Patent Application
20090004041
1. Field of the Invention
This invention relates to a steel alloy that provides high strength and high fracture toughness, and in particular to such an alloy that also provides improved fatigue resistance.
2. Description of Related Art
Steel alloys having a good combination of high strength and high fracture toughness are known. Examples of such alloys are described and claimed in U.S. Pat. No. 5,087,415; U.S. Pat. No. 5,268,044; U.S. Pat. No. 5,866,066; and U.S. Patent Application Publication No. 2007/0113931; the entire disclosures of which are incorporated herein by reference. The alloys described in those documents were originally designed for structural aerospace uses such as aircraft landing gear components, aircraft structural members, such as braces, beams, and struts. The known alloys have been made with a process that includes the use of a rare earth treatment to control the size and shape of inclusions such as sulfides, oxides, and oxysulfides that would otherwise adversely affect the strength and toughness of the alloys in the structural applications for which they are designed.
Since the original development of the above-mentioned alloys, a new application of the alloys has arisen, namely, in rotating shafts for jet engines. In the course of developing rotating shafts made from these alloys, it has been found that the fatigue life of the rare earth treated alloys, particularly the low cycle fatigue life, leaves something to be desired. However, the combination of high strength and high fracture toughness provided by these alloys is still highly desirable for the rotating shaft application. Therefore, it would be desirable to have a steel alloy that provides the combination of high strength and high fracture toughness provided by the known alloys, together with improved fatigue life for the rotating shaft application.
The present invention relates to an age-hardenable, martensitic steel alloy having a composition and microstructure that is designed to provide a significant improvement in the fatigue life of a finished component, such as a rotating shaft of the type suitable for use in a jet engine or gas turbine. Summarized in the table below are the Broad, Intermediate, and Preferred chemistries of the material according to the present invention. The values set forth in the table are given in weight percent.
The balance of the alloy is iron and usual impurities, including additional elements in amounts which do not detract from the desired combination of properties. The alloy according to this invention is further characterized by a dispersion of small inclusions having a size distribution of about 0.4 μm to about 7 μm in major dimension in the alloy matrix. Preferably, the median inclusion size is at least about 1.6 μm. The composition of the inclusions contains essentially no rare earth elements such as cerium and lanthanum.
In accordance with another aspect of the present invention, there is provided a method of improving the low cycle fatigue life of a high strength, high toughness, age-hardenable martensitic steel alloy. The method includes the step of melting an age-hardenable martensitic steel alloy having the weight percent composition set forth above. The method further includes the step of adding calcium to the molten alloy to combine with available sulfur and oxygen in the molten alloy to form inclusions that are removable from said alloy. The method also includes the steps of processing the alloy to remove at least a portion of the inclusions from the alloy and then solidifying the refined alloy, whereby the solidified alloy contains a limited dispersion of such inclusions in the alloy matrix. The retained inclusions have a size distribution of about 0.4 μm to about 7 μm in major dimension and a median size of at least about 1.6 μm. The alloy can be mechanically worked and machined to form a useful product such as a shaft for a rotating machine.
The foregoing tabulation is provided as a convenient summary and is not intended to restrict the lower and upper values of the ranges of the individual elements of the alloy of this invention for use solely in combination with each other, or to restrict the broad, intermediate or preferred ranges of the elements for use solely in combination with each other. Thus, one or more of the broad, intermediate, and preferred ranges can be used with one or more of the other ranges for the remaining elements. In addition, a broad, intermediate, or preferred minimum or maximum for an element can be used with the maximum or minimum for that element from one of the remaining ranges. Here and throughout this application percent (%) means percent by weight unless otherwise indicated.
The foregoing summary as well as the following detailed description will be better understood when read in conjunction with the drawings, wherein:
The alloy according to the prevent invention contains at least about 0.2%, better yet, at least about 0.20%, and preferably at least about 0.21% carbon because it contributes to the good hardness capability and high tensile strength of the alloy primarily by combining with other elements such as chromium and molybdenum to form carbides during heat treatment. Too much carbon adversely affects the fracture toughness of this alloy. Accordingly, carbon is limited to not more than about 0.36%, better yet, to not more than about 0.33%, and preferably to not more than about 0.27%.
Cobalt contributes to the hardness and strength of this alloy and benefits the ratio of yield strength to tensile strength (Y.S./U.T.S.). Therefore, at least about 8%, better yet at least about 10%, and preferably at least about 11% cobalt is present in this alloy. Above about 22% cobalt the fracture toughness and the ductile-to-brittle transition temperature of the alloy are adversely affected. Preferably, not more than about 17%, and better yet not more than about 14% cobalt is present in this alloy.
Cobalt and carbon are critically balanced in this alloy to provide the unique combination of high strength and high fracture toughness that is characteristic of the alloy. Thus, to ensure good fracture toughness, carbon and cobalt are preferably balanced in accordance with the following relationship:
To ensure that the alloy provides the desired high strength and hardness, carbon and cobalt are preferably balanced such that:
Chromium contributes to the good hardenability and hardness capability of this alloy and benefits the desired low ductile-brittle transition temperature of the alloy. Therefore, at least about 1.3%, better yet at least about 2%, and preferably at least about 2.9% chromium is present. Above about 4% chromium the alloy is susceptible to rapid overaging such that the unique combination of high tensile strength and high fracture toughness is not attainable with the preferred age-hardening heat treatment. Preferably, chromium is limited to not more than about 3.5%, and better yet to not more than about 3.3%. When the alloy contains more than about 3% chromium, the amount of carbon present in the alloy is adjusted upwardly in order to ensure that the alloy provides the desired high tensile strength.
At least about 0.75% and preferably at least about 1.0% molybdenum is present in this alloy because it benefits the desired low ductile-brittle transition temperature of the alloy. Above about 2.7% molybdenum, the fracture toughness of the alloy is adversely affected. Preferably, molybdenum is limited to not more than about 1.75%, and better yet to not more than about 1.5%. When more than about 1.5% molybdenum is present in this alloy the % carbon and/or % cobalt must be adjusted downwardly in order to ensure that the alloy provides the desired high fracture toughness. Accordingly, when the alloy contains more than about 1.5% molybdenum, the % carbon is not more than the median % carbon for a given % cobalt as defined by equations a) and b) or a) and c).
Nickel contributes to the hardenability of this alloy such that the alloy can be hardened with or without rapid quenching techniques. Nickel benefits the fracture toughness and stress corrosion cracking resistance provided by this alloy and contributes to the desired low ductile-to-brittle transition temperature. Accordingly, at least about 10%, better yet, at least about 10.5%, and preferably at least about 11.0% nickel is present. Above about 15% nickel the fracture toughness and impact toughness of the alloy can be adversely affected because the solubility of carbon in the alloy is reduced which may result in carbide precipitation in the grain boundaries when the alloy is cooled at a slow rate, such as when air cooled following forging. Preferably, nickel is limited to not more than about 13.0%, and better yet to not more than about 12.0%.
Other elements can be present in this alloy in amounts which do not detract from the desired properties. Not more than about 0.20% manganese can be present because manganese adversely affects the fracture toughness of the alloy. Preferably, manganese is restricted to about 0.15% max. and better yet to about 0.10% max. For best results the alloy contains not more than about 0.05% manganese. Up to about 0.1% silicon, up to about 0.01% aluminum, and up to about 0.02% titanium can be present as residuals from small additions for deoxidizing the alloy.
A small but effective amount of calcium is present in this alloy to provide sulfide and oxide inclusion shape control which benefits the fracture toughness and the low cycle fatigue life of the alloy. It is believed that the use of calcium in this alloy benefits the low cycle fatigue (LCF) life because it beneficially affects the size and distribution of inclusions that form in the alloy matrix during processing. Unlike the alloys described in U.S. Pat. No. 5,087,415; U.S. Pat. No. 5,268,044; U.S. Pat. No. 5,866,066; and U.S. Patent Application Publication No. 2007/0113931, the present invention avoids the use of rare earth treatments. Therefore, the alloy or product made in accordance with the present invention may contain only trace amounts of such elements as cerium, lathanum, and other rare earths.
The balance of the alloy according to the present invention is essentially iron except for the usual impurities found in commercial grades of alloys intended for similar service or use. The levels of such elements must be controlled so as not to adversely affect the desired properties of this alloy. For example, phosphorus is limited to not more than about 0.01%, preferably not more than about 0.008%. Sulfur adversely affects the fracture toughness provided by this alloy. Accordingly, sulfur is restricted to about 0.0040% max., better yet to about 0.0025% max., and preferably to about 0.0020% max. Best results are obtained when the alloy contains not more than about 0.001% sulfur. Tramp elements such as lead, tin, arsenic and antimony are limited to about 0.003% max. each, better yet to about 0.002% max. each, and preferably to about 0.001% max each. Oxygen is limited to not more than about 20 parts per million (ppm) and nitrogen to not more than about 40 ppm.
The alloy of the present invention is readily melted using conventional vacuum or inert gas melting techniques. For best results, as when additional refining is desired, a multiple melting practice is preferred. The preferred practice is to melt a heat in a vacuum induction furnace (VIM) and cast the heat in the form of an electrode. The alloying addition for sulfide shape control referred to above is preferably made before the molten VIM heat is cast. Preferably, the electrode is then remelted in a vacuum arc furnace (VAR) and recast into one or more ingots. It is expected that the alloy would contain not more than about 0.001% calcium after VAR processing. Prior to VAR the electrode ingot is preferably stress relieved at about 1250° F. for 4-16 hours and air cooled. After VAR the ingot(s) is(are) preferably homogenized at about 2150-2250° F. for 6-24 hours. This alloy can also be prepared using powder metallurgy techniques, for example, VIM followed by inert gas atomization.
The alloy can be hot worked from about 2250° F. to about 1500° F. The preferred hot working practice is to forge an ingot from about 2150-2250° F. to obtain at least a 30% reduction in cross sectional area. The ingot is then reheated to about 1800° F. and further forged to obtain at least another 30% reduction in cross sectional area. It will be appreciated that the alloy can also be hot worked using a single reduction step for some product forms.
The alloy according to the present invention is austenitized and age hardened as follows. Austenitizing of the alloy is carried out by heating the alloy at about 1550-1800° F. for about 1 hour plus about 5 minutes per inch of thickness and then quenching in oil. The hardenability of this alloy is good enough to permit air cooling or vacuum heat treatment with inert gas quenching, both of which have a slower cooling rate than oil quenching. Whatever quenching technique is used, the quench rate is preferably rapid enough to cool the alloy from the austenitizing temperature to about 150° F. in not more than about 2 h. When this alloy is to be oil quenched, however, it is preferably austenitized at about 1550-1600° F., whereas when the alloy is to be vacuum treated or air hardened it is preferably austenitized at about 1575-1650° F. After austenitizing, the alloy is preferably cold treated as by deep chilling at about −100 to −320° F. for ½ to 1 hour and then warmed in air. Age hardening of this alloy is preferably conducted by heating the alloy at about 850-950° F. for about 5 hours followed by cooling in air.
In order to demonstrate the improvement in fatigue life provided by the alloy according to the present invention relative to the known alloy, comparative testing was conducted. Samples for testing were obtained from a heat of a known rare-earth treated (RE Treated) material and from a calcium treated (Ca Treated) heat of the alloy according to the present invention. The weigh percent chemistries of the two heats are set forth in Table 1 below.
The balance of each composition is iron and the usual impurities.
Longitudinal and transverse sections of the sample materials were obtained. Standard tensile and fracture toughness specimens were prepared from each of the sections. The tensile and fracture toughness specimens were heat treated by heating at 1625° F. for one hour and then cooled in air. The test specimens were then deep chilled at −100° F. for one hour, followed by warming in air. The specimens were then age hardened by heating at 900° F. for five hours and then air cooled. The results of the tensile land fracture toughness testing are shown in Table 2 below including the 0.2% offset yield strength and the ultimate tensile strength in ksi, the percent elongation, the percent reduction in area, and the KIc fracture toughness in ksi√in.
The data presented in Table 2 show that the Ca-treated alloy according to the presented provides tensile properties and fracture toughness that is at least as good as the known rare earth treated alloy. There was no adverse effect on the tensile and fracture toughness properties that resulted from the calcium treatment.
Transverse blanks measuring ¾ in. square×4½ in. long were cut from each of the heats. The blanks were heat treated using the same heat treatment used for the tensile and fracture toughness specimens as described above. The heat treated blanks were low stress ground to form axial-axial fatigue test specimens.
Smooth fatigue test specimens (Kt=1.0) for room temperature axial-axial fatigue testing were prepared in accordance with ASTM E466-96. Servo-Hydraulic test equipment was used with a 20 Hz sinusoidal waveform to cycle the specimens from zero stress to each of three maximum tensile stress levels of 1400 MPa, 1200 MPa, and 1100 MPa (203 ksi, 174 ksi, and 160 ksi, respectively). Six samples of each heat were tested at each stress level. This axial-axial fatigue test was therefore performed under conditions that gave R=0 and Kt=1. For economic reasons, testing of a specimen was discontinued at 1,728,000 cycles (24 hours) if the specimen did not fail by that time.
The results of the fatigue tests for the rare earth treated specimens are shown in Table IIIA and the results for the calcium treated specimens are shown in Table IIIB.
The data presented in Tables IIIA and IIIB are graphed in
Material from the samples that broke during the fatigue testing were analyzed in a scanning electron microscope (SEM) to characterize the inclusions formed in each alloy by size and composition. The SEM examined 191 inclusions in the samples from the calcium treated heat and 156 inclusions in the rare earth treated material. Table IV presents the sizes and compositions of the inclusions observed in the rare earth treated alloy. Table V presents the sizes and compositions of the inclusions observed in the calcium treated material.
La, S, Ce, Fe, Cr, Co, Ni, P, C, O
La, S, Ce, Fe, Cr, Co, Ni, P, C, O
La, S, Ce, Fe, Cr, Co, Ni, O, C
La, S, Ce, Fe, Cr, Co, Ni, O, C
La, Ce, S, Fe, Cr, Co, Ni, O and C
S, La, Ce, Fe, Cr, P, Co, Ni, As, C, W
La, Ce, S, Fe, Cr, Co, Ni, O and C
S, La, Ce, Fe, Cr, P, Co, Ni, As, C, Mo
Ce, P, Fe, La, S, Cr, Co, As, Ni, C
La, Fe, Ce, Cr, Co, Ni, S
S, Fe, La, Ce, Cr, Co, Ni, As, W, P, C
La, S, Ce, Fe, Cr, Co, Ni, P, O, As, C, W
La, S, Ce, Fe, Cr, Co, Ni, P, O, As, C, W
S, La, Ce, Fe, Cr, O, Co, Ni, C and As
S, La, Ce, Fe, Cr, Co, Ni, C and As
Ce, Fe, Co, Ni, La, Ce and C
La, S, Ce, Cr, Fe, O, C, Co, Ni, Sr, Ga
La, S, Ce, Cr, Fe, O, C, Co, Ni, Sr, Ga
S, La, Ce, Fe, Cr, O, Co, C, Ni, Sr
La, S, Ce, Cr, Fe, O, C, Co, Ni, Sr, Ga
La, S, Ce, Cr, Fe, O, C, Co, Ni, Sr, Ga
S, La, Ce, Fe, Cr, O, Co, C, Ni, Sr
La, S, Ce, Cr, Fe, O, C, Co, Ni, Sr, Ga
La, S, Ce, Cr, Fe, O, C, Co, Ni, Sr, Ga
La, S, Ce, Cr, Fe, O, C, Co, Ni, Sr, Ga
La, S, Ce, Cr, Fe, O, C, Co, Ni, Sr, Ga
Ce, P, Fe, La, Cr, O, Co, Ni, C, As
Fe, Ce, P, La, Cr, Co, Ni, O, C
Ce, P, Fe, La, Cr, O, Co, Ni, C, As
Ce, P, La, Fe, O, Cr, C, Co, As, Si, Ni
La, S, Ce, Fe, Cr, Co, O, Ni, C, W, P
La, S, Ce, Fe, Cr, Co, Ni, O, W, C
Ce, P, La, Fe, O, Cr, C, Co, As, Si, Ni
Ce, P, La, Fe, O, Cr, C, Co, As, Si, Ni
La, S, Ce, Fe, Cr, Co, O, Ni, C, W, P
Fe, Ce, P, La, Cr, Co, Ni, O, C
La, S, Ce, Fe, Cr, Co, Ni, O, W, C
La, S, Ce, Fe, Cr, Co, Ni, O, W, C
Ca, Fe, Ti, As, Co, Ni, Cr, O, S, P
Ca, Fe, As, Co, O, Ti, Cr, Ni and S
Ca, As, Fe, O
S, Ca, Fe, Co, Ni, Cr, Ti
S, Ca, Fe, Co, Ni, Cr, Ti
Ca, Fe, Ti, As, Co, Ni, Cr, O, S, P
Ti, S, Ca, Fe, Co, Cr, Ni
Al, O and Ti plus a trace of Fe
S, Ca, Fe, Co, Ni and Cr
S, Ca, Fe, Co, Ni and Cr
Ca, Fe, S, Co, Ti, Cr, Ni
S, Ca, Fe, Co, Ni, Cr
Ca, Fe, Co, As, Ni, Cr, S, O, Ti
Ti, S, Fe, Ca, Zr, Cr, Co, Ni, As
S, Ca, Fe, Co, Ni, Cr
Ca, Fe, S, Co, Ni, Cr, As, O
S, Ca, Fe, Co, Ni, Cr
S, Ca, Fe, Co, Ni, Cr
S, Ca, Fe, Co, Ni, Cr
Ca, Fe, S, Co, Ti, Cr, Ni
S, Ca, Fe, Co, Ni, Cr
Ca, Fe, As, Cr, Ni, O, S, W
Ca, Fe, S, Co, Ti, Cr, Ni
S, Ca, Fe, Co, Ni, Cr
Ca, Fe, S, Co, Ti, Cr, Ni
S, Ca, Fe, Co, Ni, Cr
S, Ca, Fe, Co, Ni, Cr
S, Ca, Fe, Co, Ni, Cr
Ca, Fe, S, Co, Ni, Cr, As, O
Ca, O and Fe plus traces of S, Si and Al
Ca, O and Fe plus traces of S, Si and Al
Ca, O and Fe plus traces of S, Si and Al
Ca, O and Fe plus traces of S, Si and Al
Ca, O and Fe plus traces of S, Si and Al
Ca, O and Fe plus traces of S, Si and Al
Ca, Fe, S, Co, Cr, Ni, O, Si
Ca, Fe, S, Co, Cr, Ni, O, Si
Ca, O and Fe plus traces of S, Si and Al
Ca, Fe, S, Co, Cr, Ni, O, Si
Ca, O and Fe plus traces of S, Si and Al
Ca, Fe, S, Co, Cr, Ni, O, Si
Ca, O and Fe plus traces of S, Si and Al
Ca, Fe, S, Co, Cr, Ni, O, Si
Ca, O and Fe plus traces of S, Si and Al
Ca, O and Fe plus traces of S, Si and Al
Ca, Fe, Co, Ni, Cr, Mo, Ti, O
Ca, O and Fe plus traces of S, Si and Al
Ca, Fe, Co, Cr, Ni,
O and S
Ca, Fe, S, Co, Cr, Ni, O, Si
Ca, O and Fe plus traces of S, Si and Al
Ca, Fe, S, Co, Cr, Ni, O, S
Ca, O and Fe plus traces of S, Si and Al
Ca, Fe, Co, Cr, Ni, O and S
Ca, O and Fe plus traces of S, Si and Al
Ca, S, Fe and O
Ca, O and Fe plus traces of S, Si and Al
Ca, S, Fe and O
Ca, O and Fe plus traces of S, Si and Al
Ca, S, Fe and O
Ca, O and Fe
Ca, O and Fe
Ti, Fe, Mo, Zr, Cr, Ni, C
Ti, Mo, Fe, Cr, Co, Ca, Zr, C, Ni
Ca, S, Fe, Zr, Ce, Ti, Cr, Cr, Ni, Mg, Al
S, Fe, Ca, Co, Ni, Cr, Mg
Ca, S, Fe, Zr, Ce, Ti, Cr, Cr, Ni, Mg, Al
S, Ca, Fe, Co, Cr and Ni
Ca, S, Fe, Co, Ni, Cr, Si
Ca, Fe, As, Co, S, Cr, Ni, Si
Ca, Fe, As, Co, S, Cr, Ni, Si
Ca, S, Fe and Co
Ca, Fe, As, Co, S, Cr, Ni, Si
Ca, S, Fe, Co, Ni, Cr, Si
Ca, S, Fe, Co, Ni, Cr, Si
Ca, Fe, As, Co, S, Ni, Cr, Zr, Si, Ti, O
Ca, S, Fe, Co, Ni, Cr, Si
Ca, Fe, S, Co, Cr, Ni, As, Si, Ti, Zr
Ca, Fe, As, Co, S, Cr, Ni, Si
Ca, S, Fe, Co, Ni, Cr, Si
Ti, Ca, Mo, Fe and Zr
Ti, Fe, Mo, Co, Cr and Ni
The data in Tables IV and V show that the type of desulfurization treatment greatly affects the composition of the inclusions. Most of the inclusions in the rare earth treated alloy contain the rare earth elements Ce and La. In contrast, most of the inclusions in the calcium treated alloy contain Ca, but essentially no rare earth elements such as Ce and La. Both the rare earth treatment and the calcium treatment were effective in gettering tramp elements such as P, S, and As.
The inclusion size data set forth in Tables IV and V fit a Weibull distribution as shown in
The terms and expressions which have been employed herein are use as terms of description and not of limitation. There is no intention in the use of such terms and expressions to exclude any equivalents of the features described or any portions thereof. It is to be recognized, however, that various modifications are possible within the scope of the invention described and claimed.
This application claims the benefit of U.S. Provisional Application No. 60/946,244 filed Jun. 26, 2007.
| Number | Date | Country | |
|---|---|---|---|
| 60946244 | Jun 2007 | US |
