Cyclic deformation behavior of linear friction welded Ti6Al4V joints

Author's Accepted Manuscript

Cyclic deformation behavior of linear friction welded Ti6Al4V joints G.D. Wen, T.J. Ma, W.Y. Li, J.L. Li, H.Z. Guo, D.L. Chen

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S0921-5093(14)00023-9 http://dx.doi.org/10.1016/j.msea.2014.01.006 MSA30662

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Materials Science & Engineering A

Received date: 16 September 2013 Revised date: 29 December 2013 Accepted date: 4 January 2014 Cite this article as: G.D. Wen, T.J. Ma, W.Y. Li, J.L. Li, H.Z. Guo, D.L. Chen, Cyclic deformation behavior of linear friction welded Ti6Al4V joints, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2014.01.006 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Cyclic deformation behavior of linear friction welded Ti6Al4V joints G.D. Wen1, 2, T.J. Ma1, W.Y. Li1*, J.L. Li1, H.Z. Guo1, D.L. Chen2 1

State Key Laboratory of Solidification Processing, Shaanxi Key Laboratory of Friction

Welding Technologies, Northwestern Polytechnical University, 127 Youyi Road, Xi’an 710072, Shaanxi, PR China,

2

Department of Mechanical and Industrial Engineering, Ryerson University, 350 Victoria Street, Toronto, Ontario M5B 2K3, Canada

Abstract

This study was aimed at characterizing microstructural change and fatigue properties of linear friction welded (LFWed) Ti6Al4V joints with special attention to the relationship between the microstructure and cyclic deformation behavior. The welding process resulted in a remarkable microstructure change across the LFWed joint. The microstructure of the joint consisted of fine subgrains in the weld zone (WZ) and elongated grains in the thermomechanically affected zone (TMAZ), leading to a significantly higher hardness in the WZ. Both the base metal (BM) and joint exhibited essentially symmetrical hysteresis loops and equivalent fatigue life, while the cyclic strain hardening exponent of the welded joint was lower. Cyclic stabilization appeared at lower strain amplitudes up to 0.6% in both the BM and joint, however, cyclic softening occurred at higher strain amplitudes. Fatigue failure was observed to occur in the BM, and fatigue crack

*

Corresponding author. Tel: +86 29 88495226; Fax: +86 29 88492642. E-mail address: [email protected] (W.Y. Li);

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initiated from the specimen surface or near-surface defect. Fatigue crack propagation was basically characterized by fatigue striations, together with some secondary cracks.

Keywords: Titanium alloys; Linear friction welding; Low cycle fatigue; Microstructure.

1.

Introduction

Titanium alloys have been successfully used in many critical structural applications in the aerospace industry, because of their low density, high strength, and excellent fatigue and superior corrosion resistance [1-3]. Among titanium alloys, Ti6Al4V has a bimodal microstructure, which is essentially ductile [4] and has been widely used in low- and high-temperature applications [5]. There are many welding methods to join titanium alloys, such as electron beam welding (EBW) [6], laser beam welding (LBW) [7], linear friction welding (LFW) [8], friction stir welding (FSW) [9], gas tungsten arc welding (GTAW) [10] or tungsten inert gas (TIG) welding [11], among which LFW has drawn particular attention due to its advantages. It is an advanced process involving metallurgically sound solid-state joining through the reciprocating movement of one component relative to another under an axial force to generate frictional heat input. LFW has been used to join non-symmetric components on their flat faces, e.g., a turbine blade can be bonded to a turbine disk by simply rubbing the blade on the disk to form a blisk (integrally bladed disk) [8]. To date, most of the LFW development has been driven by the desire of aeroengine industry’s to fabricate integrally bladed titanium alloy disks giving a lower weight and to improve the performance over existing slotted blade/disk assemblies. In addition, as a solid-state process, LFW can eliminate solidification defects, e.g., air hole associated with the conventional

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fusion welding, thus providing defect-free welds having good service properties. At the same time, narrow welding zone and fine grains can be achieved due to the severe deformation during LFW [12]. Therefore, LFW has been acknowledged as a key technology of aero-engine blisk manufacture and maintenance. In many applications, structures with welded joints often inevitably involve dynamic/alternating loading, which would result in the occurrence of worrisome fatigue failure. Therefore, it is necessary to evaluate the microstructure, tensile properties and fatigue resistance of the linear friction welded (LFWed) joints so as to guarantee the safety and durability of aircraft.

Extensive studies on the LFWed titanium alloys have been reported, including the microstructure, hardness, tensile strength, and impact toughness by changing welding parameters and via finite element simulations [12-21]. For example, Li et al. [12] and Wanjara et al. [13] studied the microstructure and hardness of the LFWed Ti6Al4V joint. Chen et al. [14] investigated the formation mechanism of fine grains in the LFWed Ti6Al4V joint. Vairis et al. [15,16] focused on the effect of power and strain rate in the LFW process for Ti6Al4V alloy. Frankel et al. [17] and Preuss et al. [18] reported the effect of heat treatment on the residual stresses of the LFWed titanium alloy joints. Ma et al. [21] examined the impact toughness of the LFWed Ti6Al4V joint. However, studies on the stress-controlled or strain-controlled low cycle fatigue behavior of titanium joints remain quite limited to date [22-27]. Mohandas et al. [22] studied the stresscontrolled low cycle fatigue (LCF) behavior of friction welds and electron beam welds of the α-β titanium alloy (Ti-6.5Al-3.3Mo-1.6Zr-0.3Si) at ambient temperature, and reported that the friction welds exhibit substantially superior LCF behavior as compared to EBWed joints. Fu et al. [25] reported strain controlled low cycle fatigue tests for EBWed Ti6Al4V titanium alloy joint

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with triangular waveform, strain ratio of Rε=-1, and strain amplitude from 0.25% to 0.6%. Wang et al. [26,27] reported the cyclic deformation of EBWed dissimilar titanium alloy joint with triangular waveform, strain ratio of Rε=-1, at a constant strain rate of 1×10-2 s-1, and strain amplitude from 0.2% to 1.2%. To the authors’ knowledge, no information on the straincontrolled fatigue behavior of LFWed Ti6Al4V joints is available in the open literature. It is unknown what the effect of LFW on the LCF life of the Ti6Al4V alloy is, if the cyclic hardening or softening appears, and where (in the weld zone (WZ) or base metal (BM)) the failure occurs. The objectives of the present study were, therefore, to evaluate the LCF behavior of LFWed Ti6Al4V joints in relation to the microstructural characteristics.

2.

Material and Experimental Procedure

The material used in the present study was a forged Ti6Al4V titanium alloy with a chemical composition (in wt.%) of 6 Al, 4 V, 0.3 Fe, 0.1 C, 0.05 N, 0.015 H, 0.2 O, and balance Ti. Specimens were machined into blocks with dimensions of 70×18×11 mm, where the weld interface was 18×11mm with the oscillation along the direction of 18 mm long. The welding parameters applied were set as the optimal values according to our previous studies: a frequency of oscillation 35 Hz, an oscillation amplitude of 3.5 mm, a friction time of 4s and a friction pressure of 45 MPa.

Metallographic samples were cut from the LFWed joint perpendicular to the oscillation direction, then ground, polished and etched using Keller’s reagent (12 ml HF, 36 ml HNO3 and 42 ml H2O). Microstructure examinations were performed via optical microscope and scanning electron

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microscope (SEM) (JSM-6380LV) having energy-dispersive X-ray spectroscopy (EDS) and three-dimensional (3D) fractographic analysis capacity. Microhardness was determined across the weld using a computerized Buehler hardness tester with a load of 500g and a dwell time of 15s at an interval of 0.025 mm in WZ. Fatigue specimens with a gauge length of 12 mm and a width of 3 mm were machined perpendicularly to the oscillation direction using electrodischarge machining (EDM). The gauge area was ground by using SiC papers up to #600 to remove the EDM cutting marks and other surface irregularities so as to achieve an equivalent surface condition. Total strain-controlled pull-push type of fatigue tests was performed in air at room temperature using a computerized Instron 8801 fatigue testing system at different strain amplitudes up to 1.2%. A triangular waveform with a strain ratio of R =-1 was applied at a constant strain rate of 1×10-2 s-1, where R is the minimum peak strain divided by the maximum peak strain in the strain-controlled fatigue tests. The strain-controlled testing at low strain amplitudes was carried on until 10,000 cycles, after which it was changed to load control at 50 Hz. At least two specimens were tested at each strain amplitude. Fatigue crack initiation site and crack propagation mechanisms were examined on the fracture surfaces of failed samples via SEM.

3.

Results and Discussion

3.1

Microstructure

Fig. 1 shows the typical appearance of the LFWed joint. It is seen that the flash has extruded on all four sides due to the material softened at elevated temperatures arising from the friction heat

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and under a fairly high pressure of 45 MPa during LFW. The microstructure of the BM is shown in Fig. 2. It is seen that the BM consists of a combination of equiaxed α grains and inter-granular α+β lamellae. After LFW, a remarkable microstructural change occurred in the WZ and thermomechanically affected zone (TMAZ) as shown in Fig. 3(a). It can be clearly seen that the width of WZ is about 500 µm, and is composed of mainly subgrain structures, ranging from approximately 2-5 µm wide and 10-15 µm long as shown in Fig. 3(a) and (b). In addition, the dispersed fine recrystallized globular α grains in the WZ were observed as indicated by arrows in Fig. 3(b), which are associated with the dynamic recrystallization and the rapid heating and cooling during LFW [8,12,13]. Similar results have been reported by Li et al. [12] and Wanjara et al. [13]. The TMAZ is visibly narrower with elongated grains as shown in Fig. 3(a), which suggests the presence of large plastic deformation in TMAZ, and the plastic strain increases with increasing distance from the BM to TMAZ. Furthermore, a small amount of recrystallized grains can be seen inside of the elongated grains in the TMAZ (Fig. 3(c)).

3.2

Microhardness

Vickers microhardness profile across the LFWed joint is shown in Fig. 4. It is seen that a characteristic symmetrical hardness profile across the weld was obtained with an average hardness of about 340 HV for the BM. The highest hardness appeared in the WZ, which is attributed to the considerably finer grains stemming from the occurrence of dynamic recrystallization during LFW, as shown in Fig. 3(a). Generally, the grain size dependence of the strength or hardness in polycrystals could be described in terms to Hell-Petch type relationship ( σ = σ o + kd −1 / 2 ) [28-30], where σ is the yield strength, σ o is the friction stress for the

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movement of dislocation, and d is the average grain size. The slope k is sometimes referred to as the “Hell-Petch strength coefficient”. In this study, the grain size in the WZ is considerably smaller than that of the BM as shown in Fig. 3(a). Therefore, the highest Vickers hardness in WZ is expected based on the Hell-Petch relationship. The similar result about effect of grain size on the hardness of LFWed Ti6Al4V joint was reported by Li et al. [12]. In addition to the effect of grain size, it has also been reported that the presence of strain hardening enhanced the Vickers hardness in titanium alloys [31,32], which is attributed to the increase in the density of dislocations generated by the severe plastic deformation during LFW.

3.3

Hysteresis loops

Fig. 5 shows the hysteresis loops of the first cycle and the mid-life cycle at a total strain amplitude of 1.2%. It is seen that the first and mid-life hysteresis loops of both BM and joint are basically symmetrical. Similar symmetrical hysteresis loops was also observed in cast or semisolid processed magnesium alloys [33,34], rare-earth element containing extruded magnesium alloys [35,36], TiAl intermetallic compound [37], titanium alloys [38,39], and in EBWed titanium alloy joints [25-27]. This is, however, in sharp contrast to the hysteresis loops of the extruded or rolled Mg alloys due to the presence of strong crystallographic texture [40-42]. Furthermore, unlike the case of the EBWed titanium alloy joint [26], the Young’s modulus of the present LFWed titanium joint was almost the same as that of BM as shown in Fig. 5(a), which is attributed to the absence of coarse and soft β phase and martensite α′′ [26,43-46].

3.4

Cyclic deformation response

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Fig. 6 shows the evolution of cyclic stress amplitude as a function of the number of cycles at different strain amplitudes for the Ti6Al4V BM and LFWed joint. It is seen that with increasing total strain amplitude applied, the cyclic stress amplitude increases and fatigue life decreases in both BM and joint. Similar cyclic deformation characteristics have been reported for titanium alloy under strain-controlled low cycle fatigue tests [26,27,47]. In addition, at lower strain amplitudes (0.2-0.6%) both BM and joint exhibited basically constant or stable stress amplitude. As the strain amplitude increases (0.8-1.2%) cyclic softening occurred in both BM and joint, which corresponds well to the increasing plastic strain amplitude (Fig. 7). At the same time, a tendency of initial cyclic stabilization appeared in both BM and joint at the higher strain amplitude of 0.8-1.2% as shown in Fig. 6. As seen from Figs. 6 and 7, as the strain amplitude decreases the cyclic stabilization stage becomes increasingly longer, or the extent of cyclic softening decreases, which may be attributed to the rearrangement and partial annihilation of high density dislocations during the high strain cyclic deformation [48]. Indeed, a strain amplitude of 0.6% appears to be a critical value, at or below which cyclic stabilization or saturation remains in the entire cyclic deformation process, which has also been observed in the EBWed titanium alloy joint [26].

3.5

Fatigue life and fatigue parameters

The fatigue life (i.e., the number of cycles to failure, Nf) as a function of the applied total strain amplitude ( Δε t / 2 ) of the joint is plotted in Fig. 8, in conjunction with the experimental data reported in the literature [25,26,49-51]. All the titanium alloys show a similar trend of increasing

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fatigue life with decreasing strain amplitude. Especially, the fatigue life of LFWed titanium joint is longer than that of EBWed Ti6Al4V joint [25], which indicates that the welding parameters selected in the present study are more appropriate and the quality of the LFWed joints is sound. Cyclic deformation behavior is normally considered to be related to the portion of the plastic strain amplitude. This relationship can be express as: n′

⎛ Δε ⎞ Δσ = K ′⎜⎜ p ⎟⎟ 2 ⎝ 2 ⎠ ,

where

(1)

Δε p Δσ is the mid-life stress amplitude, is the mid-life plastic strain amplitude, n' is the 2 2

cyclic strain-hardening exponent and K' is the cyclic strength coefficient. Based on Basquin equation and Coffin-Manson relation, the total strain amplitude could be expressed as elastic strain amplitude and plastic strain amplitude [40,41,52], i.e., Δε t Δε e Δε p σ ′f (2 N f ) c = + = + ε ′f (2 N f ) 2 2 2 E , b

(2)

where E is the Young’s modulus, Nf is the fatigue life or number of cycles to failure (the term of 2Nf is referred to as the number of reversals to failure), σ ′f is the fatigue strength coefficient, b is the fatigue strength exponent, ε ′f is the fatigue ductility coefficient, and c is the fatigue ductility exponent. Fig. 9 shows the elastic, plastic, and total strain amplitudes plotted as a function of the number of reversals to failure taken from the mid-life cycles in a double-log scale. It is seen that with increasing strain amplitude, all the amplitudes increase linearly and the fatigue life decreases in this plot. However, the plastic strain amplitude shows a steeper line than the other two amplitudes. The fatigue life parameters evaluated on the basis of Eqs (1) and (2) are summarized in Table 1 along with the results reported in the literature [48,53-55]. It is seen

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that the obtained fatigue parameters are well within the range of other titanium alloys reported in the literature.

3.6

Fractography

The failure of all the LFWed joint samples was observed to occur in the BM during fatigue tests. A typical macroscopic image is shown in Fig. 10, where the failure location is about 3 mm away from the WZ. This is due to the fact that the LFW led to a much higher hardness in WZ as shown in Fig. 4. In addition, unlike the EBWed Ti6Al4V/Ti17 dissimilar joints [26] and fiber laser welded DP980 dual-phase steel [56,57] or diode laser welded DP600 dual-phase steel [58-60], no soft zone in the TMAZ or heat-affected zone was present in the present LFW. This suggests that the welding parameters selected in the present study are appropriate and the quality of LFWed joints is sound. Fracture surfaces of the fatigued specimens were also examined using SEM. Fig. 11 shows an overall view of fracture surfaces of the LFWed joint tested at different strain amplitudes, i.e., 0.4% and 1.0%, containing fatigue crack initiation, propagation, and final fast fracture regions. It is seen from these low magnification images that fatigue crack initiated from the specimen surface or near-surface defect, and the river line patterns appeared in the joint which are irregular and broken and flow along the crack propagation direction. Furthermore, the size of the crack propagation area is mainly associated with the applied strain amplitude in the joint. The higher the applied strain amplitude is, the smaller the propagation area is. As shown in Fig. 11(a), the propagation area of samples tested at a lower strain amplitude of 0.4% has a larger propagation area, compared with the samples tested at a higher strain amplitude of 1.0% (Fig. 11(b)), due to the lower peak stress. Multiple fatigue crack initiation sites were observed at a

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higher strain amplitude of 1.0% as shown in Fig. 11. Similar multiple crack initiation sites in a fiber laser welded AZ31B-H24 magnesium alloy were also observed at a higher stress level in [61,62]. Fig. 12 shows typical SEM micrographs taken near the crack initiation sites (the red dashed box indicated in Fig. 11) and in the crack propagation zone of the fatigued samples tested at strain amplitudes of 0.4% and 1.0% at a higher magnification. It is seen that a mix of cleavage-like features and fatigue striations appears on the fracture surface in the near-initiation area. Fatigue crack propagation is basically characterized by fatigue striations which are perpendicular to the crack propagation direction, in conjunction with some secondary cracks. A better view of the characteristic fatigue striations in the propagation area on the fracture surface of the joint fatigued at a strain amplitude of 0.4% could be seen from a three-dimensional image as shown in Fig. 13, where some secondary cracks are also visible, as indicated by arrows in Fig. 13. It is known that the fatigue striations normally occur by a repeated plastic bluntingsharpening process in face-centered cubic (fcc) materials arising from the slip of dislocations in the plastic zone ahead of the fatigue crack tip [63]. The formation of the fatigue striations in the hexagonal close-packed (hcp) titanium alloy is expected to be related to both dislocation slip and twinning in the plastic zone during fatigue crack propagation [40,41]. Further studies in this aspect are needed.

4.

Conclusions

Strain-controlled low cycle fatigue tests were conducted on LFWed Ti6Al4V titanium alloy joints, and cyclic deformation characteristics, fatigue parameters and fracture mode of the joints were evaluated. The following conclusions can be drawn from this investigation:

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1. The microstructure across the LFWed joint exhibited a marked change, mainly consisting of fine subgrains in the WZ, and elongated α in TMAZ. 2. A symmetrical microhardness profile across the welded joint was observed with a highest value in the WZ stemming from the formation of fine subgrains. 3. The hysteresis loops of both BM and joint were essentially symmetrical, unlike those of extruded magnesium alloys also with a hcp crystal structure reported in the literature. 4. During cyclic deformation, both Ti6Al4V BM and welded joint exhibited cyclic stabilization behavior in the entire deformation process at lower strain amplitudes up to 0.6%, while cyclic softening occurred after initial cyclic stabilization at higher strain amplitudes. The stage of initial cyclic stabilization was observed to be gradually shortened with increasing strain amplitude at high strain amplitudes. 5. The strain-controlled fatigue resistance of both BM and joint was observed to be equivalent within the experimental scatter, suggesting that sound LFWed joints were made with proper welding parameters. 6. The fatigue failure of the LFWed joint occurred in the BM. Fatigue crack initiated from the specimen surface or near-surface defect. Fatigue crack propagation was basically characterized by the characteristic fatigue striations along with some secondary cracks.

Acknowledgements

The authors would like to thank the financial support received from the National Natural Science Foundation of China (51005180), the Fok Ying-Tong Education Foundation for Young Teachers

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in the Higher Education Institutions of China (131052), the Fundamental Research Fund of Northwestern Polytechnical University (JC201233) and the 111 Project (B08040). One of the authors (D.L. Chen) is also grateful for the financial support by the Natural Sciences and Engineering Research Council of Canada (NSERC), Premier’s Research Excellence Award (PREA), NSERC-Discovery Accelerator Supplement (DAS) Award, Canada Foundation for Innovation (CFI), and Ryerson Research Chair (RRC) program. The authors would also like to thank Messrs. Q. Li, A. Machin, J. Amankrah and R. Churaman for easy access to the laboratory facilities of Ryerson University and their assistance in the experiments.

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Table Captions

Table 1 Fatigue parameters for the BM and LFWed joint in comparison with those of other titanium alloys reported in the literature.

Figure Captions

Fig. 1

Typical appearance of the specimens after linear friction welding.

Fig. 2

Microstructure of Ti6Al4V BM, (a) OM image, and (b) SEM micrograph.

Fig. 3

Microstructure change across a LFWed joint, (a) overall view of the cross-section, (b) WZ of the joint, and (c) TMAZ of the joint.

Fig. 4

Profile of microhardness across a LFWed joint.

Fig. 5

Typical stress-strain hysteresis loops of the (a) first and (b) mid-life cycle at a total strain amplitude of 1.2% and strain ratio of Rε=-1 for the Ti6Al4V BM and LFWed joint, respectively.

Fig. 6

Stress amplitude as a function of the number of cycles at different total strain amplitudes, (a) Ti6Al4V BM and (b) LFWed joint.

Fig. 7

Plastic strain amplitude vs. the number of cycles at different total strain amplitudes, (a) Ti6Al4V BM and (b) LFWed joint.

17

Fig. 8

Total strain amplitude as a function of the number of cycles to failure for the Ti6Al4V and LFWed joint, in comparison with the data reported in the literature for various Ti alloys.

Fig. 9

Strain amplitudes at the mid-life as a function of the number of reversals to failure obtained for the LFWed joint tested at room temperature.

Fig. 10

An overall view of fracture location in a LFWed joint.

Fig. 11

An overall view of fracture surfaces of the LFWed joint fatigued at different strain amplitudes, (a) 0.4% and (b) 1.0%.

Fig. 12

SEM images of fracture surfaces of the LFWed joint fatigued at different strain amplitudes of 0.4% ((a) and (b)) and 1.0% ((c) and (d)), where (b) and (d) were taken at a high magnification of (a) and (c), respectively.

Fig. 13

A typical three-dimensional image taken in the propagation area on the fracture surface of the LFWed joint fatigued at a strain amplitude of 0.4%.

Table 1. Fatigue parameters for the BM and LFWed joint in comparison with those of other titanium alloys reported in the literature.

Cyclic strain hardening exponent, n'

Cyclic strength coefficient, K' , MPa

Fatigue strength coefficient, σ ′f MPa

Fatigue strength exponent, b

795

0.11

1556

1342

800

0.11

1553

1267

Cyclic yield strength, σ ′y , MPa

Ti-6Al-4V BM LFWed joints

Fatigue parameters

18

Fatigue ductility coefficient,

Fatigue ductility exponent, c

-0.08

0.18

-0.67

-0.08

0.12

-0.67

ε ′f

Ti-13Nb-13Zr arc melted [53] LT26A asreceived [48] LT26A β treated [48] IMI834 ST-A [54] Ti-27Al-15Nb ST [55]

-

-

-

1154

-0.082

0.913

-0.723

-

-

-

1400

-0.073

0.066

-0.487

-

-

-

1440

-0.07

0.465

-0.83

850

-

-

1200

-0.04

0.049

-0.8

-

-

-

766.7

-0.03

0.004

-0.132

19

Figure(s)

Table 1. Fatigue parameters for the BM and LFWed joint in comparison with those of other titanium alloys reported in the literature.

Fatigue parameters

Cyclic yield strength,

 y , MPa

Cyclic strain hardening exponent, n'

Cyclic strength coefficient, K' , MPa

Fatigue strength coefficient,

 f MPa

Fatigue strength exponent, b

Fatigue ductility coefficient,

 f

Fatigue ductility exponent, c

Ti-6Al-4V BM

795

0.11

1556

1342

-0.08

0.18

-0.67

LFWed joints

800

0.11

1553

1267

-0.08

0.12

-0.67

-

-

-

1154

-0.082

0.913

-0.723

-

-

-

1400

-0.073

0.066

-0.487

-

-

-

1440

-0.07

0.465

-0.83

850

-

-

1200

-0.04

0.049

-0.8

-

-

-

766.7

-0.03

0.004

-0.132

Ti-13Nb-13Zr arc melted [53] LT26A as-received [48] LT26A β treated [48] IMI834 ST-A [54] Ti-27Al-15Nb ST [55]

10mm

(a)

(b)

β





β 100µm 40µm

BM

TMAZ

TMAZ

WZ

BM

(a)

c

b

200 µm

(b)

(c)

450 BM

TMAZ

WZ

TMAZ

BM

Hardness, HV

400

350

300

250

-8

-6 -4 -2 0 2 4 6 Distance from the weld centreline, mm

8

1200

(a)

800 Stress, MPa

400 0 -400 -800

Ti6Al4V BM first cycle Joint first cycle

-1200 -1.5 1200

-1

-0.5

0 0.5 Strain, %

1

1.5

(b)

Stress, MPa

800 400 0

-400 -800

Ti6Al4V BM mid-life cycle Joint mid-life cycle

-1200 -1.5

-1

-0.5

0 0.5 Strain, %

1

1.5

Stress amplitude, MPa

1000

(a)

800 600

1.2% 1.0% 0.8% 0.6% 0.4% 0.2%

400 200 0 1E+0

1000 Stress amplitude, MPa

Ti6Al4V BM

(b)

1E+1 1E+2 1E+3 1E+4 Number of cycles, N Joint

800 600

1E+5

1.2% 1.0% 0.8% 0.6% 0.4% 0.2%

400 200 0 1E+0

1E+1 1E+2 1E+3 1E+4 Number of cycles, N

1E+5

Plastic strain amplitude, %

0.6

(a)

0.5 0.4 0.3

1.2% 1.0% 0.8% 0.6% 0.4% 0.2%

0.2 0.1 0

-0.1 1E+0 0.6 Plastic strain amplitude, %

Ti6Al4V BM

(b)

1E+1 1E+2 1E+3 1E+4 Number of cycles, N Joint

0.4

1E+5

1.2% 1.0% 0.8% 0.6% 0.4% 0.2%

0.2

0

-0.2 1E+0

1E+1 1E+2 1E+3 1E+4 Number of cycles, N

1E+5

1.4

Ti6Al4V LFWed joint Ti6Al4V BM Ti6Al4V EBWed joint [25] Ti6Al4V/Ti17 EBWed joint [26] Ti17 BM [26] IMI 834 alloy [49] 829 alloy [50] 685 alloy [51]

1.2

0.8 0.6 0.4 0.2 0.0 1E+1

1E+3 1E+5 1E+7 Number of cycles to failure, Nf

-1

1E+9

Total

∆εt/2 = -0.13(2Nf) - 1.67 R² = 0.86

Elastic

-2 Log (D/2)

Dt/2, %

1.0

Plastic

∆εe/2 = -0.08(2Nf) - 1.91 R² = 0.91

-3 -4

∆εp/2 = -0.67(2Nf) - 0.92 R² = 0.81

-5 -6 1

2

3

4 5 Log(2Nf)

6

7

8

(a)

(b)

(a)

(b)

(c)

(d)