Strengthening mechanisms and microstructures in Ti-5553 alloys

The properties of metastable beta titanium alloys depend on several strengthening mechanisms such as grain size, solid solution strengthening, and precipitation hardening.

All of these mechanisms can be manipulated during processing ultimately leading to particular microstructures. A deep understanding of the phase transformations and microstructural modifications is necessary to improve and modify the mechanical properties for superior performance ⁴⁶.

As with other metastable beta titanium alloys, strengthening is primarily achieved by the precipitation of the alpha phase in the parent beta ⁴⁷. The volume fraction, location, and morphology of the precipitates strongly influences the mechanical properties of the

49 alloy⁴⁸. Figure 17 shows a schematic of a pseudo-binary phase diagram for titanium alloys. The β-transus temperature is found between 850-880 ˚C. If the alloy is heated above this temperature, a single phase solid solution of the β phase is obtained. If it is subsequently cooled down from the β-transus, α is precipitated in the α-β field ⁴⁷.

Figure 17. Schematic for a binary phase diagram for Ti alloys ⁴⁷

The β to α transformation is considered to be a displacive transformation, bearing in mind that a chemical composition change occurs with a long range migration of interstitial and substitutional atoms ⁴⁶. It has been shown that depending on the Moeq, aging at a given temperature after quenching from the β phase field to obtain a solid solution leads to a varied amount of α precipitates as well as different types of morphologies ⁴⁶. If aged at a relatively high temperature, α precipitates heterogeneously at grain boundaries into a globular morphology ⁴⁶. If aging takes place at lower temperatures, the precipitate is distributed more homogenously and tends to have a plate shaped morphology with a basket weave orientation. When slowly cooled, the plate like structure is coarsened which

50 typically improves the damage tolerance characteristics ⁴⁷. This principle is illustrated in Figure 18 which shows some microstructures of Ti-5553 with different α-morphologies at various aging temperatures after homogenization above the β transus and ageing time of 3.6 ks. It is possible to combine the two ageing temperature regions (both high temperature and low temperature) to generate a bimodal α microstructure to bring about a combined effect on the mechanical properties. Such a micrograph is shown in Figure 18 d) in which large primary α platelets are present within a β matrix hardened by fine secondary α platelets ⁴⁶

The microhardness of the Ti-5553 alloys after aging at different temperatures is shown in Figure 19. The hardness is the lowest for the as quenched β state and increases until a maximum hardness is reached at 400 ˚C, after which it then decreases ⁴⁶. This effect is directly related to the microstructures as seen in Figure 18. The coarse α-morphologies that are scarcely dispersed in the β matrix have a relatively poor hardening efficiency.

The drop below 400˚C is due to an insufficient amount of energy that is required to precipitate out the α-phase in the time scale that was considered ⁴⁶.

51

(a) 800 ˚C (b) 700 ˚C

(c) 500˚C (d) two step heat treatment for bimodal structure

Figure 18. Scanning electron micrographs of microstructures resulting from various ageing temperatures for Ti-5553 ⁴⁶

Figure 19 (b) depicts the tensile properties of the Ti-5553 alloy as a function of ageing temperature ⁴⁶. Although maximum strength of the alloy is reached at approximately 400˚C, the elongation/ductility is respectively lower. It is seen that ductility becomes an issue for the lowest ageing temperatures ⁴⁶. The bimodal structure may present a method

52 to recover ductility in aged β alloys, although the results in this study indicated otherwise

46.

(a) (b)

Figure 19. (a) microhardness and (b) tensile properties of Ti-5553 alloy as a function of ageing temperature

The direct transformation of the β-phase into an equilibrium α-phase occurs at relatively high temperatures. At lower temperatures, this decomposition becomes sluggish and other transformation products can be formed such as a hexagonal or trigonal phase called the omega phase. There are three types of omega precipitates as described by Ankem; (1) an athermal omega phase, ωa,, which forms by a displacive mechanism during quenching and generally lends very high strength but with very low ductility, (2) an isothermal omega phase, ωi,, which generally forms in the temperature of ~200-500˚C and (3) a stress induced omega phase ωs49

. The stress induced omega phase, ωs, has been observed after dynamic compressive deformation at room temperature of a metastable β Ti-20V alloy ⁴⁹. It should be noted that the presence of the omega phase is typically undesirable since it is known to cause severe embrittlement of the alloy concerned ⁹.

53 2.3.4 Weld Defects in Titanium Alloys

Weldability is determined by the capability of the alloy to produce a weld that is free of discontinuities or defects. Defects can arise from deviations from a qualified welding procedure with respect to material properties, laser beam characteristics, and processing parameters ¹³. Some of these defects are associated with the geometry of the weld bead and can be assessed by simple visual inspection methods, while others are concealed in the fusion zone or heat affected zone and can only be detected by non-destructive methods or metallurgical examination ³⁹. Defects that may be encountered during welding titanium alloys is outlined below:

- Porosity

- Embrittlement

- Macrosegregation

- Solidification cracking

- Contamination cracking

Porosity

During the solidification of the weld pool, gases may be entrapped in the molten metal resulting in porosity⁵⁰. The source of gases involved in pore formation can be the environment, shielding gas, the filler wire, and/or the base metal. There are several theories that try to explain the gas evolution process from the base metal. The most popular theory attributes the gas evolution to the rejection of dissolved gases which is brought about by a decrease in solubility, when transforming from liquid to solid⁵⁰.

54 Another theory attributes the gas evolution to be the result of chemical reactions that liberate gaseous products⁵⁰.

Many authors single out hydrogen as the likely culprit. It is believed that the main cause for porosity is the hydrogen dissolved in the weld pool and the internal hydrogen contained in the filler wire and base metal as well as that picked up from the environment during welding ⁵⁰. Putting aside cause of gas evolution, the bubbles nucleate from the growing solid interface as the molten metal solidifies. The small gas bubbles then proceed to coalesce into larger ones as they rise in the molten metal ⁵⁰. If the weld metal completely solidifies before they have sufficient time to float and escape to the top, they are entrapped as porosity ⁵⁰. When welding speeds are high such as that in laser welding, the cooling rates tend to be high, resulting in insufficient time for bubbles to nucleate and grow. When welding speeds are slow, the bubbles are allowed sufficient time to escape

50. Both of these extremes during welding are beneficial in reducing the amount of porosity, thereby enhancing weld quality. Several researchers have shown that porosity can cause cracks in titanium alloys that lowers the fatigue resistance of the alloys so it is important that they be minimized ^50-51.

Avoiding porosity begins with protecting the weld pool from exposure to oxygen and/or hydrogen during welding. These means that the joint must be adequately cleaned and degreased before welding and that adequate shielding methods must be utilized as was discussed in the previous sections.

55 Embrittlement

Embrittlement is caused by weld metal contamination usually by gas absorption. Recall that titanium has a high affinity for interstitial elements such as oxygen, nitrogen, and hydrogen at temperatures above 500˚C, if not adequately protected from the environment by proper shielding methods.

Titanium can contain a surface layer of α grains that are commonly referred to as an alpha-case which is enriched of oxygen ⁵². The alpha-case is formed when the molten metal picks up oxygen as a solid solution. As the molten metal solidifies, oxygen diffuses further below the surface at high temperature producing alpha case whose thickness depends on the cooling rate ⁵². The alpha-case which is hard and brittle due to its high oxygen content reduces wear resistance ⁵².

Hydrogen also poses a threat during welding. In fact, the dangers of hydrogen pickup are of greater importance than that of oxidation because hydrogen does not create a visible surface condition that can be used as a check against excess hydrogen ³⁹. The current specifications of hydrogen limit the content to a maximum of 100 to 200 ppm. When the content is above this limit, hydrogen can embrittle the alloy thereby reducing its impact strength and notch tensile strength and also cause delayed cracking ³⁹.

A number of researchers have studied the role of hydrogen cracking in titanium welds

53-54. Tal-Gutelmacher investigated the phenomena of hydrogen cracking in Ti-6Al-4V and a metastable Beta-21S alloy ⁵⁴. It was found that the prior microstructure plays an

56 important role in their behaviour under hydrogen containing environments ⁵⁴. The main mechanism for hydrogen cracking in the Ti-6Al-4V alloy was due to the formation and rupture of brittle hydride phases ⁵⁴. It was found that the severity of hydrogen degradation depends on the amount and distribution of the beta phases in the microstructure ⁵⁴. In contrast, the metastable Beta-21A alloy exhibited a fair resistance to hydrogen⁵⁴. It should be noted that this fair resistance to hydrogen was only in the aged condition. In the mill-annealed condition (Figure 20 a), consisting predominantly of the beta phase, the hydrogen absorption was high. It was only after ageing to obtain an acicular alpha phase in the beta matrix, was the hydrogen absorption lower. This can be attributed to the BCC structure of the beta phase which has a higher diffusion rate than the alpha HCP phase due to a diffusion coefficient that is almost 2 orders of magnitude lower than that of the BCC structure ⁸.

Figure 20. Microstructure of Timetal Beta-21S alloy (a) mill-annealed condition and (b) ageing at 538˚C for 8 hr ⁵⁴

(a)

(b)

57 Macrosegregation

Macrosegregation in titanium alloys is defined as segregation that extends over distances of several grain diameters and it may be an issue during fusion welding ³⁹. It is often seen in the form of transverse solute bands ³⁹. They are usually attributed to thermal variations in the weld pool which change the solid-liquid interface velocity. Several studies have shown that the bands result from vanadium and aluminum segregation. The effect of macrosegregation has been reported to be more prominent in the more heavily beta-stabilized alloys ³⁹. This is schematically shown in Figure 21 for a gas tungsten arc weld fusion zone produced between the alpha-beta alloy Ti-6Al-4V and the metastable beta alloy Ti-15V-3Cr-3Al-3Sn ³⁹. The transverse solute bands shown with the smaller arrow found to be depleted of vanadium and chromium (both beta stabilizers) transformed to martensite while the more heavily stabilized regions were retained as the beta phase during weld cooling ³⁹.

It should be noted that the degradation of weld integrity, structure, or properties due to macrosegregation is often not significant in titanium alloys as it may be in other alloy systems due to the limited extent of segregation of common alloying elements and the diffusional homogenization when cooling through the beta phase field ³⁹.

Solidfication Cracking

Solidification cracking occurs during the solidification phase of the weld bead in certain susceptible alloys ¹³. Titanium alloys are generally not considered to be susceptible to fusion zone solidification cracking. However, under severe conditions of restraint,

58 cracking along the columnar beta grains may occur ²⁵. Solid forms in the weld metal as interlocking dendrites during cooling which is accompanied by the rejection of solutes ¹³. This solute rich liquid has a low solidification temperature which tends to remain at the grain boundaries lowering the cohesion between adjacent grains. The presence of tensile stresses can lead to solidification cracks ¹³. Furthermore, liquation cracking is also not an issue for conventional titanium alloys in the HAZ due to the absence of second phase precipitate particles ²⁵.

(a) (b)

Figure 21. Dissimilar welds between Ti-6Al-4V and Ti-15V-3Cr-3Al-3Sn sheets showing martensitically transformed solute bands (small arrows) for (a) light optical microscope

and (b) SEM ³⁹

Contamination Cracking

The presence of interstitial elements such as oxygen, nitrogen above a certain amount may cause stresses during welding that further result in the generation of cracks ²⁵. For example, oxygen levels on the order of 3000 ppm in the weld have been known to result in transverse cracks. The HCP alpha phase is particularly sensitive to contamination cracking since the presence of interstitial atoms tend to distort the HCP crystal lattice ²⁵.

59 These types of cracks may be avoided by minimizing the exposure time between the molten metal and interstitial elements, by properly degreasing the joint prior to welding and by providing adequate shielding with the use of the trailing and backup shields ²⁵.

2.3.5. Studying the mechanical properties of weldments

Typically, the tensile mechanical properties of weldments are assessed using a conventional transverse tensile configuration for which the weld line is perpendicular to the loading direction and/or a longitudinal tensile configuration for which the weld line is parallel to the loading direction ⁵⁵. By using this method, one can determine the global tensile behaviour of the weld. Although being a useful technique to determine the global properties of the weld, this technique gives us information on the weakest region of the weld and overlooks the discrepancies in structure and ultimately the strength and ductility among the base metal, heat affected zone, and fusion zone which may be vastly different in laser welds ⁵⁶.

Different techniques exist to study the local properties of weldments. If the weld zones are sufficiently large enough, micro-tensile specimens may be extracted for tensile testing. This gives direct information on the mechanical properties of the weld ⁵⁷. In many situations, the different zones in laser welds are not large enough and other techniques must be employed. Several researchers have used digital image correlation (DIC) to study inhomogenous strains across friction stir and laser welds. In this technique, digital images are obtained of the weld prior to and during deformation ^58-59. A random speckled pattern applied on the specimen surface with sufficient contrast allows

60 for correlation of the images and the displacement is then determined between features on the specimen surface ^58-59.

2.2.6. Summary

The laser welding of titanium alloys is increasingly being considered for developing near net shape structural components. Processing parameters for the laser welding of titanium alloy Ti-6Al-4V have clearly been identified for producing repeatable and high quality joints on a variety of material thicknesses and joint configurations. Less information is available on metastable beta titanium alloys especially the recently emerged titanium alloy Ti-5553. As this alloys gains acceptance in the aerospace industry, it can be predicted that the laser welding process will play a crucial role in its wider implementation hence it is necessary to address its weldability by producing high quality defect-free welds with good mechanical performance.

61 Chapter 3. Effect of Defocusing Distance and Weld Speed on Laser

Welding of Ti-5553

3.1. Introduction

There are many advantages associated with laser welding such as a low and precise heat input, ease of automation, and rapid processing rates. In addition, laser welding offers additional flexibility when welding titanium alloys since it offers the possibility of welding either autogenously or with filler material in the form of a wire or powder. For autogenous continuous wave Nd:YAG laser welding, the power, defocusing distance, and welding speed are amongst the most important parameters that influence how the energy is applied to the joint and ultimately, the quality of the welds. Hence, it is crucial that these parameters be optimized to fit within an optimum processing window that satisfies the aerospace specification tolerances. The work presented in this chapter deals with the effect of defocusing distance and welding speed on the welding quality of Ti-5553 autogenous welds.

3.2. Experimental Procedure

Ti-5553 material, received in ingot form, was sectioned to obtain weld coupons of 76 mm in length x 38 mm in width x 3.1 mm in thickness. Prior to welding, the weld coupons were (1) solution treated at 815.5°C for 45 min. in vacuum followed by an argon quench and (2) aged at 621°C for 8 hrs in argon partial pressure followed by an argon quench.

The faying surfaces and neighbourhood of all the specimens were brushed and then cleaned with methanol to remove surface oxides and other contaminants prior to

62 clamping. Butt joints were welded in the direction along the length of each coupon using a 4 kW CW Nd: YAG laser system (manufactured by TRUMPF, Germany) equipped with an ABB robot and magnetic holding fixture system. A collimation lens of 200 mm, focal lens of 150 mm, and a fiber diameter of 0.6 mm were employed to produce a laser beam with a spot diameter of approximately 0.45 mm. Titanium, in its molten state and at temperatures above 300°C, is reactive with most atmospheric gases such as oxygen, nitrogen, carbon, and hydrogen ^{8, 60}. Therefore, it is necessary to take adequate measures to shield the weld region until the weldment is cooled below the reactivity temperature.

High purity argon at a flow rate of 23.6 l/min was used to shield the top surface of the work-piece. The trail on the top surface and the bottom of the work-piece was shielded using helium at a flow rate of 66.1 l/min, which according to AWS C7.2 minimizes the HAZ size during welding of titanium alloys. The laser power (P) used was kept constant at 4 kW while the defocusing distance (∆z) and welding speed (v) were adjusted to obtain fully penetrated autogenous butt welds. Table 5 shows the processing parameters used.

After welding, the specimens were examined by radiography to detect any cracks and/or porosity in the weld regions. The surface quality of all laser welds was visually evaluated and macroscopically recorded using an Olympus SZ40 stereoscope. For the examination of the transverse section and microstructure of the welds, two metallurgical specimens were cut from each joint and then mounted using cold-setting epoxy resin. The specimens were prepared using standard metallographic techniques and final polishing was performed using 0.04 micron colloidal silica with 10% hydrogen peroxide to produce a mirror-like finish. Etching to reveal the microstructure was accomplished using Kroll’s

63 reagent (1-3 mL HF + 2-6 mL HNO3 + 100 mL H2O) for ~ 10 seconds. Microstructural examination was carried out using an inverted optical microscope (Olympus GX71) equipped with digital image analysis software (AnalySIS Five). The microhardness profiles across the welded joints were measured at a testing load of 500 g and a dwell time of 15 seconds using a Vickers microindentation machine (Struers Duramin A300) with an automated testing cycle. For each weld condition, three hardness profiles across the weld joint near the top, center, and root height of the joint were made with an indent interval of 0.3 mm. Three tensile specimens per weld having a standard sub-size geometry of 25 mm in gage length, 6 mm in width and 3.1 mm in thickness were machined in accordance with ASTM E8M-01. All specimens were tested at room temperature using a 250 kN MTS 810 tensile machine equipped with an Aramis 3D deformation measurement system. The Aramis system is a non-contact optical system that automatically measures strains along the gage length of the tensile sample during the testing. The system comprises of two CCD cameras capable of recoding 15 frames per second (fps), a trigger box and a high performance PC system. It is noteworthy that the functionality of the Aramis deformation system depends on the quality of the speckle

63 reagent (1-3 mL HF + 2-6 mL HNO3 + 100 mL H2O) for ~ 10 seconds. Microstructural examination was carried out using an inverted optical microscope (Olympus GX71) equipped with digital image analysis software (AnalySIS Five). The microhardness profiles across the welded joints were measured at a testing load of 500 g and a dwell time of 15 seconds using a Vickers microindentation machine (Struers Duramin A300) with an automated testing cycle. For each weld condition, three hardness profiles across the weld joint near the top, center, and root height of the joint were made with an indent interval of 0.3 mm. Three tensile specimens per weld having a standard sub-size geometry of 25 mm in gage length, 6 mm in width and 3.1 mm in thickness were machined in accordance with ASTM E8M-01. All specimens were tested at room temperature using a 250 kN MTS 810 tensile machine equipped with an Aramis 3D deformation measurement system. The Aramis system is a non-contact optical system that automatically measures strains along the gage length of the tensile sample during the testing. The system comprises of two CCD cameras capable of recoding 15 frames per second (fps), a trigger box and a high performance PC system. It is noteworthy that the functionality of the Aramis deformation system depends on the quality of the speckle