TECHNICAL FIELD OF THE INVENTION

The present invention relates to a method of producing a high speed steel alloy containing, in percent by weight (wt.%):

C 1 .00-1 .10, N 0.003-0.025, Cr 3.80-4.40, Mo 3.90-4.50, W 0- 1 .0, Co 0-0.99, V 1 .8-2.2, Nb 0-0.30, Mn 0.20-0.40, Si 1 .40- 1 .55, Ni 0-0.50, and Cu 0-0.50, the balance being Fe and wherein the content of normally occurring impurities is less than 1 .0 wt.%, , and wherein said method comprises the following steps: providing a melt of said alloy, casting said melt followed by solidification thereof, hot forming the alloy into a predetermined body, soft annealing the solidified alloy, and hardening said body of the alloy at a hardening temperature T in the range of 1 100°C-1200°C for a predetermined time t which is in the range of t1 -t2, wherein t1 is a time which is sufficient for carbide-forming elements of the alloy to be dissolved in an austenitic phase presented by the alloy, and, after said hardening step, it comprises the further step of tempering said cast alloy member.

BACKGROUND AND PRIOR ART

High speed steels (HSS) are steels being used especially in tools for different types of machining, such as drilling, milling and sawing, but other applications are also conceivable, such as for example in tools for hot-working, such as dies for extrusion of aluminium profiles and rollers for hot-rolling, in advanced machine elements and press rollers, i.e. tools for stamping of patterns or profiles in metals etc. Another application of such steels is in cold-working tools, for example thread rolling. Low cost non-coated cutting tools, mainly drills, are products that preferably may be made of a high speed steel. A high toughness will be required for such material.

Other important properties of such a high speed steel are a high hardness and wear or abrasive resistance as well as an easiness to be machined after soft annealing for manufacturing tools out of tool blanks of such a steel. A further property often required is a good grindability.

The high speed steel most frequently occurring on the market today is the so-called M2, which may have compositions differing slightly, but mainly has the following composition in weight-%: C 0.90, Cr 4.2, Mo 5.0, W 6.4 and V 2.0. Cr is used for obtaining an appropriate hardening capacity of the steel, whereas the alloying elements Mo, W and V are used together with the carbon forming metal carbides necessary for obtaining the hardness and wear and abrasive resistance aimed at.

High speed steel having lower contents of expensive alloying element than what is used in M2, but still having mechanical properties comparable to those of M2 are being developed. The high speed steel produced by means of the method suggested in this application is such a steel. Different alloying elements, such as Mo, W, V and Nb are used in known low alloyed high speed steels for forming metal carbides in the steel for obtaining a desired high toughness and abrasive resistance as well as a high strength and hardness of the steel.

During hardening of a cast alloy having a composition of a high speed steel with low content of alloying elements as defined hereinabove and hereinafter and at a hardening temperature as defined hereinabove and hereinafter, the holding time at the hardening temperature is long enough to guarantee that the carbides formed in the alloy during cooling after casting and during a subsequent soft annealing are dissolved in the austenitic phase of the alloy. The requested hardness of the alloy will be achieved as the alloy, after being held at the hardening temperature, is cooled rapidly enough to form a martensitic structure and after a tempering at a suitable tempering temperature.

However, high speed steel having the composition defined hereinabove and hereinafter have a tendency to have a suppressed impact toughness when being heat treated in accordance with teachings of prior art. Such prior art may be represented by W02009/082328 A1 , submitted by the present applicant and disclosing steel having compositions similar to the steel of the present application.

It is therefore an object of the present invention to present a method for production of a high speed steel alloy as defined hereinabove or hereinafter that results in a high speed steel having an improved impact toughness compared to if it had been heat treated in accordance with teachings of prior art.

SUMMARY OF THE INVENTION

The object of the present invention is achieved by means of the initially defined method, which is characterized in that t2 is below a time at which a medium austenite grain size of the alloy, as measured with the Snyder-Graff method, is such that the Snyder- Graff intercept grain size number (SG) is at least 13. The measurement is in accordance with ASTM E 1 12’’Standard test methods for determining average grain size”.

The present inventors have realized that, in particular when the high speed steel has relatively low contents of W and, in particular, Mo, an accelerated growth rate of the austenitic grains in the alloy is obtained. Accordingly, Mo and W seem to have a growth rate-suppressing effect on the austenitic grains. A large austenitic grain size upon hardening has been observed to result in a reduced impact toughness of the hardened alloy. By controlling the time at hardening temperature, and not letting it be too long, an improved impact toughness is thus achieved. The hardening is preferably ended by cooling said body from said temperature T such that at least a partly martensitic structure is obtained.

According to one embodiment, t2 is below a time at which a medium austenite grain size of the alloy, as measured with the Snyder-Graff method, is such that the Snyder-Graff intercept grain size number (SG) is at least 14.

According to one embodiment, t2 < minutes.

minutes or

t2<25 minutes, whichever is the lowest.

According to one embodiment, 1 100°C<T<1 180°C. According to another embodiment 1 150°C<T<1200°C, and according to yet another embodiment, 1 150°C<T<1 180°C.

According to one embodiment, (MO < 4.5,

Mo and W being the contents of molybdenum and tungsten expressed in weight percent.

According to an alternative embodiment, (MO + ) < 4.0.

According to one embodiment, the content of Co is less than 0.50 wt.%.

According to one embodiment, the content of W is less than 0.50 wt.%. According to one embodiment, the content of Mo is 3.90-4.10 wt.%. According to one embodiment, the method according to the invention is further characterised in that, after said hardening step, it comprises the further step of tempering said cast alloy member.

According to one embodiment, said tempering is carried out at a temperature of 500°C-600°C, for 0.5-2 hours 2-4 times. Tempering is controlled such that a fully martensitic structure is obtained in the body formed by the alloy.

According to one embodiment, the hardening temperature T is in the range of 1 140°C - 1 160°C, and the tempering is carried out 3 times at a temperature in the range of 535°C - 545°C. According to one embodiment, the duration length of each tempering sequence (time at tempering temperature) is 45-75 minutes, followed by cooling down to less than 25°C. According to another embodiment, the duration length of each tempering sequence (time at tempering temperature) is 55-65 minutes. The hardening time may be longer than 30 minutes.

Further features and advantages will be presented in the following detailed description of embodiments.

BRIEF DESCRIPTION OF THE DRAWINGS

Fig. 1 is a diagram showing Snyder-Graff intercept grain size number (SG) versus time t at hardening temperature T for three different values on T for a high speed steel with a composition according to the method of the present invention,

Fig. 2 is a diagram showing impact toughness, as measured according to standard SEP1314, versus austenite grain size expressed as Snyder-Graff intercept grain size number (SG) for a high speed steel with a composition according to the method of the present invention and hardened at 1 180°C, and Fig. 3 is a diagram showing hardness and impact toughness for samples hardened at 1 150°C, versus number of tempering steps.

DETAILED DESCRIPTION OF EMBODIMENTS

For reasons to be disclosed hereinafter, the high speed steel comprises the following alloying element in the amounts that are specified here and in the appended claims: Carbon (C) should exist at a content of 1 .00-1 .10 weight-% for resulting in about 3 atom-% in the austenite at a typical hardening temperature, such as 1 180°C, which is favourable for giving the material a hardness in the hardened and tempered condition that is suitable for its purposes. Carbon contributes to the formation of an adequate amount of primarily precipitated MC-carbides, which may be of the type M6C and MC as disclosed further below. These carbides are important for obtaining a desired hardness and wear and abrasive resistance. Nitrogen (N) may partially replace carbon and has the same function as the carbon while forming M-nitrides and carbon nitrides. It should not be present in a content above 0.025 weight- %, since this may result in production of large vanadium nitrides already in the melt.

Chromium (Cr) should exist in the steel at a content of at least 3.8 weight-% in order to, when dissolved in the matrix of the steel, contribute to the steel achieving adequate hardness and toughness after hardening and tempering. Chromium can also contribute to the resistance to wear of the steel by being included in primarily precipitated hard phase particles, mainly M6C- carbides. Chromium shall not be present in a content above 4.40 weight-%, since that would only result in extra alloying element costs without adding anything to the hardness of the steel.

Molybdenum (Mo) is used for forming M6C-carbide contributing to hardness and the resistance of wear of the steel. The content should be at least 3.9 weight-% for obtaining sufficient contribution to wear resistance and hardness of the steel, but it is expensive and should not be above 4.50 weight.

Tungsten (W) form M6C-carbides contributing to the wear resistance of the steel. However, tungsten shall not be present in a content above 1 .0 weight-%, preferably not above 0.5 weight- %, since the relationship of the content of Mo/W shall be high, such as at least above 3 for enabling Si to contribute to the hardness of the steel and partially replacing Mo. Vanadium (V) is used for forming MC-carbides contributing to resistance to wear and hardness of the material. MC-carbides are harder than M6C-carbide, so that it is better to have MC-carbides of a certain size than M6C-carbides of that size. However, the content of V may not be above 2.2 weight-%, since that would result in formation of large carbides reducing the easiness to machine the material after soft annealing, and reducing the grindability and the toughness of the material. Too high amounts of V also involve a risk of formation of MC-carbides already in the cast making the manufacturing process more difficult.

Niobium (Nb) may partially replace vanadium to some extent and has substantially the same behaviour as vanadium with respect to formation of MC-carbides and the properties thereof. However, V may be preferred, since it results in easier handling of scrap of the alloy than does Nb. The content of Nb should not be above 0.3 weight-%.

Silicon (Si) should be present at a content of at least 1 .40 weight- % for contributing to the hardness and abrasive resistance of the steel. However, higher contents are desired for the ability of Si to replace Mo, so that the content of the more costly Mo may be lowered and by that costs may be saved. The content of Si should not exceed 1.55 weight-% since the hardness after soft annealing will then be too high for making it comfortable to machine the material. Another effect of Si is that it destabilizes M2C, which may be present in the cast, in favour of M6C-carbides for transforming the M2C into M6C and MC when the cast is heat treated. Si is a ferrite stabiliser. Manganese (Mn), is an austenite stabiliser. The steel alloy comprises 0.20-0.40 weight-% Mn. If the content of Mn is too low, Fe will form FeS which ends up in the grain boundaries, thereby making the material brittle. Mn in combination with Si also improves de-oxidation during the production of the steel and result in a steel with less oxide inclusions.

Nickel (Ni) is a strong austenite stabiliser. It may be present in the steel but in order to avoid remaining austenite after hardening and tempering, the amount of Ni should not be above 0.5 weight- %, preferably not above 0.3 weight-%. If copper is present in the steel, Ni + Cu should not be above 0.7 weight-%, preferably not above 0.5 weight-%.

Copper (Cu) may be present in the steel in amounts up to 0.5 weight-%, preferably not more than 0.3 weight-%.

EXPERIMENTAL RESULTS

An alloy having the following final composition was molten, cast and permitted to solidify into an ingot.

Impurity levels of aluminium (Al), titanium (Ti), lead (Pb) and tin (Sn) were also present in the steel. The total content of such impurities was below 0.1 weight %.

Test samples having the shape of rods with a diameter of 6-13 mm were formed from the ingot through a process that included forging and rolling of the ingot to rods having a diameter of 6.5- 13.5 mm, final drawing of the rod down to final dimension and, finally, cutting thereof.

Before final drawing the rods or threads was then soft annealed at 880°C during a time period of 2 hours, followed by controlled cooling to 700 °C with a cooling rate of approximately 10 °C/minute, and thereafter free cooling from 700 °C to room temperature.

Samples were then subjected to hardening step at 1 100°C, 1 150°C and 1 180°C. For each hardening temperature, samples were held at the hardening temperature for different times, in this case 2 minutes, 20 minutes and 60 minutes.

The samples were cooled from the respective hardening temperature with an approximate cooling rate of 7 °C/second. A partly austenitic and partly martensitic structure was obtained as a result thereof.

The hardened samples were then subjected to tempering, which consisted of heating the samples to a tempering temperature of 550 °C, holding the sample at said temperature for 1 hour, and repeating this procedure one time (two times in total).

For each sample, the austenite grain size was measured by means of the Snyder-Graff method, and the austenite grain size was expressed by the Snyder-Graff intercept grain size number (SG). A higher SG number indicates a smaller grain size. The measurements were performed in accordance with ASTM E 1 12 ’’Standard test methods for determining average grain size”. Results are shown in fig. 1 . As can be seen, for each hardening temperature, the austenite grain size increased (as shown by a lower SG number) almost linearly with increasing hold time at the respective temperature. Then, impact toughness was measured in accordance with SEP1314 for samples hardened at 1 180°C, and subsequently tempered and showing different austenite grain size as a result of different hold times at the hardening temperature. The results are shown in fig. 2. As can be seen, a remarkable improvement of impact toughness was recognised for samples having a Snyder- Graff intercept grain size number of approximately 14 compared to those having a Snyder-Graff intercept grain size number of approximately 12. Samples having a lower Snyder-Graff intercept grain size number, i.e. a larger austenite grain size, showed a remarkably lower impact toughness.

Samples from alloys having different contents of Mo and W than the alloy used for the test samples above have been used in order to see if there is a correlation between the content of the these strong formers of M6C and MC carbides and the austenite grain growth as a function of time and temperature. Tests corresponding to the tests described herein above have thus been carried out and show that there is such a correlation. The correlation between maximum time at a given hardening temperature and the content of Mo and W can be written as follows: minutes

Fig. 3 shows the test results for samples that have been hardened at 1 150°C for 60 minutes, quenched down to <50°C and then subjected to tempering, which consisted of heating the samples to 520°C, 540°C and 560°C respectively, holding the sample at said temperature for 60 minutes, cooling down to <25°C, and repeating this procedure different number of times. It can be seen from fig. 3 that an advantageous combination of hardness and impact toughness was achieved for samples having a tempering temperature of 540°C and subjected to three tempering sequences. The grain size was 13.8 (Snyder-Graff intercept grain size number).