Microstructures and mechanical properties of friction stir welded joints ...

Microstructures and mechanical properties of friction stir welded joints of Zr–Ti alloy Y. J. Li 1,2, R. D. Fu*1,2, D. X. Du1,2, L. J. Jing1,2, D. L. Sang1,2 and X. Y. Zhang1,2 Zirconium–titanium alloy joints were successfully produced by friction stir welding. Unlike the (azb) dual phase microstructure in base metal, only the b phase existed in the region in which temperature exceeded the b transient point for the as welded joint. Accordingly, the hardness in these regions exhibited integral decrement and uniform distribution features. The thermal simulation further showed that hardness variation was mainly determined by phase composition. Microstructure development in the nugget zone was mainly governed by continuous dynamic recrystallisation. Satisfactory ultimate tensile strength and elongation equal to the base metal were achieved in the as welded joint. Tensile fracture occurred at the heat affected zone near the retreating side of the joint. The fracture surface of the joint exhibited a mixing feature with quasicleavage facets and small dimples. Keywords: Zr–Ti alloy, Friction stir welding, Recrystallisation, Phase transformation, Hardness, Tensile properties

Introduction Zr and Zr alloys have been widely used in the nuclear industry and in fields that require high corrosion resistance.1–3 A series of Zr–Ti alloys has recently been developed for the aerospace industry for their excellent anti-irradiation property.4–7 The microstructures of Zr– Ti alloys undergo a series of transformation products from the a (a9) phase to the b phase with increasing Zr content. Alloy ductility is gradually enhanced, whereas tensile strength remains at a high level. For example, the 47Zr–45Ti–5Al–3V alloy has comparable strength with and higher plasticity than Zr based bulk metallic glass.4 Zr alloy welding remains challenging. In the past decades, various welding processes, such as gas tungsten arc welding,8 gas metal arc welding,9 electron beam welding10 and resistance welding8, have been tested for Zr and Zr alloy welding. The problems associated with these welding methods include composition contamination, formation of brittle and coarse microstructure, larger weld distortion and residual stress. Friction stir welding (FSW), which was invented at The Weld Institute in 1991, is a solid state welding process, which has been successfully applied to join various metal materials with low melting temperatures11 and several steels.12–16 Studies on the FSW of Ti and Zr alloys are relatively few, and most existing works mainly focus on Ti alloys,17–28 with only a few focusing on Zr alloys.29 The composition proportion of Zr to Ti in current Zr–Ti alloys is nearly 1 : 1. Thus, this alloy is no longer a type of simple Zr or Ti based alloy. Consequently, the 1 State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao 066004, China 2 College of Materials Science and Engineering, Yanshan University, Qinhuangdao 066004, China

*Corresponding author, [email protected]

weldability of this Zr–Ti alloy should be evaluated. We attempted to weld this Zr–Ti alloy using FSW. Detailed discussions were presented on the relationships between the mechanical properties and microstructures of the joint.

Experimental Zr–Ti alloy with a 3 mm thick gauge was used in this work. The alloy chemical composition was as follows: Zr–41Ti–4Al–3V (wt-%). An FSW tool made of tungsten–rhenium (W–25Re) alloy consisted of a shoulder with a 16 mm diameter and an unthreaded pin with 3 mm length. The FSW tool had rotational and travel speeds of 700 rev min21 and 50 mm min21 respectively. The tilt angle was set at 2u from the vertical axis. Argon shielding was employed around the tool during FSW to avoid surface oxidation. The temperature variation at the heat affected zone (HAZ) was monitored with the use of a thermocouple fixed at a distance of 10 mm from the welding centre. After FSW, the Vickers hardness profile was measured on a cross-section perpendicular to the welding direction with FM-ARS9000. Tensile tests were performed on an Instron 5900 at a constant strain rate of 10–4 s–1 at room temperature. The fracture surface morphologies of the tensile specimens were examined using a HITACHI S-3400 scanning electron microscope (SEM). The phase composition in the joints was determined with the use of a D/max-2500/PC X-ray diffraction. After the tensile tests, the microstructures of the different zones in the as welded joints and fractured samples were observed by optical microscopy (OM). Orientation imaging microscopy (OIM) was performed using a HITACHI S-3400 equipped with a TSL OIM electron backscatter diffraction system. The OM specimens were mechanically ground and etched in a solution

ß 2014 Institute of Materials, Minerals and Mining Published by Maney on behalf of the Institute Received 14 May 2014; accepted 19 June 2014 DOI 10.1179/1362171814Y.0000000229

Science and Technology of Welding and Joining

2014

VOL

19

NO

7

588

Li et al.

Microstructures and properties of FSW joints of Zr–Ti alloy

a macrograph of Zr–Ti alloy FSW joint surface; b low magnification overview of transversal cross-section of Zr–Ti alloy FSW joint; c composite OIM of region on advancing side of joint; d (011) pole figure in ND–RD reference frame of different zones in Zr–Ti alloy FSW joint 1 Macrograph and microstructural observations of Zr–Ti alloy FSW joint

of 2?5% nitric acid, 1?5% hydrochloric acid, 1% hydrofluoric acid and 90% water. The OIM analysis samples were electrochemically polished in a solution consisting of 75 mL perchloric acid, 25 mL glycerol and 425 mL ethyl alcohol at room temperature with a voltage of 18 V for 8 s. To determine the reason behind the decrease in FSW joint hardness, a physical simulation test with a similar welding thermal cycle in FSW was performed on a Gleeble 3500 instrument.

Experimental results Microstructures A smooth appearance of the Zr–Ti alloy FSW joint is observed in Fig. 1a. An overview of the transverse crosssection of this joint without internal defects is presented in Fig. 1b. Four zones, namely, nugget zone (NZ), thermomechanical affected zone (TMAZ), HAZ and base material (BM), are shown in Fig. 1b. Unlike the absence of TMAZ or HAZ in the FSW joints of Ti alloys17–28 and Zr alloys,29 conspicuous TMAZ and HAZ can be found in the Zr–Ti FSW joint. The OIM map of the region marked with a black rectangle in Fig. 1b is also presented in Fig. 1c. This map shows the distribution of grain orientations and sizes in this region. Low angle boundaries (LABs) and high angle boundaries (HABs) defined by a 15u criterion are depicted as white and black lines respectively. Figure 1b shows that the BM microstructure exhibits a mixing feature of large and small grains resulting from incomplete dynamic microstructural evolution during prior hot forging. The HAZ exhibits grain features similar to those of BM in terms of insignificant grain growth during FSW. Figure 1c shows that the NZ grains

are refined to an average size of ,25 mm, which can be attributed to the dynamic recrystallisation (DRX) for the extensive thermal–mechanical deformation in NZ.17,27 Thermomechanical affected zone exists between NZ and HAZ. This region can be further divided into two parts: TMAZ I and TMAZ II. TMAZ I is composed of elongated grains and large original grains that are subdivided into smaller subgrains. Although relatively more LABs exist in TMAZ II than in NZ, the grains are homogeneously refined. The slightly elongated feature of the grains is caused by shear deformation. The (011) pole figures reflecting the microtexture features in HAZ, TMAZ and NZ are shown in Fig. 1f. The texture shows a major (001) [100] component with a maximum intensity of 10?9 in HAZ. This component was formed during prior hot forging. The variation in the maximum intensity from HAZ to TMAZ II confirms that (001) [100] textures were gradually randomised with the enhanced effects of the strain generated by the FSW tool. Although the material flowing around the FSW tool appears to be complicated, a weak (21210) [21122] shear texture component within the range of TMAZ II to NZ can be distinguished in Fig. 1f. The (21210) [21122] shear texture components in NZ became significantly stronger despite the similar maximum intensity from TMAZ II to NZ. This condition is consistent with the viewpoint that the predominant deformation mode in NZ is simple shear.18,19,29,30 Another important microstructure feature in NZ is band structure (BS), which has been observed in FSW joints of 304 stainless steel15,16 and copper.31 The BS details in the FSW joint of the studied Zr–Ti alloy are shown in Fig. 2. As shown in Fig. 2a, BS presents a lamellar structure, as determined by OM observation.


2014

VOL

19

NO

7

589

Li et al.


a OM micrograph; b SEM image 2 Band structure in Zr–Ti alloy FSW joint

The results of further BS observation through SEM are shown in Fig. 2b. Apart from the absence of internal defects in BS, the lamellar feature in Fig. 2b is actually an undular concave–convex resulting from the difference in etching resistance.

Mechanical properties The hardness distribution along the weld cross-section centreline is shown in Fig. 3. The hardness distribution exhibits uniform decrement between two HAZs of the joint. Hardness decreases from ,445 HV in BM to ,325 HV in the low hardness region of the joint. The NZ hardness remained constant despite the significant refinement of grains. A thermal simulation test with the same thermal history in FSW was conducted to investigate further the nature of the hardness decrease in FSW joint. Figure 3 shows that the distribution and values of the low hardness region in the FSW joint and of the simulated specimen exhibit similar features. This observation confirms that the hardness in the Zr–Ti alloy weld joint was mainly determined by the thermal cycle. The joint hardness is thus related to the phase composition rather than the grain size or substructures. The ultimate tensile strength (UTS) and elongation of the FSW joint were 1056 MPa and 15?83% respectively. These values are nearly equal to the BM properties (UTS, 1150 MPa; elongation, 17?18%). A satisfactory tensile property of the Zr–Ti alloy joint was clearly

obtained through FSW. The tensile fractured specimens of BM and FSW joint are presented in Fig. 4a and b respectively. Unlike the case of fracture in NZ of TC4 alloy FSW joints,20,21 the fracture of Zr–Ti alloy FSW joint occurred at HAZ. Negligible necking was found at the fracture position in terms of the relative uniform tensile deformation in the BM gauge range, whereas obvious necking was observed in the HAZ fracture position at the retreating side of the joint. The detailed observations of the deformed microstructures near the fracture surface of the joint are shown in Fig. 4c–e. The high proof strength resulting from NZ grain refinement caused strain localisation to occur at HAZ rather than NZ. The fracture surfaces of the FSW joint and BM are presented in Fig. 5 and exhibit similar features with quasi-cleavage facets and small dimples. Nevertheless, the sizes of cleavage facets and dimples in the FSW joint specimen seem larger than that in BM. This observation can be attributed to the microstructure differences at the fractured position.

Discussion Microstructure and hardness Most early researchers found no obvious HAZ in Ti alloy17,19,22,23 and Zr alloy29 FSW joints on the basis of the absence of evident grain growth in these zones.

3 Hardness proﬁle along transverse direction in FSW joint and along simulated specimen


2014

VOL

19

NO

7

590

Li et al.


a macrograph of BM; b macrograph of FSW joint; c low magnification overview of transversal cross-section of FSW joint; d NZ; e fracture location 4 Macrograph and microstructural observations of fracture location

Other researchers observed a narrow HAZ with microstructure similar to that of BM, along with an increment in the b phase volume faction.21,24 The narrowness or absence of HAZ is related to the lower thermoconductivity of Ti and Zr alloys under low welded heat input.

In TMAZ, the absence21,24–26 or formation of narrow TMAZ17,19,23,27,29 in the Ti alloy and Zr alloy FSW joints is related not only to the weld heat input and thermoconductivity of alloys but also to the crystal structures. For example, the TMAZ in Ti based alloy

5 a, b SEM images of tensile specimens with BM specimen and c, d FSW as weld joint


2014

VOL

19

NO

7

591

Li et al.

6 Evolution of speciﬁc grain boundary area statistics from HAZ to NZ

FSW joints was masked by allotropic transformation.26 The absence of TMAZ in pure titanium FSW joints can be explained by the relatively low stacking fault energy of the hexagonal close packed (hcp) structure a-Ti. In this case, TMAZ with a recovered subgrain structure inside the elongated grains could not be observed.25 Nevertheless, a negligibly narrow TMAZ was produced under low welded heat input in most of Ti alloys because of limited plastic deformations. By contrast, clear and wide HAZ and TMAZ can be distinguished in the studied Zr–Ti alloy FSW joint. Aside from the effects of high weld heat input, the microstructural evolution mechanism of this Zr–Ti alloy significantly affects microstructure evolutions in these regions. As shown in Fig. 1c, dynamic recovery occurred at TMAZ I under slight plastic deformation and low temperature in the b phase field. More subgrain structures were observed inside the elongated grains. The original grain boundaries in TMAZ I also become wavy and tend to form small equiaxed grains. This observation indicates an initial stage of the discontinuous DRX (DDRX).19 From the misorientation angle distributions of TMAZ I in Fig. 6, the mechanism of deformation induced sub-boundaries can be confirmed by the slightly increased LABs. The total length of HABs in TMAZ I also increased slightly. In particular, the HAB increase is a build-up of a wide peak centred near 45u rather than a progressive movement of the low angle peak towards high angle misorientation. This condition probably reflects the DDRX predominance during the microstructure evolution in TMAZ I.19 Although some larger grains with substructures are retained in TMAZ II, incomplete DRX has occurred for the increased strain and temperature. Based on the misorientation angle distributions in Fig. 6, the grain refinement in TMAZ II can be attributed to the continuous DRX (CDRX), which shows that the microstructure formation in TMAZ II results from the gradual accumulation of misorientation by the deformation induced boundaries. This typical CDRX feature has been found in high stacking fault energy metals, such as aluminium and b-Ti, during large deformation at high temperatures.17 The complete CDRX occurrence can also be demonstrated by the clear decrease in LABs and continued accumulation of


NZ misorientation. Thus, CDRX plays a major role in the microstructure evolution of NZ and TMAZ II. The microstructure evolutions in these Zr–Ti alloy FSW joints are also affected by the transformation dynamics. The a phase with a hcp structure is stable below 813 K. The b phase with a body centred cubic structure is stable above 993 K. A two-phase (azb) field exists between the critical temperatures. The Zr and Ti combination in the binary alloy can also significantly increase the b phase stability. Therefore, the b phase transition temperature decreases gradually with the increase in Zr content in the Zr–Ti alloys. When the Zr content in the alloys increases to ,50%, the single b phase can be reserved under water cooling conditions.6 Considering the effect of V (i.e. b phase stabilised element) in the alloy, the b phase stability is further enhanced. Even in the case of a thin plate, only a single b phase in the alloy exists during normalisation because of the relatively faster cooling rate. Therefore, the single b phase can be easily obtained in the joint because the cooling rate of the shielding gas is faster than the critical cooling rate. The X-ray diffraction results of phase composition analysis are shown in Fig. 7, which confirm the above estimation. Results show that BM is composed of b phase and a little of a phase, but only b phase exists in HAZ and NZ. These results also correspond to the uneven optical contrast in Fig. 1b. The single b phase region is comparatively lighter, and BM is darker. From the welding thermal cycle measured at HAZ in Fig. 8, the peak temperature of the welding thermal cycle during FSW has exceeded the b phase transient point of this Zr–Ti alloy because of the tool’s relatively high rotation speed. Consequently, the a phase in BM is redissolved into the b phase matrix at a high temperature during the welding thermal cycle heating stage. Given the b phase stability in the alloy, the transformation from b to a phase can also be restricted by the cold shielding gas during the cooling stage of the welding thermal cycle. Therefore, the single b phase microstructures are preserved at room temperature in these zones. As shown in Fig. 7, the microstructure in the low hardness region of the thermally simulated specimen also exhibits a single b phase. This condition further proves that the phase transformation occurring in the low hardness zones of the Zr–Ti alloy is mainly determined by the thermal effect. Zr or Ti alloy hardness largely depends on the microstructure. For example, Zhang et al.26 reported that the NZ with fully a/b lamellar microstructures has a higher hardness value than that of BM for the plate FSW joint of Ti–6Al–4V alloys. However, Liu et al.20 and Zhou et al.21 reported that the mean hardness value of the NZ is lower than the BM for the mill annealed Ti– 6Al–4V FSW joints. The NZ microstructure is characterised by a fully lamellar microstructure with dislocation free or low dislocation density a colonies. In this case, the NZ hardness decrease is attributed to the reduction of dislocation density during the recrystallisation processes. In comparison, the present Zr–Ti alloy shows a single b phase microstructure in the low hardness region of the joint. Although the grain sizes in NZ are significantly refined, the NZ hardness does not clearly increase. The a phase with an hcp structure also has higher strength and hardness relative to the b phase.


2014

VOL

19

NO

7

592

Li et al.


7 X-ray diffraction patterns of different Zr–Ti alloy FSW joint positions and simulated specimen

Reference to Figs. 1b and 3, the low hardness region length (,325 HV) is equal to the single b phase zone (comparatively lighter). Together with the phase composition (Fig. 7) and hardness distribution (Fig. 3) of the simulated specimen, the hardness of the Zr–Ti alloy FSW joint is determined by the phase composition, but not the grain size or substructures.

Tensile and fracture Generally, tensile deformation and fracture occur at the lowest hardness regions in FSW joints. For example, a Ti–6Al–4V alloy FSW joint has been reported to have lower tensile strength than BM.21,26 The decrease in joint strength can be attributed to the decrease in hardness induced by dislocation density reduction and microstructural coarsening. In this study, the appearance of a tensile specimen in Fig. 4a shows that the tensile deformation in the BM specimen is homogeneous because of the microstructure’s uniform distribution. The high UTS of BM is mainly related to the solid solution strengthening by a high proportion of alloy elements32 and to the dispersed strengthening attributed

to some a phases in b phase grains. The acceptable plasticity is also determined by the b phase microstructure in the Zr–Ti alloy.6 Figure 4b confirms that tensile deformation is mainly concentrated in the FSW joint soft regions. A further observation of the deformation zone in the FSW joint’s fractured specimen is shown in Fig. 4c. Given the difference in phase composition and grain size in the soft regions, the FSW joint deformation is not homogeneous. Many slip bands are formed in the single b phase region, whereas none are formed in BM during tensile deformation. This condition indicates that the phase composition has a primary effect on the nonuniform tensile deformation of the joint. Although a soft region exists in the FSW joint, the UTS of the joint is also nearly equal to the BM. This condition can be attributed to the fact that the grain size in low hardness regions is important in the tensile deformation process. Based on the Hall–Petch relationship, the yield strength of the zones with refined grains must be higher than that with coarser ones. This observation can be verified by comparing the slip bands

8 Temperature variation at HAZ


2014

VOL

19

NO

7

593

Li et al.

formed in NZ and HAZ near to the fracture surface as shown in Fig. 4. In comparison, many slip bands are found in the entire grains in HAZ. This observation indicates that tensile deformation is concentrated in softened HAZ with coarser grains. However, the higher tensile strength of the joints than BM should be explained in terms of the restrain effect. This condition indicates that tensile deformation, which occurs in soft regions of the joint, can be restrained by adjacent regions with higher strength. The stress state in strain localisation has transformed from uni- to multiaxial stress. As a result, the FSW joint exhibits higher strength. This phenomenon has also been observed in stainless steel FSW joints.14 The soft region composed of a single b phase is expected to have higher toughness and elongation. However, the FSW joint only exhibits equivalent elongation to BM. The strain hardening behaviour of the metal in the soft region plays a key role during the subsequent tensile deformation after yield strain. The balance between softening and hardening, which occurs in the deformation zones, ultimately determines the total join elongation. Thus, the lower elongation of the entire FSW joint than BM may be attributed to the workhardening rate decrease in the soft region.

Conclusions In this study, a 3 mm thick Zr–Ti alloy plate was successfully friction stir welded. The FSW joint microstructure and mechanical properties were investigated. The following conclusions were drawn. 1. A conspicuous TMAZ and HAZ were formed in the joint. Unlike the (azb) phase in BM, only b phase existed in NZ, TMAZ and a portion of HAZ because of the dominant effect of the welding thermal cycle. 2. Based on the distribution of the grain structures from HAZ to NZ, the NZ microstructure evolution can be mainly attributed to the CDRX process. 3. The hardness within the NZ, TMAZ and a portion of HAZ were lower than that of the BM, although the grain has been refined in NZ. The most important factor determining the joint hardness is the phase composition of the current Zr–Ti alloy. 4. An acceptable tensile property of the Zr–Ti alloy FSW joint was obtained. A satisfactory UTS and elongation equal to those of BM was observed. The FSW joint fracture surface exhibited a similar mixing feature to BM with quasi-cleavage facets and small dimples.

Acknowledgement Financial support from the National Basic Research of China (grant no. 2010CB731606) is greatly acknowledged.


References 1. D. G. Franklin, G. E. Lucas and A. L. Bement: ‘Creep of zirconium alloys in nuclear reactors’, ASTM STP815, ASTM, West Conshohocken, PA, USA, 1983. 2. P. Rodriguez: Proc. Conf. on ‘Zirconium alloys for reactor components’, Bombay, India, December 1991, Department of Atomic Energy, 46–95. 3. G. Sabol, G. R. Klip, M. G. Balfour and E. Roberts: Proc. 8th Int. Symp. on ‘Zirconium in the nuclear industry’, January 1989, ASTM STP1023, West Conshohocken, PA, USA, 227–244. 4. S. X. Liang, L. X. Yin and M. Z. Ma: Mater. Sci. Eng. A, 2013, A561, 13–16. 5. R. Jing, S. X. Liang and C. Y. Liu: Mater. Des., 2013, 52, 981–986. 6. S. X. Liang, M. Z. Ma and R. Jing: Mater. Sci. Eng. A, 2012, A532, 1–5. 7. S. X. Liang, M. Z. Ma and R. Jing: Mater Sci. Eng. A, 2012, A539, 42–47. 8. R. A. Bordoni and A. M. Olmedo: J. Mater. Sci., 1981, 16, 1527– 1532. 9. R. Boyer, G. Welsch and E. W. Collings: ‘Titanium alloys, technical note 10: welding and brazing’, in ‘Materials properties handbook’, 1159; 1994, Materials Park, OH, ASM. 10. N. K. Babu, S. G. S. Raman and C. V. S. Murthy: Sci. Technol. Weld. Join., 2005, 10, 583–590. 11. B. T. Gibson, D. H. Lammlein and T. J. Prater: J. Manuf. Process, 2013, 16, 56–73. 12. Y. C. Chen, H. Fujii and T. Tsumura: J. Nucl. Mater., 2012, 420, 497–500. 13. Y. C. Chen, H. Fujii and T. Tsumura: Sci. Technol. Weld. Join., 2009, 14, 197–201. 14. Y. Miyano, H. Fujii and Y. Sun: Mater. Sci. Eng. A, 2011, A528, 2917–2921. 15. S. H. C. Park, Y. S. Sato and H. Kokawa: Sci. Technol. Weld. Join., 2005, 10, 550–556. 16. S. H. C. Park, Y. S. Sato and H. Kokawa: Scr. Mater., 2003, 49, 1175–1180. 17. L. H. Wu, D. Wang, B. L. Xiao and Z. Y. Ma: Scr. Mater., 2014, 78–79, 17–20. 18. S. Mironov, Y. Zhang and Y. S. Sato: Scr. Mater., 2008, 59, 27–30. 19. S. Mironov, Y. S. Sato and H. Kokawa: Mater. Sci. Eng. A, 2010, A527, 7498–7504. 20. H. J. Liu, L. Zhou and Q. W. Liu: Mater. Des., 2010, 31, 1650– 1655. 21. L. Zhou, H. J. Liu and Q. W. Liu: Mater. Des., 2010, 31, 2631– 2636. 22. L. Zhou, H. J. Liu and P. Liu: Scr. Mater., 2009, 61, 596–599. 23. S. Mironov, Y. S. Sato and H. Kokawa: Acta Mater., 2009, 57, 4519–4528. 24. L. Zhou, H. J. Liu and Q. W. Liu: J. Mater. Sci., 2010, 45, 39–45. 25. W. Lee, C. Lee and W. Chang: Mater. Lett., 2005, 59, 3315–3318. 26. Y. Zhang, Y. S. Sato and H. Kokawa: Mater Sci. Eng. A, 2008, A485, 448–455. 27. Y. Zhang, Y. S. Sato and H. Kokawa: Sci. Technol. Weld. Join., 2010, 15, 500–505. 28. A. Steuwer, D. G. Httingh and M. N. James: Sci. Technol. Weld. Join., 2012, 17, 525–533. 29. Y. S. Sato, Y. Nagahama and S. Mironov: Scr. Mater., 2012, 67, 241–244. 30. R. W. Fonda and K. E. Knipling: Sci. Technol. Weld. Join., 2011, 16, 288–294. 31. K. Savolainen, T. Saukkonen and H. Ha¨nninen: Sci. Technol. Weld. Join., 2012, 17, 111–115. 32. H. C. Hsu, S. C. Wu and Y. C. Sung: J. Alloys Compd, 2009, 488, 279–283.


2014

VOL

19

NO

7

594