the nitrided zone and at the interface with substrate willreduce adhesion and lead to chipping of this layer in the caseof X160 Cr MoV12 / D2. Steels are heated at only 500◦C(930°F) for this treatment, thenTENASTEEL as well as X160 Cr Mo V12 / D2 save their matrix hardness.The cutting, forming tools, as well as molds elements for aluminum andplastic injection are frequently covered by titanium nitride which reducesto a significant degree the coefficient of friction, and very largely improvesthe abrasive and adhesive wear resistance. These coatings,obtained by va-por condensation on the surface of the substrate make it possible to form


a metal deposit, which will grapple to the heat-treated surfaces. For thistype of coating, the preparation of the surface of the substrate is an essentialstep. It will make it possible to solve possible problems of adherence by acleaning and an activation of surface. For these PVD (Physical vapor depo-sition) and CVD (Chemical Vapor Deposition) coatings the main advantageof the TENASTEEL compared to the X160 Cr Mo V12 / D2 is due to thesmoothness and the distribution of carbides. Indeed, the presence of largechromium carbides to the interface between the substrate and the coatingdecreases the adhesion of this one.

CONCLUSION

The aim of the new cold work tool steel grade with improved toughness,TENASTEEL, is to take the place of X160CrMoV12 / D2, currently mostwidespread on the market in spite of big problems of rupture,of damageby chipping or adhesion due to its too low toughness. Its chemical com-position was adapted in order to decrease the volume fraction of large pri-mary chromium carbides and improve the toughness of steel. An increase inmolybdenum content compensates this decrease of the carbonand chromiumcontent to preserve a good wear resistance because of finer and dispersedsecondary carbides contribution. Moreover, an addition oftitanium refinesthe structure. After austenitization at 1030 or 1050◦C, temperatures com-patible with the current practices for the processing of theother grades, theTENASTEEL will be tempered twice between 525 and 575◦Cto obtain astandard hardness ranging between 58 and 62 HRC. For particular applica-tions requiring a still increased toughness, tempering at higher temperaturecould be practised in order to decrease the hardness of steel. At equal levelof hardness, the toughness of TENASTEEL is twice better thanof D2 for awear resistance comparable. Moreover, the higher tempering temperaturesand the thinner carbides and structure confer to TENASTEEL avery goodaptitude for the surface coating :

to obtain homogeneous layers in thickness and in morphology;

to reduce heating time ;

to improve coating adherence ;

to save the matrix hardness.


Lastly, weldability, polishability and machinability of the TENASTEEL arealso higher than that of the D2. For example, compared with grades X160Cr Mo V12 / D2, TENASTEEL allows an increasing of the tool lifeduringmachining :

> 30% in softened condition

> 75% in hardened condition

The size and the dispersion of the carbides can explain thesegood prop-erties.

REFERENCES

[1] R. EBNER, H. LEITNER, F. JEGLITSCH, D. CALISKANOGLU "Tool Steels in thenext century" 5th International Conference on Tooling – Loeben 1999.

[2] H. BERNS "New Materials Processes Experiences for Tooling" International EuropeanConference on Tooling Materials Interlaken 1992.

[3] S. CORRE, C. LE CALVEZ, P. MABELLY, F. CHENOU, J. BEGUINOT"Tool Steelsin the next century" 5th Interntional Conference on Tooling– Loeben 1999.

[4] H. JESPERSON "Tool Steels in the next century" 5th Interntional Conference on Tooling– Loeben 1999.

[5] D. YOKOI, N. TSUJII European Patent Application EP 0930 374 A1.

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OPTIMIZING MICROSTRUCTURE FOR HIGH TOUGHNESS