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    • 专利摘要:
    A low-chromium hot-work tool steel consisting of (in wt-%): C 0.08-0.40 N 0.015-0.30 C +N 0.30-0.50 Cr 1-4 Mo 1.5-3 V 0.8-1.3 Mn 0.5-2 Si 0.1-0.5 optionally Ni <3 Co =5 B < 0.01 Fe balance apart from impurities, and A process for making a low-chromium hot-work tool steel article having increased tempering resistance.
    公开号:SE1150200A1
    申请号:SE1150200
    申请日:2011-03-04
    公开日:2012-09-05
    发明作者:Joergen Andersson；Henrik Jesperson；Andr N Hans-Olof；Lars-Erik Svensson
    申请人:Uddeholms Ab；
    IPC主号:
    专利说明:

    ol5 1 (23) HOT-WORK TOOL STEEL ANT) A PROCESS FOR MAKING A HOT-WORKTOOL STEEL DESCRIPTION TECHNICAL FIELDThe present invention relates to a process for making a low-chromium hot-Work toolsteel article. ltalso relates to a low-chromium hot-Work tool steel suitable for making the article. i BACKGROUND ART , The term “hot-Work tools” is applied to a great number of different kinds of tools for theWorking or fonning of metals at comparatively high temperatures, for example tools fordie casting, such as dies, inserts and cores, inlet parts, nozzles, ejector elements, pistons,pressure chambers, etc.; tools for extrusion tooling, such as dies, die holders, liners,pressure pads and stems, spindles, etc.; tools for hot-pressing, such as tools for hot-pressing of aluminium, magnesium, copper, copper alloys and steel; moulds for plastics, such as moulds for injection moulding, compression moulding and extrusion; together With various other kinds of tools such as tools for hot shearing, shrink-rings/collars and “ wearing parts intended for use in Work at high temperatures. Low-alloyed hot-Work toolsteel is used in small to medium sized tools in applications Where the demands ontempering resistance and thermal fatigue are high. Tempering resistance is the ability ofa hot-Work tool steel to keep its hardness at an elevated temperature for prolonged time.Hot-Work tool steels are developed for strength and hardness during prolonged exposureto elevated temperatures and generally use a substantial amount of carbide forming alloys.
    Another type of tool steels is the high speed steels, Which are used for cutting toolsWhere strength and hardness must be retained at temperatures up to or exceeding 760 °C. To reduce the amount of tungsten and chromium required, e. g. 18 and 4 wt-%,respectively, variants using molybdenum (5-10 wt-%) Were developed. High speedsteel differs from hot-Work steel in composition and price and cannot be used as a substitute for hot-Work steel.
    US 5,2'32,660 discloses a loW-alloy steel especially useful in the fabrication of hot Work . implements such as die blocks, dies, and other equipment for forging and other hot '10 2 (23) forming operations, and a method of manufacturing such steels. The low alloy steel hasa medium carbon content, and less than about 5 % Cr and less than about 1 % Mn, and it has hardenability approaching unity derived from the presence of N in an amount of from about 100 to about 400 ppm. It also discloses a method of manufacturing such a steel, which preferably contains the processing step in the late stages of manufacture of 'bubbling nitrogen gas through the molten steel. The preferred concentration of nitrogen is about 200 ppm. If the nitrogen concentration drops below about 100 ppm, thedisclosed desired objectives of through hardness and increased wear resistance will notbe met. On the other hand, nitrogen concentrations of greater than about 400 ppm arestated to require excessive amounts of aluminum, titanium, manganese and chromium inorder not to adversely affect the quality of the steel. Austenitizing and tempering of the steel are not mentioned.
    SUMMARY OF THE IN VENTION One object of the present invention is to provide a process for making a low-chromiumhot-work tool steel article having increased tempering resistance.
    In accordance with the present invention this object is achieved in that the processcomprises the steps of: a) incorporating nitrogen in a low-chromium hot-work tool steel composition having achromium content of at most 4 wt-% to forrn a steel composition having a nitrogencontent of 0.015 to 0.30 vvt-%; i b) forming a steel article from the steel composition; c) austenitizing the steel article obtained in step b) at a temperature of at most 1200 °Cfor a time on the order of half an hour followed by quenching; and d) tempering the quenched steel article at least twice at a temperature between 500 and700 °C for a time on the order of 2 hours.
    In a creep resistant steel having a high chromium content, i.e. 9-12 wt-% , it is possibleto dissolve vanadium carbide-nitrides already at relatively low temperatures, i.e. 1020-1050 °C. However, if the chromium content is low, less than about 4-5 wt-%, primary vanadium carbide-nitrides will be formed in the melt, and they are virtually impossible to dissolve afterwards.
    In the steel of the present invention, the total amount of carbon and nitrogen shall beregulated to 0.30 5 (C+N) í 0.50, preferably 0,36 3 (C+N) 5 0.44. The nominal contentshall be in the order of 0.40 Wt-%. At the same time, it is advantageous to regulate the 3 (23) nitrogen content to between 0.015 and 0.15 N, preferably 0.015 - 0.10 and carbon maypreferably be regulated to at least 0.20 wt-%.
    When the nitrogen content is balanced to about 0.05 to 0,10 wt-% vanadiumcarbonitrides will form, Which will be partly dissolved during the austenitizing step andthen precipitated during the tempering step as partícles of nanometer size. The thermalstability of vanadium carbonitrides is better than that of vanadium carbides, andconsequently the tempering resistance of the low-chromium hot-work tool steel articleWill be much improved. Further, by tempering at least twice, the tempering curve(showing hardness as a function of tempering temperature) will have a higher, secondary peak. ln the most preferred embodiment of the invention the nitrogen content preferably is onthe order of 0.05 wt-%. This value gives a better performance than higher values. Anitro gen content on the order of 0.05 Wt-% gives a higher potential for secondaryhardening during quenching than higher contents do, thus giving the steel a highhardness. However, an amount in the order of 0.10 wt-% has shown to give a shift ofthe secondary hardening peak to somewhat higher tempering temperatures which ispositive. Additionally, the performed tests and modelling calculations indicate that anincreased austenitizing temperature is required in connection With increased nitrogen contents.
    Chromium promotes the hardenability and corrosion resistance of steels. At too lowcontents the corrosion resistance will be adversely affected. A minimum chromiumcontent in the steel, therefore, is set to 1 wt-%. The maximum content is set to 4 Wt-% inorder to avoid undesired formation of chromium rich carbides/carbonitrides, e.g.
    M23C6.. The chromium content preferably shall not exceed 3 wt-%, and even morepreferred preferably not exceed 2.6 wt-%. In the most preferred embodiment of theinvention, the chromium content is l.5-1 .7 wt-%. A low chromium content delays theprecipitation of chromium carbides in the microstructure in favour of the morethermally stable vanadium-rich carbo-nitride. Thus the recovery is slowed down in the material and the tempering resistance becomes improved The steel shall contain vanadium in an amount of at least°0.8 wt-% in order to provide asufficient precipitation potential and thus an adequate tempering resistance and desiredhigh temperature strength properties. In order to avoid excessive formation of M(C,N) precipitates which would increase the risk of large undissolved precipitates remaining in 10_ 302 4 (23) the matrix after heat treatment and further risk a depletion of carbon and nitro gen in thematrix, the upper limit of vanadium is 1.3 wt-%. Preferably vanadium is between 1.0and 1.3 wt-%.
    Silicon shall be present in the steel in an amount of between 0.2 - 0.5 wt-%, preferably0.2 - 0.4 wt-%. By keeping the content of silicon low it is possible to obtain an initialprecipitation of meta-stable MgC carbides. These carbides will act as a reservoir forcarbon for subsequent precipitation of the desired M(C,N) particles. Also, precipitationof undesired chromiurn-rich M23C6 particles in the grain boundaries and lattice boundaries is avoided.
    Manganese is present in order to give the steel an adequate hardenability, particularlygiven the relatively low content of chromium and molybdenum in the steel. The content of manganese in the steel is between 0,5 and 2 wt-%, preferably between 1.0 and 2.0 wt- .%.
    Molybdenum shall be present in the steel in an amount of between 1.5 and 3 wt-%,preferably 2.2 - 2.8 wt-%, in order to provide a secondary hardening during temperingand to give a contribution to the hardenability. Part of the molybdenum may besubstituted for tungsten in a manner known per sebut the steel shall preferably not contain any intentionally added amounts of tungsten, i.e. shall not contain tungsten in _ amounts exceeding impurity level, because of certain drawbacks related to the presence of that element.
    Nickel and cobalt are elements that may be included in the steel in amounts up to 3, wt-% and 4 wt-% respectively. Cobalt may increase the hardness at high temperatureswhich may be advantageous for some applications of the steel. Nickel may increase the corrosion resistance, hardenability and toughness of the steel.
    In principle, austenitizing may be carried out at a temperature between the softannealing temperature 820 °C and the maximum austenitizing temperature 1200 °C, butthe austenitizing of the steel article preferably is carried out at a temperature on the ,order of 1050 - 1150 °C, preferably at 1100 °C. In-house tests indicate that higher iaustenitizing temperatures shift the tempering hardness to higher temperatures, i.e. thesecondary hardening peak will be shifted to higher temperatures, which means that thedesired hardness will be reached at a higher initial tempering temperature. Thus the 5 (23) material will obtain an improved tempering resistance and the Work temperature of the tools could be elevated.
    The tempering of the quenched steel article preferably is carried out at least twice at aretention time of 2 hours at a temperature between 550 and 680 °C. ln the mostpreferred embodimentof the steel composition, the tempering is carried out at atemperature between 600 and 650 °C, preferably between 625 and 650 °C.
    Nitro gen contents in the range of 0.05 - 0.10 wt-% may be obtained by incorporatingthe nitrogen by conventional casting methods to form a melt, casting the melt to form aningot, and homogenizing the ingot by heat treatment. Nitrogen additions Will producelarge primary Vanadium-rich M(C,N) precipitates, which in turn will give the materialuneven hardness. However, the large primary carbo-nitrides will not occur if thenitrogen content is lowered and there is a homo genizing heat treatment prior to a subsequent forging.
    In a variant of the steel, higher nitrogen contents than indicated for the preferredembodiment is also conceivable. ln this Variant, nitrogen may amount to up to 0.30 wt~%. To obtain higher nitrogen contents, conventional casting methods are insufficient.Instead, the nitro gen could be incorporated by first manufacturing a steel powder ofessentially the desired composition, except for the nitrogen, then nitriding this powderin solid state by nitrogen containing fluid, e.g. nitrogen gas, thereafter hot pressing thepowder isostatically at a temperature on the orderof 150 °C and a pressure on theorder of 76 MPa to form an ingot. By manufacturing the tool steel by powder metallurgy, the problem of large primary carbide occurrence is avoided.
    The ingot is preferably forged at a temperature on the order of 1270 °C, and then softannealed at a temperature on the order of 820 °C, followed by cooling at a rate of 10 °Cper hour to a temperature of 650 °C and then free cooling in air to make it ready for austenitizing. 6 (23) After the incorporation of the nitrogen the composition of the steel article suitably isessentially (in Wt-%): C _ 0.08~0.40, preferably at least 0.20N 0015-030, preferably 0.05 4 0.10C +'N 0.30 4 0.50, preferably 0.36 - 0.44Cr 1 - 4, preferably max 3 Mo 1.5 -3 V 0.8 - 1.3, preferably 1.0 - 1.3 Mn 05-2 Si 0.2-0.5, preferably 0.2 - 0.4 Ni i <3 Co <4 Fe and incidental impurities remainder Such steel article has a much improved tempering resistance perrnitting a longer articlelife in hot-Work applications. As already indicated above, the nitro gen contentpreferably is on the order of 0.05 Wt-% and the chromium content preferably is 1.5-l .7Wt-%.
    Another object of the present invention is to provide a low-chromium hot-Work toolsteel suitable for making a steel article having increased tempering resistance. The steelarticle of the present invention shall preferably also satisfy some of the following demands : - good tempering resistance,- good high-temperature strength,- good thermal conductivity, - not have an unacceptably large coefficient of heat expansion.
    In accordance With the present invention this object is achieved in that the low- chromium hot-Work tool steel consists essentially of (in Wt-%): 7 (23) C 008-040, preferably at least 0.20N 0.015-0.30, preferably 0.05 ~ 0.10C + N 0.30 - 0.50, preferably 0.36 - 0.44Cr 1 - 4, preferably max 3 Mo 1.5 -3 V 0.8 - 1.3, preferably 1.0 - 1.3 Mn 0.5-2 v Si . 0.2~0.5, preferably 0.2 - 0.4 Ni <3 ' Co <4 Fe and incidental impurities remainder In the tool steel, the nitrogen content preferably is on the order of 0.05 wt-% and thechromium content preferably is l.5-l .7 wt-% .
    BRIEF DESCRIPTION OF THE DRAWINGSIn the following, the invention will be described in more detail with reference to preferred embodiments and the appended drawings.
    Fig. 1 is a diagram showing hardness Vs. tempering temperature of an exemplary prior art low-chromium hot-work tool steel containing no nitrogen.
    F ig. 2 is a diagram showing hardness of prior art steels (contents in wt-%) Cr 15, Mo 1,C 0.6 and Cr 15, Mo l, C029, N 0.35 at different tempering temperatures.
    Fig. 3 is a diagram illustrating the effect of low chromium content on the stability ofM(C,N) in austenite.
    Fig. 4 is a diagram showing the mole fraction of MÖC, M(C,N) and the bcc matrix as a function of temperature. (Balance phase: austenitic matrix.) Fig. 5 is a diagram showing the amount of M(C,N) phase and meta-stable MZC as function of temperature. (Balance phase: ferrite.) Fig. 6 is a diagram showing hardness Vs. tempering temperature curves for trial alloysN0.05, N0.l0 and N0.30 so, s (23) F ig. 7 is a back-scattered SEM image showing small undissolved M(C,N) precipitatesand a globular mixed oXide-sulphide particle in N0.05ï Fig. 8 is a back-scattered SEM image revealing undissolved, primary M(C,N) at former austenite grain boundaries in alloy-N0.10.
    Fig. 9 is a back-scattered SEM image depicting primary particles in soft annealedN0.10. ^ ' Fig. 10 is a back-scattered SEM image revealing an even distribution of undissolved.M(C,N) particles in N0.30.
    Fig. 11 is a back~scattered SEM image revealing some clusters of undissolved M(C,N)found in N0.30.
    MODE(S) FOR CARRYIN G OUT THE INVENTIONMolybdenum and vanadium medium alloyed hot-Work tool steels have good resistanceto thermal fatigue, softening and high-temperature creep. An exemplary nominal chemical composition of such a prior art steel is presented in Table 1 (Wt-%). i Tablelc cr M0 v Mn si Fe0.38 2.6 2.3 0.9 0.75 0.3 92.8 lt has been suggested that the steel of Table 1 owes its high-temperature properties tothe precipitation of nanometre-sized vanadium carbides during tempering. These hardcarbides of MC type (2900 HV) give a secondary hardening, of the material. Figure 1presents a tempering curve (hardness vs. tempering temperature) for the exemplary priorart tool steel. The samples Were austenitized at 1030 °C, and then tempered two times atdifferent temperatures; from 200 °C up to 700 °C for a tempering time of 2 + 2 hours.
    As can be seen, in the interval 500 to 650 “PC there is a pronounced secondary hardeningpeak at 550 °C. Later Work has also shown that there is a significant precipitation of themeta-stable molybdenum-richMgC in the exemplary prior art tool steel during tempering at 625 °C, Which contributes to the secondary hardening effect.
    The ability of a hot-Work tool steel to keep its hardness at an elevated temperature for prolonged time, the tempering resistance, can normally be connected to the initial 9 (23) tempering temperature; if the material is held at a temperature Well below the initialtempering temperature it will not soften. At holding temperatures closer to or above the initial tempering temperature the softening will be more pronounced.
    If the secondary hardening peak could be shifted to higher temperatures, this wouldmean that the desired hardness (eg. 44-46 HRC) could be reached at a higher initialtempering temperature. Thus the material would get an improved tempering resistance, and the Work temperature of the tools could be raised.
    Earlier Work on high-chromium steels suggests that When nitrogen is added to the steel,it is possible to achieve higher hardness during tempering. Samples of Cr 15, Mo 1, C0.6 and Cr 15, Mo 1, C 0.29, N 0.35 were solution treated at 1050 °C followed by Waterquenching and cooling to liquid nitrogen, and then they Were tempered at differenttemperatures for 2 hours. As can be seen in Fig. 2, the peak hardness becamesignificantly higher When adding nitro gen. The initial hardness of the martensite is - lower for the nitrogen containing steel, but during tempering this steel achieves a higher i hardness than the steel containing no nitro gen.
    The explanation for this is that nitrogen makes the chromium more homogeneouslydistributed in the matrix, due to increased solubility of chromium in the austenitic ' iphase. After quenching the martensitic phase inherited the evenly distributed chromiumfromthe austenite, and during tempering a very finely distributed precipitation of chromium nitrides takes place, thus giving a stronger hardening effect in the material.
    Furthermore, the substitution of nitrogen for part of the carbon is used to achieve ahigher hardness of the martensitic steel matrix. The nitrogen addition initially causes alarger amount of retained austenite. However, this austenite can later be transformed tomartensite by cold Work, and it is possible to achieve hardness as high as» 608 HRC in i this manner.
    A low chromium content appears to have a positive effect on the tempering resistance.A comparison of two different hot-Work tools steels With 1-.5 and 5.0 Wt-% chromiumshows that the lower chromium content delays the precipitation of chromium carbides .in the microstructure in favour of the more thermally stable vanadium-rich MC. Thusthe recovery is slowed down in the material and the tempering resistance becomes improved. 10 (23) However, studies made on a typical creep resistant 9-12 wt-% chromium steelcontaining 0.06 wt-% N indicate that low chromium contents stabilize the MX (X being ~C + N) particles dramatically, see Fig. 3. lf the austenitizing were to be performed at 1100 °C, then all of the M(C,N) particles Would be dissolved in the steel containing 10Wt % chromium. If the chromium content Were lowered to 2.5 Wt % (ef. the exemplarylow-chromium tool steelof Fig. 1), then large-amounts of M(C,N) Would still be presentin the austenite. Apparently, the consequence of a low chromium content is that onlysmall amounts of interstitials Will be dissolved into the austenite during austenitizing treatment.
    According to the present invention a low-chromium hot-work tool steel article having increased tempering resistance is made by carrying out the following process steps: a) incorporating nitrogen in a low-chromium hot-Work tool steel composition having achromium content of at most 4 percent by Weight to form a steel composition havinga nitrogen content of 0.015 to 0.30 percent by Weight; b) forming a steel article from the steel composition; c) austenitizing the steel article obtained in step b) at a temperature of at most 1200 °Cfor a time on the order of half an hour followed by quenching; and d) tempering the quenched steel article at least twice at a temperature between 500 and700 °C for a time on the order of 2 hours.
    Considering the conventional understanding in the present technical field, these resultsare surprising since the prevalent teaching is that lowering of the chromium content willresult in a reduced hardenability and difficulties to dissolve primary M(C,N) particles In a creep resistant steel having a high chromium content, i.e. 9-12 % by weight, it ispossible to dissolve vanadium carbo-nitrides already at relatively low temperatures, i.e.1020-1050 °C. However, if the chromium content is low, less than about 4-5 % byWeight, primary vanadium carbo-nitrides Will be formed in the melt, and they are virtually impossible to dissolve afterwards.
    The inventors have found that when the nitro gen content is balanced to about 0.015 to0.30 Wt-% in a low-chromium steel, vanadium carbo-nitrides Will form, Which will bepartly dissolved during the austenitizing step and then precipitated duringlthe temperingstep as particles of nanometer size. The particles are in the order of about l um to about10 um. In some cases, Where the nitrogen content is low, typically at 0.05 wt-% the average size of the particles are less than 1 um. The thermal stability of vanadium .30 p, 11 (23) carbo-nitrides is better than that of vanadium carbides, and consequently the temperingresistance of the low-chromium hot-work tool steel article will be much improved.Further, by tempering at least twice, the tempering curve (showing hardness asa function of tempering temperature) will have a higher, secondary peak.
    In a preferred embodiment of the steel, the nitrogen content preferably is on the order of0.05 percent by Weight. This value gives a better performance than higher values. Anitrogen content on the order of 0.05 percent by weight gives a higher potential for secondary hardening during quenching than higher contents do.
    In the preferred embodiment, the chromium content preferably is l.5-l .7 percent byweight. A low chromium content delays the precipitation of chromium carbides in themicrostructure in favour of the more thermally stable vanadium-rich carbo-nitrides.Thus the recovery is slowed down in the material and the tempering resistance becomes improved.
    In principle, austenitizing may be carried out at a temperature between the softannealing temperature 820 °C and the maximum austenitizing temperature 1200 °C. In apreferred embodiment, i.e. in a composition having a nitro gen content in the order of0.05 percent by weight and a chromium content in the order of 1.5 to 1.7 percent byweight, the austenitizing of the steel article preferably is carried out at a temperature onthe order of 1050 - 1.150 °C, preferably at 1100 °C. In-house tests indicate that higheraustenitizing temperatures shift the tempering hardness to higher temperatures, i.e. thesecondary hardening peak will be shifted to higher temperatures, which means that thedesired hardness will be reached at a higher initial tempering temperature. Thus thematerial will get an improved tempering resistance and the work temperature of the tools will be raised.
    The tempering of the quenched steel article preferably is carried out at least twice at aretention time of 2 hours at a temperature between 550 and 680 °C. ln the mostpreferred embodiment of the steel composition, the tempering is carried out at atemperature between 600 and 650 °C, preferably between 625 and 650 °C.
    Nitrogen contents in the range .of 0.05 - 0.10 percent by weight may be obtained byincorporating the nitrogen by conventional casting methods to form a melt, casting themelt to form an ingot, and homogenizing the ingot by heat treatment. Nitro gen additions will produce large primary vanadium-rich M(C,N) precipitates, which in turn will give 12 (23) the material uneven hardness. However, large primary carbo-nitrides will not occur ifthe nitrogen content is lowered and there is a homogenizing heat treatment prior to a subsequent forging.
    In the preferred embodiment of the invention the nitrogen content preferably is on theorder of 0.05 wt-%. This value gives a better performance than higher values. Anitrogen content on the order of 0.05 wt-%gives a higher potential for secondaryhardeningduring quenching than higher contents do, thus giving the steel a highhardness. However, an amount in the order of 0.10 wt-% has shown to give a shift ofthe secondary hardening peak to somewhat higher tempering temperatures which ispositive. Additionally, the performed tests and modelling calculations indicate that anincreased austenitizing temperature is required in connection with increased nitrogen contents. ln a variant of the steel, higher nitro gen contents than indicated for the preferredembodiment is also conceivable. In this variant, nitrogen may amount to up to 0.30 Wt-%. To obtain higher nitro gen contents, conventional casting methods are insufficient.instead, the nitrogen then is incorporated preferably by first manufacturing a steelpowder of essentially the desired composition, except for the nitro gen, then nitridingthis powder in solid state by nitrogen gas, thereafter hot pressing the powderisostatically at a temperature on the order of .l 150 °C and a pressure on the order of 76MPa to form an ingot. By manufacturing the tool steel by powder metallurgy, the problem of primary carbide occurrence is avoided.
    The ingotis preferably forged at a temperature on the order of 1270 °C, and then softannealed at a temperature on the order of 820 °C, followed by cooling at a rate of 10 °Cper hour to a temperature of 650 °C and then free cooling in air to make it ready for austenitizing. 13 (23) After the incorporation of the nitrogen the composition of the steel article suitably isessentially in Weight-% (Wt-%): CNC+N Co i Fe and incidental impurities Such steel article has a much improved tempering resistance permitting a longer article i 008-040, preferably at least 0.200.015-0.30,preferab1y 0.05 _ 0.100.30 - 0.50, preferably 0.36 - 0.441 - 4, preferably max 3 1.5 -3 0.8 -1.3, preferably 1.0 -1.30.5~2 0.2~0.5, preferably 0.2 - 0.4 <3 <4 I remainder life in hot Work applications. As already indicated above, the nitro gen content preferably is on the order of 0.05 wt-% and the chromium content preferably is 1.5-l .7 Wt-%.
    Another object of the present invention is to provide a low-chromium hot-Work tool steel suitable for making a steel article having increased ternpering resistance.
    In accordance With the present invention this object is achieved in that the low- chromium hot-Work tool steel consists essentially of (in wt-%): CNC+N Co Fe and incidental impurities 0.08-0.40, preferably at least 0.200015-030, preferably 0.05 - 0.100.30 - 0.50, preferably 0.36 - 0.441 - 4, preferably max 3 1.5 -3 0.8 - 1.3,_preferably 1.0 - 1.30.5-2 0.2-0.5, preferably 0.2 ~ 0.4 <3 <4 remainder 715 14 (23) In the tool steel, the nitrogen content preferably is on the order of 0.05 wt-% and thechromium content preferably is 1.5-1.7 wt-%.
    EXAMPLE 1 In Table 2 below the chemical compositions in percent by weight of three differentalloys N0.05; N0.10 and N0.30. are presented. N0.05 designates a material having anitrogen content of 0.05 wt-%, and so on. Note that these are the actual cornpositions ofthe trial ingots.
    The aim was to keep the level of all alloyíng elements except carbon and nitrogenconstant. Compared to the standard low chromium hot-work tool steel of Table 1,chromium was also slightly decreased. There was a small decrease in rnolybdenum'content and an increase in manganese content. For carbon and nitrogen, the aim was tohave a constant sum of around 0.40 wt-% of these elements, and this Was relatively wellachieved.
    Table 2 Material C N Cr Mo V Mn Si Fe N0.05 0.38 0.05 1.70 2.77 1.20 1.09 0.30 92.5N0.10 0.27 0.10 1.53 2.32 1.20 1.85 0.26 92.5NO.30 0.08 0.32 1.51 2.20 1.20 1.88 0.29 '92.5 The tempering stage concerns mainly meta-stable phases, and previous electron Imicroscopy work has shown that they exist in standard low chromium hot-work toolsteel at tempering temperature intervals, i.e. 400 to 700 °C. These carbide phases aremainly vanadium-rich MC (FCC) and molybdenum-rich MgC (HCP). Some amount ofchromium-rich M7C3 has also .been found in the standard low chromium hot-work toolsteel.
    The following calculations were made in order to decide whether or not these nitrogencontainingalloys were possible to harden, i.e. if enough alloying elements could bedissolved into the austenitic matrix at the austenitizing temperature, so that martensitewould form during quenching. The interesting temperature interval thus was betweenthe soft annealing temperature, 820 °C and the set practically usable maximumaustenitizing temperature, 1200 °C. i 15 (23) The results of these equilibrium calculations are presented in Fig. 4. Here the molefraction of MÖC, M(C,N) and the bcc matrix is shown as a function of temperature. Thebalance phase is austenite. The full curves represent N0.05; the dashed *curves representN0.10 and the dotted curves represent N03 0. Note the high content of M(C,N) in theN0.30 alloy even up to 1200 °C. As expected, the bcc phase is unstable above 850°C. ltis interesting to see that the slope of the equilibrium curve, representing the amount ofM(C,N), decreases as the nitrogen content increases. This means that it is more difficultto dissolve M(C,N) in N0.30 compared to N0.05 _ Thus, it is expected that the amount ofcarbon, nitrogen and vanadium would be lower in the .N0.30 matrix after austenitizing at1100 °C than in the N0.05 matrix.
    Since the molybdenum-rich MÖC phase only dissolves carbon and no nitrogen, it suffersfrom the lower carbon content in N0.l0 and N0.30, thus the amount of MGC decreaseswith decreasing carbon content. It should also be noted that all M6C is dissolved at the austenitizing temperatures used.
    The calculations performed in the tempering temperature region were only done in orderto estimate the potential for secondary precipitation in N0.05, N0.10 and N0.3 0. Theequilibria found can at best show what phases would be present in the material after asufficiently long time. Previous work has shown that in practice there is some auto-tempering in the standard low chromium hot-work tool steel. This means that M3C (cementíte) will precipitate after the austenitizing process.
    The results from the calculations in the tempering temperature region are presented inFig. 5. The full curves represent N0.05; the dashed curves represent N0.10 and thedotted curves represent N0.30. Secondary hardening normally takes place between 500and 650 °C, and in this temperature interval there is no big difference between N0.05and NO. 1 0 regarding the amount of M(C,N). N0.30 on the other hand has a high andalmost constant amount of M(C,N), probably due to the high vanadium and nitrogen contents.
    The higher carbon content in N0.05 produces more MgC phase in equilibrium with thematrix compared to N0. 1 0. In N0.30 there is much less MgC.
    Based upon the previous calculations it should be possible to estimate the potential forsecondary precipitation in these alloys after austenitizing at a certain temperature. Thispotential depends on the difference in amount of M(C,N) phase and MgC phase between 10. 16 (23) the meta-stable equilibrium at tempering temperature and the equilibrium ataustenitizing temperature. In Table 3, these differences are presented as the secondaryprecipitatíon potentialvfor the three different alloys. The values are given in mole percent.
    Table 3Phases, mole percentAlloy N0.05. M(C,N) MgC TotalTempering, 625 °C 2.1 2.8Austenitizing, 1100 °C 1.1 0.0Precipitation Potential 1.0 2.8 - 3.8Alloy N0.10Tempering, 625 °C 1.8 2.3Austenitizing, 1100 °C 1.2 0.0Precipitation Potential 0.6 2.3 2.9Alloy N0.30Tempering, 625 °C 2.6 1.2Austenitizing, 1100 °C 2.4 0.0Precipitation Potential 0.2 1.2 1.4 The results presented in Table 3 indicate that N0.05 Would have the best hardeningresponse due to the low amount of M(C,N) phase present at 1100 °C, i.e. a lot ofalloying elements can be dissolved into the austenitic matrix. It also indicates that N0.05 has the best potential for a good secondary hardening during tempering at 625 °C.
    EXAMPLE 2The two alloys N0.05 and N0.lO were conventionally cast as small ingots of 50 kg.NO. 10' was the first trial and there was no homo genizing treatment done on this ingot before the forging process. The second trial, N0.05, a homogenizing treatment at i l300°C for 15 hours was applied before forging. The third alloy, N0.30 had a too high nitrogen content to be manufactured by conventional casting. Therefore this alloy was eproduced using powder metallurgy. First the steel powder Was manufactured and thenthis powder was nitrided in solid state by pressurized Nz-gas. The powder was then hotisostatically pressed (HIP) at 1150 °C With the pressure of 76 MPa. 17 (23) All three ingots Were forged at 1270 °C and then samples Were cut out With thedimensions: 15X15x8 mm. The samples Where heat treated by first soft annealing at820 °C; the sequence for Cooling after annealing is 10 °C per hour to 650 °C and thenfree Cooling in air. After soft annealing, N0.05 Was austenitized at 1100 °C for 30minutes. ln order to cornpensate for the poorer potential for precipitation, N0.10 Wasaustenitized at 1150 °C for 30 minutes, and N0.30 Was austenitized at 1200 °C for 30minutes. Nine samples from each of the three alloys Were tempered at followingtemperatures: 450, 525, 550, 575, 600, 625, 650, 675 and 700 °C. The soaking time Wastwo hours and it Was a double tempering, i.e. the total tempering time Was four hours.After heat treatment, the hardness of the samples Was measured. Scamiing electronmicroscopy (SEM) Was performed in order to further investigate the morphology,distribution and size of the undissolved particles in the samples. The SEM instrumentused Was a FEl Quanta 600 F.
    Hardness measurements The results from the hardness measurements are presented in Fig. 6. As can be seen, allthree alloys have a secondary hardening peak in the temperature interval 500 to 650 °C.All tempering Was done for 2 + 2 hours. N0.05 has the highest hardness in the as-quenched condition (53HRC), While N0.10 and N0.30 had someWhat loWer hardness.HoWever, all threealloys are regarded as hardenable. The hardness curve of N0.05 isvery similar to that of the standard low-chromium hot-Work tool steel With a maximumof around 54 HRC as shown in Fig. 1.
    The secondary hardening peak of NO. 1 0 seems to be someWhat shifted to a highertemperature With peak hardness at 600 °C. The peak hardness for both N0.05 and N0.30Was at 550 °C.
    Scanning electron mícroscopy The undissolved M(C,N) particles in the conventionally east N0.05, the alloy With theloWest nitrogen content, have an average size smaller than 1 um. This is comparableWith ordinary undissolved carbides in steel. Another phase that is easily found in N0.05is the mixture of aluminium-oxide and manganese-sulphide, see Fig. 7, Which is a SEMimage (back-scattered) showing small undissolved M(C,N) precipitates 2 and a globularmixed oxide-sulphide particle 1 in N0.05. The sample Was austenitized at 1100 °C for30 min and tempered at 625 °C for 2 + 2 hours. 20' 1893) The reason for the many non-metallic inclusions in N0.05 (and NO.10) is that all trial ingots were manufactured and cast in open atmosphere.
    The most common size of the M(C,N) particles in NO.10 is between 5 and 10 umEquivalent Circle Diameter (ECD) after austenitizing at 1150 “C for 30 minutes andtempering at 625 °C for 2 + 2 hours. Larger, primary carbides 3 (precipitated in themelt) are frequently found in former austenite grain boundaries, see Fig. 8, which is aback-scattered SEM image revealing undissolved, primary M(C,N) at former austenitegrain boundaries in alloy NO.10. The sample was austenitized at 1150 °C for 30 min andtempered at 625 °C for 2 + 2 hours.
    F ig. 9 is a detail SEM mícrograph of primary M(C,N) particles 4 in NO.10. They werediscovered automatically in SEM using the lNCA Feature software from OxfordInstruments. Their sharp edges indicated that they had precipitated from the melt. Thewhite areas in the image are molybdenum-rich MóC particles 5. Note that in this casethe sample was soft annealed NO.10.
    In the powder metallurgically manufactured N0.3 0, the undissolved M(C,N) particles 6had a size distribution (ECD) between 1 to 5 um With the most common size 2 um, thus the particles were small even though the nitro gen content was high. The particles were Vhomogeneously distributed in the microstructure, see Fig. 10. However, as shown in iFig. 11, some clusters 7 of M(C,N) were found.
    The chemical composition of the undissolved particles of the M(C,N) phase in all threealloys was measured by EDS and the result is presented in Table 4, Which shows thechemical composition of the M(C,N) particles in alloys N0.05, NO.10 and N0.30. Thevalues are given in atomic percent. Note that, even though the accuracy in EDSregarding light elements such as carbon and nitro gen is not so high, one can see that thebalance of carbon and nitrogen in the M(C,N) phase is what can be expected based uponthe nominal compositions. The i values given in the table are the ones given in theINCA program (Oxford instruments). Some of the iron recorded probably comes from the surrounding matrix, especially for the alloy N0.05.
    Table 4Alloy _ C N V Fe Cr MoN0.05 39.41 ~ 15.4: 42.13; 1.18: -0.34 0.14 0.29 0.08 19 (23) N0.10 26.4 i 27.6 32.4 i0.33 11.0 i0.17 1.3 i0.l0 1.0 i0.28 i0.42 i 0.12N0.30 12.1 i 41.9 21.4 i0.2 21.5 i0.5 2.1 i0.09 0.38 i4 0.24 i0.32 0.1 INDUSTRIAL APPLICABILITY The process and the loW-chromium hot-Work tool steel of the present invention are applícable Where it is desired to get hot-Work steel tools, Which can be utílized at increased temperatures for an extended period of time.
    权利要求:
    Claims (14)
    [1] 1. : . A process for making a low-chromium hot-work tool steel article having increased temperíng resistance, comprising a) incorporating nitrogen in a low-chromium hot-work tool steel compositionhaving a chromium content of at most 4. wt-% to form a steel compositionhaving a nitrogen content of 0.015 to 0.30 Wt-%; b) forming a steel article from the steel composition; c) austenitizing the steel article obtained in step b) at a temperature of at most1200 °C for a time on the order of half an hour followed by quenching; and d) temperíng the quenched steel article for a time of 2 hours at least twice at a temperature between 500 and 700 °C. . A process as claimed in claim 1, wherein the nitrogen content *is 0.05 to 0.10, preferably 0.05 wt-%. . A process as claimed in claim l or 2, wherein the steel composition contains carbon in an amount of at least 0,20 wt-%. . A process as claimed in any one of claims l~3, wherein the chromium content is at least 1 wt-%, preferably 1-3 wt-%, and even more preferred 1.5-1.7 wt~%. . A process as claimed in any one of claims 1-4, further comprising carrying out the austenitizing of the steel article at a temperature of 1050 ~ 1150 °C, preferably at1100~ 1150 °C. . process as claimed in any one of claims l~5, further comprising carrying out the temperíng of the quenched steel article at a temperature of 550 - 680 °C, preferably600 - 650 °C and even more preferred at 625 - 650 °C.' . 'A process as claimed in any one of claims 1-6, further comprising incorporating the nitrogen by conventional casting methods to form a melt, casting theimelt to form an ingot, and homo genizing the ingot by heat treatment. . A process as claimed in any one of claims l-6, further comprising incorporating the nitrogen by first manufacturing a steel powder of essentially the desired 10. 11. 21 (23) cornposition, except for the nitrogen” then nitriding this powder in solid state by nitrogen gas, and thereafter hot pressing the powder to form an ingot. A process as claimed in clairn 7 or 8, further cornprising the steps of hornogenizing, forging and soft annealing the ingot before austenitizing. . A process as claimed in claim l, Whereín the compositíon of the steel article after the incorporation of the nitrogen essentially is (in Wt-%): c 008-040-N 0015-030c +N 030-050Cr 1-4 Mo 15-3 V 08-13Mn ' i 0.5-2 si 0.2-0.5 Ni i <3 Co i <4 Fe and incidental impurítíes remainder A process as claimed in claim l, Wherein the cornposítion of the steel article after the incorporation of the nitrogen essentially is (in Wt-%): C at least 020-040N fl 0,05 - 0.10 c + N 036 - 044 cr 15-17 Mo 2-3 V 1.0 - 1.3 Mn l-Z si ' 02-04 Ni ' <3 Co <4 Fe and incidental impurities remainder 22 (23) 12. A loW-chromium hot-Work tool steel consisting essentially of (in Wt-%): C 0.08-0.40N 0.0l5-0.30C +N ' 030-050Cr _ 1-4 Mo l.5-3 V 0.8-l.3Mn 0.5-2 Si 0.2-0.5 Ni <3 Co <4 Fe and incidental irnpurities remainder 13. A low-chromium hot-Work tool steel as claimed in claim 12, Wherein the nitrogencontent is 0.05 - 0.10 Wt-% and the carbon content is at least 0.20 Wt-%. 14. A loW-chrornium hot-Work tool steel as claimed in claim 12 or 13, Wherein tlie chromium content is 1.5-1.7 Wt~%.
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    同族专利:
    公开号 | 公开日
    EP2681340B1|2015-04-15|
    DK2681340T3|2015-06-29|
    US20140056749A1|2014-02-27|
    TWI535863B|2016-06-01|
    KR20140015445A|2014-02-06|
    CN103703150A|2014-04-02|
    US20160115573A1|2016-04-28|
    PL2681340T3|2015-10-30|
    RU2013142584A|2015-04-10|
    EP2681340A1|2014-01-08|
    CN103703150B|2015-12-23|
    ES2540905T3|2015-07-14|
    CA2828962C|2018-11-06|
    SI2681340T1|2015-10-30|
    SE536596C2|2014-03-18|
    TW201303043A|2013-01-16|
    KR102012950B1|2019-08-21|
    KR20150047636A|2015-05-04|
    JP2014512456A|2014-05-22|
    KR20170105138A|2017-09-18|
    WO2012119925A1|2012-09-13|
    BR112013022606A2|2016-12-06|
    PT2681340E|2015-08-25|
    JP5837945B2|2015-12-24|
    CA2828962A1|2012-09-13|
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    法律状态:
    优先权:
    申请号 | 申请日 | 专利标题
    SE1150200A|SE536596C2|2011-03-04|2011-03-04|Hot work steel and a process for producing a hot work steel|SE1150200A| SE536596C2|2011-03-04|2011-03-04|Hot work steel and a process for producing a hot work steel|
    KR1020157009651A| KR102012950B1|2011-03-04|2012-03-01|Hot-work tool steel and a process for making a hot-work tool steel|
    CA2828962A| CA2828962C|2011-03-04|2012-03-01|Hot-work tool steel and a process for making a hot-work tool steel|
    KR1020137026324A| KR20140015445A|2011-03-04|2012-03-01|Hot-work tool steel and a process for making a hot-work tool steel|
    KR1020177025271A| KR20170105138A|2011-03-04|2012-03-01|Hot-work tool steel and a process for making a hot-work tool steel|
    US14/002,967| US20140056749A1|2011-03-04|2012-03-01|Hot-work tool steel and a process for making a hot-work tool steel|
    BR112013022606A| BR112013022606A2|2011-03-04|2012-03-01|hot tool steel and process for manufacturing it|
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    PL12707998T| PL2681340T3|2011-03-04|2012-03-01|Hot-work tool steel and a process for making a hot-work tool steel|
    ES12707998.6T| ES2540905T3|2011-03-04|2012-03-01|Hot work tool steel and a process for manufacturing a hot work tool steel|
    RU2013142584/02A| RU2575527C2|2011-03-04|2012-03-01|Tool steel for work at high temperatures and method of manufacturing of tool steel for work at high temperatures|
    PT127079986T| PT2681340E|2011-03-04|2012-03-01|Hot-work tool steel and a process for making a hot-work tool steel|
    CN201280021117.7A| CN103703150B|2011-03-04|2012-03-01|The method of hot working tool steel and manufacture hot working tool steel|
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    PCT/EP2012/053563| WO2012119925A1|2011-03-04|2012-03-01|Hot-work tool steel and a process for making a hot-work tool steel|
    SI201230252T| SI2681340T1|2011-03-04|2012-03-01|Hot-work tool steel and a process for making a hot-work tool steel|
    EP12707998.6A| EP2681340B1|2011-03-04|2012-03-01|Hot-work tool steel and a process for making a hot-work tool steel|
    TW101106919A| TWI535863B|2011-03-04|2012-03-02|Hot-work tool steel and a process for making a hot-work tool steel|
    US14/989,469| US20160115573A1|2011-03-04|2016-01-06|Hot-work tool steel and a process for making a hot-work tool steel|
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