Technical Field

The present disclosure relates to metastable beta titanium alloys and especially to methods of processing such alloys and articles produced by said methods.

Background

Titanium alloys are attractive engineering materials due to their excellent mechanical properties, high strength to weight ratios and good corrosion resistance. These alloys are generally categorized based on the phases present in their microstructures. The stable phase of pure titanium at temperatures below 882 °C (the beta transus temperature, T _p of pure titanium) is a hexagonal close packed (hep) crystal structure known as alpha (a) phase. Above the beta transus temperature, beta (β) phase, a body centred cubic (bec) crystal structure, is the stable phase of pure titanium. Alloying elements generally alter the T _p and in titanium alloys a mixture of both phases is commonly seen in the microstructure based on the type and amount of the alloying elements as well as the processing parameters used to form the alloys. For example, in metastable beta alloys, chemical compositions are designed in such a way to retain single beta phase at room temperature upon quenching from the beta phase field.

An aging treatment is commonly carried out in the final stages of manufacturing process in order to adjust and optimise the required mechanical properties of the titanium alloy product. In the aging treatment alpha particles precipitate within the beta grains.

In titanium alloys, a wide range of values of strength, ductility and fracture toughness can be achieved by carefully designing a microstructure that contains both phases with specific volume fractions and morphologies. In metastable beta titanium alloys, beta grain size is one of the microstructure parameters that influence the final mechanical properties. In general, an average large beta grain size has a detrimental effect on the ductility and fatigue life of the alloy and therefore it is desirable to control its size during the manufacturing process. The grain size of as-cast metastable beta titanium alloys is often quite large and can be in the order of millimetres due to the nature of the casting process for such alloys. Therefore, extensive thermomechanical processing is often employed to homogenise and break down the as-cast microstructure during the manufacturing process.

However, high temperature processing of titanium alloys is generally acknowledged as a difficult process due to the alloys' relatively high strength, high solubility of oxygen at elevated temperatures and because of the complex alpha-beta microstructures that can be produced during hot deformation. Like other metals and alloys, the hot deformation mechanism of these alloys depends on their chemical composition, temperature and strain rate as well as the total strain. Dynamic recovery may be the main mechanism for energy restoration of the beta matrix during hot deformation of metastable beta titanium alloys. However, both continuous dynamic recrystallization in the presence of alpha precipitates and discontinuous dynamic recrystallization close to the beta grain boundaries have been observed. The former usually is attained after large amount of strains and the latter causes non-uniform grain size structure. Therefore, the processing window for refining beta grains is generally narrow and its optimization is a challenge in the titanium industry.

In conventional titanium alloy fabrication processes, breaking down the as cast ingot microstructure is typically carried out in a primary fabrication stage where the cast ingot is converted into general mill products such as billets, bars, or plates for example. The working temperature of the primary fabrication process starts well above the beta transus temperature (T _p ) and finishes either at the same temperature (beta isothermal forging) or at a temperature below the beta transus temperature followed by a beta annealing treatment to facilitate recrystallization of the beta grains.

During the fabrication process it is necessary to provide steps to deal with the potential porosities and inclusions that have formed in the ingot during the casting process. Such porosities and inclusions can be the site of cracking in the alloy either during the fabrication process or during use of the finished article. Generally, this problem is dealt with by subjecting the as cast ingot to a substantial amount of deformation usually up to about 75% reduction in at least one dimension of the cast ingot. It has also been generally understood that when producing such high levels of deformation in metastable beta titanium alloys, the temperature should be kept high enough (no more than 50°C below the beta transus temperature, but preferably over the beta transus temperature) to minimise cracking as a result of the deformation process itself and to minimise the formation of a nonhomogeneous microstructure. Conventional understanding is that if deformation is carried out at low temperatures with high stresses on metastable beta titanium alloys, strain induced porosities (SIP) or cracks can form in the alloy.

Deformation is also conventionally carried out at temperatures higher than 50°C below the beta transus temperature because holding the workpiece at lower temperatures leads to non-uniform alpha precipitations, especially when the ingot is directly cooled from a previous deformation stage above the beta transus temperature. If these non-uniform alpha precipitations occur, then deforming of such microstructure causes non uniform recrystallization of beta grains and abnormal grain growth during the deformation process and/or in any subsequent heat treatment stages resulting in the alloy having a non-uniform beta grain size structure.

There remains an ongoing desire to provide improved methods of processing metastable beta titanium alloys.

Summary of the Disclosure

The present disclosure relates to improved methods of processing metastable beta titanium alloys and to articles produced from those methods.

According to an embodiment of the present disclosure, there is provided a method of processing a metastable beta titanium alloy comprising:

a) controllably heating the alloy which has a substantially homogenised beta phase to an aging temperature that is close to but below the alloy's beta transus temperature (Tp);

b) aging the alloy at the aging temperature; and

c) deforming the aged alloy at a deformation temperature that is equal or less than the aging temperature.

According to another embodiment of the present disclosure, there is provided a method of processing a metastable beta titanium alloy comprising:

a) solution treating the alloy at a solution treatment temperature that is above the alloy's beta transus temperature (T _p ) before cooling;

b) aging the solution treated alloy by heating the alloy at a heating rate of 10°C per minute or lower to an aging temperature that is 90 - 1 10°C less than the T _p and holding the alloy at the aging temperature;

c) deforming the aged alloy at a deformation temperature that is equal to or below the aging temperature; and

d) annealing the deformed alloy by heating the alloy to an annealing temperature that is at or above the T _p .

According to a further embodiment of the present disclosure, there is provided a method of processing a metastable beta titanium alloy comprising:

a) controllably heating the alloy which has a substantially homogenised beta phase to an aging temperature that is close to but below the alloy's beta transus temperature (Tp); and

b) aging the alloy at the aging temperature, wherein the aged alloy has a substantially homogenised alpha phase within the beta phase grains. According to a further embodiment of the present disclosure, there is provided a metastable beta phase titanium alloy produced by a method as described in any one of the above mentioned embodiments.

According to a further embodiment of the present disclosure, there is provided a metastable beta phase titanium alloy produced by a method as described in any one of the above mentioned embodiments.

Brief Description of the Figures

Embodiments of the present disclosure will now be described, by way of example only, with reference to the accompany Figures, in which:

Figure 1 is a time against temperature graph showing schematically a method of processing metastable beta titanium alloys according to an embodiment of the present disclosure that includes a solution treatment step, an aging treatment step, a deformation step and an annealing treatment step;

Figure 2 is an optical image of the microstructure of a solution treated Ti-5Al-5Mo-5V-3Cr (ΤΊ5553) alloy;

Figure 3 is an SEM image of the sample shown in Figure 2 after aging at 770 °C for 300 minutes;

Figure 4 is an SEM image of the sample shown in Figure 3 after rolling at 700 °C to achieve a 36% reduction in thickness in one pass;

Figure 5 is an optical image superimposed by EBSD maps for the rolled sample of Figure 4 after annealing at 860 °C for 10 minutes;

Figure 6 is the misorientation angle distributions (correlated) obtained from each of EBSD maps shown in Figure 5 along with random (MacKenzie) distributions (shown in dashed lines);

Figure 7 is a graph of the volume fraction of alpha phase and recrystallized beta during annealing treatment of the rolled aged sample; Figure 8 is the stress strain curves of a ΤΊ5553 alloy which is compressed at 720 °C with different strain rates; and

Figure 9 is an optical image of a ΤΊ5553 alloy after being uniaxial compressed (forged) at strain rates of 10 ^"3 and 10 s ^"1 and subsequently annealed at 860 °C for 10 minutes.

Detailed Disclosure of Embodiments

Disclosed in some embodiments is a method of processing a metastable beta titanium alloy comprising:

a) controllably heating the alloy which has a substantially homogenised beta phase to an aging temperature that is close to but below the alloy's beta transus temperature (T _P );

b) holding the alloy at the aging temperature to age the alloy; and

c) deforming the aged alloy at a deformation temperature that is at or below the aging temperature.

It is understood that the beta transus temperature (T _p ) is the temperature at which the alloy will be completely transformed to the beta phase in an equilibrium condition. It is to be appreciated that the beta transus temperature will be different for different metastable beta titanium alloys.

The present inventors have surprisingly found that the above described method of processing a metastable beta titanium alloy provides a process in which substantially less deformation at much lower temperatures can be carried out to uniformly fragment and refine the microstructure. This means that much lower levels of work (energy) are applied to the alloy in order to obtain the final shape with desirable mechanical properties

This deformation process is enabled by the aging process described in steps (a) and (b).

The aging process described in this embodiment is not equivalent to the aging treatments that are conventionally carried out for the purpose of improving the mechanical properties of the alloy. In particular the aging temperature is substantially higher and the aging time is longer than conventional aging treatments used for improving a titanium alloy's mechanical properties.

The aging process described above obtains relatively coarse and homogenous alpha precipitates inside the large beta grains. Thus, heating to a higher aging temperature and holding for longer time may be referred to as "over aging". In some circumstances, this particular "over aging" treatment also causes long alpha precipitations along the beta grain boundaries.

The present inventors have also found that if the aging temperature is too high, then the alpha precipitations dissolve, which is not desirable for further processing steps.

Disclosed in some embodiments is a method of processing a metastable beta titanium alloy comprising:

a) controllably heating the alloy which has a substantially homogenised beta phase to an aging temperature that is close to but below the alloy's beta transus temperature (T _p ); and

b) aging the alloy at the aging temperature,

wherein the aged alloy has a substantially homogenised alpha phase within the beta phase grains.

In an embodiment, the above methods also comprise annealing the deformed alloy to cause recrystallization of the beta phase.

The present inventors have found that the homogeneous alpha precipitates that are formed during the aging process play an important role in the controlling of the beta grains recrystallization and their grain growth during deformation and the subsequent annealing treatment. Alpha precipitates that are responsible for beta phase fragmentation advantageously prevent beta recrystallization or any abnormal grain growth during the deformation or the subsequent annealing treatment. Beta recrystallization is delayed until the massive dissolution of alpha precipitates occurs. Without wishing to be bound by theory, the inventors attribute this result to the alpha precipitates hindering dislocation movement and causing local beta lattice fragmentation by a Particle Stimulated Nucleation (PSN) mechanism. Further, the alpha precipitates prevent any non-uniform beta recrystallization or abnormal grain growth due to the Zener pinning effect. Onset of beta recrystallization only occurs after massive dissolution of alpha precipitates during the annealing treatment. As a result, a cast microstructure with a grain size on an order of millimetres can be readily refined by the above described methods to a fully recrystallised uniform microstructure with an average beta grain size of approximately 100 μιη or less.

It is to be appreciated that a titanium fabrication process can involve a number of additional processing steps before, between or after the aging, deforming and optionally the annealing steps. Such steps may, in some embodiments be forming steps such as machining, grinding, pickling, rolling, wire drawing and/or welding for example. In some embodiments, the method comprises any one or more of these forming steps.

In an embodiment, the method also comprises, prior to step a), solution treating the alloy to substantially homogenise the beta phase in the alloy.

In an embodiment, solution treating the alloy comprises heating the alloy to a solution treatment temperature that is above the alloy's beta transus temperature (T _p ).

In an embodiment, controllably heating the alloy comprises heating the alloy sufficiently slowly to produce a microstructure having a substantially uniform dispersion of alpha precipitates.

In an embodiment, controllably heating comprises heating at a rate of 10°C per minute or lower.

In an embodiment, the aging temperature is 90 - 1 10°C less than the T _p .

In an embodiment, step b) comprises aging the alloy at the aging temperature for at least 5 hours.

In an embodiment, the deformation temperature is between the aging temperature and 100°C below the aging temperature.

In an embodiment, the deformation temperature is 90 - 210°C less than the T _p .

Disclosed in some embodiments is a method of processing a metastable beta titanium alloy comprising:

a) solution treating the alloy at a solution treatment temperature that is above the alloy's beta transus temperature (T _p ) before cooling;

b) aging the solution treated alloy by heating the alloy at a heating rate of 10°C per minute or lower to an aging temperature that is 90 - 1 10°C less than the T _p and holding the alloy at the aging temperature;

c) deforming the aged alloy at a deformation temperature that is equal to or below the aging temperature; and

d) annealing the deformed alloy by heating the alloy to an annealing temperature that is at or above the T _p . In an embodiment, the solution treatment temperature is at least 50°C above the T _p .

In an embodiment, solution treating the alloy comprises holding the alloy at the solution treatment temperature for at least 30 minutes.

In an embodiment, cooling the alloy after solution treating the alloy comprises cooling the alloy to ambient temperature at a rate equal to or faster than air cooling.

In an embodiment, aging the solution treated alloy comprises holding the alloy at the aging temperature for at least 5 hours.

In an embodiment, the deformation temperature is between the aging temperature and 100°C below the aging temperature.

In an embodiment, the deformation temperature is 90 - 210°C less than the T _p .

In an embodiment, the method also comprises cooling the deformed alloy to ambient temperature prior to annealing, the cooling rate being equal to or faster than air cooling.

In an embodiment, annealing the deformed alloy comprises heating the alloy at a rate of at least 50°C per minute when heating from the aging temperature to the annealing temperature.

In an embodiment, annealing the deformed alloy comprises holding the alloy at the annealing temperature for at least 10 minutes.

In an embodiment, the alloy is cooled to ambient temperature between each of steps a) to d).

In any of the above embodiments, deforming comprises rolling or forging.

In an embodiment, deforming the aged alloy comprises isothermal deformation.

Disclosed in some embodiments is an article formed from a metastable beta titanium alloy, the article produced by a method as described in any one of the above embodiments.

The article may be a sheet, bar, tube, plate, wire or a forging for example. Disclosed in some embodiments is a metastable beta phase titanium alloy produced by a method as described in any one of the above embodiments.

In an embodiment, the alloy has a substantially homogenous microstructure.

In an embodiment, the alloy has a beta grain size of less than 100μιη.

Referring to Figure 1 , a schematic for a heat treatment regime that provides a method 10 of processing a metastable beta titanium alloy is illustrated. The method 10 comprises a solution treatment step 1 1 for homogenising the beta phase of the alloy from the non-uniform grain size and chemical composition of the as cast alloy, an aging treatment step 12, a deformation step 13 such as rolling or forging and an annealing treatment step for recrystallising the aged and deformed alloy. The dashed line 15 indicates the beta transus temperature of the alloy being processed.

The solution treatment step 1 1 comprises heating 20, usually from ambient temperature, to a solution treatment temperature that is at least 50°C above the beta transus temperature. Solutionizing occurs by holding 21 the alloy at the solution treatment temperature for at least 30 minutes. Subsequently, the alloy is cooled 22 to ambient temperature. The cooling rate needs to be at least as fast as air cooling in order to minimise the formation of alpha phase in the alloy.

The aging treatment step 12 comprises heating 23 the solution treated alloy at a controlled rate that is sufficiently slow enough to stabilise the alpha precipitations as they form and maintain the homogeneity of the alloy microstructure. The rate of heating for the treatment step is not faster than 10 °C per minute, preferably about 6 °C per minute. In the aging treatment step 12, the alloy is heated to a temperature which is 90 - 1 10°C less than the beta transus temperature and aged by holding 24 at this temperature for a minimum of 5 hours. The aged alloy is subsequently cooled 25 to ambient temperature.

The deformation step 13 comprises heating 26 to a deformation temperature that is between the aging temperature and 100°C below the aging temperature (i.e. between 90 and 200°C below the beta transus temperature). The alloy is deformed 27 at this temperature. In one embodiment, the deformation 27 comprises rolling the alloy in a single pass to achieve 10 - 50% reduction in the alloy thickness. In another embodiment of the present disclosure, the deformation 27 comprises forging (uniaxial deformation) the alloy to cause 10 - 50% reduction in dimensions of the alloy at a strain rate between 10 ^"3 to 10 s ^' After deforming the alloy, the deformed alloy is cooled 28 to ambient temperature at a cooling rate that is equivalent to or faster than air cooling.

To reform the beta phase in the alloy, the annealing treatment step 14 comprises heating 29 the deformed alloy from ambient temperature to an annealing temperature that is greater than the beta transus temperature. The heating rate to the annealing temperature is not critical up to the previous aging temperature. Above the aging temperature, the heating rate needs to be at least 50 °C per minute. The alloy is subsequently annealed by holding 30 the alloy at the annealing temperature for at least 10 minutes to cause recrystallisation of the beta phase. The annealed alloy is cooled 31 to ambient temperature by air cooling. Beta grains after processing by the above described method is substantially refined with uniform size across the thickness of the alloy.

The method may also comprise one or more forming steps such as machining, grinding, pickling, rolling, wire drawing and/or welding for example. The one or more forming steps may be carried out before, between or after the steps of solution treatment 1 1 , aging 12, deforming 13 and annealing 14.

Example 1

The following processes were carried out on a Ti-5AI-5Mo-5v-3Cr (Ti5553) alloy. Ti5553 is a metastable beta titanium alloy that has been employed in landing gear of commercial planes. The chemical composition of the alloy used is given in Table 1. Ti5553 has a beta transus temperature (T _p ) of -856° C.

Table 1 : Chemical Composition of Ti5553 Alloy as cast, in weight %

This alloy which is a modified version of a Russian alloy VT22 presents deeper hardenability and wider processing window compared with other metastable beta titanium alloys such as Ti-10V-2Fe-3AI due to its low sensitivity to the heat treatments variables. The alloy was received as an as-cast ingot that was 120 mm in diameter. A plate of 12 mm in thickness was sliced from the ingot. The plate was subjected to beta solution treatment in a tube furnace at 900 °C for 30 minutes before being cooled to ambient temperature. After solution treatment, the plate aged at 770 °C for 300 minutes with a heating rate of ~6 °C per minute. Hot rolling was carried out at 700 °C. Prior to rolling, the sample was preheated for -15 minutes at a temperature 30 °C above the rolling temperature. During the rolling, the temperature of the sample was monitored by inserting a thermocouple into the sample. Thickness of the sample was reduced from 1 1 .7 mm to 7.5 mm (~ 36% reduction or ~ 0.44 true strain) in one pass. Annealing was conducted on the hot rolled sample in a fluid bed furnace at 860 °C. The heating rate was approximately 600 °C per minute. Water quenching was employed at the end of all hot rolling and heat treatment processes. In order to minimize oxidation during thermo-mechanical treatments, all samples were coated with a glass based coating (Acheson's Deltaglaze FB-412) and an Argon gas atmosphere was employed during annealing.

Figure 2 shows the typical microstructure of the solution treated sample. The microstructure is taken from the plate thickness and as can be seen it possesses a single beta phase, with grains larger than 1 mm. Microstructure is etched with a solution of 6% HN0 ₃ + 2% HF in water for 3 minutes. Bulk texture measurements from the solution treated sample by electron back scattered diffraction (EBSD) method reveal no particular texture for the grains.

The typical microstructure of the plate after the aging treatment is shown in Figure 3. The image is taken by scanning electron microscopy (SEM) using angular selective backscatter (ASB) detector. Large quantities of alpha precipitates are uniformly formed inside the beta grains as well as along the beta grain boundaries. The relatively slow heating rate (~6 °C per minute) used here during aging treatment provides a uniform dispersion of alpha precipitates across the interior of the beta grains, probably due to an omega assisted alpha nucleation mechanism. The aging temperature employed (770 °C) is almost 100 °C over the range of aging temperature conventionally used for this alloy. As a result, the alpha precipitates are relatively coarse and possess a larger interparticle spacing than what is commonly seen in beta alloys aged for optimum properties. The volume fraction of the alpha phase, F _v , was estimated using point counting and alpha interparticle spacing, A, is calculated from λ=(1- F _V )/N _L equation where N _L is the number of alpha particles intercepting a line of unit length. The alpha volume fraction is 0.35 ± 0.02 and the interparticle spacing of alpha precipitates, A, is 1.35 μιη.

Figure 4 shows a typical microstructure of a metastable beta titanium alloy after the deforming step. The image is taken by SEM using ASB detector. As shown, beta phase is fragmented into small subgrains among the alpha precipitates.

Figure 5 shows a typical optical microstructure of a metastable beta titanium alloy at the end of the annealing treatment. Microstructure is etched with a solution of 6% HN0 ₃ + 2% HF in water for 3 minutes. Three EBSD maps in which each one is collected from the interior of a parent beta grain are superimposed on the optical image. The crystallographic data of the maps are collected from an EBSD detector in an SEM machine that was equipped with the HKL Technology CHANNEL 5 software system. Beta grain boundaries larger than 15 degrees are shown as black lines in the EBSD maps. As can be seen the annealed sample is almost fully recrystallized. The recrystallized grains are uniform over the structure and equiaxed in shape with an average linear intercept grain size of 100 μιη. After etching, the primary beta grain boundaries are still visible under the optical microscope and it is clear that there is interagranular nucleation of beta recrystallized grains. The corresponding correlated misorientation angle distributions for each EBSD map are also shown in Figure 6. Random (Mackenzie) distribution is also displayed on each plot (dashed lines) in order to assess the degree of randomness of each boundary misorientation data. The crystallographic orientations of the recrystallized grains are also shown by means of inverse pole figures (IPF) along the normal direction (ND). As can be seen the misorientation angles of the recrystallized grains are well spread over different angles and are close to the random distribution.

Beta recrystallization and alpha phase fractions for the hot rolled aged sample that has been annealed at 860 °C is shown in Figure 7. The fraction of recrystallized area as well as the fraction of alpha phase have been quantified by EBSD mapping. Figure 7 shows that less than -5% of the alpha precipitates have remained in the microstructure after 90 seconds of annealing. At this point the recrystallization of beta is seen to begin. It appears that the initiation of beta recrystallization is triggered by the disappearance of the alpha phase. This may indicate that beta grain boundaries migration is hindered due to the presence of closely spaced alpha particles - i.e. a Zener pinning effect -so beta recrystallization is delayed until massive alpha dissolution occurs.

Example 2

The same alloy that was used in Example 1 was used in uniaxial compression testing. Cylindrical samples of 15 mm height and 10 mm diameter were cut from the ingot after solution treatment at 900 °C for 30 minutes in a tube furnace. The samples were then aged at 770 °C for 300 minutes with a heating rate of ~6 °C per minute. Uniaxial compression testing was carried out at 720 °C with different strain rates ranging from 10 ^"3 to 10 s ^"1 using a servohydraulic thermomechanical treatment simulator (Servotest, 500 kN). Annealing was conducted on the deformed samples in a fluid bed furnace at 860 °C. The heating rate was approximately 600 °C per minute. Water quenching was employed at the end of all uniaxial deforming and heat treatment processes. In order to minimize oxidation during thermomechanical treatments, all samples were coated with a glass based coating (Acheson's Deltaglaze FB-412) and an Argon gas atmosphere was employed during annealing. The stress strain curves of the samples deformed at 720 °C at different strain rates are shown in Figure 8. The stress strain curves of samples that have been solution treated at 900 °C followed by quenching to 720 °C holding for 10 minutes then deforming with strain rates between 10 ^"3 to 10 s ^"1 are also depicted in Figure 8. The results show that at a given true strain between 0.1 - 0.7, increasing strain rate causes the flow stress to increase. In addition, it is seen that for each given strain rate, at early stage of deformation (up to ~ 0.3 true strains) the flow stresses of the samples containing alpha precipitates are higher than the solution treated samples (i.e. specimens that have been quenched after solution treatment to 720 °C soaked for 10 minutes then deformed). However, with further increasing of strain, the flow stresses of the aged samples gradually decrease and remain close to the values obtained for the solution treated samples. This indicates that the flow stress increment due to aging treatment is not too large to affect the formability of the material during the forging process.

Figure 9 shows the optical microstructure of the samples that have been compressed to cause up to 50% in reduction in thickness (-0.7 true strain) at strain rates of 10 ^"3 and 10 s ^"1 and subsequently annealed at 860 °C for 10 minutes. The images show the whole vertical cross section of the deformed samples where the compression axis is illustrated by a vertical line. The microstructure is etched with a solution of 6% HN0 ₃ + 2% HF in water for 3 minutes. The results show massive grain refinement of the deformed samples after the annealing treatment. The dead metal zones close to the top and bottom of the samples resulted from the contact between the workpiece and the die.

The above described method of processing a metastable beta titanium alloy enables a substantial refinement of the as cast microstructure. The rate of deformation (strain rate) has only a minor effect on the final grain size. Grain refinement is almost uniform across the cross section of the workpiece except at the region close to the tooling due to the plastic flow localization. This effect is noticeable for the samples deformed at higher strain rates.

In the claims which follow and in the preceding description of the invention, except where the context requires otherwise due to express language or necessary implication, the word "comprise" or variations such as "comprises" or "comprising" is used in an inclusive sense, i.e. to specify the presence of the stated features but not to preclude the presence or addition of further features in various embodiments of the invention.