US3705827A - Nickel-iron base alloys and heat treatment therefor - Google Patents

Nickel-iron base alloys and heat treatment therefor Download PDF

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US3705827A
US3705827A US142635A US3705827DA US3705827A US 3705827 A US3705827 A US 3705827A US 142635 A US142635 A US 142635A US 3705827D A US3705827D A US 3705827DA US 3705827 A US3705827 A US 3705827A
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temperature
nickel
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alloy
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Donald R Muzyka
Donald K Schlosser
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Carpenter Technology Corp
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22F—CHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/10—Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of nickel or cobalt or alloys based thereon

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  • This invention relates to nickel-iron base alloys with or Without chromium and cobalt in which columbium, ttanium and aluminum take part in the main precipitation hardening and strengthening reaction which is controlled by forming and heat treating so as to provide a unique combination of room temperature and elevated temperature properties.
  • a stabilization aging treatment has been used at a temperature intermediate the solution treating temperature and the precipitation aging temperature.
  • Nig (Cb, Ti, Al) the precise composition of which may vary, depending upon process and composition variables.
  • the form of the precipitate is diticult to determine, but may be gamma prime which is a face centered cubic (FCC) structure, gamma double prime which is a body centered tetragonal (BCT) structure or a combination of the two.
  • FCC face centered cubic
  • BCT body centered tetragonal
  • a further important object of the present invention is to provide a nickel-iron base alloy which responds to the heat treatment of the present invention and which has improved elevated temperature tensile and stress rupture properties.
  • FIGS. 1A and 2A are light micrographs having a magnilication of 500x and FIGS. 1B and 2B are electron micrographs having a magnification of 7700 all prepared from specimens formed and heat treated in accordance with the present invention;
  • FIGS. '3A/4A and FIGS. 3Bf4B are respectively corresponding light and electron micrographs prepared from specimens of the same composition as those shown in FIGS. l and 2, and formed and treated in the same way except that they were solution treated at a higher temperature outside the scope of the present invention.
  • the process of the present invention is applicable to a narrowly defined class of alloys in which the predominant elements are nickel and iron that coact to provide an austenitic or gamma face centered cubic microstructure.
  • nickel can be present from about 30% to and at least about 30% iron should be present.
  • the strengthening elements colnmbium, titanium and aluminum react with some of the nickel to form one or more strengthening phases brought out as an intragranular precipitate by age or precipitation hardening.
  • the composition of those phases is generalized as Ni3(Cb.Ti,Al) and may be gamma prime which has a face centered cubic structure, gamma double prime which has a body centered tetragonal structure, or both phases may be present.
  • the composition of these alloys is such that the Ni3(Cb,Ti,Al) intragranular precipitate has a solvus temperature which is considerably below that of the same r similar phase found in commercial nickel base super alloys now in use containing little or no iron. Its solvus temperature is also substantially lower than that of additional precipitated phases formed of the elements Ni, Cb, and Ti designated as eta phase for NiSTi and delta phase for NigCib.
  • Ni3'1 ⁇ i in the form of eta phase has a close packed hexagonal crystal structure and is thus distinguishable from the face centered cubic structure typical of the gamma prime form of NiaTi by means of well ikuown X- ray diffraction techniques.
  • the orthorhombic delta phase formed by NiCb as a grain boundary precipitate can be distinguished from the intragranular body centered tetragonal gamma double prime phase of NiaCb brought out by aging. Hitherto, eta phase and/or delta phase not only appeared along the grain boundaries where they could be beneficial, but also within the grain structure where their presence was not desired. As will be more fully pointed out hereinafter, eta phase and/or delta phase are brought out along the grain boundaries so as to provide improved properties, and no effective amount appears within the grains to detract from the over-al1 properties.
  • the only other essential elements in the compositions that respond to the present process are the elements columbium, titanium and a small but necessary amount of aluminum.
  • aluminum is not one of the main strengthening elements, some small amount is necessary, and the minimum amount required should ⁇ be increased as the columbiurn content tends toward the lower end of the range stated.
  • the minimum amount of aluminum preferably should not be less than about L' 0.l5% and better yet not less than about 0.20%.
  • boron can be beneficially used as for example in compositions t containing little or no chromium while when significant amounts of chromium are present, then preferably boron is limited to no more than about 0.01% and better yet to no more than about 0.006%.
  • Optional elements that can be present include up to about chromium to impart stainless properties, up to about 20% cobalt which may be used in addition to or in place of some of the nickel particularly in controlled expansion-type alloys.
  • up to about 1% vanadium may be used to beneficially affect hot workability of.
  • up to about 0.1% zirconium may be included for its beneficial effect on ductility
  • up lo 2% hafnium may be included as a solid solution strengthener and as a carbide former.
  • solid solution strengtheners as up to about 3% molybdenum and up to about 3% tungsten may also be present.
  • carbon up to about 0.1%, up to about 1%, preferably no more than 0.50% manganese, and up to about 0.50% silicon may be present in the nickel-iron base alloys with which the present invention is concerned.
  • Phosphorus and sulfur are not desirable additions, and preferably each does not exceed about 0.020%.
  • Such alloys are readily melted and cast as ingots using conventional techniques: however, for best results, a multiple melting practice is preferred. For example, a heat can be first melted and cast as an ingot under vacuum in an induction furnace, and that ingot is then used as a consumable electrode and remelted under vacuum.
  • hot working is preferably carried out so as to provide a fine grain structure of at least A.S.T.M. 4 or finer.
  • forging from a furnace temperature of about 2000 F. to 2100 F. to at least about reduction in cross-sectional area is adequate, but reductions of as much as to 90% provide better results.
  • the nal forging operations be carried out at least in part within about F., above or below, of the effective solvus temperature of the eta and delta phases of the composition. This ensures the desired grain structure of no coarser than A.S.T.M. 4.
  • the starting grain structure for heat treatment which has been found to give best results is at least as fine as A.S.T.M. 8 or finer.
  • the solution treatment of the present process serves several functions including the usual one of putting back into solution the intragranular strengthening gamma prime and double prime phases brought out during hot working. This is to avoid banding or other non-uniform distributions of the phases which, as a practical matter, usually cannot be avoided during hot working.
  • Another and important function of the solution treatment of the present invention is the formation of the eta phase and/ or delta phase precipitate along the grain boundaries.
  • the solution treatment is carried out at a temperature and for time long enough to provide the amount and also the distribution of the eta and delta phases which favors the attainment of elevated temperature tensile and stress rupture ductility.
  • the solution-treating temperature is readily determined empirically using the following guidelines. Hot worked specimens are heated at about 25 F. increments from about 1400 F. up to determine the solvus temperature of the intragranular gamma prime and double prime precipitates. Further testing is carried out in the 25 F. increments above the solvus temperature of the gamma prime and double prime precipitates to the temperature at which the eta phase and delta phase effective solvus temperature is found. That is the temperature at which enough of the eta and delta phases have been put back into solution so that what remains is no longer effective to prevent grain growth and other objectionable effects.
  • the effective solvus temperature is somewhat below the equilibrium solvus temperature, the latter being the temperature at which those phases are entirely taken into solution.
  • the effective solvus temperature is readily identified by examining the microstructure of the solution treated specimens because of the grain growth from the as-forged condition which occurs as soon as the solutiontreating temperature used is above the effective solvus temperature.
  • the best solution-treating temperature for use can be readily veried by aging tensile and stress rupture specimens in the usual way, that is, below the gamma prime/gamma double prime solvus and observing the effect of the various solution-treating temperatures on the elevated temperature tensile and stress rupture properties.
  • the term hot working it is not intended to exclude warm working, which includes working the metal while it is below its recrystallization temperature, or other thermo-mechanical procedures.
  • the effective solvus temperature of the eta and delta phases is the recrystallization temperature of the alloy for most practical purposes but in some instances following extreme warm working the recrystallization temperature can be lower.
  • EXAMPLE 1 As an example of a preferred embodiment of the present invention, an experimental vacuum induction heat was prepared of a hitherto known composition falling within an intermediate range of up to about 0.06%, preferably about 0.01% to 0.05% carbon, up to about 0.35% manganese, up to about 0.35% silicon, no more than 0.020% phosphorus or sulfur, about 114.5% to 17.5% chromium, about 39% to 44% nickel, up to about 1% molybdenum, up to about 1% cobalt, about 2.5% to 3.3% columbium, about 1.5% to 2% titanium, about 0.15% to 0.40% aluminum, about 0.001% to 0.01%,
  • 1800 F.A indicates solution treatmerit for 1 hour at 1800 F. followed by a stabilization treatment at 1550l F. for 3 hours before aging.
  • the 1800 F.-B solution treatment plus stabilization treatment was the same except that the stabilization treatment was carried out at ⁇ 1650 F. for l hour.
  • 1800 F.-C indicates the same solution treatment followed by a stabilization treatment at 1650 F. for 4 hours
  • 1800 F.-D indicates that a stabilization treatment at 1650 F. for 8 hours was used. The effect of different aging treatments is also indicated.
  • Example 1 had the following was then heated for 8 hours at one of 4 primary aging analysis in weight percent: temperatures ranging from 1325 F. to 1400 F. in 25 F. increments. This was followed by cooling at the rate C b 0.027 Manglnesg 008 of 100 F./hr. to a final aging temperature of either Silicon 0:10 1150c F. or 1200 F.
  • the 2% inch square ingot was homogenized and then forged from a temperature of 2000o F. to 2 in. sq., reheated to 2000 F. and forged to 1% in. sq., reheated to 2000 F. for 1 hour and forged to 11/2 in. sq., reheated to 2000 F. for 1 hour and then forged to 7/8-in. sq.
  • the as-forged grain structure was A.S.T.M. 9-10 Blanks for forming test specimens were cut from the thus forged bar stock which were heat treated and then machined and tested.
  • the combination smooth/notch stress-rupture specimens utiiized were standard A.S.T.M. specimens having a 0.178-in. diameter, a 0.712-in. long smooth gage section and a notch stress concentration factor (Kt) of 3.8. Stress rupture tests at 1200 F. under a load of 100,000 p.s.i. (100 k.s.i.) were carried out, and the stress rupture life in hours (R.L., hrs.) the percent elongation El.) and the percent reduction in area (percent R.A.) are indicated in Table I below.
  • test specimens were subjected to various heat treatments including solution treatment for 1 hour at temperatures ranging in 25 F. increments from 1600 F. to 1800 F. In each instance, solution treatment was for 1 hour followed by cooling in air.
  • FIGS. 1A, 2A, 3A and 4A and the corresponding electron micrographs, 7700 magnification were prepared from the tested stress rupture specimens of Test No. 13, No. 7, No. 4 and No. 2 respectively, with the area shown extending longitudinally with respect to the test specimen axes. Inasmuch as the primary and final aging temperatures for these four tests were the same, the results can be directly compared to show the critical effect of variations of the solution-treating temperature from l650 F.
  • a further important feature of the present invention resides in the fact that it makes possible the provision of controlled expansion-type alloys characterized by high strength and ductility for use at high temperatures.
  • a further intermediate range to which the process of the present invention is applicable consists essentially in weight percent of about Broad Preferred Carbon 10.1 0. 01-005 l 0.50 l 0. 20 l 0.50 l 0. Z) l 0.020 l 0. 20 l 0. 020 l 0.0m l0. l 0.5 l 0.5 1 0.5 -40 315-39 13-17 14. 516. 5 2. 5-6 2. 75-3. 2 1-3 1. 65-1. 85 0. 1-2 0. 85-1. 15 l 0. 030 0. 0015-0. 020
  • the balance of the composition is iron and incidental impurities which are preferably kept low which is facilitated by using a multiple vacuum melting practice such as that previously described hereinabove. While the broad range for columbium is indicated as about 2.5% to 6%, the better practice is to limit columbium to no more than about 3.5%. In the case of aluminum, the larger amounts tend to provide better properties, and thus an intermediate broad range of 0.5% to 1.5% is preferred. While the absence of boron can be tolerated, when elevated temperature stress rupture ductility is not desired, a small but definite amount of at least about 0.003% boron is required to obtain the outstanding stress rupture properties of the present invention. For best results 0.005% to 0.015% is used.
  • the composition is balanced within the foregoing intermediate range so as to satisfy the following two equations in which percents are also by weight:
  • the composition is balanced within the preferred range so as to provide a value of alphal of 4.()X10-6 to 4.5X10-/ F. and a Curie temperature range of 760 F. to 860 F.
  • EXAMPLE 2 As an example of the preferred controlled expansion alloy of the present invention, an experimental vacuum induction heat having the following composition in weight percent was prepared:
  • the ingot was forged to bar stock which was then used for making test blanks which were heat treated, machined and tested.
  • 1200 F. tensile specimens having a 0.252-in. diameter and a 1.0-in. long gage section were also prepared. Solution treatment of all of the specimens was for 1 hour at the temperature indicated followed by aging at 1325D F. for 8 hours, then cooling at the rate of F. per hour to 1l50 F., holding that temperature for 8 hours and then cooling in air.
  • the effect of solution-treating temperatures from l550 F. to 1700 F. on the stress rupture properties of the alloy of Example 2 measured at 1150 F. under a load of 110,000 p.s.i. are set forth in the following table.
  • optimum elevated temperature properties are provided by a solution-treating temperature of about 1575 F. to 1625 F., the effect being mainly on the .2% yield strength (.2% Y.S.) rather than on the ultimate tensile strength (U.T.S.).
  • Example 2 For comparison purposes, a heat was prepared as described in connection with Example 2 having an equiva lent analysis except that the boron content was 0.0022% and containing 0.031% carbon, 37.73% nickel, 16.19% cobalt, 3.02% columbium, 1.74% titanium, 1.00% aluminum, and the balance iron except for inconsequential impurities.
  • Two stress rupture specimens were prepared which were solution treated at 1625 F. for 1 hour followed by primary aging at 1325 F. and final aging at 1150 F. as was described in connection with Example 2. When subjected to a stress of 110 k.s.i. at 1150 F., both specimens failed at the notch one after only 1.2 hours and the other after only 2.1 hours.
  • EXAMPLE 3 As a further example of a controlled expansion alloy and its heat treatment in accordance with the present invention, a vacuum induction heat was prepared as described in connection with Example 2 having the following composition in weight percent:
  • the alloy of Example 3 corresponds to that of said Pat. No. 3,157,495.
  • the unique stress rupture ductility provided by the present invention is clearly brought out by the data in Table IV. It may be noted that the difference in the aluminum content of about 0.5% between the compositions of Example 2 and Example 3 is believed to account for the difference in optimum solution-treating temperatures between the two alloys, that of Example 2 being higher because of the greater amount of aluminum.
  • columbium when the element columbium is referred to, it is to be understood as including a certain amount of tantalum ranging from about 1% to 20% of the amount of columbium. That amount of tantalum usually is present in commercial supplies of columbium customarily used for alloying purposes. Further, additional amounts of columbium can be replaced if desired by tantalum. Thus columbium is to be read as including tantalum or as the combined columbium plus tantalum content of the composition.
  • a nickel-iron base alloy which consists essentially by weight of about Carbon Up to 0.1%. Manganese Up to 1%. Silicon Up to 0.5%. Chromium Up to 20%. Molybdenum Up to 3%. Tungsten do Nickel 30% to 50% Cobalt Up to 20% Columbium 2.5% to 6%. Titanium 1% to 3%. Aluminum 0.1% to 2%. Vanadium Up to 1%. Zirconium Up to 0.1%. Hafnium Up to 2% Boron Up to 0.030%.
  • the iron being at least about 30% and which is strengthened by precipitation hardening to bring out gamma prime and gamma double prime phases, the steps of solution treating said alloy at a temperature which is above the effective solvus temperature of the gamma prime and double prime phases and below the effective solvus temperature of eta and delta phases so as to form a precipitate in the grain boundaries of said alloy made up of said eta and delta phases, and then aging said alloy at a temperature below the effective solvus temperature of said gamma prime and gamma double prime phases to form a tine dispersion thereof of said gamma prime and gamma double prime phases within the grains of said alloy.
  • said alloy being strengthened by precipitation hardening to bring out gamma prime and gamma double prime phases and which includes the steps of solution treating said alloy at a temperature which is above the effective solvus temperature of the gamma prime and double prime phases and below the effective solvus temperature of eta and delta phases so as to form a precipitate n the grain boundaries of said alloy made up of said eta and delta phases, and then aging said alloy at a temperature below the effective solvus temperature of said gamma prime and gamma double prime phases to form a line dispersion of said gamma prime and gamma double prime phases within the grains of said alloy.
  • said alloy being strengthened by precipitation hardening to bring out gamma prime and gamma double prime phases and which includes the steps of solution treating said alloy at a temperature which is about the effective solvus temperature of the gamma prime and double prime phases and below the effective solvus temperature of eta and delta phases so as to form a precipitate in the grain boundaries of said alloy made up of said eta and delta phases, and then aging said alloy at a temperature below the effective solvus temperature of said gamma prime and gamma double prime phases to form a fine dispersion of said gamma prime and gamma double prime phases within the grains of said alloy.
  • said alloy consists essentially of about 0.01% to 0.05% carbon, about 0.20% max. manganese, about 0.20% max. silicon, about 0.5% max. chromium, about 0.5% max. molybdenum, about 36% to 39% nickel, about 14.5% to 16.5% cobalt, about 2.75% to 3.2% Columbium, about 1.65% to 1.85% titanium, about 0.85% to 1.15% aluminum, about 0.005% to 0.020% boron, and the balance iron, and includes forging said alloy before said solution treatment, said forging at least in part being carried out within about 100 F. of the effective solvus temperature of the eta. and delta phases of said alloy so as to establish a grain size which is substantially no coarser than about A.S.T.M. 8.

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US4006011A (en) * 1972-09-27 1977-02-01 Carpenter Technology Corporation Controlled expansion alloy
US4066447A (en) * 1976-07-08 1978-01-03 Huntington Alloys, Inc. Low expansion superalloy
US4165997A (en) * 1977-03-24 1979-08-28 Huntington Alloys, Inc. Intermediate temperature service alloy
US4190437A (en) * 1977-12-08 1980-02-26 Special Metals Corporation Low thermal expansion nickel-iron base alloy
US4200459A (en) * 1977-12-14 1980-04-29 Huntington Alloys, Inc. Heat resistant low expansion alloy
US4225364A (en) * 1978-06-22 1980-09-30 The United States Of America As Represented By The United States Department Of Energy High strength nickel-chromium-iron austenitic alloy
US4231795A (en) * 1978-06-22 1980-11-04 The United States Of America As Represented By The United States Department Of Energy High weldability nickel-base superalloy
US4236943A (en) * 1978-06-22 1980-12-02 The United States Of America As Represented By The United States Department Of Energy Precipitation hardenable iron-nickel-chromium alloy having good swelling resistance and low neutron absorbence
US4255186A (en) * 1978-01-19 1981-03-10 Creusot-Loire Iron-containing alloys resistant to seawater corrosion
US4487743A (en) * 1982-08-20 1984-12-11 Huntington Alloys, Inc. Controlled expansion alloy
US4578130A (en) * 1979-07-27 1986-03-25 The United States Of America As Represented By The United States Department Of Energy Iron-nickel-chromium alloy having improved swelling resistance and low neutron absorbence
US4685978A (en) * 1982-08-20 1987-08-11 Huntington Alloys Inc. Heat treatments of controlled expansion alloy
US4711826A (en) * 1986-01-27 1987-12-08 Olin Corporation Iron-nickel alloys having improved glass sealing properties
US4755240A (en) * 1986-05-12 1988-07-05 Exxon Production Research Company Nickel base precipitation hardened alloys having improved resistance stress corrosion cracking
US4816216A (en) * 1985-11-29 1989-03-28 Olin Corporation Interdiffusion resistant Fe--Ni alloys having improved glass sealing
US4872924A (en) * 1986-09-12 1989-10-10 Hitachi, Ltd. Method of producing shadow mask of color cathode ray tube
US4905074A (en) * 1985-11-29 1990-02-27 Olin Corporation Interdiffusion resistant Fe-Ni alloys having improved glass sealing property
US4908069A (en) * 1987-10-19 1990-03-13 Sps Technologies, Inc. Alloys containing gamma prime phase and process for forming same
US5059257A (en) * 1989-06-09 1991-10-22 Carpenter Technology Corporation Heat treatment of precipitation hardenable nickel and nickel-iron alloys
US5169463A (en) * 1987-10-19 1992-12-08 Sps Technologies, Inc. Alloys containing gamma prime phase and particles and process for forming same
US5439640A (en) * 1993-09-03 1995-08-08 Inco Alloys International, Inc. Controlled thermal expansion superalloy
US5534085A (en) * 1994-04-26 1996-07-09 United Technologies Corporation Low temperature forging process for Fe-Ni-Co low expansion alloys and product thereof
DE19542920A1 (de) * 1995-11-17 1997-05-22 Asea Brown Boveri Eisen-Nickel-Superlegierung vom Typ IN 706
US5688471A (en) * 1995-08-25 1997-11-18 Inco Alloys International, Inc. High strength low thermal expansion alloy
US20040156737A1 (en) * 2003-02-06 2004-08-12 Rakowski James M. Austenitic stainless steels including molybdenum
US20060081315A1 (en) * 2001-09-18 2006-04-20 Honda Giken Kogyo Kabushiki Kaisha Method for producing Ni based alloy and forging die
US20070029017A1 (en) * 2003-10-06 2007-02-08 Ati Properties, Inc Nickel-base alloys and methods of heat treating nickel-base alloys
US20070044875A1 (en) * 2005-08-24 2007-03-01 Ati Properties, Inc. Nickel alloy and method of direct aging heat treatment
US20080163963A1 (en) * 2007-01-08 2008-07-10 Ling Yang Heat Treatment Method and Components Treated According to the Method
US20080213099A1 (en) * 2006-08-25 2008-09-04 Shinya Imano Ni-Fe BASED FORGING SUPERALLOY EXCELLENT IN HIGH-TEMPERATURE STRENGTH AND HIGH-TEMPERATURE DUCTILITY, METHOD OF MANUFACTURING THE SAME, AND STEAM TURBINE ROTOR
US20080257457A1 (en) * 2007-04-19 2008-10-23 Ati Properties, Inc. Nickel-base alloys and articles made therefrom
US20100276041A1 (en) * 2007-01-08 2010-11-04 Ling Yang Heat Treatment Method and Components Treated According to the Method
US20130071633A1 (en) * 2010-05-28 2013-03-21 Gosakan Aravamudan Short Fiber Reinforced Artificial Stone Laminate
US20130213531A1 (en) * 2008-04-28 2013-08-22 Canon Kabushiki Kaisha Method for producing alloy
US10184166B2 (en) 2016-06-30 2019-01-22 General Electric Company Methods for preparing superalloy articles and related articles
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US3705827A (en) 1971-05-12 1972-12-12 Carpenter Technology Corp Nickel-iron base alloys and heat treatment therefor
US4225363A (en) * 1978-06-22 1980-09-30 The United States Of America As Represented By The United States Department Of Energy Method for heat treating iron-nickel-chromium alloy
GB2054647B (en) * 1979-07-27 1983-10-26 Westinghouse Electric Corp Iron-nickel-chromium alloys
US4445943A (en) * 1981-09-17 1984-05-01 Huntington Alloys, Inc. Heat treatments of low expansion alloys
US4445944A (en) * 1981-09-17 1984-05-01 Huntington Alloys, Inc. Heat treatments of low expansion alloys
US4713576A (en) * 1985-04-24 1987-12-15 Hitachi, Ltd. Color picture tube with shadow mask
US4888253A (en) * 1985-12-30 1989-12-19 United Technologies Corporation High strength cast+HIP nickel base superalloy
EP0433072B1 (en) * 1989-12-15 1994-11-09 Inco Alloys International, Inc. Oxidation resistant low expansion superalloys
EP0588657B1 (en) * 1992-09-18 1998-04-15 Inco Alloys International, Inc. Controlled thermal expansion superalloy
DE69615977T3 (de) * 1995-08-25 2010-05-06 Inco Alloys International, Inc., Huntington Hochfeste Legierung mit niedrigem Ausdehnungskoeffizient
JP6337514B2 (ja) * 2013-05-21 2018-06-06 大同特殊鋼株式会社 析出硬化型Ｆｅ−Ｎｉ合金及びその製造方法
AT518456B1 (de) * 2016-04-14 2017-12-15 Lkr Leichtmetallkompetenzzentrum Ranshofen Gmbh Verfahren zur Herstellung eines Objektes für optische Anwendungen aus einer Aluminiumbasislegierung und entsprechend hergestelltes Objekt

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US3705827A (en) 1971-05-12 1972-12-12 Carpenter Technology Corp Nickel-iron base alloys and heat treatment therefor

Cited By (50)

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Publication number Priority date Publication date Assignee Title
US4006011A (en) * 1972-09-27 1977-02-01 Carpenter Technology Corporation Controlled expansion alloy
US4066447A (en) * 1976-07-08 1978-01-03 Huntington Alloys, Inc. Low expansion superalloy
US4144102A (en) * 1976-07-08 1979-03-13 The International Nickel Company, Inc. Production of low expansion superalloy products
US4165997A (en) * 1977-03-24 1979-08-28 Huntington Alloys, Inc. Intermediate temperature service alloy
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FR2139424A5 (ja) 1973-01-05
GB1372605A (en) 1974-10-30
CA969842A (en) 1975-06-24
DE2223114C3 (de) 1978-12-14
JPS5243763B1 (ja) 1977-11-01
DE2223114B2 (de) 1978-04-20
DE2264997A1 (de) 1976-03-04
DE2264997C2 (de) 1983-10-20
GB1372606A (en) 1974-10-30
DE2223114A1 (de) 1972-11-23

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