# Patent

## Title
Thermomechanical processing of alpha-beta titanium alloys

## Abstract

One embodiment of a method of refining alpha-phase grain size in an alpha-beta titanium alloy comprises working an alpha-beta titanium alloy at a first working temperature within a first temperature range in the alpha-beta phase field of the alpha-beta titanium alloy. The alloy is slow cooled from the first working temperature. On completion of working at and slow cooling from the first working temperature, the alloy comprises a primary globularized alpha-phase particle microstructure. The alloy is worked at a second working temperature within a second temperature range in the alpha-beta phase field. The second working temperature is lower than the first working temperature. The is worked at a third working temperature in a third temperature range in the alpha-beta phase field. The third working temperature is lower than the second working temperature. After working at the third working temperature, the titanium alloy comprises a desired refined alpha-phase grain size.

## Background
STATEMENT REGARDING FEDERALLY SPONSORED RESEARCH OR DEVELOPMENT
This invention was made with United States government support under NIST Contract Number 70NANB7H7038, awarded by the National Institute of Standards and Technology (NIST), United States Department of Commerce. The United States government may have certain rights in the invention.


BACKGROUND OF THE TECHNOLOGY
Field of the Technology
The present disclosure relates to methods for processing alpha-beta titanium alloys. More specifically, the disclosure is directed to methods for processing alpha-beta titanium alloys to promote a fine grain, superfine grain, or ultrafine grain microstructure.
Description of the Background of the Technology
Alpha-beta titanium alloys having fine grain (FG), superfine grain (SFG), or ultrafine grain (UFG) microstructure have been shown to exhibit a number of beneficial properties such as, for example, improved formability, lower forming flow-stress (which is beneficial for creep forming), and higher yield stress at ambient to moderate service temperatures.
As used herein, when referring to the microstructure of titanium alloys: the term “fine grain” refers to alpha grain sizes in the range of 15 μm down to greater than 5 μm; the term “superfine grain” refers to alpha grain sizes of 5 μm down to greater than 1.0 μm; and the term “ultrafine grain” refers to alpha grain sizes of 1.0 μm or less.
Known commercial methods of forging titanium and titanium alloys to produce coarse grain or fine grain microstructures employ strain rates of 0.03 s−1 to 0.10 s−1 using multiple reheats and forging steps.
Known methods intended for the manufacture of fine grain, very fine grain, or ultrafine grain microstructures apply a multi-axis forging (MAF) process at an ultra-slow strain rate of 0.001 s−1 or slower (see, for example, G. Salishchev, et. al., Materials Science Forum, Vol. 584-586, pp. 783-788 (2008)). The generic MAF process is described in, for example, C. Desrayaud, et. al, Journal of Materials Processing Technology, 172, pp. 152-156 (2006). In addition to the MAF process, it is known that an equal channel angle extrusion (ECAE) otherwise referred to as equal channel angle pressing (ECAP) process can be used to attain fine grain, very fine grain, or ultrafine grain microstructures in titanium and titanium alloys. A description of an ECAP process is found, for example in V. M. Segal, USSR Patent No. 575892 (1977), and for Titanium and Ti-6-4, in S. L. Semiatin and D. P. DeLo, Materials and Design, Vol. 21, pp 311-322 (2000), However, the ECAP process also requires very low strain rates and very low temperatures in isothermal or near-isothermal conditions. By using such high force processes such as MAF and ECAP, any starting microstructure can eventually be transformed into an ultrafine grained microstructure. However, for economic reasons that are further described herein, only laboratory-scale MAF and ECAP processing is currently conducted.
The key to grain refinement in the ultra-slow strain rate MAF and the ECAP processes is the ability to continually operate in a regime of dynamic recrystallization that is a result of the ultra-slow strain rates used, i.e., 0.001 s−1 or slower. During dynamic recrystallization, grains simultaneously nucleate, grow, and accumulate dislocations. The generation of dislocations within the newly nucleated grains continually reduces the driving force for grain growth, and grain nucleation is energetically favorable. The ultra-slow strain rate MAF and the ECAP processes use dynamic recrystallization to continually recrystallize grains during the forging process.
A method of processing titanium alloys for grain refinement is disclosed in International Patent Publication No. WO 98/17836 (the “WO '836 Publication”), which is incorporated by reference in its entirety herein. The method in the WO '836 Publication discloses heating and deforming an alloy to form fine-grained microstructure as a result of dynamic recrystallization.
Relatively uniform billets of ultrafine grain Ti-6-4 alloy (UNS R56400) can be produced using the ultra-slow strain rate MAF or ECAP processes, but the cumulative time taken to perform the MAF or ECAP steps can be excessive in a commercial setting. In addition, conventional large scale, commercially available open die press forging equipment may not have the capability to achieve the ultra-slow strain rates required in such embodiments and, therefore, custom forging equipment may be required for carrying out production-scale ultra-slow strain rate MAF or ECAP.
It is generally known that finer lamellar starting microstructures require less strain to produce globularized fine to ultrafine microstructures. However, while it has been possible to make laboratory-scale quantities of fine to ultrafine alpha-grain size titanium and titanium alloys by using isothermal or near-isothermal conditions, scaling up the laboratory-scale process may be problematic due to yield losses. Also, industrial-scale isothermal processing proves to be cost prohibitive due to the expense of operating the equipment. High yield techniques involving non-isothermal, open die processes prove difficult because of the very slow required forging speeds, which requires long periods of equipment usage, and because of cooling-related cracking, which reduces yield. Also, as-quenched, lamellar alpha structures exhibit low ductility, especially at low processing temperatures.
It is generally known that alpha-beta titanium alloys in which the microstructure is formed of globularized alpha-phase particles exhibit better ductility than alpha-beta titanium alloys having lamellar alpha microstructures. However, forging alpha-beta titanium alloys with globularized alpha-phase particles does not produce significant particle refinement. For example, once alpha-phase particles have coarsened to a certain size, for example, 10 μm or greater, it is nearly impossible using conventional techniques to reduce the size of such particles during subsequent thermomechanical processing, as observed by optical metallography.
One process for refining the microstructure of titanium alloys is disclosed in European Patent No. 1 546 429 B1 (the “EP '429 patent”), which is incorporated by reference herein in its entirety. In the process of the EP '429 patent, once alpha-phase has been globularized at high temperature, the alloy is quenched to create secondary alpha phase in the form of thin lamellar alpha-phase between relatively coarse globular alpha-phase particles. Subsequent forging at a temperature lower than the first alpha processing leads to globularization of the fine alpha lamellae into fine alpha-phase particles. The resulting microstructure is a mix of coarse and fine alpha-phase particles. Because of the coarse alpha-phase particles, the microstructure resulting from methods disclosed in the EP '429 patent does not lend itself to further grain refinement into a microstructure fully formed of ultrafine to fine alpha-phase grains.
U.S. Patent Publication No. 2012-0060981 A1 (the “U.S. '981 Publication”), which is incorporated by reference herein in its entirety, discloses an industrial scale-up to impart redundant work by means of multiple upset and draw forging steps (the “MUD Process”). The U.S. '981 Publication discloses starting structures comprising lamellar alpha structures generated by quenching from the beta-phase field of titanium or a titanium alloy. The MUD Process is performed at low temperatures to inhibit excessive particle growth during the sequence of alternate deformation and reheat steps. The lamellar starting stock exhibits low ductility at the low temperatures used and, scale-up for open-die forgings may be problematic with respect to yield.
It would be advantageous to provide a process for producing titanium alloys having fine, very fine, or ultrafine grain microstructure that accommodates higher strain rates, reduces necessary processing time, and/or eliminates the need for custom forging equipment.
SUMMARY
According to one non-limiting aspect of the present disclosure, a method of refining alpha-phase grain size in an alpha-beta titanium alloy comprises working an alpha-beta titanium alloy at a first working temperature within a first temperature range. The first temperature range is in an alpha-beta phase field of the alpha-beta titanium alloy. The alpha-beta titanium alloy is slow cooled from the first working temperature. On completion of working at and slow cooling from the first working temperature, the alpha-beta titanium alloy comprises a primary globularized alpha-phase particle microstructure. The alpha-beta titanium alloy subsequently is worked at a second working temperature within a second temperature range. The second working temperature is lower than the first working temperature and also is in the alpha-beta phase field of the alpha-beta titanium alloy.
In a non-limiting embodiment, subsequent to working at the second working temperature, the alpha-beta titanium alloy is worked at a third working temperature in a final temperature range. The third working temperature is lower than the second working temperature, and the third temperature range is in the alpha-beta phase field of the alpha-beta titanium alloy. After working the alpha-beta titanium alloy at the third working temperature, a desired refined alpha-phase grain size is attained.
In another non-limiting embodiment, after working the alpha-beta titanium alloy at the second working temperature, and prior to working the alpha-beta titanium alloy at the third working temperature, the alpha-beta titanium alloy is worked at one or more progressively lower fourth working temperatures. Each of the one or more progressively lower fourth working temperatures is lower than the second working temperature. Each of the one or more progressively lower fourth working temperatures is within one of a fourth temperature range and the third temperature range. Each of the fourth working temperatures is lower than the immediately preceding fourth working temperature. In a non-limiting embodiment, at least one of working the alpha-beta titanium alloy at the first temperature, working the alpha-beta titanium alloy at the second temperature, working the alpha-beta titanium alloy at the third temperature, and working the alpha-beta titanium alloy at one or more progressively lower fourth working temperatures comprises at least one open die press forging step. In another non-limiting embodiment, at least one of working the alpha-beta titanium alloy at the first temperature, working the alpha-beta titanium alloy at the second temperature, working the alpha-beta titanium alloy at the third temperature, and working the alpha-beta titanium alloy at one or more progressively lower fourth working temperatures comprises a plurality of open die press forging steps, the method further comprising reheating the alpha-beta titanium alloy intermediate two successive press forging steps.
According to another aspect of the present disclosure, a non-limiting embodiment of a method of refining alpha-phase grain size in an alpha-beta titanium alloy comprises forging an alpha-beta titanium alloy at a first forging temperature within a first forging temperature range. Forging the alpha-beta titanium alloy at the first forging temperature comprises at least one pass of both upset forging and draw forging. The first forging temperature range comprises a temperature range spanning 300° F. below the beta transus temperature of the alpha-beta titanium alloy up to a temperature 30° F. less than the beta transus temperature of the alpha-beta titanium alloy. After forging the alpha-beta titanium alloy at the first forging temperature, the alpha-beta titanium alloy is slow cooled from the first forging temperature.
The alpha-beta titanium alloy is forged at a second forging temperature within a second forging temperature range. Forging the alpha-beta titanium alloy at the second forging temperature comprises at least one pass of both upset forging and draw forging. The second forging temperature range is 600 F below the beta transus temperature of the alpha-beta titanium alloy up to 350° F. below the beta transus temperature of the alpha-beta titanium alloy, and the second forging temperature is lower than the first forging temperature.
The alpha-beta titanium alloy is forged at a third forging temperature within a third forging temperature range. Forging the alpha-beta titanium alloy at the third forging temperature comprises radial forging. The third forging temperature range is 1000° F. and 1400° F., and the final forging temperature is lower than the second forging temperature.
In a non-limiting embodiment, after forging the alpha-beta titanium alloy at the second forging temperature, and prior to forging the alpha-beta titanium alloy at the third forging temperature, the alpha-beta titanium alloy may be annealed.
In a non-limiting embodiment, after forging the alpha-beta titanium alloy at the second forging temperature, and prior to forging the alpha-beta titanium alloy at the third forging temperature, the alpha-beta titanium alloy is forged at one or more progressively lower fourth forging temperatures. The one or more progressively lower fourth forging temperatures are lower than the second forging temperature. Each of the one or more progressively lower fourth forging temperatures is within one of the second temperature range and the third temperature range. Each of the progressively lower fourth working temperatures is lower than the immediately preceding fourth working temperature.
According to another aspect of the present disclosure, a non-limiting embodiment of a method of refining alpha-phase grain size in an alpha-beta titanium alloy comprises forging an alpha-beta titanium alloy comprising a globularized alpha-phase particle microstructure at an initial forging temperature within a initial forging temperature range. Forging the alpha-beta titanium alloy at the initial forging temperature comprises at least one pass of both upset forging and draw forging. The initial forging temperature range is 500° F. below the beta transus temperature of the alpha-beta titanium alloy to 350° F. below the beta transus temperature of the alpha-beta titanium alloy.
The workpiece is forged at a final forging temperature within a final forging temperature range. Forging the workpiece at the final forging temperature comprises radial forging. The final forging temperature range is 1000° F. to 1400° F. The final forging temperature is lower than the initial forging temperature.



BRIEF DESCRIPTION OF THE DRAWINGS
The features and advantages of articles and methods described herein may be better understood by reference to the accompanying drawings in which:
FIG. 1 is a flow diagram of a non-limiting embodiment of a method of refining alpha-phase grain size in an alpha-beta titanium alloy according to the present disclosure;
FIG. 2 is a schematic illustration of the microstructure of alpha-beta titanium alloys after processing steps according to a non-limiting embodiment of the method of the present disclosure;
FIG. 3 is a backscattered electron (BSE) micrograph of the microstructure of a forged and slow cooled alpha-beta phase titanium alloy workpiece according to a non-limiting embodiment of the method of the present disclosure;
FIG. 4 is a BSE micrograph of the microstructure of a forged and slow cooled alpha-beta phase titanium alloy according to a non-limiting embodiment of the method of the present disclosure;
FIG. 5 is an electron backscattered diffraction (EBSD) micrograph of a forged and slow cooled alpha-beta phase titanium alloy according to a non-limiting embodiment of the method of the present disclosure;
FIG. 6A is a BSE micrograph of the microstructure of a forged and slow cooled alpha-beta phase titanium alloy according to a non-limiting embodiment of the present disclosure, and FIG. 6B is a BSE micrograph of the microstructure of a forged and slow cooled alpha-beta phase titanium alloy according to the non-limiting embodiment of FIG. 6A that was further forged and annealed according to a non-limiting embodiment of the method of the present disclosure;
FIG. 7 is an EBSD micrograph of a forged and slow cooled alpha-beta phase titanium alloy that was further forged and annealed according to a non-limiting embodiment of the method of the present disclosure;
FIG. 8 is an EBSD micrograph of a forged and slow cooled alpha-beta phase titanium alloy that was further forged and annealed according to a non-limiting embodiment of the method of the present disclosure;
FIG. 9A is an EBSD micrograph of the sample of Example 2 that is a forged and slow cooled alpha-beta phase titanium alloy that was further forged and annealed according to a non-limiting embodiment of the method of the present disclosure;
FIG. 9B is a plot showing the concentration of grains having a particular grain size in the sample of Example 2 shown in FIG. 9A;
FIG. 9C is a plot of the distribution of disorientation of the alpha-phase grain boundaries of the sample of Example 2 shown in FIG. 9A;
FIGS. 10A and 10B are BSE micrographs of respectively the first and second forged and annealed samples;
FIG. 11 is an EBSD micrographs of the first sample of Example 3;
FIG. 12 is an EBSD micrographs of the second sample of Example 3;
FIG. 13A is an EBSD micrograph of the second sample of Example 3;
FIG. 13B is a plot of the relative amount of alpha grains in the sample of Example 3 having particular grain sizes;
FIG. 13C is a plot of the distribution of disorientation of the alpha-phase grain boundaries in the sample of Example 3;
FIG. 14A is an EBSD micrograph of the second sample of Example 3;
FIG. 14B is a plot of the relative amount of alpha grains in the sample of Example 3 having particular grain sizes;
FIG. 14C is a plot of the distribution of disorientation of the alpha-phase grain boundaries in the sample of Example 3;
FIG. 15 is a BSE micrograph of the microstructure of a forged and slow cooled alpha-beta phase titanium alloy that was further forged according to a non-limiting embodiment of the method of the present disclosure;
FIG. 16 is an EBSD micrograph of a forged and slow cooled alpha-beta phase titanium alloy that was further forged according to a non-limiting embodiment of the method of the present disclosure;
FIG. 17A is an EBSD micrograph of the sample of Example 4 that is a forged and slow cooled alpha-beta phase titanium alloy that was further forged according to a non-limiting embodiment of the method of the present disclosure;
FIG. 17B is a plot showing the concentration of grains having a particular grain size in the sample of Example 4 shown in FIG. 17A;
FIG. 17C is a plot of the distribution of disorientation of the alpha-phase grain boundaries of the sample of Example 4 shown in FIG. 17A;
FIG. 18 is an EBSD micrograph of a forged and slow cooled alpha-beta phase titanium alloy that was further forged according to a non-limiting embodiment of the method of the present disclosure;
FIG. 19A is an EBSD micrograph of the sample of Example 4 that is a forged and slow cooled alpha-beta phase titanium alloy that was further forged according to a non-limiting embodiment of the method of the present disclosure;
FIG. 19B is a plot showing the concentration of grains having a particular grain size in the sample of Example 4 shown in FIG. 19A; and
FIG. 19C is a plot of the distribution of disorientation of the alpha-phase grain boundaries of the sample of Example 4 shown in FIG. 19A;



The reader will appreciate the foregoing details, as well as others, upon considering the following detailed description of certain non-limiting embodiments according to the present disclosure.
DETAILED DESCRIPTION OF CERTAIN NON-LIMITING EMBODIMENTS
It is to be understood that certain descriptions of the embodiments described herein have been simplified to illustrate only those elements, features, and aspects that are relevant to a clear understanding of the disclosed embodiments, while eliminating, for purposes of clarity, other elements, features, and aspects. Persons having ordinary skill in the art, upon considering the present description of the disclosed embodiments, will recognize that other elements and/or features may be desirable in a particular implementation or application of the disclosed embodiments. However, because such other elements and/or features may be readily ascertained and implemented by persons having ordinary skill in the art upon considering the present description of the disclosed embodiments, and are therefore not necessary for a complete understanding of the disclosed embodiments, a description of such elements and/or features is not provided herein. As such, it is to be understood that the description set forth herein is merely exemplary and illustrative of the disclosed embodiments and is not intended to limit the scope of the invention as defined solely by the claims.
Also, any numerical range recited herein is intended to include all sub-ranges subsumed therein. For example, a range of “1 to 10” is intended to include all sub-ranges between (and including) the recited minimum value of 1 and the recited maximum value of 10, that is, having a minimum value equal to or greater than 1 and a maximum value of equal to or less than 10. Any maximum numerical limitation recited herein is intended to include all lower numerical limitations subsumed therein and any minimum numerical limitation recited herein is intended to include all higher numerical limitations subsumed therein. Accordingly, Applicants reserve the right to amend the present disclosure, including the claims, to expressly recite any sub-range subsumed within the ranges expressly recited herein. All such ranges are intended to be inherently disclosed herein such that amending to expressly recite any such sub-ranges would comply with the requirements of 35 U.S.C. §112, first paragraph, and 35 U.S.C. §132(a).
The grammatical articles “one”, “a”, “an”, and “the”, as used herein, are intended to include “at least one” or “one or more”, unless otherwise indicated. Thus, the articles are used herein to refer to one or more than one (i.e., to at least one) of the grammatical objects of the article. By way of example, “a component” means one or more components, and thus, possibly, more than one component is contemplated and may be employed or used in an implementation of the described embodiments.
All percentages and ratios are calculated based on the total weight of the alloy composition, unless otherwise indicated.
Any patent, publication, or other disclosure material that is said to be incorporated, in whole or in part, by reference herein is incorporated herein only to the extent that the incorporated material does not conflict with existing definitions, statements, or other disclosure material set forth in this disclosure. As such, and to the extent necessary, the disclosure as set forth herein supersedes any conflicting material incorporated herein by reference. Any material, or portion thereof, that is said to be incorporated by reference herein, but which conflicts with existing definitions, statements, or other disclosure material set forth herein is only incorporated to the extent that no conflict arises between that incorporated material and the existing disclosure material.
The present disclosure includes descriptions of various embodiments. It is to be understood that all embodiments described herein are exemplary, illustrative, and non-limiting. Thus, the invention is not limited by the description of the various exemplary, illustrative, and non-limiting embodiments. Rather, the invention is defined solely by the claims, which may be amended to recite any features expressly or inherently described in or otherwise expressly or inherently supported by the present disclosure.
According to an aspect of this disclosure, FIG. 1 is a flow chart illustrating several non-limiting embodiments of a method 100 of refining alpha-phase grain size in an alpha-beta titanium alloy according to the present disclosure. FIG. 2 is a schematic illustration of a microstructure 200 that results from processing steps according to the present disclosure. In a non-limiting embodiment according to the present disclosure, a method 100 of refining alpha-phase grain size in an alpha-beta titanium alloy comprises providing 102 an alpha-beta titanium alloy comprising a lamellar alpha-phase microstructure 202. A person having ordinary skill in the arts knows that a lamellar alpha-phase microstructure 202 is obtained by beta heat treating an alpha-beta titanium alloy followed by quenching. In a non-limiting embodiment, an alpha-beta titanium alloy is beta heat treated and quenched 104 in order to provide a lamellar alpha-phase microstructure 202. In a non limiting embodiment, beta heat treating the alloy further comprises working the alloy at the beta heat treating temperature. In yet another non-limiting embodiment, working the alloy at the beta heat treating temperature comprises one or more of roll forging, swaging, cogging, open-die forging, impression-die forging, press forging, automatic hot forging, radial forging, upset forging, draw forging, and multiaxis forging.
Still referring to FIGS. 1 and 2, a non-limiting embodiment of a method 100 for refining alpha-phase grain size in an alpha-beta titanium alloy comprises working 106 the alloy at a first working temperature within a first temperature range. It will be recognized that the alloy may be forged one or more times in the first temperature range, and may be forged at one or more temperatures in the first temperature range. In a non-limiting embodiment, when the alloy is worked more than once in the first temperature range, the alloy is first worked at a lower temperature in the first temperature range and then subsequently worked at a higher temperature in the first temperature range. In another non-limiting embodiment, when the alloy is worked more than once in the first temperature range, the alloy is first worked at a higher temperature in the first temperature range and then subsequently worked at a lower temperature in the first temperature range. The first temperature range is in the alpha-beta phase field of the alpha-beta titanium alloy. In a non-limiting embodiment, the first temperature range is a temperature range that results in a microstructure comprising primary globular alpha phase particles. The phrase “primary globular alpha-phase particles”, as used herein, refers to generally equiaxed particles comprising the close-packed hexagonal alpha-phase allotrope of titanium metal that forms after working at the first working temperature according to the present disclosure, or that forms from any other thermomechanical process known now or hereafter to a person having ordinary skill in the art. In a non-limiting embodiment, the first temperature range is in the higher domain of the alpha-beta phase field. In a specific non-limiting embodiment, the first temperature range is 300° F. below the beta transus up to a temperature 30° F. below a beta transus temperature of the alloy. It will be recognized that working 104 the alloy at temperatures within the first temperature range, which may be relatively high in the alpha-beta phase field, produces a microstructure 204 comprising primary globular alpha-phase particles.
The term “working”, as used herein, refers to thermomechanical working or thermomechanical processing (“TMP”). “Thermomechanical working” is defined herein as generally covering a variety of metal forming processes combining controlled thermal and deformation treatments to obtain synergistic effects, such as, for example, and without limitation, improvement in strength, without loss of toughness. This definition of thermomechanical working is consistent with the meaning ascribed in, for example, ASM Materials Engineering Dictionary, J. R. Davis, ed., ASM International (1992), p. 480. Also, as used herein, the terms “forging”, “open die press forging”, “upset forging”, “draw forging”, and “radial forging” refer to forms of thermomechanical working. The term “open die press forging”, as used herein, refers to the forging of metal or metal alloy between dies, in which the material flow is not completely restricted, by mechanical or hydraulic pressure, accompanied with a single work stroke of the press for each die session. This definition of open press die forging is consistent with the meaning ascribed in, for example, ASM Materials Engineering Dictionary, J. R. Davis, ed., ASM International (1992), pp. 298 and 343. The term “radial forging”, as used herein, refers to a process using two or more moving anvils or dies for producing forgings with constant or varying diameters along their length. This definition of radial forging is consistent with the meaning ascribed in, for example, ASM Materials Engineering Dictionary, J. R. Davis, ed., ASM International (1992), p. 354. The term “upset forging”, as used herein, refers to open-die forging a workpiece such that a length of the workpiece generally decreases and the cross-section of the workpiece generally increases. The term “draw forging”, as used herein, refers to open-die forging a workpiece such that a length of the workpiece generally increases and the cross-section of the workpiece generally decreases. Those having ordinary skill in the metallurgical arts will readily understand the meanings of these several terms.
In a non-limiting embodiment of the methods according to the present disclosure the alpha-beta titanium alloy is selected from a Ti-6Al-4V alloy (UNS R56400), a Ti-6Al-4V ELI alloy (UNS R56401), a Ti-6Al-2Sn-4Zr-2Mo alloy (UNS R54620), a Ti-6Al-2Sn-4Zr-6Mo alloy (UNS R56260), and a Ti-4Al-2.5V-1.5Fe alloy (UNS 54250; ATI 425® alloy). In another non-limiting embodiment of the methods according to the present disclosure the alpha-beta titanium alloy is selected from Ti-6Al-4V alloy (UNS R56400) and Ti-6Al-4V ELI alloy (UNS R56401). In a specific non-limiting embodiment of the methods according to the present disclosure the alpha-beta titanium alloy is a Ti-4Al-2.5V-1.5Fe alloy (UNS 54250).
After working 106 the alloy at the first working temperature in the first temperature range, the alloy is slow cooled 108 from the first working temperature. By slow cooling the alloy from the first working temperature, the microstructure comprising primary globular alpha-phase is maintained and is not transformed into secondary lamellar alpha-phases, as occurs after fast cooling, or quenching, as disclosed in the EP '429 patent, discussed above. It is believed that a microstructure formed of globularized alpha-phase particles exhibits better ductility at lower forging temperatures than a microstructure comprising lamellar alpha-phase.
The terms “slow cooled” and “slow cooling”, as used herein, refer to cooling the workpiece at a cooling rate of no greater than 5° F. per minute. In a non-limiting embodiment, slow cooling 108 comprises furnace cooling at a preprogrammed ramp-down rate of no greater than 5° F. per minute. It will be recognized that slow cooling according to the present disclosure may comprise slow cooling to ambient temperature or slow cooling to a lower working temperature at which the alloy will be further worked. In a non-limiting embodiment, slow cooling comprises transferring the alpha-beta titanium alloy from a furnace chamber at the first working temperature to a furnace chamber at a second working temperature. In a specific non-limiting embodiment, when the diameter of the workpiece is greater than to or equal 12 inches, and it is ensured that the workpiece has sufficient thermal inertia, slow cooling comprises transferring the alpha-beta titanium alloy from a furnace chamber at the first working temperature to a furnace chamber at a second working temperature. The second working temperature is described hereinbelow.
Before slow cooling 108, in a non-limiting embodiment, the alloy may be heat treated 110 at a heat treating temperature in the first temperature range. In a specific non-limiting embodiment of heat treating 110, the heat treating temperature range spans a temperature range from 1600° F. up to a temperature that is 30° F. less than a beta transus temperature of the alloy. In a non-limiting embodiment, heat treating 110 comprises heating to the heat treating temperature, and holding the workpiece at the heat treating temperature. In a non-limiting embodiment of heat treating 110, the workpiece is held at the heat treating temperature for a heat treating time of 1 hour to 48 hours. It is believed that heat treating helps to complete the globularization of the primary alpha-phase particles. In a non-limiting embodiment, after slow cooling 108 or heat treating 110 the microstructure of an alpha-beta titanium alloy comprises at least 60 percent by volume alpha-phase fraction, wherein the alpha-phase comprises or consists of globular primary alpha-phase particles.
It is recognized that a microstructure of an alpha-beta titanium alloy including a microstructure comprising globular primary alpha-phase particles may be formed by a different process than described above. In such a case, a non-limiting embodiment of the present disclosure comprises providing 112 an alpha-beta titanium alloy comprising a microstructure comprising or consisting of globular primary alpha-phase particles.
In non-limiting embodiments, after working 106 the alloy at the first working temperature and slow cooling 108 the alloy, or after heat treating 110 and slow cooling 108 the alloy, the alloy is worked 114 one or more times at a second working temperature within a second temperature range, and may be forged at one or more temperatures in the second temperature range. In a non-limiting embodiment, when the alloy is worked more than once in the second temperature range, the alloy is first worked at a lower temperature in the second temperature range and then subsequently worked at a higher temperature in the second temperature range. It is believed that when the workpiece is first worked at a lower temperature in the second temperature range and then subsequently worked at a higher temperature in the second temperature range, recrystallization is enhanced. In another non-limiting embodiment, when the alloy is worked more than once in the first temperature range, the alloy is first worked at a higher temperature in the first temperature range and then subsequently worked at a lower temperature in the first temperature range. The second working temperature is lower than the first working temperature, and the second temperature range is in the alpha-beta phase field of the alpha-beta titanium alloy. In a specific non-limiting embodiment the second temperature range is 600° F. to 350° F. below the beta transus. and may be forged at one or more temperatures in the first temperature range.
In a non-limiting embodiment, after working 114 the alloy at the second working temperature, the alloy is cooled from the second working temperature. After working 114 at the second working temperature, the alloy can be cooled at any cooling rate, including, but not limited to, cooling rates that are provided by any of furnace cooling, air cooling, and liquid quenching, as know to a person having ordinary skill in the art. It will be recognized that cooling may comprise cooling to ambient temperature or to the next working temperature at which the workpiece will be further worked, such as one of the third working temperature or a progressively lower fourth working temperature, as described below. It will also be recognized that, in a non-limiting embodiment, if a desired degree of grain refinement is achieved after the alloy is worked at the second working temperature, further working of the alloy is not required.
In non-limiting embodiments, after working 114 the alloy at the second working temperature, the alloy is worked 116 at a third working temperature, or worked one or more times at one or more third working temperatures. In a non-limiting embodiment, a third working temperature may be a final working temperature within a third working temperature range. The third working temperature is lower than the second working temperature, and the third temperature range is in the alpha-beta phase field of the alpha-beta titanium alloy. In a specific non-limiting embodiment, the third temperature range is 1000° F. to 1400° F. In a non-limiting embodiment, after working 116 the alloy at the third working temperature, a desired refined alpha-phase grain size is attained. After working 116 at the third working temperature, the alloy can be cooled at any cooling rate, including, but not limited to, cooling rates that are provided by any of furnace cooling, air cooling, and liquid quenching, as know to a person having ordinary skill in the art.
Still referring to FIGS. 1 and 2, while not being held to any particular theory, it is believed that by working 106 an alpha-beta titanium alloy at a relatively high temperature in the alpha-beta phase field, and possibly heat treating 110, followed by slow cooling 108, the microstructure is transformed from one comprising primarily of an alpha-phase lamellar microstructure 202 to a globularized alpha-phase particle microstructure 204. It will be recognized the certain amounts of beta-phase titanium, i.e. the body-centered cubic phase allotrope of titanium, may be present between the alpha-phase lamella or between the primary alpha phase particles. The amount of beta-phase titanium present in the alpha-beta titanium alloy after any working and cooling steps is primarily dependent on the concentration of beta-phase stabilizing elements present in a specific alpha-beta titanium alloy, which is well understood by a person having ordinary skill in the art. It is noted that the lamellar alpha-phase microstructure 202, which is subsequently transformed into primary globularized alpha-particles 204, can be produced by beta heat treating and quenching 104 the alloy prior to working the alloy at the first working temperature and quenching, as described hereinabove.
The globularized alpha-phase microstructure 204 serves as a starting stock for subsequent lower-temperature working. Globularized alpha-phase microstructure 204 has generally better ductility than a lamellar alpha-phase microstructure 202. While the strain required to recrystallize and refine globular alpha-phase particles may be greater than the strain needed to globularize lamellar alpha-phase microstructures, the alpha-phase globular particle microstructure 204 also exhibits far better ductility, especially when working at low temperatures. In a non-limiting embodiment herein in which working comprises forging, the better ductility is observed even at moderate forging die speeds. In other words, the gains in forging strain allowed by better ductility at moderate die speeds of the globularized alpha-phase microstructure 204 exceed the strain requirements for refining the alpha-phase grain size, e.g., low die speeds, and may result in better yields and lower press times.
While still not being held to any particular theory, it is further believed that because the globularized alpha-phase particle microstructure 204 has higher ductility than a lamellar alpha-phase microstructure 202, it is possible to refine the alpha-phase grain size using sequences of lower temperature working according to the present disclosure (steps 114 and 116, for example) to trigger waves of controlled recrystallization and grain growth within the globular alpha-phase particles 204,206. In the end, in alpha-beta titanium alloys processed according to non-limiting embodiments herein, the primary alpha-phase particles produced in the globularization achieved by the first working 106 and cooling steps 108 are not fine or ultrafine themselves, but rather comprise or consist of a large number of recrystallized fine to ultrafine alpha-phase grains 208.
Still referring to FIG. 1, a non-limiting embodiment of refining alpha-phase grains according to the present disclosure comprises an optional annealing or reheating 118 after working 114 the alloy at the second working temperature, and prior to working 116 the alloy at the third working temperature. Optional annealing 118 comprises heating the alloy to an annealing temperature in an annealing temperature range spanning 500° F. below the beta transus temperature of the alpha-beta titanium alloy up to 250° F. below the beta transus temperature of the alpha-beta titanium alloy for an annealing time of 30 minutes to 12 hours. It will be recognized shorter times can be applied when choosing higher temperatures, and longer annealing times can be applied when choosing lower temperatures. It is believed that annealing increases recrystallization, albeit at the cost of some grain coarsening, and which ultimately assists in the alpha-phase grain refinement.
In non-limiting embodiments, the alloy may be reheated to a working temperature before any step of working the alloy. In an embodiment, any of the working steps may comprise multiple working steps, such as for example, multiple draw forging steps, multiple upset forging steps, any combination of upset forging and draw forging, any combination of multiple upset forging and multiple draw forging, and radial forging. In any method of refining alpha-phase grain size according to the present disclosure, the alloy may be reheated to a working temperature intermediate any of the working or forging steps at that working temperature. In a non-limiting embodiment, reheating to a working temperature comprises heating the alloy to the desired working temperature and holding the alloy at temperature for 30 minutes to 6 hours. It will be recognized that when the workpiece is taken out of the furnace for an extended time, such as 30 minutes or more, for intermediate conditioning, such as cutting the ends, for example, the reheating can be extended to more than 6 hours, such as to 12 hours, or however long a skilled practitioner knows that the entire workpiece is reheated to the desired working temperature. In a non-limiting embodiment, reheating to a working temperature comprises heating the alloy to the desired working temperature and holding the alloy at temperature for 30 minutes to 12 hours.
After working 114 at the second working temperature, the alloy is worked 116 at the third working temperature, which may be a final working step, as described hereinabove. In a non-limiting embodiment, working 116 at the third temperature comprises radial forging. When previous working steps comprise open-end press forging, open end press forging imparts more strain to a central region of the workpiece, as disclosed in co-pending U.S. application Ser. No. 13/792,285, which is incorporated by reference herein in its entirety. It is noted that radial forging provides better final size control, and imparts more strain to the surface region of an alloy workpiece, so that the strain in the surface region of the forged workpiece may be comparable to the strain in the central region of the forged workpiece.
According to another aspect of the present disclosure, non-limiting embodiments of a method of refining alpha-phase grain size in an alpha-beta titanium alloy comprises forging an alpha-beta titanium alloy at a first forging temperature, or forging more than once at one or more forging temperatures within a first forging temperature range. Forging the alloy at the first forging temperature, or at one or more first forging temperatures comprises at least one pass of both upset forging and draw forging. The first forging temperature range comprises a temperature range spanning 300° F. below the beta transus up to a temperature 30° F. below a beta transus temperature of the alloy. After forging the alloy at the first forging temperature and possibly annealing it, the alloy is slow cooled from the first forging temperature.
The alloy is forged once or more than once at a second forging temperature, or at one or more second forging temperatures, within a second forging temperature range. Forging the alloy at the second forging temperature comprises at least one pass of both upset forging and draw forging. The second forging temperature range is 600° F. to 350° F. below the beta transus.
The alloy is forged once or more than once at a third forging temperature, or at one or more third forging temperatures within a third forging temperature range. In a non-limiting embodiment, the third forging operation is a final forging operation within a third forging temperature range. In a non-limiting embodiment, forging the alloy at the third forging temperature comprises radial forging. The third forging temperature range comprises a temperature range spanning 1000° F. and 1400° F., and the third forging temperature is lower than the second forging temperature.
In a non-limiting embodiment, after forging the alloy at the second forging temperature, and prior to forging the alloy at the third forging temperature, the alloy is forged at one or more progressively lower fourth forging temperatures. The one or more progressively lower fourth forging temperatures are lower than the second forging temperature. Each of the fourth working temperatures is lower than the immediately preceding fourth working temperature, if any.
In a non-limiting embodiment, the high alpha-beta field forging operations, i.e., forging at the first forging temperature, results in a range of primary globularized alpha-phase particles sizes from 15 μm to 40 μm. The second forging process starts with multiple forge, reheats and anneal operations, such as one to three upsets and draws, between 500° F. to 350° F. below the beta transus, followed by multiple forge, reheats and anneal operations, such as one to three upsets and draws, between 550° F. to 400° F. below the beta transus. In a non-limiting embodiment, the workpiece may be reheated intermediate any forging step. In a non-limiting embodiment, at any reheat step in the second forging process, the alloy may be annealed between 500° F. and 250° F. below the beta transus for an annealing time of 30 minutes to 12 hours, shorter times being applied when choosing higher temperatures and longer times being applied when choosing lower temperatures, as would be recognized by a skilled practitioner. In a non-limiting embodiment, the alloy may be forged down in size at temperatures of between 600° F. to 450° F. below the beta transus temperature of the alpha-beta titanium alloy. Vee dies for forging may be used at this point, along with lubricating compounds, such as, for example, boron nitride or graphite sheets. In a non-limiting embodiment, the alloy is radial forged either in one series of 2 to 6 reductions performed at 1100° F. to 1400° F., or in multiple series of 2 to 6 reductions and reheats with temperatures starting at no more than 1400° F. and decreasing for each new reheat down to no less than 1000° F.
According to another aspect of the present disclosure, a non-limiting embodiment of a method of refining alpha-phase grain size in an alpha-beta titanium alloy comprises forging an alpha-beta titanium alloy comprising a globularized alpha-phase particle microstructure at an initial forging temperature within a initial forging temperature range. Forging the alloy at the initial forging temperature comprises at least one pass of both upset forging and draw forging. The initial forging temperature range is 500° F. to 350° F. below the beta transus temperature of the alpha-beta titanium alloy.
The alloy is forged at a final forging temperature within a final forging temperature range. Forging the workpiece at the final forging temperature comprises radial forging. The final forging temperature range is 600° F. to 450° F. below the beta transus. The final forging temperature is lower than each of the one or more progressively lower forging temperatures.
The examples that follow are intended to further describe certain non-limiting embodiments, without restricting the scope of the present invention. Persons having ordinary skill in the art will appreciate that variations of the following examples are possible within the scope of the invention, which is defined solely by the claims.
Example 1
A workpiece comprising Ti-6Al-4V alloy was heated and forged in the first working temperature range according to usual methods to those familiar in the art of forming a substantially globularized primary alpha microstructure. The workpiece was then heated to a temperature of 1800° F., which is in the first forging temperature range, for 18 hours (as per box 110 in FIG. 1). Then it was slow cooled in the furnace at −100° F. per hour or between 1.5 and 2° F. per minute down to 1200° F. and then air cooled to ambient temperature. Backscattered electron (BSE) micrographs of the microstructure of the forged and slow cooled alloy are presented in FIGS. 3 and 4.
In the BSE micrographs of FIGS. 3 and 4, it is observed that after forging at a relatively high temperature in the alpha-beta phase field, followed by slow cooling, the microstructure comprises primary globularized alpha-phase particles interspersed with beta-phase. In the micrographs, levels of grey shading are related to the average atomic number, thereby indicating chemical composition variables, and also vary locally based on crystal orientation. The light-colored areas in the micrographs are beta phase that is rich in vanadium. Due to the relatively higher atomic number of vanadium, the beta phase appears as a lighter shade of grey. The darker-colored areas are globularized alpha phase. FIG. 5 is an electron backscattered diffraction (EBSD) micrograph of the same alloy sample showing the diffraction pattern quality. Again, the light-colored areas are beta-phase as it exhibited sharper diffraction patterns in these experiments, and the dark-colored areas are alpha-phase as it exhibited less sharp diffraction patterns. It was observed that forging an alpha-beta titanium alloy at a relatively high temperature in the alpha-beta phase field, followed by slow cooling, results in a microstructure that comprises primary globularized alpha-phase particles interspersed with beta-phase.
Example 2
Two workpieces in the shape of 4′′ cubes of Ti-6-4 material produced using similar method as for Example 1 was heated to 1300° F. and forged through two cycles (6 hits to 3.5′′ height) of rather rapid, open-die multi-axis forging operated at strain rates of about 0.1 to 1/s to reach a center strain of at least 3. Fifteen second holds were made between hits to allow for some dissipation of adiabatic heating. The workpieces were subsequently annealed at 1450° F. for almost 1 hour and then moved to a furnace at 1300° F. to be soaked for about 20 minutes. The first workpiece was finally air cooled. The second workpiece was forged again through two cycles (6 hits to 3.5′′ height) of rather rapid, open-die multi-axis forging operated at strain rates of about 0.1 to 1/s to impart a center strain of at least 3, viz. a total strain of 6. Fifteen second holds were made as well between hits to allow for some dissipation of adiabatic heating. FIGS. 6A and 6B are BSE micrographs of the first and second samples, respectively, after they underwent processing. Again, grey shading levels are related to the average atomic number, thereby indicating chemical composition variations, and also variations locally with respect to crystal orientation. In this sample shown in FIGS. 6A and 6B, light-colored regions are beta phase, while the darker-colored regions are globular alpha-phase particles. Variation of the grey levels inside the globularized alpha-phase particle reveals crystal orientation changes, such as the presence of sub-grains and recrystallized grains.
FIGS. 7 and 8 are EBSD micrographs of respectively the first and second samples of Example 2. The grey levels in this micrograph represent the quality of the EBSD diffraction patterns. In these EBSD micrographs, the light areas are beta-phase and the dark areas are alpha-phase. Some of these areas appear darker and shaded with substructures: these are the unrecrystallized, strained areas within the original or primary alpha particles. They are surrounded by the small, strain-free recrystallized alpha grains that nucleated and grew at the periphery of those alpha particles. The lightest small grains are recrystallized beta grains interspersed between alpha particles. It is seen in the micrographs of FIGS. 7 and 8 that by forging the globularized material like that of the sample of Example 1, the primary globularized alpha-phase particles are beginning to recrystallize into finer alpha-phase grains within the original or primary globularized particles.
FIG. 9A is an EBSD micrograph of the second sample of Example 2. The grey shading levels in the micrograph represent alpha grain sizes, and the grey shading levels of the grain boundaries are indicative of their disorientation. FIG. 9B is a plot of the relative amount of alpha grains in the sample having particular grain sizes, and FIG. 9C is a plot of the distribution of disorientation of the alpha-phase grain boundaries in the sample. As can be determined from FIG. 9B, a larger number of the alpha-grains achieved on forging the globularized sample of Example 1 and then annealing at 1450° F. then forging again are superfine, i.e., 1-5 μm in diameter and they are overall finer than the first sample of example 2, right after the anneal at 1450° F. that allowed some grain growth and intermediate, static progression of recrystallization.
Example 3
Two workpieces shaped as a 4′′ cube of ATI 425 alloy material produced using similar method as for Example 1 was heated to 1300° F. and forged through one cycle (3 hits to 3.5′′ height) of rather rapid, open-die multi-axis forging operated at strain rates of about 0.1 to 1/s to reach a center strain of at least 1.5. Fifteen second holds were made between hits to allow for some dissipation of adiabatic heating. The workpieces were subsequently annealed at 1400° F. for 1 hour and then moved to a furnace at 1300° F. to be soaked for 30 minutes. The first workpiece was finally air cooled. The second workpiece was forged again through one cycle (3 hits to 3.5′′ height) of rather rapid, open-die multi-axis forging operated at strain rates of about 0.1 to 1/s to impart a center strain of at least 1.5, viz. a total strain of 3. Fifteen second holds were made as well between hits to allow for some dissipation of adiabatic heating.
FIGS. 10A and 10B are BSE micrographs of respectively the first and second forged and annealed samples. Again, grey shading levels are related to the average atomic number, thereby indicating chemical composition variations, and also variations locally with respect to crystal orientation. In this sample shown in FIG. 10A and FIG. 10B, light-colored regions are beta phase, while the darker-colored regions are globular alpha-phase particles. Variation of the grey levels inside the globularized alpha-phase particle reveals crystal orientation changes, such as the presence of sub-grains and recrystallized grains.
FIGS. 11 and 12 are EBSD micrographs of respectively the first and second samples of Example 3. The grey levels in this micrograph represent the quality of the EBSD diffraction patterns. In these EBSD micrographs, the light areas are beta-phase and the dark areas are alpha-phase. Some of these areas appear darker and shaded with substructures: these are the unrecrystallized, strained areas within the original or primary alpha particles. They are surrounded by the small, strain-free recrystallized alpha grains that nucleated and grew at the periphery of those alpha particles. The lightest small grains are recrystallized beta grains interspersed between alpha particles. It is seen in the micrographs of FIGS. 11 and 12 that by forging the globularized material like that of the sample of Example 1, the primary globularized alpha-phase particles are beginning to recrystallize into finer alpha-phase grains within the original or primary globularized particles.
FIG. 13A is an EBSD micrograph of the first sample of Example 3. The grey shading levels in the micrograph represent alpha grain sizes, and the grey shading levels of the grain boundaries are indicative of their disorientation. FIG. 13B is a plot of the relative amount of alpha grains in the sample having particular grain sizes, and FIG. 13C is a plot of the distribution of disorientation of the alpha-phase grain boundaries in the sample. As can be determined from FIG. 13B, the alpha-grains achieved on forging the globularized sample of Example 1 and then annealing at 1400° F. recrystallized and grew again during the anneal resulting in a wide alpha grain size distribution in which most grains are fine, i.e., 5-15 μm in diameter.
FIG. 14A is an EBSD micrograph of the second sample of Example 3 The grey shading levels in the micrograph represent alpha grain sizes, and the grey shading levels of the grain boundaries are indicative of their disorientation, FIG. 14B is a plot of the relative amount of alpha grains in the sample having particular grain sizes, and FIG. 14C is a plot of the distribution of disorientation of the alpha-phase grain boundaries in the sample. As can be determined from FIG. 14B, a number of the alpha-grains achieved on forging the globularized sample of Example 1 and then annealing at 1400° F. then forging again are superfine, i.e., 1-5 μm in diameter. The coarser unrecrystallized grains are remnants of the grains that grew the most during the anneal. It shows that anneal time and temperature must be chosen carefully to be fully beneficial, i.e. allow an increase in recrystallized fraction without excessive grain growth.
Example 4
A 10′′ diameter workpiece of Ti-6-4 material produced using similar method as for Example 1 was further forged through four upsets and draws performed at temperatures between 1450° F. and 1300° F. decomposed as first a series of draws and reheats at 1450° F. down to 7.5′′ diameter, then second, two similar upset-and-draws sequences made of an about 20% upset at 1450° F. and draws back to 7.5′′ diameter at 1300° F., then third, draws down to 5.5′′ diameter at 1300° F., then fourth, two similar upset-and-draws sequences made of an about 20% upset at 1400° F. and draws back to 5.0′′ diameter at 1300° F., and finally draws down to 4′′ at 1300° F.
FIG. 15 is a BSE micrograph of the resulting alloy. Again, grey shading levels are related to the average atomic number, thereby indicating chemical composition variations, and also variations locally with respect to crystal orientation. In the sample, light-colored regions are beta phase, and darker-colored regions are globular alpha-phase particles. Variation of the grey shading levels within globularized alpha-phase particles reveals crystal orientation changes, such as the presence of sub-grains and recrystallized grains.
FIG. 16 is an EBSD micrograph of the sample of Example 4. The grey levels in this micrograph represent the quality of the EBSD diffraction patterns. It is seen in the micrograph of FIG. 16 that by forging the globularized sample of Example 1, the primary globularized alpha-phase particles recrystallize into finer alpha-phase grains within the original or primary globularized particles. The recrystallization transformation is almost complete as only few remaining unrecrystallized areas can be seen.
FIG. 17A is an EBSD micrograph of the sample of Example 4. The grey shading levels in this micrograph represent grain sizes, and the grey shading levels of the grain boundaries are indicative of their disorientation. FIG. 17B is a plot showing the relative concentration of grains with particular grain sizes, and FIG. 17C is a plot of the distribution of disorientation of the alpha-phase grain boundaries. It may be determined from FIG. 17B that after forging the globularized sample of Example 1 and conducting the additional forging through 4 upsets and draws at temperature between 1450° F. and 1300° F., the alpha-phase grains are superfine (1 μm to 5 μm diameter).
Example 5
A full-scale billet of Ti-6-4 was quenched after some forging operations performed in the beta field. This workpiece was further forged through a total of 5 upsets and draws in the following approach: The first two upsets and draws were performed in the first temperature range to start the lamellae break down and globularization process, keeping its size in the range of about 22′′ to about 32′′ and a length or height range of about 40′′ to 75′′. It was then annealed at 1750° F. for 6 hours and furnace cooled down to 1400° F. at −100° F. per hour, with the aim of obtaining a microstructure similar to that of the sample of Example 1. It was then forged through 2 upsets and draws with reheats between 1400° F. and 1350° F., keeping its size in range of about 22′′ to about 32′′ with a length or height of about 40′′ to 75′′. Then another upset and draws was performed with reheats between 1300° F. and 1400° F., in a size range of about 20′′ to about 30′′ and a length or height range of about 40′′ to 70′′. Subsequent draws down to about 14′′ diameter were performed with reheats between 1300° F. and 1400° F. This included some V-die forging steps. Finally the piece was radially forged in a temperature range of 1300° F. to 1400° F. down to about 10′′ diameter. Throughout this process, intermediate conditioning and end-cutting steps were inserted to prevent crack propagation.
FIG. 18 is an EBSD micrograph of the resulting sample. The grey shading levels in this micrograph represent the quality of the EBSD diffraction patterns. It is seen in the micrograph of FIG. 18 that by forging first in the high alpha-beta field, slow cool, and then in the low alpha-beta field, the primary globularized alpha-phase particles begin to recrystallize into finer alpha-phase grains within the original or primary globularized particles. It is noted that only three upsets and draws were performed in the low alpha-beta field as opposed to Example 3 where four such upsets and draws had been carried out in that temperature range. In the present case, this resulted in lower recrystallization fraction. An additional sequence of upset and draws would have brought the microstructure to be very similar to that of Example 3. Also, an intermediate anneal during the low alpha-beta series of upsets and draws (box 118 of FIG. 1) would have improved the recrystallized fraction.
FIG. 19A is an EBSD micrograph of the sample of Example 5. The grey shading levels in this micrograph represent grain sizes, and the grey shading levels of the grain boundaries are indicative of their disorientation. FIG. 19B is a plot of the relative concentration of grains with particular grain sizes, and FIG. 19C is a plot of the orientation of the alpha-phase grains. It may be determined from FIG. 19B that after forging the globularized sample of Example 1, with additional forging through 5 upsets and draws and an anneal performed at 1750° F. to 1300° F., the alpha-phase grains are considered to be fine (5 μm to 15 μm) to superfine (1 μm to 5 μm diameter).
It will be understood that the present description illustrates those aspects of the invention relevant to a clear understanding of the invention. Certain aspects that would be apparent to those of ordinary skill in the art and that, therefore, would not facilitate a better understanding of the invention have not been presented in order to simplify the present description. Although only a limited number of embodiments of the present invention are necessarily described herein, one of ordinary skill in the art will, upon considering the foregoing description, recognize that many modifications and variations of the invention may be employed. All such variations and modifications of the invention are intended to be covered by the foregoing description and the following claims.

## Claims

1. A method of refining alpha-phase grain size in an alpha-beta titanium alloy, the method comprising:
working an alpha-beta titanium alloy at a first working temperature within a first temperature range, wherein the first temperature range is from a temperature 300° F. below a beta transus temperature of the alpha-beta titanium alloy to a temperature 30° F. below the beta transus temperature;
slow cooling the alpha-beta titanium alloy from the first working temperature, wherein on completion of working at the first working temperature and the slow cooling from the first working temperature, the alpha-beta titanium alloy comprises a primary globularized alpha-phase particle microstructure;
wherein the slow cooling occurs prior to any further working of the alpha-beta titanium alloy and comprises cooling the workpiece at a cooling rate no greater than 5° F. per minute;

working the alpha-beta titanium alloy at a second working temperature within a second temperature range, wherein the second temperature range is from a temperature 600° F. below the beta transus temperature to a temperature 350° F. below the beta transus temperature; and
working the alpha-beta titanium alloy at a third working temperature in a third temperature range, wherein the third working temperature is lower than the second working temperature, wherein the third temperature range is 1000° F. to 1400° F., and wherein after working at the third working temperature, the alpha-beta titanium alloy comprises a desired refined alpha-phase grain size.

2. The method according to claim 1, wherein the alpha-beta titanium alloy is selected from Ti-6Al-4V alloy (UNS R56400), Ti-6Al-4V ELI alloy (UNS R56401), a Ti-6Al-2Sn-4Zr-2Mo alloy (UNS R54620), a Ti-6Al-2Sn-4Zr-6Mo alloy (UNS R56260), and a Ti-4Al-2.5V-1.5Fe alloy (UNS 54250).

3. The method according to claim 1, wherein the alpha-beta titanium alloy is selected from Ti-6Al-4V alloy (UNS R56400) and Ti-6Al-4V ELI alloy (UNS R56401).

4. The method according to claim 1, wherein the alpha-beta titanium alloy is a Ti-4Al-2.5V-1.5Fe alloy (UNS 54250).

5. The method according to claim 1, wherein the slow cooling comprises furnace cooling.

6. The method according to claim 1, wherein the slow cooling comprises transferring the alpha-beta titanium alloy from a furnace chamber at the first working temperature to a furnace chamber at the second working temperature.

7. The method according to claim 1, further comprising, before the slow cooling the alpha-beta titanium alloy from the first working temperature:
heat treating the alpha-beta titanium alloy at a heat treating temperature in a heat treating temperature range that is from a temperature 300° F. below a beta transus temperature of the alpha-beta titanium alloy up to a temperature 30° F. below the beta transus temperature of the alpha-beta titanium alloy; and
holding the alpha-beta titanium alloy at the heat treating temperature.

8. The method according to claim 7, wherein holding the alpha-beta titanium alloy at the heat treating temperature comprises holding the alpha-beta titanium alloy at the heat treating temperature for 1 hour to 48 hours.

9. The method according to claim 1, further comprising, after working the alpha-beta titanium alloy at the second working temperature, annealing the alpha-beta titanium alloy.

10. The method according to claim 9, wherein annealing the alpha-beta titanium alloy comprises heating the alpha-beta titanium alloy at a temperature in an annealing temperature range of 500° F. below the beta transus temperature to 250° F. below the beta transus temperature for 30 minutes to 12 hours.

11. The method according to claim 1, wherein at least one of working the alpha-beta titanium alloy at the first temperature, working the alpha-beta titanium alloy at the second temperature, and working the alpha-beta titanium alloy at the third temperature comprises open die press forging.

12. The method according to claim 11, wherein each of the open die press forgings comprises upset forging.

13. The method according to claim 11, wherein each of the open die press forgings comprises draw forging.

14. The method according to claim 11, wherein each of the open die press forgings comprises at least one of upset forging and draw forging.

15. The method according to claim 11, wherein working the alpha-beta titanium alloy at the third working temperature comprises radial forging the alpha-beta titanium alloy.

16. The method according to claim 1, wherein at least one of working the alpha-beta titanium alloy at the first working temperature, working the alpha-beta titanium alloy at the second working temperature, and working the alpha-beta titanium alloy at the third working temperature comprises a plurality of open die press forgings, the method further comprising reheating the alpha-beta titanium alloy intermediate two successive press forgings.

17. The method according to claim 16, wherein reheating the alpha-beta titanium alloy comprises heating the alpha-beta titanium alloy to a previous working temperature and holding the alpha-beta titanium alloy at the previous working temperature for 30 minutes to 12 hours.

18. The method according to claim 1, further comprising:
beta heat treating the alpha-beta titanium alloy at a beta heat treating temperature prior to working the alpha-beta titanium alloy at the first working temperature;
wherein the beta heat treating temperature is within a temperature range from a beta transus temperature of the alpha-beta titanium alloy to a temperature 300° F. greater than the beta transus temperature of the alpha-beta titanium alloy; and

quenching the alpha-beta titanium alloy.

19. The method according to claim 18, wherein beta heat treating the alpha-beta titanium alloy further comprises working the alpha-beta titanium alloy at the beta heat treating temperature.

20. The method according to claim 19, wherein working the alpha-beta titanium alloy at the beta heat treating temperature comprises one or more of roll forging, swaging, cogging, open-die forging, impression-die forging, press forging, automatic hot forging, radial forging, upset forging, draw forging, and multiaxis forging.