Effects of annealing on microstructure and properties of linear friction ...

Effects of annealing on microstructure and properties of linear friction welded dissimilar titanium joints T. J. Ma, X. Li, B. Zhong and W. Y. Li* Linear friction welds of Ti–6Al–4V (TC4, according to Chinese classification) to Ti–6?5Al–3?5Mo– 1?5Zr–0?3Si (TC11, according to Chinese classification) were subject to post-weld heat treatment (PWHT) at 650uC64 h (PWHT1) and at 950uC61 hz530uC64 h (PWHT2) under air cooling. In the as welded joint, TC4 had recrystallised sufficiently compared to TC11. After PWHT1, the a grains in the TC4 weld centre zone grew to some extent and many superfine equiaxed recrystallised a grains precipitated along b boundaries on the TC11 side. The grain growth along the weld line in the weld centre zone under PWHT2 was evident. PWHT2 reduced the joint tensile strength due to the coarsening of the microstructure. The joint microhardness decreased after PWHT. Keywords: Titanium alloys, Linear friction welding, Microstructure, Post-weld heat treatment, Mechanical properties

Introduction With increasingly demanding specifications for high thrust/weight ratio aeroengines, the integrally bladed disk (blisk) has been attracting more attention due to its simplified design, reduced number of components, lower weight and elimination of leakage flow. There are a number of manufacturing methods for making blisks, such as welding, milling from solid blocks and electrochemical machining (ECM). Compared to these methods welding allows blisk blades to be manufactured to fit stress conditions, reduce material use and allow blades to be replaced when needed.1 Linear friction welding (LFW) is a self-cleaning process as it removes surface contaminants2,3 with its continuous extrusion of a plasticized layer from the rubbing interface. The localized material yielding is caused by frictional heat, which forms a flash isolating the interface of the joint from outside air during oscillation. LFW does not generate the typical solidification defects found in fusion welding due to the absence of a liquid phase,4 making it suitable to use in a wide range of materials, such as steels, intermetallics, aluminium, copper, nickel and titanium alloys. Currently, LFW is the most attractive joining process for the production of a titanium alloy blisk,2,5 as it can be used to join dissimilar alloys together and benefit from their specific properties in a blade (exposure to high cycle fatigue loads under high temperatures) and disc (exposure to low cycle fatigue loads)6 to improve the performance of the assembly.

State Key Laboratory of Solidification Processing, Shaanxi Key Laboratory of Friction Welding Technologies, Northwestern Polytechnical University, Xi’an 710072, Shaanxi, PR China *Corresponding author, email [email protected]

Nevertheless, there are few papers on LFW of dissimilar alloys available. Researchers obtained Ti– 6Al–4V joints where they had been heat treated to correspond to typical aerospace blade and disk materials.2,7,8 However, these two materials showed slight differences in mechanical properties and had the same elastic modulus. In our previous study,9 LFW of the dissimilar titanium alloys joints between Ti–6Al–4V (TC4, according to Chinese classification) and Ti– 6?5Al–3?5Mo–1?5Zr–0?3Si (TC11, according to Chinese classification) was conducted for their wide application in aeroengines. It was found that the as welded dissimilar material joint has a higher strength than the parent TC4.9 Previous experiments with LFW of Ti–5Al–2Sn–2Zr– 4Mo–4Cr (Ti17, according to Chinese classification) revealed that post-weld heat treatment (PWHT) affected positively the overall properties of the joint.10 Therefore, in this work, PWHT was also applied to dissimilar TC4 to TC11 joints under different conditions to improve the microstructure and mechanical properties of the joint.

Materials and experimental procedures The hot rolled TC4 and the TC11 annealed at 910uC after forging were received and machined to rectangular blocks with nominal dimensions of 15618645 mm (with welding interface 18615 mm) for the following LFW experiment. As shown in Figs. 1a and d, the microstructure of the TC4 and TC11 parent metals consists of a bimodal azb structure showing prior a grains in a transformed b matrix (dark grey). TC4 has elongated a grains, while TC11 presents equiaxed a grains. The lab scale LFW machine (type XMH-160) built at the Northwestern Polytechnical University (China) was used for welding. The process parameters of welding were selected based on our preliminary experimental work11,12 with a forging pressure of 110 MPa. One of the welding joints was used for microstructure examination and

ß 2014 Institute of Materials, Minerals and Mining Published by Maney on behalf of the Institute Received 23 May 2014; accepted 10 August 2014 DOI 10.1179/1362171814Y.0000000243

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Effects of annealing on LFWed dissimilar titanium joints

1 Optical micrographs of a–c TC4 and d–f TC11 parent materials in a, d as received, b, e PWHT1 and c, f PWHT2 conditions

mechanical properties measurement in the as welded condition. The others were heat treated under two different heat treatment conditions, i.e. at 650uC64 h followed by air cooling (referred to as PWHT1) and at 950uC61 hz530uC64 h followed by air cooling (referred to as PWHT2), as in literature.10,13 Both PWHT treatments were performed in a vacuum resistance furnace. The joints were cut as illustrated in Fig. 2b for metallographic specimens and tensile specimens, and the sizes of the tensile specimen are shown in Fig. 2c.

Results and discussion Microstructure characterisation Overall view

Figure 2a shows the photo of a typical linear friction welded TC4 to TC11 joint, which shows appreciable

flash from all sides of the welding interface and is similar to same material joints.7,11,14–,15,16 Figure 3a shows the overall view of cross-sectional microstructure of the joint along the direction of oscillation. The narrow weld zone can be clearly identified, with its width gradually increasing from about 190 mm in the middle to about 880 mm at the edge, indicating non-uniform temperature and stress fields at the weld zone during LFW. Zone A in Fig. 3a in the as welded condition exhibits different microstructural regions on both TC4 and TC11 sides, i.e. the weld center zones, thermomechanically affected zones (TMAZs) and parent metals as shown in Fig. 3b. This microstructural feature is different from what was observed in the similar TC47,11,12,14,15 or TC1116 joint due to the presence of a distinct weld line. It can be observed that the width of the weld zone and TMAZ of TC4 are larger than those of TC11, which can be attributed to the different thermal and mechanical

2 a photo of welded joint, b drawing of samples and c dimensions of tensile test specimen


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3 Optical micrographs of joint cross-section: a overall view; b–d high magnifications of zone A in Fig. 3a in as welded, PWHT1 and PWHT2 conditions respectively; e–g high magnifications of zone B in Fig. 3a in as welded, PWHT1 and PWHT2 conditions respectively

properties of TC4 and TC11 under rapid heating (of the order of 1000uC s21) and cooling (of the order of 250uC s21)17 conditions that prevail in LFW. As shown in Fig. 3c, the joint microstructure heat treated under PWHT1 shows no differences from the as welded condition, while the joint microstructure after PWHT2 (Fig. 3d) has been improved with the elimination of flow lines on the TC4 side and fully recrystallised grains at the weld zone, which will be discussed in the following sections. Weld centre zone

Figure 4a shows the microstructure of the weld centre zone in the as welded condition, exhibiting superfine a grains in the b matrix on the TC4 side. It has been reported that forming superfine grains in the TC4 weld was associated with recrystallisation and rapid heating and cooling during the process.7,11,12,15 Nucleation and growth of dynamic recrystallisation grains were accompanied by deformation, elongating some of the a grains along the oscillation direction as marked in Fig. 4a. Owing to insufficient recrystallisation, the metastable b phase appeared and a phase precipitated along b boundaries is limited on the TC11 side as shown in Fig. 4b, for the material is stable up to high temperatures. According to literature, it is believed

that impurity atoms tend to segregate at recrystallised grain boundaries, while the mobility of the grain boundaries could be diminished by the interaction between impurity atoms and the grain boundaries. This results in a decrease of the recrystallisation rate.18 Therefore, more alloying elements in TC11 would limit recrystallisation during the deformation process. Residual strain energy may drive static recovery and recrystallisation with the assistance of PWHT. After PWHT1, the a grains in the TC4 weld centre zone grew to a certain extent, while many superfine equiaxed recrystallised a grains precipitated along b boundaries (Fig. 4c) and the metastable b phase was decomposed to form secondary a phase (Fig. 4d) on the TC11 side. These show that the effect of PWHT1 on the weld centre zone is still limited, since the b-transus of TC4 and TC11 are 995 and 1020uC respectively,19 which are much higher than the heat treatment temperature. Study of the microstructure of the weld centre zone after PWHT2 (Fig. 4e) shows grain growth with apparent grain boundaries compared to the as welded microstructure (Fig. 4a) and with the direction of grain growth along the weld line (Fig. 4f). During the PWHT2 process, the initial elongated a grains in both TC4 and TC11 weld centre zone tended to grow without


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4 Micrographs of weld centre zones: a in as welded condition by SEM, b high magnification of weld line in Fig. 4a by SEM, c in PWHT1 condition by SEM, d high magnification of TC11 weld centre zone in Fig. 4c by SEM, e in PWHT2 condition by OM and f high magnification of the weld line in Fig. 4e

nucleation by annexing fine structures around them. a grain boundaries constantly migrated towards b phase, thus forming a texture along the oscillation direction. The previous work10 on LFW of Ti17 revealed that the impact toughness of the weld zone after PWHT could reduce by 6?7%, which may be associated to the preferred orientation microstructure in the weld zone. In addition, it was found that acicular martensite a9 is present in the outer region of the TC4 weld zone as shown in Fig. 3e. During the LFW process, the plastically deformed material next to the weld line was exposed to the highest temperature of the weld while continuously being expelled from the weld interface. In

the forging phase, most of the high temperature plastic material was extruded into the flash, which cooled quickly after welding, forming therefore martensite, supporting the theory that temperatures at the interface exceeded b-transus during LFW. It should be pointed out that previous studies have also supported the hypothesis that the weld centre zone experiences supra b-transus temperatures (.990uC) as there had been observed martensitic or Widmannstatten microstructures in LFWed a or azb Ti alloys.7,8,15 However, after PWHT1, martensite a9 had transformed to a and b phases as shown in Fig. 3f. Moreover, a and b grains had clearly grown up after PWHT2 as shown in Fig. 3g.


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5 Micrographs of a–d TMAZs in as welded, e–g PWHT1, and h–j PWHT2 conditions: a–d correspond to zones C, D, E, F in Fig. 3b respectively, e and f correspond to TMAZs near weld centre zone and near parent TC4 respectively

TMAZ

In the TMAZs on both TC4 and TC11 sides, the appreciable grain elongation shows the flow direction of plasticised material as shown in Fig. 3b. The detailed observation of the TMAZs on both sides indicates the distinct features as shown in Figs. 5a–d. A small number of recrystallised grains can be found in the TMAZ of TC4 (Fig. 5a) in the as welded condition. On the TC11 side, the recrystallised grains are not present and the elongation of grains is pronounced (Fig. 5c). This observation is explained in the same manner as the TC11 weld centre zone. In the TMAZs closer to both parent metals there is appreciable deformation as shown in Fig. 5b and d. In addition, in Fig. 5d, the extent of

deformation of b is found to be more intense than that of a, due to the difference of crystal structure between the two phases. As a result of heat treating the joint with PWHT1, the dynamically recrystallised a formed during LFW grew up to a certain degree and the lamellar a was formed in b in the TMAZ of TC4 as shown in Fig. 5e and f. While in the TMAZ of TC11, the recrystallised globular a precipitated on the grain boundaries or in b during heat treatment (Fig. 5g). It is expected that elevated heat treatment temperature or prolonged exposure time will let the recrystallised a grains on the grain boundaries be absorbed by the adjacent coarse primary a grains, while those in the b phase would become coarse as the joint


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6 Fractured tensile specimens in a as welded, b PWHT1, and c PWHT2 conditions

microstructure under PWHT2 shows (Figs. 5j). It should be pointed out that some of the a grains are found to be larger than those of TC11 parent metal (Fig. 5i), as high strain energy stored in the TMAZ promotes the fine recrystallised grains to be absorbed by primary a grains. Clearly coarse and spheroidal grains in the TC4 TMAZ under PWHT2 are apparent as shown in Fig. 5h. Microstructure of parent metal

The micrographs of the TC4 and TC11 parent metals in the PWHT1 and PWHT2 conditions are also given in Fig. 1. As shown in Figs. 1b and e, the microstructure of TC4 and TC11 parent metals was not affected by PWHT1. PWHT2 did affect the microstructure as the clearly coarse primary a and secondary a in transformed b in the TC11 side shows, and even the elimination of the original elongated features due to rolling and the formation of equiaxed microstructure in the TC4 side also shows (Figs. 1c and f). Therefore, a higher heat treatment temperature will change unexpectedly the microstructure of parent metals.

Mechanical properties The average tensile strengths of the specimens in the as welded, PWHT1 and PWHT2 conditions are 1013¡48, 990¡30 and 947¡12 MPa respectively. All specimens failed in the TC4 parent metal far away from the weld as shown in Fig. 6, which shows that the tensile strength of the joint is actually that of parent TC4 for those three conditions, while it needs further experiment to test the properties of the joint itself. With a higher heat treatment temperature, the microstructure of the weld zone and TMAZs of the joints is modified. The joint tensile strength gets lower after PWHT due to the coarsening of the TC4 parent microstructure. Figure 7 displays the change of microhardness of the joints under the three different conditions in the direction which is perpendicular to the welding interface. In all cases the microhardness of the joints has the same tendency to decrease from the interface to parent metals on both TC4 and TC11 sides. On the TC4 side, the weld zone has the highest microhardness because of the formation of superfine grains.11 On the TC11 side, although recrystallisation is insufficient, there exist large deformations with a high dislocation density which does increase hardness. As expected, after PWHT, the microhardness of the joints decreased due to further recovery and recrystallisation which reduces the high dislocation density. The two PWHTs used in this work were found to reduce peak microhardness by about 150 HV. In addition, it can be seen that the joints in both PWHT conditions have very similar profiles, and the microhardness of the joints, except for the weld zone, did not fluctuate. It indicates that grain growth in PWHT2 condition has

7 Microhardness of joints across weld line under different conditions

little effect on microhardness compared to PWHT1 condition.

Conclusions Based on the results from this study, the following conclusions can be drawn. 1. After post-weld heat treatment at 650uC64 h under air cooling (PWHT1), the a grains in the TC4 weld centre zone grew to some extent, while many superfine equiaxed recrystallised a grains precipitated along b boundaries and the metastable b phase was decomposed to form secondary a phase in the TC11 side. In the thermomechanically affected zone of TC4, the dynamically recrystallised a formed during LFW grew to a certain degree and the lamellar a was formed in b, while in that of TC11 the recrystallised globular a precipitated on the grain boundaries or in b during heat treatment. 2. After post-weld heat treatment at 950uC61 h z530uC64 under air cooling (PWHT2), flow lines on the TC4 side were eliminated. On the TC11 side, the direction of grain growth is parallel to the weld line in the weld centre zone, and some a grains in the thermomechanically affected zone became even larger than those of parent TC11. 3. All the tensile specimens failed in the TC4 parent material far away from the weld, which indicates that the tensile strength of the weld is higher than that of parent TC4. The decrease of the average tensile strength of the joint after PWHT2 is caused by the coarser microstructure of the parent TC4 metal. 4. Post-weld heat treatment in this work reduced the microhardness of the joints due to further recovery and recrystallisation. The joints in both heat treatment conditions have very similar profiles with an almost constant value except the weld zone.

Acknowledgements The authors would like to thank for financial support the Fok Ying-Tong Education Foundation for Young Teachers in the Higher Education Institutions of China (grant no. 131052), the 111 Project (grant no. B08040), the Fundamental Research Funds for the Central Universities (grant no. JC02010404) and the Research


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Fund of the State Key Laboratory of Solidification Processing (NPU, China) (grant no. 108-QP-2014).

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