Classifications

• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
  • C21D1/18—Hardening; Quenching with or without subsequent tempering
  • C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
  • C21D1/22—Martempering
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D6/00—Heat treatment of ferrous alloys
  • C21D6/004—Heat treatment of ferrous alloys containing Cr and Ni
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D6/00—Heat treatment of ferrous alloys
  • C21D6/007—Heat treatment of ferrous alloys containing Co
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D6/00—Heat treatment of ferrous alloys
  • C21D6/02—Hardening by precipitation
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D6/00—Heat treatment of ferrous alloys
  • C21D6/04—Hardening by cooling below 0 degrees Celsius
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
  • C21D9/32—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for gear wheels, worm wheels, or the like
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/52—Ferrous alloys, e.g. steel alloys containing chromium with nickel with cobalt
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
  • C21D1/18—Hardening; Quenching with or without subsequent tempering
  • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D2211/00—Microstructure comprising significant phases
  • C21D2211/008—Martensite

Definitions

• the present invention relates to a high-performance carburized gear steel that can improve the performance of rotorcraft power transmissions, due to a unique and useful combination of surface hardness and core toughness.
• the U.S. Navy estimates that a 20% increase in gear durability would provide an annual cost saving of $17 million to the Defense Logistics Agency.
• the rotorcraft industry has not adopted a new gear steel for over twenty years, and instead focused on surface processing optimizations such as laser- peening, super-finishing, and directional forging. Such processes are providing diminishing returns in durability improvements.
• the subject invention provides a solution complementary to process enhancements and enables high-performance gears having reduced size and weight which are capable of transmitting more power at increased operating temperatures.
• Carburized X53 (U.S. Patent 4,157,258) is the incumbent material in rotorcraft transmissions. Compared to X53, the subject invention places an emphasis on increasing the case strength and the core fracture toughness, as well as increasing the thermal stability up to 450 0 C to provide hot hardness in high-temperature excursions.
• Patent No. 6,464,801 also discloses case-hardened steels. However, the embodiment Al of Patent No. 6,464,801 shows limited surface-hardness, i.e., Rockwell C scale hardness (HRC) of 60- 62. Another embodiment of Patent No. 6,464,801, steel C3, shows a greater surface hardness of 69 HRC, but the core of this steel lacks toughness.
• HRC Rockwell C scale hardness
• the core fracture toughness of the steel must exceed 50 ksiVin.
• a carburized gear steel with a surface hardness of HRC of at least about 62-64 at a usable core toughness exceeding 50 ksiVin has developed a need for a carburized gear steel with a surface hardness of HRC of at least about 62-64 at a usable core toughness exceeding 50 ksiVin.
• the present invention comprises a high-performance gear steel which is especially useful for rotorcraft transmissions.
• the steel exhibits an increase in surface hardness and core fracture toughness compared to conventional carburized gear steels.
• the steel is designed for a reasonable carbide solvus temperature, which, in turn, enables gas or vacuum carburization. Upon gas quenching from the solution heat treatment temperature, the steel transforms into a predominantly lath martensitic matrix.
• an optimal strengthening dispersion of secondary M 2 C carbide precipitates, where M is Mo, Cr, W and/or V.
• the high tempering temperature of the steel enables higher operating temperatures in transmission components compared to conventional gear steels like X53 or 9310.
• the matrix composition is carefully balanced to ensure the ductile-to -brittle transition is sufficiently below room temperature.
• the designed composition also effectively limits the thermodynamic driving force for precipitation of embrittling Topologically-Close-Packed (TCP) intermetallic phases such as ⁇ and ⁇ .
• Toughness of the invented steel is further enhanced by the distribution of a fine dispersion of grain-pinning particles that are stable during carburization and solution heat treatment cycles.
• the exemplary steel of the invention is designated as C64 in the above table.
• this steel is distinct from the steels disclosed in Patent No. 6,464,801 (i.e. Al, C2, and C3).
• Inclusion of W increases the M 2 C driving force similar to Cr or Mo, and uniquely limits the thermodynamic driving force for precipitation of undesirable TCP phases.
• Mo and Cr preferentially promote ⁇ -phase more than ⁇ -phase, W provides a reverse effect.
• W by adding W, the total driving force for ⁇ - and ⁇ -phases is balanced and precipitation of either TCP phase is avoided.
• Alloy C69B is a counterexample. Although alloy C69B does include W and successfully tends to avoid the precipitation of TCP phases, insufficient Ni in the matrix places the ductile-to -brittle transition above room temperature. The Ni content is thus greater in alloys of the embodiment of the invention to place the ductile-to -brittle transition above room temperature and concurrently maximize the driving force for M 2 C, enabling the highest surface hardness at a usable toughness compared to any other known secondary-hardening steel.
• the disclosed steel Due to high surface hardness, good core toughness, and the high-temperature capability, the disclosed steel is considered especially utilitarian with respect to gears for helicopter transmissions. Other uses of the steel include vehicle gearing and armor. With respect to the constituents in the exemplary steel set forth above, the alloy preferably includes a variance in the constituents in the range of plus or minus five percent of the mean value.
• FIG. 1 is a systems design chart representing the interactions among the desired hierarchical micro structure, the required processing and the property objectives for the alloy of the invention.
• FIG. 2 schematically illustrates the time-temperature processing steps for the subject alloy.
• FIG. 3 is a graph plotting the maximum surface hardness and core fracture toughness of various steels possibly useful for power transmission gears. Typical embodiments of the claimed invention are also plotted and identified as alloy C64.
• FIG. 4 is a graph plotting the Charpy V-Notch (CVN) impact energy of alloys C64 and
• FIG. 5 is a graph depicting the hardness profile achieved for the carburized sample of alloy C64 and the alloy Al of patent 6,464,801, in solid and open circles, respectively.
• the subject matter of the invention comprises a secondary-hardening carburized gear steel with surface hardness of HRC of at least about 62-64 and core fracture toughness greater than about 50 ksiVin.
• the interactions among the desired hierarchical microstructure, the processing and the property objectives are represented by the systems design chart in Figure 1.
• An ultimate goal of this invention was to optimize the whole system by controlling each subsystem and provide the most useful combination of surface hardness, core fracture toughness, and temperature resistance.
• Failure modes in gears are generally grouped into three categories: bending fatigue, contact fatigue, and temperature-induced scoring. Bending fatigue as well as contact fatigue can be limited by a high surface-hardness.
• the steel of the invention employs efficient secondary hardening by coherent M 2 C carbides which precipitate during tempering.
• the high Co content of the steel retards dislocation recovery and reduces the density of dislocations in response to thermal exposure.
• M 2 C carbides precipitate coherently on these dislocations during tempering and provide a strong secondary hardening response, enabling a surface hardness of 62-64 HRC.
• the steel alloy of the invention also limits temperature-induced scoring. Subsurface scoring results if the alloy's contact fatigue strength drops below the applied stress at any point below the surface. To provide adequate fatigue strength and avoid subsurface scoring, typically at least about a 1 mm-deep hardened case is preferred. The steel of the invention achieves this desirable case depth via a carbon content gradient achieved during carburization.
• the steel comprises a predominantly lath martensitic matrix free of TCP-phases, and is strengthened by a fine-scale distribution of M 2 C carbides.
• the martensite-start temperature (M s ) In order to produce a predominantly lath martensitic matrix, the martensite-start temperature (M s ) must be higher than about 100 0 C at the carburized surface.
• the invention has a carefully optimized the Ni content. While Ni is desirable for cleavage resistance, it also stabilizes austenite and thus, depresses M s .
• the Ni content is chosen to place the ductile-to-brittle transition of the steel sufficiently below room temperature, preferably below -20 0 C, while maintaining a sufficiently high M s .
• the Ductile-to-Brittle Transition Temperature (DBTT) of the steel can be characterized by measuring the CVN impact energy at varying temperatures. As shown in Figure 3, while earlier prototype alloy C69B shows susceptibility to cleavage up to 150 0 C, the optimized composition of alloy C64 of the invention successfully depresses the DBTT to about -20 0 C.
• the average grain diameter must be less than about 50 ⁇ .
• the steel employs a grain- pinning dispersion of MC particles, where M may be Ti, Nb, Zr, or V, with Ti preferred.
• M may be Ti, Nb, Zr, or V
• the particle size of the grain-pinning dispersion should be refined.
• a refined size of the MC particles is achieved by designing a system wherein the particles dissolve during homogenization and subsequently precipitate during forging. The MC particles remain stable during subsequent carburization and solution heat treatment cycles.
• the resulting lath martensitic matrix is free of undesirable TCP-phases.
• TCP-phase precipitation is to be avoided during tempering because such phases can reduce the alloy ductility and toughness.
• the thermodynamic driving force for precipitation of TCP phases is limited in the invention by the contents of Cr, Mo, and W.
• IW-0.02Ti 0. IW-0.02Ti (wt%) was prepared from high purity materials. The melt was converted to a 1.5- inch-square bar. The optimum processing condition was to solutionize at 1050° 90 minutes, quench with oil, immerse in liquid nitrogen for 1 hour, warm in air to room temperature, temper at 468°C for 56 hours, and cool in air. The DBTT in this condition was between 150 0 C and 250 0 C.
• a 300-lb vacuum induction melt of Fe-17.0Co-7.0Ni-3.5Cr-l.5Mo-0.2W-0.12C (wt%) was prepared from high purity materials. Because this prototype did not include Ti, the grain- pinning dispersion of TiC particles could not form. As a result, the average grain diameter was 83 ⁇ and toughness was very low.
• the CVN impact energy of the core material from this prototype was 5 ft- Ib at an Ultimate Tensile Strength (UTS) of 238 ksi.
• 0.12C-0.03Ti (wt%) was prepared from high purity materials. This composition did include Ti, and the average grain diameter was 35 ⁇ . Toughness improved substantially.
• the CVN impact energy of the core material from this prototype was 23 ft-lb at a UTS of 238 ksi.
• the corresponding processing condition was to carburize and concurrently solutionize the steel at 927°C for 8 hours, quench with oil, immerse in liquid nitrogen for 1 hour, temper at 496°C for 8 hours, and cool in air.
• the fracture toughness in this condition was 100 ksiVin.
• the DBTT in this condition was around room temperature.
• 0.03Ti-0.02V (wt%) was prepared from high purity materials.
• Half of the melt was converted to a 6.5-inch-diameter barstock, while the other half was converted to a 4.5-inch-diameter barstock.
• the optimum processing condition was to carburize the steel at 927°C for 3 hours, cool in air, solutionize at 1000 0 C for 40 minutes, quench with oil, immerse in liquid nitrogen for 2 hours, warm in air to room temperature, temper at 496°C for 8 hours, and cool in air.
• the average grain diameter in this condition was 27 ⁇ and the fracture toughness was 85 ksiVin at a UTS of 228 ksi.

Landscapes

• Chemical & Material Sciences (AREA)
• Engineering & Computer Science (AREA)
• Materials Engineering (AREA)
• Mechanical Engineering (AREA)
• Metallurgy (AREA)
• Organic Chemistry (AREA)
• Physics & Mathematics (AREA)
• Thermal Sciences (AREA)
• Crystallography & Structural Chemistry (AREA)
• Solid-Phase Diffusion Into Metallic Material Surfaces (AREA)
• Heat Treatment Of Articles (AREA)
• Heat Treatment Of Steel (AREA)
• Gears, Cams (AREA)

Priority Applications (5)

Applications Claiming Priority (4)

Publications (2)

Family

ID=40547343

Family Applications (1)

Country Status (7)

Families Citing this family (31)

Citations (3)

Family Cites Families (10)

• 2008
  • 2008-08-20 US US12/194,964 patent/US8801872B2/en active Active
  • 2008-08-22 EP EP08843244.8A patent/EP2181199B1/en active Active
  • 2008-08-22 WO PCT/US2008/073966 patent/WO2009055133A2/en not_active Ceased
  • 2008-08-22 BR BRPI0815648A patent/BRPI0815648B1/pt not_active IP Right Cessation
  • 2008-08-22 CN CN2008801045341A patent/CN101784681B/zh not_active Expired - Fee Related
  • 2008-08-22 JP JP2010522059A patent/JP5588869B2/ja active Active
  • 2008-08-22 CA CA2695472A patent/CA2695472C/en not_active Expired - Fee Related

Patent Citations (3)

Also Published As

Similar Documents

Legal Events