Patent Application
20130000799
Most α/β titanium alloys show superplasticity, i.e., elongation larger than 500%, at sub-transus temperatures when deformed with slower strain rates. The temperature and the strain rate at which superplasticity occurs vary depending on alloy composition and microstructure(1). An optimum temperature for superplastic forming (SPF) ranges from 1832° F. (1000° C.) to as low as 1382° F. (750° C.) in α/β titanium alloys(2). SPF temperatures and beta transus temperatures show a fairly good correlation if other conditions are the same(2).
On the production side, there are significant benefits arising from lowering SPF temperatures. For example, lowering the SPF temperature can result in a reduction in die costs, extended life and the potential to use less expensive steel dies(7). Additionally, the formation of an oxygen enriched layer (alpha case) is suppressed. Reduced scaling and alpha case formation can improve yields and eliminate the need for chemical milling. In addition, the lower temperatures may suppress grain growth thus maintaining the advantage of finer grains after SPF operations(8,9).
Grain size or particle size is one of the most influential factors for SPF since grain boundary sliding is a predominant mechanism in superplastic deformation. Materials with a finer grain size decrease the stress required for grain boundary sliding as well as SPF temperatures(2-4). The effectiveness of finer grains in lowering SPF temperatures was previously reported in Ti-6Al-4V and other alloys(5,6).
There are two approaches for improving superplastic formability of titanium alloys. The first approach is to develop a thermo-mechanical processing that creates fine grains as small as 1 to 2 μm or less to enhance grain boundary sliding. Deformation at lower temperature than conventional hot rolling or forging was studied and an SPF process was developed for Ti-64(5,6).
The second approach is to develop a new alloy system that shows superplasticity at a lower temperature with a higher strain rate. There are several material factors that enhance superplasticity at lower temperatures(1), such as (a) alpha grain size, (b) volume fraction and morphology of two phases, and (c) faster diffusion to accelerate grain boundary sliding(11,16). Therefore, an alloy having a lower beta transus has a potential to exhibit low temperature superplasticity. A good example of an alloy is SP700 (Ti-4.5Al-3V-2Mo-2Fe) that exhibits superplasticity at temperatures as low as 1400° F. (760° C.)(8).
Ti-6Al-4V (Ti-64) is the most common alloy in practical applications since the alloy has been well-characterized. However, Ti-64 is not considered the best alloy for SPF since the alloy requires higher temperature, typically higher than 1607° F. (875° C.), with slow strain rates to maximize SPF. SPF at a higher temperature with a lower strain rate results in shorter die life, excessive alpha case and lower productivity.
Ti-54M, developed at Titanium Metals Corporation, exhibits equivalent mechanical properties to Ti-6Al-4V in most product forms. Ti-54M shows superior machinability, forgeability, lower flow stress and higher ductility to Ti6Al-4V(10). In addition, it has been reported that Ti-54M has superior superplasticity compared to Ti-6Al-4V, which is the most common alloy in this application(2). This result is due partly to chemical composition of the alloy as well as a finer grain size which is a critical factor that enhances superplasticity of titanium materials.(21)
The conventional processing method of titanium alloys is shown in
A method for manufacturing thin sheets of high strength titanium alloys (primarily for Ti6Al-4V) was previously studied by VSMPO in U.S. Pat. No. 7,708,845 and is shown in FIG. 2B.(22) U.S. Pat. No. 7,708,845 requires hot rolling at very low temperatures to obtain fine grains to achieve low temperature superplasticity. The method disclosed in U.S. Pat. No. 7,708,845 can be achieved with rolling mills with very high power, which often lacks flexibility to meet the requirement of a small lot with a variety of gages.(22) The process described in U.S. Pat. No. 7,708,845 is given in the figure as a comparison. In U.S. Pat. No. 7,708,845, rolling is performed at very low temperatures, which may cause excessive mill load, therefore limit the applicability.
Thus, there is a need in the industry to provide a new method for manufacturing titanium alloys that has greater applicability compared to the conventional and prior art methods.
The present disclosure is directed to a method of manufacturing titanium alloy sheets that are capable of low temperature SPF operations. The present method is achieved by the combination of a specified alloy chemistry and sheet rolling process. The method includes the steps of (a) forging a titanium slab to sheet bar, intermediate gage of plates; (b) heating the sheet bar to a temperature higher than beta transus, followed by cooling; (c) heating the sheet bar, then hot rolling to an intermediate gage; (d) heating the intermediate gage, then hot rolling to a final gage; (e) annealing the final gage, followed by cooling; and (f) grinding the annealed sheets, followed by pickling.
In a preferred embodiment (shown in
In one embodiment, the titanium alloy is Ti-54M, which has been previously described in U.S. Pat. No. 6,786,985 by Kosaka et al. entitled “Alpha-Beta Ti—Al—V—Mo—Fe Alloy”, which is incorporated herein in its entirety as if fully set forth in this specification.
The present disclosure is directed to a method of manufacturing titanium alloy sheets that are capable of low temperature SPF operations. The present method is achieved by the combination of a specified alloy chemistry and sheet rolling process. The method includes the steps of
In a preferred embodiment, the sheet bar of step (a) has a thickness from about 0.2″ (0.51 cm) to about 1.5″ (3.8 cm) depending on the finish sheet gages. In variations of this embodiment, the sheet bar of step (a) can be about 0.2″, about 0. 3″, about 0.4″, about 0.5″, about 0.6″, about 0.7″, about 0.8″, about 0.9″, about 1.0″, about 1.1″, about 1.2″, about 1.3″, about 1.4″, about 1.5″, or any increment in between. The thickness of the sheet bar in step (a) is typically chosen based on the thickness of the desired final gage.
In a preferred embodiment, the heating of the sheet bar in step (b) is performed at a temperature between about 100° F. (37.8° C.) to about 250° F. (121° C.) higher than beta transus. In a variation of this embodiment, the heating step is performed at a temperature between about 125° F. (51.7° C.) to about 225 ° F. (107° C.) higher than beta transus. In other variations the heating step is performed at a temperature between about 150° F. (65.6° C.), about 200° F. (93.3° C.) higher than beta transus. In a specific embodiment, the heating step is performed at a temperature at about 175° F. (79.4° C.) higher than beta transus.
In a preferred embodiment, the heating of the sheet bar in step (b) is heated for about 15 to about 30 minutes. In a variation of this embodiment, the sheet bar is heated for about 20 minutes. In another variation of this embodiment, the sheet bar is heated for about 25 minutes.
The cooling in step (b) can be performed at ambient atmosphere, by increasing argon pressure, or by water cooling. In a preferred embodiment, the cooling in step (b) is performed by fan air cooling or faster. Depending on the sheet bar gage, water quench may be used for thick sheet bar (generally above about 0.5″ thick). Fan cool may be sufficient for thinner sheet bar (generally less than about 0.5″ thick). If cooling rate is too slow, structure with thick alpha laths will be formed after cooling, which will prevent material from developing fine grains during intermediate and finishing rolling.
In a preferred embodiment, the heating of the sheet bar in step (c) is performed at a temperature between about 1400° F. (760° C.) to about 1550° F. (843° C.). In a variation of this embodiment, the heating step is performed at a temperature between about 1450° F. (788° C.) to about 1500° F. (816° C.). In a specific embodiment, the heating step is performed at a temperature at about 1475° F. (802° C.).
If the heating temperature is too high, grain coarsening can occur resulting in coarse grain structure even after hot rolling. If the heating temperature is too low, flow stress of material increases resulting overload of rolling mill. Hot rolling is preferably performed with a cascade rolling method without reheat after each pass. Steel pack can be, but does not have to be, used for this intermediate hot rolling. However, reheat can be done, if necessary.
In a preferred embodiment, the sheet bar in step (c) is heated for about 30 minutes to about 1 hour. In variations of this embodiment, the sheet bar is heated for about 40 minutes to about 50 minutes. In another variation of this embodiment, the sheet bar is heated for about 45 minutes.
In a preferred embodiment, the intermediate gage (formed in step c) has a thickness from about 0.10″ (0.3 cm) to about 0.60″ (1.5 cm). In variations of this embodiment, the intermediate gage has a thickness of about 0.10″, about 0.20″, about 0.30″, about 0.40″, about 0.50″, about 0.60″ or any increment in between. The thickness of the intermediate gage is typically chosen based on the thickness of the desired final gage.
The reduction in step (c) is defined as (Ho−Hf)/Ho*100, wherein Ho is the gage of input plate and Hf is a gage of finished gage. In a preferred embodiment, the hot rolling of step (c) has a total reduction controlled between about 40% to about 80%. In variations of this embodiment, the hot rolling step (c) has a total reduction controlled between about 60% to about 70%. In other variations of this embodiment, the hot rolling step (c) has a total reduction controlled at about 40%, 45%, 50%, about 55%, about 60%, about 65%, about 70%, about 75%, or about 80%.
Following the heating and rolling in step (c), the intermediate gage can proceed directly to the finishing hot rolling step (step d) or it can be cooled by a number of methods prior to proceeding. For example, the intermediate gage can be cooled using ambient atmosphere, by increasing argon pressure, or by water cooling. In a preferred embodiment, the cooling is performed by ambient atmosphere.
In a preferred embodiment, the heating of the intermediate gage in step (d) is performed at a temperature between about 1400° F. (760° C.) to about 1550° F. (843° C.). In a variation of this embodiment, the heating step is performed at a temperature between about 1450° F. (788° C.) to about 1500° F. (816° C.). In a specific embodiment, the heating step is performed at a temperature at about 1475° F. (802° C.).
If the heating temperature is too high, grain coarsening takes place resulting coarse grain structure. If the heating temperature is too low, flow stress of materials increases resulting overload of rolling mill. Final hot rolling should be performed with a cascade rolling method without reheat after each pass. In a preferred embodiment, the hot rolling of step (d) is performed with a rolling direction perpendicular to the rolling direction of step (c). In a preferred embodiment, the hot rolling of step (d) utilizes a steel pack in order to avoid excessive heat loss during rolling.
In a preferred embodiment, the intermediate gage in step (d) is heated for about 30 minutes to about 3 hours. In variations of this embodiment, the sheet bar is heated for about 1 hour to about 2 hours. In another variation of this embodiment, the sheet bar is heated for about 1 hour and 30 minutes.
In a preferred embodiment, the final gage (formed in step d) has a thickness from about 0.01″ (0.025 cm) to about 0.20″ (0.51 cm). In variations of this embodiment, the final gage has a thickness of about 0.025″ to about 0.15″. In other variations of this embodiment, the final gage has a thickness of about 0.05″ to about 0.1″. In still other variations of this embodiment, the final gage has a thickness of about 0.010″, about 0.020″, about 0.030″, about 0.040″, about 0.050″, about 0.060″, about 0.070″, about 0.080″, about 0.090″, about 0.100″, about 0.110″, about 0.120″, about 0.130″, about 0.140″, about 0.150″, about 0.160″, about 0.170″, about 0.180″, about 0.190″, about 0.200″, or any increment in between. The thickness of the final desired gage is typically chosen according to the ultimate application of the alloy.
The reduction in step (d) is defined as (Ho−Hf)/Ho*100, wherein Ho is the gage of input plate and Hf is a gage of finished gage. In a preferred embodiment, the hot rolling step of (d) has a total reduction controlled between about 40% to about 75%. In variations of this embodiment, the hot rolling step (d) has a total reduction controlled between about 50% to about 60%. In other variations of this embodiment, the hot rolling step (c) has a total reduction controlled at about 45%, about 50%, about 55%, about 60%, about 65%, about 70%, or about 75%.
Following the heating and rolling in step (d), the final gage can proceed directly to the annealing step (step e) or it can be cooled by a number of methods prior to proceeding. For example, the final gage can be cooled using ambient atmosphere, by increasing argon pressure, or by water cooling. In a preferred embodiment, the cooling is performed by ambient atmosphere.
In a preferred embodiment, the heating of the final gage in step (e) is performed at a temperature between about 1300° F. (704° C.) to about 1550° F. (843° C.). In a variation of this embodiment, the heating step is performed at a temperature between about 1350° F. (732° C.) to about 1500° F. (816° C.). In another variation of this embodiment, the heating step is performed at a temperature between about 1400° F. (760° C.) to about 1450° F. (788° C.). In yet another variation of this embodiment, the heating step is performed at a temperature between about 1300° F. (704° C.) to about 1400° F. (760° C.). In a specific embodiment, the heating step is performed at a temperature at about 1425° F. (774° C.).
If annealing temperature is too low, stress from hot rolling will not be relieved and rolled microstructure will not fully be recovered.
In a preferred embodiment, the heating of the final gage in step (e) is heated for about 30 minutes to about 1 hour. In a variation of this embodiment, the sheet bar is heated for about 40 minutes to about 50 minutes. In another variation of this embodiment, the sheet bar is heated for about 45 minutes.
The cooling in step (e) can be performed at ambient atmosphere, by increasing argon pressure, or by water cooling. In a preferred embodiment, the cooling in step (e) is performed by ambient atmosphere.
The grinding of the annealed gage in step (f) is performed by any suitable grinder. In a preferred embodiment, the grinding is performed by a sheet grinder.
In a preferred embodiment, the annealed gage in step (f) is pickled to remove oxides and alpha case formed during thermo-mechanical processing after the grinding step.
In a preferred embodiment, the titanium alloy is Ti-54M, which has been previously described in U.S. Pat. No. 6,786,985 by Kosaka et al. entitled “Alpha-Beta Ti—Al—V—Mo—Fe Alloy”, which is incorporated herein in its entirety as if fully set forth in this specification.
Superplastic forming (SPF) properties of Ti-54M (Ti-5Al-4V-0.6Mo-0.4Fe) sheet were investigated. A total elongation of Ti-54M exceeded 500% at temperatures between 750° C. and 850° C. at a strain rate of 10−3/S. Values of strain rate sensitivity (m-value) measured by jump strain rate tests were 0.45 to about 0.6 in a temperature range of 730° C. to 900° C. at a strain rate of 5×10−4/S or 1×10−4/S. Flow stress of the alloy was 20% to about 40% lower than that of Ti-6Al-4V(Ti-64) mill annealed sheet. The observed microstructure after the tests revealed the indication of grain boundary sliding in a wide range of temperatures and strain rates.
A piece of Ti-54M production slab was used for the experiment. Two Ti-54M sheets 0.375″ (0.95 cm) were produced using different thermo-mechanical processing procedures, denoted by Process A and Process B, in a laboratory facility. A Ti-64 production sheet sample 0.375″ (0.95 cm) was evaluated for comparison. Chemical compositions of the materials are shown in Table 1. As can be seen, Ti-54M contained a higher concentration of beta stabilizer with a lower Al content compared to Ti-64. Room temperature tensile properties of a typical Ti-54M sheet are shown in Table 2.
Throughout this example “Process A” and “Process B” signify the method performed according to the standard/known process. The processing history for the production of Ti-54M sheets in this example is set forth in Table 1.
Two kinds of tests were conducted to evaluate SPF capability of the sheet materials. Elevated temperature tensile tests were performed at a strain rate of 1×10−3/S until failure with sheet specimens with a gage length of 7.6-mm. Strain rate sensitivity tests to measure m-values were performed in accordance with ASTM E2448-06. Strain rates of the tests were 5×10−4/S and 1×10−4/S at temperatures between 732° C. and 899° C. Microstructures of the cross-section of the reduced section were observed after the tests.
Uniaxial tension tests were conducted at a strain rate of 1×10−3/S in an Argon gas environment at temperatures from 677° C. to 899° C.
True stress-true strain curves obtained by jump strain rate tests for Ti-54M Process A material at a strain rate of 5×10−4/S are shown in
Measurement of Strain Rate Sensitivity (m-value)
The true stress-true strain curves obtained by the jump strain rate tests showed three types of flow curves due to the difference of dynamic restoration process. Flow softening was observed in the tests at lower temperature or higher strain rate. Steady flow curves were obtained in the tests at intermediate temperatures. Flow hardening or strain hardening was seen in the tests at higher temperature with slower strain rate. Microstructures of the reduced section after the test were observed on the tested specimens.
The present work revealed that the flow stress of Ti-54M was significantly lower than that of Ti-64. A primary contributor of lower flow stress is considered to be the effect of Fe that accelerates diffusion leading to lower flow stress, which is evident from the equation for strain rate given by Mukherjee et. al.(23). In addition, lower Al content is another contributor to lower flow stress as Al strengthens both alpha and beta phases at elevated temperatures.
The present results indicated that there was a significant difference in the flow stress between Process A and Process B materials. It is commonly understood that grain size is one of the most influential factors on superplastic formability, which is also shown in the aforementioned equation. The characterization of Ti-54M materials revealed that Process B sheet has slightly smaller primary alpha grains, however, the volume fraction of primary alpha phase in these two materials was very close. An attempt was made to quantify grain boundary length of microstructures shown in
Ti-54M exhibited superplastic forming capability at a temperature range between 730° C. to 900° C. Values of strain rate sensitivity were measured between 0.45 to 0.60 at a strain rate of 5×10−4/S and 1×10−4/S. Flow stress of the alloy was approximately 20% to about 40% lower than that of Ti-64 mill annealed sheet. The morphology of alpha phase and grain boundary density as well as constituents of transformed beta phase had a significant influence on the flow stress levels and the flow curves of superplastic forming in Ti-54M.
Ti-54M exhibits superior machinability in most machining conditions and strength comparable to that of Ti-64. The flow stress of the alloy is typically about 20% to about 40% lower than that of mill-annealed Ti-64 under similar test conditions, which is believed to be one of the contributors to its superior machinability. SPF properties of this alloy were investigated and a total elongation exceeding 500% was observed at temperatures between 750° C. and 850° C. at a strain rate of 10−3/S. A steady flow behavior, which indicates the occurrence of superplasticity, was observed at a temperature as low as 790° C. at a strain rate of 5×10−4/S. It is well understood that grain size is one of the critical factors that influences superplasticity. Fine grain Ti-54M sheets with about 2 to about 3 μm grain size, produced in a laboratory facility, demonstrated that SPF would be possible at temperatures as low as 700° C. The following results report superplastic behavior of fine grain Ti-54M compared with Ti-64 and discuss metallurgical factors that control low temperature superplasticity.
A piece of Ti-54M production slab was used for making sheets in the laboratory. The chemical composition of the material was the same as in Example 1: Ti-4.94% Al-3.83% V-0.55% Mo-0.45% Fe-0.15% O (β transus: 950° C.). Ti-54M sheets with a gage of 0.375″ (0.95 cm) were produced using two different thermo-mechanical processing routes to obtain different microstructures.
Throughout this example, standard grain (SG) signifies that the Ti-54M sheets were process according the standard/known process as discussed in Example 1, Process A. Fine grain (FG) signifies that the Ti-54M sheets were processed according to the embodiments of the present disclosure. Specifically, Fine Grain (FG) sheets were produced with the thermo-mechanical processing routes as shown in Table 4.
Two kinds of tests were conducted to evaluate SPF capability of the sheet materials. Elevated temperature tensile tests were performed at a strain rate of 1×10−3/S until failure with sheet specimens of gage length was 7.6-mm. Strain rate sensitivity tests to measure m-values were performed in accordance with ASTM E2448-06. Strain rates of the tests were selected between 1×10−4/S and 1×10−3/S at temperatures between 1250° F. (677° C.) and 1650° F. (899° C.) in argon gas. Microstructures of the cross-section of the reduced section were assessed after the tests.
Flow Curve and Strain Rate Sensitivity (m-value)
Flow stress and strain rate sensitivity (m-value) were measured on Ti-54M (FG) and Ti-54M (SG) at various test conditions. Flow curves tested at 5×10−4/S are shown in
Ti-54M (SG) material exhibited stable flow behavior at 787° C. and 815° C., where grain boundary sliding is considered to be a predominant mechanism of plastic deformation. In practical superplastic forming operations, the best results are expected at this temperature range. A similar flow behavior was obtained by Ti-54M (FG) material, however, the temperature range that showed a flatter flow curve was observed between 704° C. and about 760° C., and the flow behavior was stable over a wider temperature range.
Strain rate sensitivity (m-value) obtained for Ti-54M (FG) material at various temperatures and strain rates is given in
Flow stress is one of the factors that limit SPF operations since the superplastic forming of higher stress materials may require operations with higher gas pressures or at higher temperatures.
Microstructure after Superplastic Deformation
Microstructures of the reduced sections after the deformation of a true strain=1 are given in
Comparison of SPF Properties with Ti-6Al-4V
It is useful to compare SPF characteristics of Ti-54M and Ti-64, since Ti-64, being the most common alloy for SPF applications, can be considered as a baseline.
Fine grain Ti-54M material exhibited a capability of superplastic forming at temperatures as low as 700° C., which is nearly 100° C. lower than standard grain Ti-54M, and almost 200° C. lower than that of Ti-64. It is useful to discuss metallurgical factors that control superplastic forming behavior of α/β titanium alloys focusing on Ti-54M and Ti-6Al-4V.
Alloy System
Beta transus may be important for two reasons. Primary α grains tend to become smaller with decrease in β transus, since the optimum hot working temperature to manufacture alloy sheets reduces in line with β transus. The temperature that shows approximately 50%/50% of α and β phases will also be proportional to the β transus of the material. Lower SPF temperature of Ti-54M is thus due in part to the lower β transus compared with Ti-64.
Effect of Alloying Elements
Ti-54M contains elevated levels of Mo and Fe and a reduced level of Al compared with Ti-64. The addition of Mo to titanium is known to be effective for grain refinement as Mo is a slow diffuser in α and β phases. On the other hand, Fe is known to be a fast diffuser in both α and β phases(11). Diffusivity of Fe in titanium is faster than self diffusion of Ti by an order of magnitude. A predominant mechanism of superplasticity in α/β titanium alloys is considered to be grain boundary sliding, specifically at grain boundaries of α and β grains. Dislocation climb is an important mechanism to accommodate the strains during grain boundary sliding. As dislocation climb is a thermal activation process, the diffusion of substitutional elements in β phase has a critical role in superplastic deformation. Unusually fast diffusion of Fe is believed to play an important role in accelerating diffusion in β phase, resulting in an enhanced dislocation climb in the beta phase and the activity of dislocation sources and sinks at α/β grain boundaries(11-13).
As demonstrated for Ti-64, finer grain size is an effective way to achieve lower temperature superplasticity(3-6). Ultra-fine grains of Ti-64, typically primary α grains finer than 1 μm, can lower the SPF temperature more than 200° C.(6). The present work demonstrated that a similar grain size effect occurred in Ti-54M.
In addition to lowering SPF temperature in Ti-54M, lower flow stress was measured, particularly in fine grain Ti-54M. Flow stress of fine grain Ti-54M was as low as ¼ of that of fine grain Ti-64 at superplastic conditions, i.e. slow strain rate. The results indicate that grain boundary sliding of Ti-54M was easier than that of Ti-64 when other conditions are the same. Since β phase is more deformable than α phase, flow stress of β phase and mobility of α/β grain boundary may determine overall flow stress of the material. Assuming a sphere for α grain shape, a total surface area of grains can be expressed by A=NπD2, where A is the total surface area of grains; D is a diameter of average α grains; and N is the number of grains in a unit volume. When α grain diameter is different between two materials, and two materials have different average grain sizes, DL and DS, the number of α grains in a unit volume is expressed in Equation (1), where NL and NS are the number of α grains of coarse α material and finer α materials, respectively.
NS=(DL/DS)3NL (Equation 1)
A total α grain boundary area, AS will be given in Equation (2).
AS=π(DS)2NS=(DL/DS)AL (Equation 2)
Equation (2) shows that a total α grain boundary area is inversely proportional to α grain size. Therefore, there will be approximately 4 times of α grain boundary area that can work as sink sources of dislocations in the fine grain Ti-54M compared with standard grain Ti-54M. Significantly larger grain boundary area due to finer grain size will be responsible for lower temperature SPF and low flow stress of fine grain Ti-54M.
Practically, it is also important to consider the effect of prior thermal cycles on the grain growth of primary alpha grains prior to superplastic forming. Diffusion bonding is the most likely heat cycle the materials would receive prior to a multi-sheet superplastic forming operations(14,15) resulting in a certain amount of grain growth. Therefore, the improved superplastic performance arising from the presence of a significant amount of Fe in Ti-54M and the use of Mo to reduce grain growth results in robust SPF performance irrespective of the prior thermal cycle.
Ti-54M has superior superplastic forming properties to that of Ti-64. Fine grain Ti-54M has an SPF capability as low as 700° C.
In addition to low temperature superplasticity, fine grain Ti-54M (FG) possesses significantly lower flow stress compared with standard grain Ti-54M and Ti-64. Superior superplastic capability of Ti-54M is explained by its lower beta transus and chemical composition. Finer grain size will be an additional contributor to low temperature superplasticity.
Ti-54M sheets were produced in the production facility using the disclosed process to produce finer grain sheets. Two sheet bars from the same heat of Ti-54M (Ti-5.07Al-4.03V-0.74Mo-0.53Fe-0.160) were used for the manufacture of 0.180″ and 0.100″ gage sheets. One sheet bar from other heat of Ti-54M (Ti-5.10Al-4.04V-0.77Mo-0.52Fe-0.150) was used for producing the 0.040″ gage sheet material. All sheet bars were beta quenched followed by subsequent rolling operations to the final sheet gage. The sheets were then ground and pickled to remove any alpha case or oxide layer. Detailed process procedure is presented in Table 3.
The resulting microstructure from the final gage material is shown in
Ti-54M (Ti-4.91Al-3.97V-0.51Mo-0.45Fe-0.150) sheet bar of 0.25″ thick was used for making fine grain sheets in a laboratory at three different rolling temperatures as shown in Table 8. Each final gage sheet is annealed at three different temperatures to determine the optimum rolling-annealing condition for the manufacture of Ti-54M fine grain sheets. Metallography samples were excised off of each sheet and average alpha size estimated according to ASTM standards.
Finally,
Additionally, comparing
It appears to be the general trend that as the rolling temperature and/or the annealing temperature is increased, average alpha grains coarsen.
To exemplify the benefits of Ti-54M over Ti-64 and the present invention over the prior art, a process simulation was performed using measured flow stress of two materials (Ti-54M and Ti-64) that are geometrically same dimensions and rolled on a mill whose maximum limit on separating forces is 2500 m. tonnes.
It is evident that separating forces on the rolling mill will increase to unusually high values with lower rolling temperatures, such as temperatures below 1400° F. Therefore, a rolling mill with very high capacities would be required to perform rolling at such low temperatures.
It will be appreciated by persons skilled in the art that the present invention is not limited to what has been particularly shown and described in this specification. Rather, the scope of the present invention is defined by the claims which follow. It should further be understood that the above description is only representative of illustrative examples of embodiments. For the reader's convenience, the above description has focused on a representative sample of possible embodiments, a sample that teaches the principles of the present invention. Other embodiments may result from a different combination of portions of different embodiments.
The description has not attempted to exhaustively enumerate all possible variations. The alternate embodiments may not have been presented for a specific portion of the invention, and may result from a different combination of described portions, or that other undescribed alternate embodiments may be available for a portion, is not to be considered a disclaimer of those alternate embodiments. It will be appreciated that many of those undescribed embodiments are within the literal scope of the following claims, and others are equivalent. Furthermore, all references, publications, U.S. patents, and U.S. Patent Application Publications cited throughout this specification are incorporated by reference as if fully set forth in this specification.
It should be understood that all elemental/compositional percentages (%) are in “weight percent”. Also, it should be understood that the term “inches” has been abbreviated with the quote symbol (″) throughout the application.
This application claims priority under 35 U.S.C. §119(e) to U.S. Provisional Patent Application No. 61/498,447 which was filed on Jun. 17, 2011, the entirety of which is incorporated by reference as if fully set forth in this specification.
| Number | Date | Country | |
|---|---|---|---|
| 61498447 | Jun 2011 | US |
