US3723193A - Process for producing a fine-grained 316 stainless steel tubing containing a uniformly distributed intragranular carbide phase - Google Patents

Abstract

A FINE-GRAINED METAL ALLOY HAVING A SECOND PHASE RANDOMLY-DISTRIBUTED THROUGHOUT THE STRUCTURE AS FINE DISCRETE PARTICLES IS PREPARED BY INTIALLY HEATING THE ALLOY TO A SUFFICIENTLY HIGH TEMPERATURE TO PLACE ESSENTIALLY ALL OF THE SECOND PHASE FORMING MATERIAL IN THE ALLOY INTO SOLUTION AND THEREAFTER SUBJECTING THE ALLOY TO REPETITIVE HEAT TREATMENSTS, INTERPERSED WITH COLD WORK, TO CAUSE RECRYSTALLIZATION AND PARTIAL PRECIPITATION OF SAID SECOND PHASE MATERIAL AT THE LOWEST TEMPERATURE COMMENSURATE WITH THE DEGREE OF COLD WORK. IN ONE EMBODIMENT, AUSTENITIC STAINLESS STEEL TYPE 316 TUBING HAVING A GRAIN SIZE OF ASTM-14 (2.4-2.8 MICRONS) AND A FINE RANDOMLY DISTRIBUTED CARBIDE PRECIPITATE (100-1000 ANGSTROMS) WAS PREPARED. DUE TO IMPROVED METALLURIGICAL PROPERTIES THESE STAINLESS STEELS ARE QUITE SUITABLE AS A CLADDING FOR LIQUID METAL FAST BREEDER REACTOR FUEL ELEMENTS.

Description

March 27, 1973 PROCESS FOR PRODUCING A FINE-GRAINED 316 STAINLESS STEEL TUBING CONTAINING A UNIFORMLY DISTRIBUTED Filed Oct. 27, 1970 TE M P E RAT UR E (C) s. A. RVEIMANN ET AL 3,723,193

INTRAGRANULAR CARBIDE PHASE 2 Sheets-Sheet 1 CARBON (WT.

Fig.1.

INVENTORS. George A. Reimann B William R. Martin ATTORNEY.

March 27, 1973 G. A. REIMANN T 3,723,193

' PROCESS FOR PRODUCING A FINE-GRAINED 316 STAINLESS STEEL TUBING CONTAINING A UNIFORMLY DISTRIBUTED INTRAGRANULAR CARBIDE PHASE Filed Oct. 27, 1970 2 Sheets-Sheet 2 INVEN'IORS. George A. Reimann By William R. Martin ATTORNEY.

US. Cl. 148-123 2 Claims ABSTRACT OF THE DISCLOSURE A fine-grained metal alloy having a second phase randomly-distributed throughout the structure as fine discrete particles is prepared by initially heating the alloy to a sufiiciently high temperature to place essentially all of the second phase forming material in the alloy into solution and thereafter subjecting the alloy to repetitive heat treatments, interspersed with cold Work, to cause recrystallization and partial precipitation of said second phase material at the lowest temperature commensurate with the degree of cold work.

In one embodiment, austenitic stainless steel Type 316 tubing having a grain size of ASTM-14 (2.4-2.8 microns) and a fine randomly-distributed carbide precipitate (100-1000 angstroms) was prepared. Due to improved metallurgical properties these stainless steels are quite suitable as a cladding for liquid metal fast breeder reactor fuel elements.

BACKGROUND OF THE INVENTION This invention relates generally to methods for preparing metal alloys having improved metallurgical properties and, more particularly, to a method for preparing metal alloys which have improved high temperature ductility and resistance to neutron embrittlement. The invention described herein was made in the course of, or under, a contract with the United States Atomic Energy Commission.

The Liquid Metal Fast Breeder Reactor (LMFBR) is now under active consideration. The most critical fuelelement problem is development of a reliable tubular cladding. Various structural materials, such as stainless steel or certain nickel-base alloys, are being evaluated as LMFBR claddings. Of these, stainless steel tubing, in particular Type 316, appears to be the prime candidate as the fuel tubes for the LMFBR. The significant effects in LMBFR cladding in reactor service are reduced ductility and swelling, the latter being caused by radiationproduced voids.

It has heretofore been recognized that reducing the grain size improves the ductility and reduces swelling. It may also be desirable to have a finely-distributed precipitate throughout the alloy structure to enhance metallurgical properties thereof. Various attempts, heretofore, have been made to improve the mechanical properties of candidate cladding materials by thermal-mechanical treatments. In British Pat. 1,124,287 there is disclosed a method for preparing austenitic stainless steel tubing which has improved high temperature ductility. The tubing is cold worked (planetary swaging) and annealed at a temperature within the range of 800 to 900 C. with the working and annealing steps being repeated at least once. Tubing having a final grain size of 4-6 microns is prepared thereby. While the grain size was quite fine, the process provided a coarse carbide precipitate.

Another method is described in Brookhaven National Laboratory Report 50161 (T-526) dated Jan. l-Dec. 31, 1968. There, austenitic stainless steel having a grain size of about 5 microns and a finely-distributed precipitate nited States Patent is prepared. In order to achieve such a fine-grained structure, however, the steel had to be heavily cold worked (-9'0% reduction in area) at room temperature followed by recrystallization at a temperature of about 750 C. Much of the carbide precipitate was concentrated at the grain boundaries by this procedure. While this method was successful in improving the ductility for sheet stock, the extremely large area reductions found necessary are not feasible for the fabrication of tubing. As far as can be ascertained, there has not been a process devised wherein such fine grain size and finely-distributed carbide precipitates were obtained for stainless steel tubing. The smallest grain size in tubing available from a commercial process is about 8-10 microns.

It is desirable and an object of this invention to provide a method for preparing metal alloys having improved metallurgical properties. It is another object of this invention to increase by improvement in the grain structure and dispersion of the carbide precipitate, the high temperature ductility of stainless steel or alloy tubing beyond that which is obtainable in commercially available tubing in order to ofiset, at least in part, the embrittling effects and to minimize swelling of these tubes in a neutron environment.

SUMMARY OF THE INVENTION We have discovered a method for improving the high temperature ductility and resistance to neutron embrittlement of metal alloys. By metal alloy it is intended to refer herein to a metal alloy having a constituent which is precipitatable as discrete particles of a second phase and which has a solubility that decreases markedly with decreasing temperature, the alloy for the temperatures of interest (350-700 C.) comprising a single phase and being capable of recrystallization by repetitive heat treatments interspersed with cold working. Suitable examples of such metal alloys include, but are not limited to, alloy steels (e.g., austenitic stainless steels) and heat treatable aluminum alloys (e.g., as Type 6061 aluminum). These metal alloys, when prepared by our invention, are characterized by the same crystal structure before and after processing, the after structure being of a fine grain size with randomly-distributed discrete particles dispersed through the structure. That is to say that where the metal alloy has a face-centered cubic structure, such as austenitic stainless steel, before being processed in accordance with our method, it will still have a facecentered cubic structure as a final product, the product being of fine grain size and having discrete carbide particles randomly-distributed throughout the structure. This method comprises the steps of:

(a) heating said metal alloy to a temperature to place essentially all of said second phase forming constituent into solution;

(b) cold working said heated metal alloy;

(c) annealing the worked metal alloy at a temperature sufficient to cause partial precipitation of said second phase forming constituent as discrete particles and complete recrystallization of said structure without dissolving any of said second phase; and

(d) repeating steps (b) and (c) until essentially no additional second phase forming constituent is precipitated on the grain boundaries of said structure during a final anneal to thereby give a fine-grained structure having a fine, randomly-distributed second phase as discrete particles throughout said structure.

Type 3 l6 stainless steel tubing (0.250 in. OD, 0.016 in. wall thickness) was produced by cold working (mandrel drawing) with recrystallization anneals in accordance with this method and had a grain size of ASTM-14 (2.42.8 microns) and fine, randomly-distributed carbide precipitate (100-1000 angstroms) throughout the structure. The grain size obtained with this method is considerably finer than the finest grained stainless steel commercially obtainable. The resulting ductility was increased from to (for commercial) to -25% for the finegrained structure when subjected to stress-rupture testing at 650 C.

BRIEF DESCRIPTION OF THE DRAWINGS FIG. 1 is a graph illustrating a typical carbon solubility curve, as function of temperature, for austenitic stainless steels.

FIG. 2 is an electron micrograph showing the microstructure of the grain structure (12,500 of Type 316 stainless steel prepared by the method of this invention.

DESCRIPTION OF THE PREFERRED EMBODIMENT While the invention is applicable to the preparation of metal alloys as hereinbefore defined in widely different sizes and shapes, the invention for illustrative purposes will be hereinafter described in detail with regard to preparing austenitic stainless steel tubing.

In accordance with this invention a stainless steel tubing is initially heated to a temperature which is sufiicient to place all of the carbon present in the steel into solution. Referring to FIG. 1, the stainless steel tubing, which has for example, a carbon content of 0.06 (C should be heated to a temperature of approximately 1000 C. (T to insure that essentially all of the carbon is in solution. While the carbon content may vary widely in stainless steel, those suitable as LMFBR cladding will not likely have a carbon content greater than 0.1 weight percent. Accordingly, for these steels an initial temperature of approximately 1000 C. is quite satisfactory in placing essentially all of the carbon into solution. Also the tubing may be held at this temperature for a period of time, such as for one hour, to further insure complete carbide dissolution.

After the carbon is placed into solution, the tubing is subjected to a series of cold working steps interspersed with low temperature annealing steps to achieve recrystallization and stepwise carbide precipitation. Generally, the subsequent recrystallization and carbide precipitation steps are conducted at the lowest temperature commensurate with the degree of cold Work, and it is critical to the successful practice of this invention that the schedule of cold working and annealing be carried out such that there is insufficient carbon remaining in solution to be deposited on the grain boundaries during the final anneal. Only when such a schedule is carried out is the extremely fine grain size (2.4-2.8 microns) and fine, randomly-distributed carbide precipitate produced which results in the improved ductility and other properties of the finished tubing.

The intermediate annealing steps are performed at a temperature which is low enough to precipitate rather than dissolve the carbides and which will cause complete recrystallization of the cold worked tubing. Referring again to FIG. 1, tubing which has, for example, a carbon content of 0.06 (C and which was heated initially to a temperature of approximately 1000 C. (T to place all of the carbon in solution is annealed after cold working at a temperature not to exceed about 780 C. (T It is critical to this invention that the temperature of 780 C. (T not be exceeded during annealing, otherwise carbides will be dissolved. If this occurs then the final product will contain more carbide precipitate at the grain boundaries than would have occurred if the procedure had been followed precisely and the grain size will be larger. The annealing tempearture (T may, as may the initial temperature (T vary somewhat for each heat of stainless steel or alloy tubing. Temperature T also will further vary depending upon the degree of cold work developed in the tubing. This latter restriction on annealing temperature T is placed because the temperature T must be sufiicient to cause complete recrystallization of the cold worked material. If it is not sufficient to cause complete recrystallization, a higher level of cold Work is necessary. Techniques such as metallography, mechanical properties, hardness and eddy current can be used to evaluate the degree of recrystallization after a given heat treatment. The degree of cold work would be increased by increments best suitable to the fabricator until complete recrystallization is achieved. Where the tubing is simply heat treated without cold working, the finished tubing has large carbide precipitates (-10,000 angstroms) that are located primarily at the grain boundaries. It is believed that the precipitate can be more beneficial to reducing the magnitude of irradiation damage if the precipitate is located within the grains. While carbides are precipitated at grain boundaries during each anneal of our method, subsequent cold working and annealing produce new grains with the precipitates from the prior anneal no longer situated at the new boundaries but, instead are within the grains.

While other methods may be employed to accomplish the cold working of stainless steel or alloy tubing, it is most convenient to employ mandrel drawing as the mode of deformation. To obtain the ultra-fine grain size, the starting stock should be of sufiicient size to permit at least 5 passes of 40% reduction each. If the starting stock has a reasonably fine grain size (ASTM-5-7) initially, achieving the desired size is less difficult. The temperature required to cause dissolution of precipitated carbides as an initial operation may result in a grain size increase in the starting material, however, and caution should be used to avoid excessive heating for prolonged periods lest the starting grain size become inordinately coarse.

After dissolution of carbon, at least 40% reduction in area (with austenitic stainless steels) is required to enable complete recrystallization at the low annealing temperatures (750780 C.) required to produce the structures described by our invention. In some cases, complete recrystallization may not occur after the first cold working step, particularly if the starting grain size is too coarse. Subsequent cold working of a mostly recrystallized structure apparently causes no harm, and the desired final structure may be obtained if partial recrystallization is limited to the initial cold working step. Accepting partial recrystallization is preferred to increasing the annealing temperature if complete recrystallization can be obtained during the second anneal.

The practical reduction limit for cold working austenitic stainless steels by mandrel drawing is about 45%, so if 100% recrystallization is insisted upon at the low annealing temperatures, an additional drawing pass of 20 to 25% reduction can be made to increase total cold work. But additional drawing passes become increasingly difiicult as the level of cold work increases, resulting in increased tooling wear and possible lubrication problems.

The high temperature ductility of the finished finegrained tubing ranged between 15 and 25% when subjected to biaxial burst testing at 650 C. conventionally produced tubing exhibits 0 to 5% ductility under the same test conditions. As a result of the extremely fine grain size and the randomly-distributed carbide precipitate these austenitic stainless steels are believed to be have improved resistance to neutron embrittlement and have reduced swelling. The overall strength at room temperature is increased While the elevated temperature creep strength is reduced. Conventionally produced tubing contains a grain size of ASTM-7-9 and exhibits longer creeprupture times because the structure is strengthened by the carbon remaining in solution. In the subject process, carbon has been precipitated from the structure by the time final size is achieved, leaving a relatively soft matrix which is less creep resistant. Thus, while the finished tubings of this invention have somewhat lower overall creep strength at elevated temperature than previously prepared stainless steel or alloy tubing, the much improved ductility and expected satisfactory resistance to neutron embrittlement render the present tubing excellent candidate cladding for LMFBR type fuels.

Having described the invention in general fashion the following examples are given to illustrate the method with greater particularlity.

EXAMPLE I A Type 316 stainless steel tubing (carbon-0.06%, chromium-17 .3 nickel-l3 .3 molybdenum-2.3 manganese-1.72%, silicon-0.40% which is similar to that required for the LMFBR, was fabricated in accordance with the method of this invention as follows.

The commercially-produced tubing 1.505 in. OD. by 1.245 in. I.D., 0.561 in?) was initially heated to a temperature of 1050 C. for 1 hour to place all carbon into solution. Thereafter, it was subjected to cold work via mandrel drawing and recrystallization anealing as specified in the following schedule. Drawing was accomplished with a standard hydraulic drawbench, using the maximum speeds consistent with the limits of the equipment and the characteristics and geometry of the material drawn. Conventional industrial lubricants were used. Mandrel drawing was employed throughout except for the final pass which was plug drawn because of required manufacturing specifications.

PROCESS SCHEDULE Crosssee- Reduction tional in area, Pass Die Mandrel area, in. percent First anneal-100 hr. at 775 C. 3 0. 700 0. 550 147 42. 8

Second anneal-1 hr. at 776 C 4.....; 5 0. 450 .0873 40. 8

Third anneal-1 hr. at 776 C. 5 0. 440 0. 360 0503 42. 3

Fourth anneal-1 hr. at 775 C. 6 0. 338 0. 275 0303 39. 7

Fifth anneal-1 hr. at 775 C.

Sixth anneal-1 hr. at 775 0.

a Figure in parenthesis is total percent reduction of two (2) passes.

The final tubing had an CD. of 0.250 in. and a Wall thickness of 0.016 in. which was specified for LMFBR fuel tubes. An electron micrograph was made of a sample of the tubing and this is illustrated in FIG. 2. The grain size was determined to be ASTM-14 (2.4-2.8 microns) and the fine carbide precipitate (200 angstroms) was randomly distributed throughout the structure. The grain size ASTM- 14 is about an order of magnitude less than the finest grained stainless steel tubing commercially available. Due to the extremely small grain size and the fine, randomlydistributed carbide precipitate, the high temperature ductility of the resultant stainless steel tubing would be greatly improved over that heretofore commercially available.

EXAMPLE II Samples of Type 316 stainless steel tubing were made from another heat of material produced by a dilferent manufacturer, which contained 0.04% carbon. This material was fabricated according to the process described in Example I in order to verify ability to obtain desired structure and grain size. The metallographic examination confirmed the formation of the fine grain size and random distribution of the fine precipitate.

What is claimed is:

1. A method of improving the resistance of 316 stainless steel tubing containing up to 0.1 weight percent carbon to neutron-irradiation-induced embrittlement which comprises heating said tubing to a temperature sufficient to dissolve all of the carbon and then subjecting the tubing to a series of cold-drawing operations to effect a reduction in area combined with intermediate anneals conducted at a temperature in the range 750-780 C. to eifect essentially complete recrystallization and produce a product having a grain size in the range 2.4-2.8 microns and a randomly distributed intragranular carbide precipitate in the range -1000 angstroms in size.

2. A method of improving the resistance of 316 stainless steel tubing consisting essentially of, in weight percent, 17.3% Cr, 13.3% Ni, 2.3% Mo, 1.72% Mn, 0.40% Si, and the balance Fe to neutron-irradiation-induced embrittlement which comprises heating said tubing to a temperature sufiicient to dissolve all of the carbon and then subjecting the tubing to a series of cold-drawing operations to effect a reduction in area combined with intermediate anneals conducted at a temperature in the range 750- 780 C. to effect essentially complete recrystallization and produce a randomly distributed intragranular carbide precipitate.

References Cited UNITED STATES PATENTS 3,347,715 10/1967 Pfeil 148-12 3,284,250 11/1966 Yeo et al. 148-12 3,473,973 10/1969 Maekawa et al. 148-123 3,573,109 3/1971' Levy et al. 148-12 WAYLAND W. STALLARD, Primary Examiner

Images