US3444011A - Low-temperature tough steel - Google Patents

Description

Filed Nov. 18, 1964 Ml CROGRAP H MICRO GRAPH MlCROGRAPH Sheet 2 of 3 ATTORNEYS y 1969 SHlNlCHlI-LNAGASHIMA ET AL 3,444,011

LOW-TEMPERATURE TOUGH STEEL Filed Nov. 18, 1964 Sheet 1 or :5

MICROGRAPH MICROGRAPH FIG. 6-8

MlCROGRAPH ATTORNEYS nited States 3,444,011 LOW-TEMPERATURE TOUGH STEEL Shinichi Nagashima, Kitakyushu, Takayuki oka, :I0 kyo,

Hiroshi Mimura, Kawasaki, and Toshlyukl Fupshrma, Kitakyushu, Japan, assignors to Yawata Iron & Steel Co., Ltd., Tokyo, Japan, a corporation of Japan Filed Nov. 18, 1964, Ser. No. 412,111 Claims priority, application Japan, Nov. 18, 1963, 38/ 61,988 Int. Cl. C220 39/44, 41/02; C21d N18 US. Cl. 148-31 6 Claims ABSTRACT OF THE DISCLOSURE This invention relates to alloy steels having a sufiicient toughness and strength even at low temperatures.

The so-called 9%-Ni steel is already widely known as an economical commercial steel to take the place of the 18-8 stainless steel. As properly heat-treated, this steel shows a toughness of about 8 to 11 kg.-m./cm. in the V-notch Charpy impact value at -196 C., the boiling point of liquid nitrogen, and has such considerable strengths as a tensile strength of 75 to 85 kg./mm. and a yield strength of 60 to 65 kg./mm. at room temperature. Although 9%-Ni steel has such toughness and strength as are mentioned above, as such large amount as about 9% of Ni which is a costly alloying element for steels is used in this steel, the price of the steel is high. Specifically in a country poor in Ni resources, even if a steel material high in the Ni content would be produced, it will be used only in an economically limited range. Therefore, in order to reduce the price and structural weight, the development of more economical tough steels is required today.

An object of the present invention is to provide an economical steel which contains such alloying elements as 0.01 to 0.15% C, 0.05 to 0.4% Si, 4.5 to 7.5% Ni, 0.5 to 3.5% Mn, N and Al, in which costly Ni is reduced to be in small amount and which, as properly heat-treated, will have a toughness and strength as high as or higher than those of the above mentioned 9%-Ni steel at loW temperatures in order to meet the above mentioned requirement.

The accompanying drawings show various embodiments of the present invention.

FIGURE 1 shows the influence of the austenitizing temperature on the impact value at 196 C.

FIGURE 2 shows the effect of Mn on the impact absorption energy of a simple composition steel.

FIGURE 3 shows a micrograph by direct observation with an electron microscope of a magnification of 50,000 after the specimen was austenitized and air-cooled.

FIGURE 4 shows the micrograph by direct observation with an electron microscope of a magnification of 50,000 where the same specimen as that of FIG. 3 was tempered.

FIGURE 5-A shows a micrograph by direct observation with an electron microscope of a magnification of 50,000 of a tempered steel austenitized in one step.

FIGURE 5-B is a micrograph showing the dispersion atent 0 3,444,011 Patented May 13, 1969 of AlN by the extraction replica method with an electron microscope of a magnification of 5,000.

FIGURE 6-A shows a micrograph by direct observation with an electron microscope of a magnification of 50,000 of a tempered steel austenitized in two steps.

FIGURE 6B is a micrograph showing the dispersion of AlN by the extraction replica method with an electron microscope of a magnification of 5,000.

The present invention shall now be detailed 'with reference to the accompanying drawings.

The steel of the present invention is characterized by having such composition and structure as 0.01 to 0.15 C, 0.05 to 0.40% Si, 4.50 to 7.50% Ni, 0.50 to 3.50% Mn, 0.001 to 0.050% N and Al in an amount sufiicient to fix N or smaller than that (less than 0.05% as acid soluble Al) in which Al may be replaced with one or more of Zr, Ti, Be, Nb, V, Hf, Ta and B, i.e., nitride forming elements, in the chemically equivalent amount of the nitride; M0 is added as required, and its amount of addition may be 0.05 to 0.50%, and the rest other than the above added alloying elements is Fe and impurities. This steel is then austenitized followed by quenching (or air-cooling) and then is tempered in order to produce a ferrite structure containing fine dispersed austenite precipitates, part or all of which may be transformed into martensite.

Thus, the fundamental metallurgical structure of the steel of this invention consists of ferrite and precipitated austenite (part or all of which has been converted into martensite). More particularly, it has a ferrite structure with austenite precipitates on old (former) martensite grain boundaries, old (former) austenite crystal grain boundaries or ferrite subgrain boundaries, produced by an appropriate heat-treating method as hereinafter described. Further, in case fine dispersing precipitants composed mostly of nitrides are added for the purpose of improving the toughness by grain refining and of increasing the strength by a dispersion hardening mechanism, the abovementioned fundamental structure will become a structure in which appropriate precipitants are additionally dispersed.

The production of the steel of the present invention shall be described in the following. In the melting and producing step for producing the steel of the present invention, smelting can be easily carried out in such steel making furnace as a converter, open hearth-furnace, electric furnace or high frequency furnace. There is no problem in particular in this regard. The molten steel containing the above mentioned alloying elements is smelted in the above mentioned furnace, is cast and hot-rolled. In the casting and hot-rolling steps, too, no specific rectriction is required. But, depending on the object, such steps may be carried out by limiting the atmosphere of the heat-treatment. When the hot-rolled steel is heattreated properly, the above described fundamental structure will be able to be obtained. However, this heat-treatment must be regulated differently as explained below depending on the contents of N and Al. That is to say, in a steel which is normally smelted in a steel making furnace and N is not positively added, the amount of Al (which may be replaced by one or more of such nitride forming elements as Zr, Ti, Be, Nb, V, Hf, Ta and B here and also hereinafter) required to fix N is added and the austenitizing followed by quenching (or air-cooling) and tempering treatments are carried out. However, the heating temperature in such case is preferably about 800 C. If it becomes higher than the grain coarsening temperature, the toughness at low temperatures will tend to reduce. Further, the tempering temperature is preferably in the range of 525 to 650 C. in which the fine dispersed austenite precipitates in the properly tempered ferrite 3 matrix. It is evident from Table 1 and FIGURE 1 that the above mentioned heating temperature is suitable.

TABLE 1 Chemical composition (in percent):

FIGURE 1 shows Charpy impact values in the case that the steel of the composition shown in Table 1 is heated at respective austenitizing temperatures for 1 hour, is air-cooled, is tempered at 600 C. for 1 hour and is then water-cooled. That is to say, it is seen that, in case the steel is heated to about 800 C., the impact value will.

be the highest and that, at a heating temperature above 850 C., the impact value will quickly reduce.

In the present invention, in case there is a possibility that more than 0.005% AlN will be formed in the steel as calculated from the N content (including the case that N is not positively added) and Al content, there will be carried out as required the following heat-treatment wherein the steel is hot-rolled, then subjected to a partial solution treatment at a temperature above 850 C. but below the crystal grain coarsening temperature, or heated to a proper temperature above the A transition point without solution treatment so that AlN may precipitate in finely dispersed manner and at the same time the austenite crystal grains may be made fine. It is then quenched or air-cooled so as to obtain martensite or mixed structure of martensite and bainite. It is tempered in the temperature range of 525 to 650 C. so that a fine austenite structure may precipitate and is quenched or air-cooled. In the present invention, the heat-treatment wherein the partial solution treatment is carried out is called a twostep austenitizing treatment and the heat-treatment including no partial solution treatment is called a one-step austenitizing treatment. Further, in case there is a possibility that an AlN content will be formed in an amount exceeding 0.005%, the object of the present invention will be attainable by the following treatment wherein the steel is hot-rolled, is subjected to a complete solution treatment at a temperature above 1200 C., is quenched, is heated for a proper time at a temperature around the A transition point so that AlN may precipitate, in finely dispersed manner, is cooled, is then heated at a temperature just above the A transition point so that the austenite grains may be made fine, is then quenched or air-cooled so as to obtain martensite or mixed structure of martensite and bainite, is then tempered at 525 to 650 C. so that a fine austenite structure may precipitate, and is quenched or air-cooled. This is called a three-step austenitizing treatment. In the present invention, the steel made to contain the required elements in the above described ranges is heat-treated in response to the composition to regulate the fine structure so that the toughness and strength at low temperatures may be increased. That is to say, in the precipitated austenite, such alloying elements as Ni, Mn, N and C in the steel are enriched more than in the average composition of the alloy and it forms an austenite whereas, in the ferrite matrix, those elements are rather less than in the average composition and especially the amounts of C and N solid-dissolved in the ferrite matrix are extremely small. The effects of both of such facts as described above serve to improve the low temperature toughness.

In the steel of the present invention, in the process of tempering the martensite structure or martensite-bainite mixed structure produced by quenching or air-cooling from a proper temperature, with the help of the effect of accelerating the diffusion by the presence of many dislocation groups therein, a fine austenite structure in which such alloying elements as N, C, Ni and Mn are enriched will precipitate in the martensite grain boundaries, austenite grain boundaries or ferrite subgrain boundaries as mentioned above, therefore the amounts of such elements in the ferrite matrix will reduce and, with the elimination and rearrangement of the dislocations within the martensite matrix, ferrite matrix containing fine subgrain groups in which the amounts of C and N are very small will be formed. Thus, free N and C existing in solid-solution which are undesirable to the toughness of the steel will be fed from the ferrite matrix into the precipitated austenite. The austenite precipitated in the grain b undaries will be in a state in which N, C, Ni and Mn are enriched as mentioned above, but its fine dispersed state and stability will be determined by the tempering temperature and time corresponding to the alloy compositions. The preferable temperature is 525 to 650 C. In case a comparatively large amount of Mn is contained, the role of free N in the tempering process will be especially important to the state of dispersion and the stability of the precipitated austenite, and so to the toughness and strength of the steel at low temperatures. Further, the nitride formed of such element chemically strongly combined with N as, for example, Al (or Zr, Ti, Be, Nb, V, Hf, Te or B) will serve as a grain refining and dispersion hardening agent for the steel of the present invention. It is already widely known that such precipitants are effective to refine the austenite grains and toughen the steel. However, the present invention includes also a heat-treating method of making the state of the formation and dispersion of the nitrides as fine as possible. That is to say, there is carried out a treatment wherein gigantic AlN formed at the time of freezing or hot-rolling an ingot is partly or completely solid-dissolved in an austenite and precipitated from supersaturated state at a comparatively low temperature.

The reasons why the contents of the respective elements in the present invention are defined to be in the above mentioned ranges shall be described.

C is useful to improve the quenchability of the steel. That is to say, it is necessary to obtain martensite structure as quenched from the austenitizing temperature and will form the dislocations of a high density in the matrix. Further, C will diffuse and be absorbed in to the austenite precipitated at the time of tempering and will increase the stability of the austenite at low temperatures. By taking these points into consideration, its lower limit is made 0.01%. On the other hand, if the content of C increases, the amount of the solid-dissolved carbon in the ferrite matrix in the tempering process will increase and will impair the toughness. By taking this fact into consideration, the upper limit is made 0.15%.

Si is an element which will improve the toughness of the steel and will increase the strength. It is also an element required for making steels. If it is less than 0.05%, the above mentioned object will not be attained. If it is added to be more than 0.4%, its toughness will tend to reduce. Therefore, Si is defined to be in this range.

Ni is an element useful for the toughness and strength of the steel. Especially it will serve to improve the toughness at the boiling point of liquid nitrogen. Further, with the help of the dislocation of a high density formed at the time of quenching or air-cooling, Ni will diffuse and be absorbed into the precipitated austenite comparatively quickly and will be able to stabilize the precipitated austenite. From these facts, it is necessary to add more than 4.5% Ni. If too much Ni is added, the cost of the steel "will become high. Therefore, by taking these points into consideration, Ni is defined to be less than 7.50%.

Mn will improve the quenchability of the steel, will stabilize the fine austenite precipitated at the time of tempering same as N, C and Ni, will increase the toughness of the ferrite matrix and will improve the toughness. However, if Mn is more than 3.5% in the steel, the toughness of the steel will be impaired. For example, in a simple series alloy steel of 0.05 to 0.1% C and 6% Ni containing 3.5% Mn, the temper brittleness at 500 to 600 C. will be so severe as to reduce toughness at low temperatures very much. This is evident from FIG- URE 2 obtained by experimenting on the effect of Mn on the impact characteristics of the simple composition steel shown in Table 2. In FIGURE 2, the solid line is of the steel which was heated at 800 C. for 1 hour, was then Water-cooled, was tempered at 600 C. for 1 hour and was then water-cooled and the dotted line is of the steel which was heated at 800 C. for 1 hour, was then air-cooled, was tempered at 600 C. for 1 hour and was then water-cooled.

That is to say, if the content of Mn is high, (Fe:Mn) C will be present stably up to high temperatures and therefore C will be retained in the ferrite. As mentioned above, this fact will have a bad influence on the toughness of the ferrite matrix at the time of tempering the steel. For such reasons, the upper limit of Mn is made 3.5%. On the other hand, if less than 0.5% Mn is added, the expected effect will not be obtainable.

M0 is an element to be added as required and is useful to reduce the temper brittleness of the simple series alloy containing Ni, Mn and C. Further, according to the observation with an electron microscope, M0 in the steel will delay the recovery of the martensite, will therefore refine the dispersed state of the austenite precipitated in the grain boundary, will accelerate the difiusion of Ni, Mn, C and N and will extend the optimum tempering temperature to a higher temperature range. In order to obtain such result, the range of 0.05 to 0.50% M0 is preferable.

In case N is present as a nitride as combined with Al (or Zr, Ti, Be, Nb, V, Hf, Ta or B), it will serve as a grain refining and dispersion hardening agent. Further, what is to be noted is that free N not fixed as nitride will contribute to the stabilization of the precipitated austenite. This fact will perform an especially important role in the process of tempering an alloy containing a. comparatively large amount of Mn. More than 0.001% N is contained in a normally smelted steel. In order to attain the above mentioned object, less than 0.05% N will be sufficient. Therefore, the range of addition of N is defined to be 0.001 to 0.05%.

Al is not only added as a deoxidizing agent but is necessary to fix the required amount of N. Its amount is diiferent depending on the setting of the ratio of N to be fixed as AlN to free N. But, in case total N exceeds 0.025%, the maximum amount of addition of Al will be made 0.05% as acid-soluble. Al is to be used for the above mentioned object. In place of or in addition to Al, there may be used one or more of Zr, Ti, Be, Nb, V, Hf, Ta and B.

Examples of the present invention are given in the following:

Example 1 A steel of the composition shown in Table 3 was hotrolled, was heated at 800 C. for 1 hour and was cooled with water or in air. The structure as air-cooled is shown in FIGURE 3 (which is a micrograph magnified by 50,000 times, by direct observation with an electron microscope). As evident from the photograph, the structure is mixed structure of martensite (A) having a high dislocation density and bainite (B) (the scattered black dots are of cementite). (C) is an old (former) austenite crystal grain boundary. The steel was then tempered at each of the temperatures of 500 to 600 C. and 625 C. for 1 hour and was then watercooled. The structure of the specimen air cooled and further tempered at 600 C. for 1 hour, is shown in FIGURE 4 (which is a micrograph magnified by 50,000 times, by direct observation with an electron microscope). That is to say, in this structure are shown ferrite matrix changed from martensite matrix (A) having fine subgrains and a fine austenite structure (B) precipitated in the old (former) martensite crystal grain boundary. The steel having such structure is very high in strength and toughness at low temperatures. The mechanical properties of this steel are shown in Tables 4-1 and 4-2. For comparison, those of the conventional 9%-Ni steel of the A.S.T.-M. specification are shown. This 9%-Ni steel was heat-treated exactly the same as in the present example of the present invention.

TABLE 3.--Chemical composition (in percent) Constituent Tested steels O Si Mn N 1 Mo Al N Steel of the present invention 0.70 0. 24 1.90 6. 26 0.22 0. 005 0.0013 9%-Ni steel of A.S.T.M. specification (Conventional steel) 0.10 0.25 0.8 9. 0 0. 01 0. 001

TABLE 4-1.TENSILE STRENGTHS AS TEMPERED AT 600 0. FOR 1 HOUR Measuring methods Tensile strength Yield strength in Kg./mm. in Kg./mm.= Elongation in percent Room Room Room tempertempertemper- Tested steels ature ---196 0. ature --196 C. ature 196 0,

Treatment A; Steel of the present invention. .0 120.3 81 .4 114.1 17.7 25 .5 Treatment B; Steel of the present invention. 87 .4 126. 6 80.1 107 .2 19 .4 30 .3 9%-Ni steel oi A.S.T.M. specification (conventional steel) 80 .1 .5 64 .7 86 .0 26 .6 28 .3

TABLE 42.2 MM. V-NOTCH CHARPY IMPACT VALUE (IN KG.-M./CM.

Tempering temperature 500 C. for 1 hour 600 C. for 1 hour 625 C. for 1 hour Measuring temperature Room Room Room tempertempertemper- Tested steels ature 0. 196 C. ature -150 0. 196 C. eture 150 0. 196 0 Treatment A; Steel of the present invention 18 .6 2 .4 0 .84 26 .5 18 .5 8 .6 18 .2 10 .1 8 .10 Treatment B; Steel of the present invention 18 .6 3 .2 1 .80 24 .5 19 .2 12 .8 18 .8 10 .5 8 .70 9%-Ni steel of A.S.T.M. specification (conventional steel) 4 .28 26 .1 15 .19 10 .7 5 .32

Note.Heat treatments denoted A and B applied to the steels of the present invention in Tables 4-1 and 42 are as follows: According to treatment A, the steel was heated at 800 C. for 1 hour, was water-cooled, was then tempered at each temperature and was cooled and according to treatment B, the steel was heated at 800 C. for 1 hour, was air-cooled, was then tempered at each temperature and was cooled.

As mentioned above, though the optimum tempering temperature of this invention is 600 C., the mechanical properties of the steel of the present invention are higher than those of the 9%-Ni steel. Especially it is a great feature of the present invention that, even by the tempering treatment at higher temperatures, reduction of the toughness is not so serious. Further, the air-cooling from the austenite temperature will give properties better than by the water-cooling, irrespective of the subsequent tempering temperature. As a result, it is shown that the steel of the present invention is a practically excellent steel.

Example 2 The steel shown in Table 5 and having had N in an amount larger than would be fixed by Al added was hotrolled, was then heated at 1200 C. for 2 hours, was quenched, was then heated at 760 C. for 1 hour and was In such case, in the calculation, about 150 ppm. of excess nitrogen was present as solid-dissolved in the steel. But the amount of N remaining in the ferrite matrix tempered at 600 C. for 1 hour was about 3 p.p.m. as measured by the internal friction method. The greater part of N had flowed out into the austenite precipitated in the tempering process, had contributed to the stabilization of the austentite and had improved the toughness of the steel. In view of the fact that the impact value of the 6-Ni fundamental series alloy of the present invention shown in FIGURE 2 is about 3 l g.-m./cm. at l96 C., the effect of free nitrogen present in the steel is more evident.

Example 3 When a steel of the composition shown in Table 7 was subjected to the respective heat-treatments shown in Table 8, such impact values as are shown in the same table were obtained.

TABLE 7.CHEMICAL COMPOSITION IN PERCENT Composition of the steel of the present invention C Ni Si Mn Al N lleat treatments Charpy impact values (with 2 mm. V-noteh) at 190 C. in

Kg.-m.lcn1. Remarks) Treatments Nos;

and water-eooled.

Heated at 600 C. for 1 hour and watercooled.

Heated at 000 C. for 1 hour and \vaterpooled.

Heated at 600 C. for 1 hour and water-cooled.

Heated at 1,200 C. for 1 hour and water-cooled. 3

Heated at 600 C. for 1 hour and water-cooled.

Heated at 1,200 C. for 1 hour and water-cooled.

Heated at 900 C. for 1 hour 5 and water-cooled.

Heated at 600 C. for 1 hour and water-cooled.

Heated at 1,350 C. for 1 hour and water-cooled.

Heated at 900 C. for 1 hour and water-cooled.

Heated at 600 0. for 1 hour and watercooled.

Heated at 1,350 0. for 1 hour and watercooled.

Heated at 700 C. for 2 hours and air-cooled.

Heated at 900 C. for hour and water-cooled.

Heated at 600 C. for 1 hour and watencooled.

{Heated at 800 C. [or 1 hour 1 Austemte onestep treating method.

. 4 Austenite two-step treating method.

5.75 Austenite three-step treating method.

water-cooled. It was then tempered at 600 C. for 1 hour and was water-cooled. An example of the mechanical properties of the obtained steel is shown in Table 6.

TABLE 5.CHEMICAL COMPOSITION (IN PERCENT) Composition of the steel of the present invention 0 Si Mn Ni Al N TABLE 6.-MEGHANICAL PROPERTIES In Table 8, the treatments Nos. 1 to 3 are by the onestep treating method, the treatments Nos. 4 and 5 are by the two-step treating method and the treatment No. 6 is by the three-step treating method. In the one-step treatment, when the heating temperature was made high as shown in the treatment No. 3, the toughness at low temperature reduced a little. But, in such case, by the two-step treatment, the toughness was increased as shown in the treatment No. 4. Further, in the two-step treatment, when ture (as in the case of the treatment No. 5 at 1350 C.),

Tensile strength Yield strength in Elongation in 2 mm. V-

in Kgjmm. KgJmm. percent notch Charpy value at 196 Room Room Room C. in Kg.- tempertempertemper- V nL/om.

ature -106 C. ature 1.)0 C. ature C.

the toughness at low temperature tended to reduce. However, a shown in the treatment No. 6, when the heating treatment around the A transition point was carried out in the middle, the steel came to havea good toughness.

FIGURE 5-A is micrograph showing a fine structure tempered at 600 C. by the direct observation with an electron microscope of a magnification of 50,000 after the one-step treatment in the treatment No. 2 in Table 8 was carried out. C represents precipitated AlN in the ferrite crystal grain boundary. D represents stable austenite islands precipitated in the crystal grain boundary. E is a dislocation network within the ferrite crystal. It is evident that the ferrite crystal grains are small and that the dislocation is in a stable arrangement and, in consequence, this structure will show a good toughness.

FIGURE 5B is a micrograph by extraction replica method with an electron microscope as taken at a lower magnification (of 5,000 times) and shows the dispersion of AlN (black dots).

FIGURE 6-A is a micrograph showing a fine structure tempered at 600 C. by the direct observation with an electron microscope of a magnification of 50,000 after the two-step treatment in the treatment No. 4 in Table 8 was carried out. C is precipitated AlN. D is stable austenite grains precipitated in the crystal grain boundary. E isa dislocation within the ferrite crystal. It seems generally that the ferrite grains are small but that the dislocation density within the ferrite crystal grain is high.

FIGURE 6-B is a micrograph by extraction replica method with an electron microscope as taken at a lower magnification (of 5,000 times) and shows the dispersion of AlN (black dots). In this two-step treating method, the dispersion of AlN is finer and the dimensions are generally smaller than in the one-step method in FIGURE 5-B and the effect of the partial solution treatment is evident.

These effects are caused mostly by the normalization of the dimensions and dispersion of AlN and the change of the austenite crystal grains by changing the solution treatment of AlN and the subsequent precipitation state.

As described above, according to the present invention, a steel of which the Ni content is comparatively low, the strength is high and the toughness at low temperatures is high can be produced by applying proper heat-treatments.

We claim:

1. A heat treated low-temperature tough steel in the tempered condition, having a finely dispersed austenitic phase in a ferritic matrix, consisting essentially of 0.01 to 0.15% C, 0.05 to 0.40% of Si, 4.50 to 7.50% Ni, 0.50 to 3.5% Mn, 0.001 to 0.050% N and less than 0.05% acidsoluble -Al, the balance being Fe; said steel being characterized by the N content being in excess of that necessary to be fixed as a nitride, wherein the excess nitrogen is substantially absorbed in the temper formed austenite phase.

2. A steel according to claim 1 wherein at least part of the dispersed austenite structure is replaced by martensite.

3. Steel according to claim 1 wherein the acid-soluble Al is replaced at least partly by at least one substance selected from the group consisting of Zr, Ti, Be, Hf, Ta and B.

4. A heat-treated low-temperature tough steel in the tempered condition having a finely dispersed austenitic phase in a ferritic matrix, conissting essentially of 0.01 to 0.15% C, 0.05 to 0.40% Si, 4.50 to 7.50% Ni, 0.50 to 2.50% Mn, 0.001 to 0.050%N, 0.05 to 0.50% "Mo and less than 0.05% acid-soluble Al, the balance being essentially Fe; said steel being characterized by the N content being in excess of that necessary to be fixed as a nitride, wherein the excess nitrogen is substantially absorbed in the temper formed austenite phase.

5. A steel according to claim 4 wherein at least part of the dispersed austenite structure is replaced by martensite.

6. Steel according to claim 4, wherein the acid-soluble Al is replaced at least partly by at least one substance selected from the group consisting of Zr, Ti, Be, Hf, Ta and B.

References Cited UNITED STATES PATENTS 2,206,370 7/1940 Scherer 123 X 2,992,148 7/1961 Yeo 148-36 2,516,125 7/1950 Kramer 75123 3,155,549 11/1964 Nakamura 75--124 3,249,426 5/1966 Nakamura 75124 CHARLES N. LOVELL, Primary Examiner.

U.S. C1. X.R.

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