US3070439A - Method for processing dispersion strengthened metals - Google Patents

Description

3,979,439 Patentecl Dec. 25, 1962 3,076,439 METHQD FOR PROCESSING DISPERSION STRENGTHENED METALS Nicholas J. Grant, Winchester, and Klaus M. Zwilsky,

Watertown, Mass, assigners to New England Materiais Laboratory, Inc, Medford, Mass, a corporation of Massachusetts No Drawing. Filed Mar. 15, 1960, Ser. No. 15,036 11 Claims. (Cl. 75206) This invention relates to a method of processing dispersion strengthened or hardened metals and, in particular to a method of hot working said metals while inhibiting the dissipation of stored energy arising from the strain deformation of such metals.

Dispersion strengthening of alloys consists of dispersing a finely divided non-metallic phase throughout a metal matrix and then storing energy of deformation in the structure. Several techniques have been employed to obtain the desired dispersion and among these, mechanical mixing and internal oxidation have thus far shown the most promise. Mechanical mixing is the fastest and least expensive way to obtain dispersion strengthened materials. Internal oxidation, by comparison, is a more expensive and time consuming process to operate on a larger scale and involves the use of fine alloy powders not readily available. tory oxide-forming metal, e.g. aluminum, is alloyed with a matrix metal, such as copper, and the alloy in particulate form subjected to an oxidation treatment to oxidize selectively aluminum to a dispersion of A1 followed by consolidation of the particles to a wrought shape. In terms of interparticle spacing, the latter technique is usually capable of providing wrought shapes exhibiting markedly improved results for smaller amounts of the non-metallic phase than wrought shapes produced from mechanically mixed powder.

From an economical viewpoint, mechanical mixing of metal and refractory oxide is preferred for the processing of common metals such as iron-base alloys.

Investigations of certain metal and refractory oxide systems have indicated that the size of the metal powder, the size of the oxide powder and the volume percent oxide are all important variables in achieving optimum properties.

With regard to dispersion strengthened alloys, stress rupture data have shown that such alloys are generally characterized by extreme flatness of slope on a log stress versus log rupture life plot. Because of this characteristic, the superiority of such alloys over conventional materials becomes markedly effective as time and/or temperature are increased. This is due to the fact that conventional materials, such as age hardenable alloys, are subject to over aging or solution of a second phase as affected by time and temperature, whereby a rapid drop-off in properties occurs. On the other hand, oxide dispersion strengthened materials have been known to be generally stable up to below the melting point of the base or matrix material.

Strengthening in these alloys has been postulated to arise from two sources. One is the effect of the dispersion, which imparts strength to the matrix metal even in the annealed condition. The other efiFect arises from the introduction of the strain energy during the working process, tag. the extrusion process, with the consequent retention of this energy by the alloy by virtue of the presence of the dispersed phase.

We have found, however, that not all matrix metals behave the same way and that in some instances, unless due care is taken during strain deformation in producing a wrought shape from the dispersion strengthened alloy, at least some of the stored energy is dissipated. We

In this process, a refracfound this to be particularly true for iron and alloys. v

It is an object of this invention to provide a method of deforming dispersion strengthened metals whereby stored energy is imparted thereto with the minimum of energy dissipation while enhancing the strength properties of said metal.

Another object is to provide a method of processing dispersion strengthened iron and iron alloys with the aim of producingmaterial of enhanced strength properties, such as improved resistance to creep.

It is a further object to provide a method for processing a dispersion strengthened metal wherein the matrix metal is characterized by a phase transformation at an elevated temperature, and wherein the deforming temperature employed in processing said metal is determined according to the phase transformation of said metal.

An additional object is to provide a method of processing a matrix metal, such as iron or an iron-base alloy, having dispersed therethrough fine particles of a refractory oxide wherein the processing temperature is determined according to the temperature of phase transformation of the matrix metal and the temperature of crystallographic transformation of the refractory oxide.

These and other objects will more clearly appear from the following disclosure.

In producing a dispersion strengthened metal, such as iron, finely divided iron is mixed with up to about 15 W0 (volume percent) of a refractory oxide, such as A1 0 the particle size of the iron being preferably below 20 microns and generally below 10 microns. We prefer for optimum results to use particle sizes below 5 microns. We find it important that the refractory oxide have a particle size less than that of the matrix metal, for example from about 30 to 250 times smaller, preferably from about 30 to times smaller, so as to obtain adequate distribution of the oxide and control the interparticle spacing between the oxide particles. The metal powder is mixed with the refractory oxide preferably in the dry state by means of a high speed dry blender, or a ball mill, for a suitable time to effect uniform mixing, e.g. ranging from about 15 to 60 minutes in the blendor and 1 to 24 hours for the ball mill. Upon completion of the mixing, the mixture is subjected to a reduction treatment with hydrogen to reduce any iron oxide present at an adequate reducing temperature, e.g. from 700 to 900 F., and then hydrostatically pressed to the desired iron-base shape, followed by sintering in hydrogen to form a compact capable of being handled. We find that the sintering temperature should preferably not exceed the transformation temperature of the matrix metal. Thus, in the case of iron, we prefer the sintering temperature be maintained in the alpha iron region, for example at 1525 F. This temperature is also important where gamma alumina is used as the dispersoid as it is below the gamma to alpha transformation temperature. Thus, severe agglomeration of the powders is substantially avoided. The sintered compact is thereafter vacuum packed in a can and the whole extruded at an elevated temperature at an extrusion ratio sufficient to achieve maximum density, e-.g. ranging from about 10 to 1 to 28 to l.

We have found that dispersion strengthened iron or iron-base alloys produced by the foregoing powder metallurgy or similar technique generally exhibit markedly improved strength properties, provided that the deformation employed in producing a dispersion strengthened metal shape of maximum density is carried out while the material is in the ferritic or alpha condition. In other words, we have found that, in producing a dispersion strengthened metal product from a matrix metal powder characterized by a phase transformation at an elevated temperature and having mixed therewith fine particles of a non-metallic dispersoid, such as alumina, the consolidation of the mixture should be conducted at an elevated temperature below the temperature at which phase transformation occurs inorder to insure optimum strength properties. Our tests have indicated that if the Working or. deforming is conducted above the phase transformation temperature of the matrix metal, the stored energy is dissipated as the worked material cools down through the, phase transformation range of the material and converts from one crystallographic structure to another; for example, as in iron from austenite to ferrite.

In order to obtain a better understanding of the invention, the following examples are given:

Iron powder of about 3 microns average particle size was mixed with varying amounts of up to about 10 v/o of alumina having an average particle size of about 0.027 micron, the ratio of particle size of iron to aluminum being a little more than .100 to l.

The alloys were prepared by dry mixing the iron matrix metal'powder and the aluminum oxide in lots of about 500 grams, the mixing being conducted in a high speed blendor, e.g. a Waring Blendor, at a speed of about 15,- 000' revolutions per minute; The mixing of each lot was carried out for about 5 minutes and then further mixed by spatulation on a sheet of clean paper for a few minutes, the procedure including the blendor and subsequent spatulation being repeated about 4 times.

The blended powder batches were thereafter subjected to a reducing treatment in dry hydrogen for a minimum of five hours at a temperature of about 800 F. to insure clean particle surfaces for subsequent consolidation of the mixture into wrought shapes. Each batch of the mixed powders was introduced into a rubber tube Supported within a perforated steel canister about two inches in diameter, one end of the rubber tube being rubber stoppered at the start. After the powder was introduced, a second rubber stopper having in communication therewith a hypodermic needle was inserted, a vacuum connection being made through the needle to remove the air from within powder mass. After completion of evacuation, the needle was removed and, the canister assembly subjected to hydrostatic pressure at about 30,000 p.s.i. to yield compacts about 1.4 inches in diameter and 3 inches long.

Using correspondingly larger amounts of powder, compacts have been prepared by the same technique up to 3 inches in diameter and 6 inches long. Depending upon the size of presses, billets of up to 20 inches diameter are envisaged.

The compacts produced as aforementioned were then subjected to sintering in dry hydrogen for a minimum of -hours at 1525 F. After that they were each canned by insertion in a mild steel can and welded vacuum tight followed .by extrusion at an elevated temperature. The extrusion ratio was about 16 to 1.

The following alloys were produced:

Table I V01. V01. Extru- Al loy No. percent percent 2 sion F81 A1203 te m? 1 3 micron iron powder. 2 0.027 micron A1203 powder.

Alloys? and 4 of the same identical composition were extruded at difierent temperatures, one at 1550 F. (in the ferritic range), the other at l9.00 F. (in the austenitic range). V

The alloys were then subjected to tensile tests at room worked No. 4.

Table II Y.S. Alloy No. Temp., (p.s.i.) Ultimate Percent F. (0.2% (p.s.i.) elong,

offset) R.T. 66, 230 8'), 800 20 1, 270 13, 700 19,100 35 .T. 68, 000 84, 500 15 1, 2l0 16, 200 21, 25 RIP. 63, 900 98,000 9 1, 230 21, 300 28, 400 14 1, 400 13, 900 18, 300 9 4 R.T. 69, 800 88, 400 8 1, 230 10, 800 16, 200 20 1, 400 6, 200 9, 000 16 5 R.T. 104, 300 113, 000 6 1, 230 26, 830 33, 300 10 Table HI Vol. 10') hr. 1200 F. Alloy No. percent rupture life, percent A1203 stress elong. (p.s.i.)

4 11,000 6 e 13, 800 s a g is, 000 2 8 6, 500 3 10 24, 000 2 For comparison purposes, an extrusion was made from the same iron powder free from the presence of A1 0 This material designated as F-l, exhibited, in the asextruded condition, a yield strength at room temperature of about 25,000 p.s.i. and at 1000 F. of about 6,600 p.s.i. while exhibiting a 100 hour rupture life (at 1200 F.) under a stress of 2,600 p.s.i. (12% elongation).

The results indicate that the properties of iron are markedly improved by the presence of A1 0 as a dispersion strengthener. The improvement is particularly noticeable with respect to strength properties at elevated temperatures. In this connection, reference is made to the 100 hour rupture properties which show an increase in rupture life stress over similarly prepared pure iron of about 4 to 9 times for alumina contents ranging from 4 to 10 v/o.

Our investigations have indicated that in order to achieve the markedly improved results, it is important that the fabrication of the material containing the dispersion strengthener be carried out below the phase transformation temperature of the matrix metal. This is illustrated by the results obtained for alloy Nos. 3 and 4, both of which contained 8 vol. percent A1 0 with the exception that No. 4 was extruded at 1900 F., i.e.

in the austenitic region of the matrix metal iron, while No. 3 was extruded at 1550 F., i.e. in the ferritic region of the matrix metal. Referring to Table II, it will be noted that No. 3 exhibited a yield strength at 1200 F. of 21,300 p.s.i. as compared to the value of 10,800 p.s.i. (about half) obtained for No. 4 which was subjected to austenitic extrusion. Similarly, at 1400 F. No. 3 exhibited a yield strength of 13,900 p.s.i. as against the much lower value of 6,200 p.s.i. for the austenitically V The same trendwas indicated with respect to the 100 hour rupture'life at. 1200" F. wherein the ferritically worked No. 3 alloy exhibited a stress of 18,000 p.s.i. as against the austenitically worked No. 4 which exhibited a stress of 6,500 p.s.i., almost one-third the former value. It is thus apparent that unless the working is conducted below the transformation temperature of the matrix metal, the elevated temperature properties of the wrought metal are adversely affected.

Tests in which MgO was used as the dispersion strengthener for ferritically worked alloys were alsocon- -changes at elevated temperatures.

ducted. Using the same particle size iron powder (3 microns) and MgO with its particle size ranging from about 0.05 to 0.1 micron, room temperature yield strengths-of 84,500; 102,000 and 129,000 p.s.i. were obtained for iron alloys containing 4, 6 and 10 v/o, re-

spectively, ranging from over 3 to over 5 times the value of that obtained for pure iron. Likewise, yield strength values at 1000 F. were obtained ranging from about 22,800 to 30,600 p.s.i., as, against 6,600 p.s.i. for pure iron similarly prepared. As for 100 hour rupture life at working of the MgO-containing alloys was likewise es- ;sential in obtaining optimum properties.

Observations in other dispersion strengthening systems ihave indicated that in order to obtain optimum properties -consistently, consideration should be given to the tendency of the dispersoid to undergo crystallographic Microstructural examinations of some dispersion hardened alloys of other systems have indicated that when an alloy is worked at a temperature above the crystallographic phase transformation of a dispersoid, for example alumina, agglomera- -tion of the dispersoid has been observed to occur as the dispersoid transforms from one crystallographic phase to another, thus partially destroying the inherently fine dispersion of the dispersoid with a consequent falling off in strength properties. In the case of alumina employed This transformation starts in excess'of 1500 F. and is believed to be complete at about 1850 F. Therefore, we prefer,

where a dispersoid is employed which crystallographically transforms at an'elevated temperature, that the temperature of consolidation of the alloy be also maintained below this transformation temperature as well as below the phase transformation of the matrix metal.

While the present invention has been described with respect to its application to iron, it is also applicable to iron-base matrix alloys such as carbon steels, certain of the ferritic stainless steels subject to gamma or other transformation, certain of the iron-base alloy steels and the like. Examples of iron-base alloys would be those single phase binary alloys which have a second metal which does not oxidize readily such as binary alloys of Fe-Cr (up to 20% Cr), FeW (up to 4% W), FeMo (up to 4%), FeNi (up to 10% Ni), FeSi (up to Si), etc. It will be appreciated that some of these alloys may have an alpha or gamma phase stable up to very high temperature in which case they may not present any problems. However, these ranges may include small amount of other elements such as carbon, manganese, or other strong austenite formers sufiicient to effect transformation. Ternary and more complex alloys, such as FeCr-Mo, FeNiW and the like would be ineluded where they are subject to phase transformation at elevated temperatures. The term iron-base alloys is meant to include solid solution alloys containing at least 50% Fe and, preferably, at least 65% Fe.

Other matrix metals and alloys to which the invention is applicable are cobalt (transforms between 800 to 900 F.), and such cobalt alloys as Co-Mo, CoNi-W, CoCrW, CoCrMo, etc., reference being made to available phase diagrams in determining adequate fabrication temperatures free from phase transformations. Still others are titanium (a phase transformation at about 1615 F.), zirconium (a phase transformation at about 1470 F.), uranium (a phase transformation at about 1115 F.) as well as other metals and alloys.

For the purposes of this invention, it is preferred that the matrix metal have a melting point above 1800" F. Where such metals are employed as the matrix, as described hereinbefore the metal in particulate form is .rnixedwith a refractory oxide powder, shaped and converted into a form capable of being easily handled during deformation working, and the whole but worked at an elevated temperature below the phase transformation temperature of the matrix metal and also preferably beloW the crystallographic transformation temperature of the dispersed refractory oxide.

Examples of refractory oxides which may be used in producing dispersion hardened alloys in accordance with the invention are SiO A1 0 MgO, CaO, BaO, SrO, BeO, ZrO TiO Th0;; and oxides of the rare earth metal group, such as oxides of cerium, lanthanum, neodymium, etc. Such refractory oxides are characterized as having a melting point above the melting point of the matrix metal, for example above 2700 F. and generally above In addition, these oxides are characterized by a negative free energy of formation at about 25 C. of over 90,000 calories per gram atom of oxygen. For example SiO has .a negative free energy of formation at 25" C. of about 96,200, A1 0 of about 125,590, MgO of about 136,130, BeO of about 139,000, etc. Where such oxides exhibit crystallographic transformation at elevated temperatures, such as SiO A1 0 ZrO TiO and others, then we prefer, depending on the crystallographic phase present, that the working temperatures employed be maintained below the transformation temperature of the oxide phase to inhibit agglomeration thereof to larger sizes. For examples, where A1 0 is used as the dispersoid and it is in the gamma form, then the temperature of fabrication of the alloy should not exceed the crystallographic transformation temperature for gamma alumina which falls within range of about 1500 F. to 1920" F., otherwise agglomerations of the alumina particles occurs. However, where the alumina used is already in the stable alpha form, the temperature of fabrication makes no difference in so far as the alpha alumina is concerned as it is in the stable high temperature form. SiO exhibits several transformations, ZrO has a monoclinic form which is stable to about 1830 F.

When using oxides of the foregoing types as dispersoids, as stated hereinabove they may range from about 1 to 15 We and preferably from about 3 to 12 v/o. Likewise, as stated hereinbefore, the particle size of the oxide should be 30 to 250 times smaller than the matrix metal powder, preferably 30 to times smaller, and also preferably range in size from about 0.01 to 0.1 micron for particle size of the matrix metal powder ranging up to about 20 microns or over the range of 1 to 20 microns.

The invention enables the production of high strength Wrought metal products characterized by structural stability and. by the ability to retain high strength properties at elevated temperatures. Examples of metal structures are heat exchangers, turbine buckets of iron or iron alloys and other metals for use in steam and gas turbines, furnace structures where resistance to creep at elevated temperatures is an important consideration, boiler tubing for carrying super-heated steam and many other applications too numerous to mention.

Although the present invention has been described in conjunction with preferred embodiments, it is to be understood that modifications and variations may be re- 7 sorted without departing from the spirit and scope of the invention as those skilled in the art will readily understand. Such modifications and variations are considered to be within the purview and scope of the intention and the appended claims.

What is claimed is:

1. A method of producing a dispersion strengthened wrought metal product from a matrix metal selected from the group consisting of iron, cobalt, iron-base and cobaltbase metals of melting point above 1800 F. characterized by a phase transformation temperature at an elevated hot working temperature and from an insoluble refractory oxide dispersion hardener characterized by a negative free energy of formation at about 25 C. of at least about 90,000 calories per gram atom'of oxygen which comprises, providing a powder mixture comprising said matrix metal and a substantially uniform dispersion of said shape atan'elevated temperature'not exceeding the temperature at which the phase transformation of said matrix metal begins, whereby dissipation of stored energy is greatly inhibited.

2. The method of claim 1, wherein the amount of dispersed refractory oxide particles ranges from about 1 to 15 volume percent of'the volume of the metal product.

3. The method of claim 2, wherein the particle size of the oxide ranges from'about.0.01 to 0.1 micron.

4. The method of claim 1, wherein the compact is hot deformed by extrusion.

5. A method of producing'a dispersion'strengthened wrought metal product from an iron-base matrix metal of melting point above 1800 F. characterized by a ferritic to austenitic transformation temperature at an elevated hot working temperature and froman insoluble refractory oxide dispersionhardener characterized by a negative free energy of-formation at about 25 C. of at least about 90,000 calories per gramatom ofoxygen which comprises, providing a powder mixture comprising said matrix metal and a substantially uniform dispersion of said refractory oxide powder, forming a compact of said mixture,hot deforming saidcompact to" a wrought metal shape at an elevated temperature not "exceeding the temperatureat which the phase transformation of said matrix metal begins, whereby dissipation of stored'energy is greatly inhibited.

6. The method of claim 5, wherein the iron-base metal is comprised substantially of iron.

7. The method of claim 5, wherein the compact is hot deformed by extrusion.

8. A method of producing a dispersion strengthened wrought metal product from an iron-base matrix metal of melting point above 1800 F. comprised substantially of iron characterized by a ferritic to austenitic transformation temperature at anelevated hot working temperature and'frorn an insoluble refractory oxide dispersion hardener characterized by a negative free energy of formation at about 25 C. of at least about 90,000 calories per gram'atom of oxygen which comprises, providing-a powder mixture of said matrix metal of average particle size ranging up to about 20 microns with about 1 to 15 volume percent of a substantially uniform dispersion of said refractory oxide powder of average particle size ranging from about 0.01 to 0.1 micron, forminga compact of said mixture, and hot deforming saidcompact to "a wrought metal shape at an elevated temperature not exceeding the temperature at which the phase transformation of said matrix metal begins, whereby dissipation of stored energy is greatly inhibited.

9. The method of claim 8, wherein the compact is-hot deformed by extrusion.

.10. The method of claim 8, wherein the refractory oxide is alumina.

11. The method of claim .10, wherein the amount of alumina ranges'from about 3 to 12 volume percent.

References Cited in the file of this patent ,UNITED STATES PATENTS 2,823,988 .Grantet'alL Feb. 18, 1958 2 ,855,659 Thomson Oct..l4, 1958 2,894,838 Gregory July 14, 1959 OTHER REFERENCES Transactions AIME, part in Journal of. Metals," February 1954, pages 247-249.

Claims (1)

1. A METHOD OF PRODUCING A DISPERSION STRENGHENED WROUGHT METAL PRODUCT FROM A MATRIX METAL SELECTED FROM THE GROUP CONSISTING OF IRON, COLBALT, IRON-BASE AND COLBALTBASE METALS OF MELTING POINT ABOVE 1800*F. CHARACTERIZED BY A PHASE TRANSFORMATION TEMPERATURE AT AN ELEVATED HOT WORKING TEMPERATURE AND FROM AN INSOLUBLE REFRACTORY OXIDE DISPERSION HARDENER CHARACTERIZED BY A NEGATIVE FREE ENERGY OF FORMATION AT ABOUT 25*C. OF AT LEAST ABOUT 90,000 CALORIES PER GRAM ATOM OF OXYGEN WHICH COMPRISES, PROVIDING A POWDER MIXTURE COMPRISING SAID MATRIXMETAL AND A SUBSTANTIALLY UNIFORM DISPERSION OF SAID REFRACTORY OXIDE POWDER, FORMING A COMPACT OF SAID MIXTURE, AND HOT DEFORMING SAID COMPACT TO A WROUGHT METAL SHAPED AT AN ELEVATED TEMPERATURE NOT EXCEEDING THE TEMPERATURE AT WHICH THE PHASE TRANSFORMATION OF SAID MATRIX METAL BEGINS, WHEREBY DISSIPATION OF STORED ENERGY IS GREATLY INHIBITED.