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Microstructure and properties of friction stir welded high strength Fe–36 wt%Ni alloy
Materials Science and Engineering A 528 (2011) 7768–7773
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and properties of friction stir welded high strength Fe–36 wt%Ni alloy Yue Zhao a,b,c,∗ , Yutaka S. Sato c , Hiroyuki Kokawa c , Aiping Wu a,b a b c
Department of Mechanical Engineering, Tsinghua University, PR China Key Laboratory for Advanced Materials Processing Technology, Ministry of Education, PR China Department of Materials Processing, Graduate School of Engineering, Tohoku University, 6-6-02 Aramaki-aza-Aoba, Aoba-ku, Sendai 980-8579, Japan
a r t i c l e
i n f o
Article history: Received 31 August 2010 Received in revised form 27 June 2011 Accepted 28 June 2011 Available online 5 July 2011 Keywords: Friction stir welding Fe–36 wt%Ni alloy Microstructure Mechanical properties Thermal expansion coefficient Electron back-scatter diffraction
a b s t r a c t High strength Fe–36 wt%Ni alloy sheets with 3 mm thickness were successfully friction stir welded and defect-free welds were acquired at travelling speed of 2 mm/s and rotational speed between 600 and 1000 rpm. Essentially, the friction stir welding process did not change the excellent thermal expansion properties of Fe–36 wt%Ni alloy. Friction stir welding resulted in a uniform coarser-grained microstructure in the stir zone with comparable fractions of low angle boundaries and 3 twin boundaries to the base material. However, approximate 10% reductions of hardness and tensile strength were found in the joint. The relationship between grain size and mechanical properties of friction stir welds was discussed. The base material has higher mechanical properties, not only because of the finer average grain size, but also attribute to its inhomogeneous grain structure. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Fe–36 wt%Ni alloy, also named Invar 36 alloy, consists of an austenite-single phase structure and is notable for its uniquely low thermal expansion coefficient (TEC) under its Curie temperature (about 230 ◦ C). Fe–36 wt%Ni alloy has been widely used in temperature-independent instruments, such as precision measuring devices which required excellent dimensional stability in a large temperature range. Recently, Fe–36 wt%Ni alloy have been increasingly used as an important, high-reliability structural material for cryogenic liquid storage and transport [1,2]. In general, conventional fusion welding processes, such as gas tungsten arc welding, are used for Fe–36 wt%Ni alloy joint. However, they tend to lead hot cracking, resulting from solidification and high thermal effect. The addition of Ti, Mn and Mo elements in the filler metal would be conductive to reduce the hot cracking sensibility of the weld [3–5], but obviously increase its TEC of the welds [4]. Friction stir welding (FSW), invented at The Welding Institute of UK in 1991, is a solid-state joining process in which a rotating tool
∗ Corresponding author at: Department of Mechanical Engineering, Tsinghua University, PR China. Fax: +86 10 62773859. E-mail address: [email protected] (Y. Zhao). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.06.082
consisting of a shoulder and a pin inserts the plates to be welded and moves along the welding direction [6,7]. FSW was initially applied to aluminum and some other low melting-temperature materials. Nowadays, with the development of welding tools, it has been expanded to higher melting-temperature materials, such as carbon steel [8–10], stainless steel [11–15] and titanium [16,17]. Compared with the fusion welding processes, FSW has much lower heat input and will reduce the sensitivity to hot creaking. Additionally, FSW would produce defect-free weld with homogeneous microstructure more easily than conventional fusion processes, resulting in good mechanical properties of the weld. Therefore, FSW might be an alternative method for Fe–36 wt%Ni alloy fusion welding processes. However, there is only one report concerning the FSW application for Fe–36 wt%Ni. Jasthi et al. [18] achieved to produce defect-free friction stir welds of normal strength (ultimate strength of 461 MPa) Fe–36 wt%Ni alloy and the results showed that its TEC, tensile strength and micro hardness were essentially matching the base material. However, the microstructural characteristics of the welds were not examined and microstructural factors governing mechanical properties were not systematically studied. Moreover, with the development of the manufacturing technology, the strength of Fe–36 wt%Ni alloy has improved from about 450 MPa to over 500 MPa. Will the FSW be applicative for the high
Y. Zhao et al. / Materials Science and Engineering A 528 (2011) 7768–7773
strength Fe–36 wt%Ni alloy and produce defect-free welds with comparable properties as the base materials has not been reported yet. The present study is first aim to test the feasibility of FSW on high strength Fe–36 wt%Ni alloy (ultimate strength of 516 MPa). Then mainly focus on the investigation of microstructure characteristics of FS welds, and discussion of the relationship between microstructure and properties.
2. Experimental procedures 3-mm thick high strength Fe–36 wt%Ni alloy sheet was used as the base material in this study. Its chemical composition (wt%) was Fe–36.0Ni–0.3Mn–0.19Si–0.03C–0.001S–0.002P. A PCBN tool and a stainless steel backing plate were used to perform the bead-onplate FSW. The PCBN tool had a convex shoulder with step-spiral pattern having a diameter of 15.3 mm and a 2 mm length tapered pin. The pin tapered from 6.9 mm at the shoulder to 4.1 mm at the pin tip. The load to the axial direction of tool was controlled to obtain a plunge depth of about 0.4 mm. The travelling speed of the welding tool changed at 2 mm/s, 3 mm/s and 5 mm/s. The rotational speed was varied among 600 rpm, 800 rpm and 1000 rpm. The defect of welds was detected by X-ray inspection and crosssectional observation. Microstructure was characterized by optical microscopy. Electron backscatter diffraction (EBSD), equipped on the Hitachi S-4300SE scanning electron microscope (SEM), was also employed to examine the orientation characteristics. Inclusions in welds were identified by energy-dispersive X-ray spectroscopy (EDS) analysis system on a SEM. Specimens for optical microscopy were cut perpendicular to the welding direction, mechanically polished with 3 m and 1 m diamond paste, and then chemically etched in a 60 vol% nitric acid + 40 vol% acetic acid solution. Specimens for EBSD were cut and mechanically polished, and finally electro-polished in a solution of 75 vol% acetic acid + 5 vol% perchloric acid + 20 vol% ethanol at 33 V for 150 s. The scanning step of EBSD data was 1 m. The average confidential indexes for each EBSD map were from 0.7 to 0.9. The experiments on f.c.c. materials have shown that when the confidential indexes are greater than 0.1, the fraction of correctly indexed pattern is over 95% [19]. Consequently, the confidence level of the present EBSD maps was high enough. The non-indexed data points and points with low confidential index were usually associated with grain boundary regions. In order to obtain more reliable map, all small grains including less than 5 pixels and data point with confidential index less than 0.1 were nullified (set to zero solutions). To eliminate spurious boundaries caused by orientation noise, a lower limit boundary misorientation was cut-off of 2◦ . The grain size was specified by diameter, which was calculated by determining the area of a grain and then assuming the grain as a circle. The average grain size was calculated by weighted average of the area fraction of grains. Linear TEC was measured by a thermal dilatometer. The dimensional change in specimen was measured during the heating process, from 50 ◦ C to 500 ◦ C, at the heating rate of 5 ◦ C/min. The cylindrical specimens for the TEC test were cut from the base material and center of the welds, with diameter of 3 mm and length of 10 mm. Vickers hardness test and transverse tensile test were employed to examine mechanical properties of the welds at room temperature. Vickers hardness profile was measured with a load of 9.8 N and a dwell time of 10 s along a 1 mm location under the top surface on the cross section of the welds. The dog-bone-like transverse tensile specimens were cut perpendicular to the welding direction with a gauge length of 30 mm. Weld was positioned in the center of gauge length.
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3. Results and discussion Defect-free welds were only acquired at the travelling speed of 2 mm/s by detection of X-ray inspection. Cross-sectional overview of defect-free and defect welds were shown in Fig. 1. Cavities were found in the advancing side (AS) of all 3 mm/s and 5 mm/s welds. These cavities penetrating through the welding direction are the typical tunnel defect which was reported due to the insufficient heat input during FSW [20]. In the present study, travelling speed was the dominate factor to make defect-free weld. So, besides the heat input which was mainly governed by rotational speed, material flow also plays an important role to the formation of FS welds. In relatively higher-travelling-speed welding process (3 mm/s, 5 mm/s), material was stirred from advancing side (AS) to retreating side (RS), but did not have enough capability return to the AS, finally leaded to cavity defect in the bottom of AS. With the decrease of welding speed, the cavities in cross-section trends to disappear, as shown in Fig. 1a, d, and e. The following microstructure and properties investigations will focus on the defect-free welds.
3.1. Microstructure of FS welded Fe–36 wt%Ni alloy As shown in Fig. 1, The stir zones of defect-free FS welds have homogeneous structure in general, except some white zones (WZ) in the top and bottom of AS. High density of inclusions and pits were observed in the WZs, as shown in Fig. 2a. Some WZs also have banded structure in the bottom AS, as shown in Fig. 2b. EDS analysis was used on as-polished WZ to clarify the composition of inclusion. Flaky-shaped inclusions were observed, as shown in Fig. 2c. The EDS spectrums reveal that WZ consists of Fe, Ni peaks as well as B, N peaks. Therefore, the inclusions could be inferred as the CBN debris arising from wearing of the PCBN tool during FSW process. The wear of PCBN tool was also observed in FSW applications to some other materials, such as stainless steel [21]. Optical microstructure in the stir zone center of defect-free FS welds and the base material were shown in Fig. 3. All of them have single austenitic equiaxed microstructure with annealing twins. However, the stir zones are relatively coarser than the base material. The grain boundary maps of base material and the stir zone center of FS welds were shown in Fig. 4. Low angle boundaries (LABs) having misorientation angles of 2–15◦ were marked by thinner pink lines; random high angle boundaries (HABs) having misorientation angles higher than 15◦ were marked by thicker black lines; and 3 twin boundaries, identified with the misorientation angles around 60◦ and misorientation axis around 1 1 1 direction, were marked by thicker gray lines. Factions fraction of LABs and 3 twin boundaries and average grain size in etch grain boundary map were summarized in Table 1. The grain structure, grain boundary characteristics and the recrystallization of FS welded Fe–36 wt%Ni alloy were studied base on the grain boundary maps as follows. Firstly, the grain size as well as grain structure of Fe–36 wt%Ni alloy changed after FSW. The grain size of stir zones is relatively bigger and its distribution is more homogeneous than that of the base material. For the base material, its average grain size is 15 m. It contains some 25–30 m big grains and many 5–10 m fine grains, the big grains are surrounded by fine grains. Accordingly, for the 600 rpm–2 mm/s FS weld, the average grain size is 20 m. Compared with the base material, it is slightly bigger, but the amount of fine grains is obvious reduced, i.e., the grain structure of FS weld becomes more homogeneous. With the increase of rotational speed, the grain structure becomes more homogeneous and coarse. The grain size is 31 m and 34 m in average for 800 rpm and 1000 rpm–2 mm/s welds, respectively.
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Fig. 1. Cross-sectional overviews of defect-free and defect FS welds.
Fig. 2. Optical microstructure (a, b) and EDS analysis of white zone in stir zone (c, d).
Table 1 Grain boundary characteristics and average grain size of the base material and stir zone center of FS welds.
LABs fraction 3 fraction Grain size
Base material
600 rpm, 2 mm/s weld
800 rpm, 2 mm/s weld
1000 rpm, 2 mm/s weld
23% 10% 15 m
25% 9% 20 m
22% 12% 31 m
19% 12% 34 m
Secondly, the LABs fraction and 3 twin boundary fraction of FS welds are comparable to those of the base material, and trends to decrease and increase with increasing of the rotational speed, respectively. The LABs fraction of base material is 23%, while LABs fractions for 600 rpm, 800 rpm and 1000 rpm–2 mm/s welds are 25%, 22% and 19%, respectively. The LABs fractions of FS welds trend to decrease with the increasing of rotational speed in general. The 3 twin boundary fraction of base material is 10%, the 3 twin boundary for 600 rpm, 800 rpm and 1000 rpm–2 mm/s welds are 9%, 12% and 12%, respectively. The 3 twin boundary of FS welds
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Fig. 3. Optical microstructure of base material and the stir zone center of FS welds.
trend to increase with the increasing of the rotational speed. As well known, LABs are always considered as forming during recovery in the deformed microstructure; 3 twin boundaries, which also named annealing twin, are believed forming with grain growth after recrystallization. So, these tendencies of LABs and 3 twin boundary fraction are reasonable, since the increasing of rotational speed results in the enhancement of heat input during welding. The greater heat input leads to more completed recovery, recrystallization even grain growth (i.e., to lower the system energy), produce a microstructure with less LABs and more 3 twin boundaries. Finally, evidence of continuous dynamic recrystallization during FSW of Fe–36 wt%Ni alloy is found in the present grain boundary map. A recovered grain “A” containing of many subgrains was observed, as shown in Fig. 4b. And some subgrain boundaries accumulate a misorientation exceeding 15◦ to high angle boundaries, which was also marked by arrowhead in the figure. This evidence is most likely associated with grain subdivision driven continuous dynamic recrystallization. Some researchers [22,23] also reported the continuous dynamic recrystallization of f.c.c. materials. Moreover, as a medium stacking-fault-energy material (163 mJ/m2 [24]), discontinuous dynamic recrystallization also probably occurred during welding process. Since it is difficult to find the transitional region between the stir zone and base material, i.e., thermal mechanical affected zone, the proof for discontinues dynamic recrystallization during FSW still requires further investigation. 3.2. TEC and mechanical properties The TEC – temperature curves of base material and FS welds were shown in Fig. 5. The curves are significantly increased
around 230 ◦ C, which related to the magnetic transformation of Fe–36 wt%Ni alloy. All welds had comparable TEC to the base material. The small amount of CBN particles did not change the good thermal expansion properties of the base material. Vickers hardness profiles of the FS welds were shown in Fig. 6. The average hardness of 600 rpm, 800 rpm and 1000 rpm–2 mm/s welds are 144 Hv , 140 Hv and 140 Hv , respectively, which is a little lower than the base material (157 Hv ). Besides, with the increasing of rotational speed, the hardness of stir zones trends to decrease. But this tendency is not significant. Transverse tensile properties of the base material and the welds were summarized in Table 2. All welds were broken in the stir zone during tensile test, indicating the stir zone had the lowest hardness in the weld. Similar to the hardness results, the welds exhibit lower ultimate tensile strength and yield strength than the base material. The elongation of FS welds is also decreased, compared with the base material. The reduction of elongation is because of the preferentially yielded stir zone, in other words, the inhomogeneity of deformation of the joint leads to the lower elongation. The strength and elongation of FS welds also have faintly trends to decrease with the increasing of rotational speed. Although the mechanical properties of FS welds are reduced compared with the base material, they are still superior to those of the conventional arc welding and previous FSW of Fe–36 wt%Ni alloy. Gottlieb reported ultimate strength of 407 MPa for gas tungsten arc welded Invar alloy with similar-to-Invar filler [3]. Jasthi reported ultimate strength of 461 MPa and yield strength of 296 MPa for FS welded Fe–36 wt%Ni alloy with PCBN tool [18].
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Fig. 4. Grain boundary maps of base material and the stir zone center of FS welds.
3.3. Relationship of microstructure characteristic and mechanical properties According to the grain boundary strengthening theory, the decrease of grain size will improve the mechanical properties. The relationship between yield strength and grain size of most polycrystalline metals could be described by Hall–Petch relationship, i.e., y = 0 + kd−1/2 , where y and d are yield strength and grain size, respectively; 0 and k are constants independent of grain size.
Table 2 Transverse tensile properties of base material and FS welds.
Fig. 5. Thermal expansion coefficient of base material and stir zones.
Ultimate strength, MPa Yield strength, MPa Elongation, %
Base material
600 rpm, 2 mm/s weld
800 rpm, 2 mm/s weld
1000 rpm, 2 mm/s weld
516 ± 5
485 ± 1
472 ± 15
475 ± 1
393 ± 17 33 ± 3
336 ± 5 19 ± 0
321 ± 26 18 ± 1
327 ± 4 19 ± 1
Y. Zhao et al. / Materials Science and Engineering A 528 (2011) 7768–7773
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strengthen of base material is probably attributed to the difference of grain structure in base material and FS welds. Base material has an inhomogeneous grain size distribution with some big grains surrounded by many fine grains. These fine grains should be helpful to strengthen the base material. 4. Conclusions Friction stir welding was successfully applied to high strength Fe–36 wt%Ni alloy and defect-free welds were acquired at travelling speed of 2 mm/s and rotational speed of 600, 800 and 1000 rpm with PCBN tool. The welds have homogeneous coarse austenite microstructure in the stir zone with comparable LABs and 3 twin boundary fraction to the base material. The outstanding TEC remains for the FS welds, while their mechanical properties are slightly reduced. The reduction of mechanical properties is probably caused by the bigger average grain size and uniform grain structure compared with the base material. Fig. 6. Microhardness of base material and FS welds.
Acknowledgements The authors are grateful to Mr. A. Honda, Prof. G. Miyamoto, Mr. Z.D. Li, Dr. S. Mironov and Dr. K. Kobayashi for technical assistance, and thank Prof. K. Maruyama and Prof. T. Furuhara for their helpful discussions. One of the authors (Y. Zhao) thanks Joint Education Program between Tohoku University and Tsinghua University. We also acknowledge financial support from the Japanese Ministry of Education, Science, Sports and Culture with a Grant-in-Aid from the Global COE program in Materials Integration International Center of Education and Research at Tohoku University. References
Fig. 7. The yield strength and grain size relationship of base material and FS welds.
Since the Hall–Petch relationship was always linear fitted by large grain size range to ensure its accuracy, it is difficult to calculate accurate Hall–Petch relationship of FS welded Fe–36 wt%Ni alloy with the existing data, for the present grain size range was limited from 20 to 34 m. Vinogradov and his colleagues [25] reported the yield strength and grain size (0.18–100 m) of equalchannel-angular-pressed (ECAP-ed) Fe–36 wt%Ni alloy and their results were linear fitted to Hall–Petch relationship. The plot of yield strength against inverse square root of grain size with base material and FS welds was shown and compared with the reported Hall-patch relationship of ECAP-ed Fe–36 wt%Ni alloy in Fig. 7. The FS welds have an about 50 MPa higher strength than the reported Hall–Petch relationship, while the deviation of the base material is about 100 MPa. All the present base material and FS welds deviated from the reported Hall–Petch relationship. It might due to the different preparation methods of base materials and specializations between FSW and ECAP process. However, it could be noticed that the base material has a much bigger deviation from the reported Hall–Petch relationship compared with the FS welds. As shown in Fig. 4, the base material has comparable fraction of 3 annealing twin boundaries and LABs of FS welds. So the
[1] M. Shiga, Curr. Opin. Sol. State Mater. Sci. 1 (1996) 340–348. [2] H. Saito, S. Chikazumi, T. Hirone, Physics and Applications of Invar alloys, Maruzen Company, LTD, Tokyo, 1978. [3] T. Gottlieb, C.S. Shira, Weld. J. 44 (1965) S116. [4] K.A. Yushchenko, N.I. Kakhovskii, T.M. Starushchenko, Met. Sci. Heat Treat. 16 (1974) 866–869. [5] G.R. Rundell, A.E. Nehrenberg, Weld. J. 45 (1966) 156–160. [6] R.S. Mishra, Z.Y. Ma, Mater. Sci. Eng. R 50 (2005) 1–78. [7] R. Nandan, T. DebRoy, H.K.D.H. Bhadeshia, Prog. Mater. Sci. 53 (2008) 980–1023. [8] Y.S. Sato, H. Yamanoi, H. Kokawa, T. Furuhara, Scripta Mater. 57 (2007) 557–560. [9] L. Cui, H. Fujii, N. Tsuji, K. Nogi, Scripta Mater. 56 (2007) 637–640. [10] A.P. Reynolds, W. Tang, M. Posada, J. DeLoach, Sci. Technol. Weld. Join 8 (2003) 455–460. [11] Y.S. Sato, T.W. Nelson, C.J. Sterling, R.J. Steel, C.O. Pettersson, Mater. Sci. Eng. A 397 (2005) 376–384. [12] R. Johnson, P.L. Threadgill, Trends in Welding Research, ASM International, Materials Park, OH, 2003, pp. 88–92. [13] Y.S. Sato, T.W. Nelson, C.J. Sterling, Acta Mater. 53 (2005) 637–645. [14] S.H.C. Park, Y.S. Sato, H. Kokawa, K. Okamoto, S. Hirano, M. Inagaki, Sci. Technol. Weld. Join 10 (2005) 550–556. [15] S.H.C. Park, Y.S. Sato, H. Kokawa, K. Okamoto, S. Hirano, M. Inagaki, Scripta Mater. 49 (2003) 1175–1180. [16] Y. Zhang, Y.S. Sato, H. Kokawa, S.H.C. Park, S. Hirano, Mater. Sci. Eng. A 488 (2008) 25–30. [17] H. Fujii, Y.F. Sun, H. Kato, K. Nakata, Mater. Sci. Eng. A 527 (2010) 3386–3391. [18] B.K. Jasthi, W.J. Arbegast, S.M. Howard, J. Mater. Eng. Perform 18 (2009) 925–934. [19] TSL. OIM Version 3.0. On-line Help, TSL, Draper, UT, 2001. [20] Y.G. Kim, H. Fujii, T. Tsumura, T. Komazaki, K. Nakata, Mater. Sci. Eng. A 415 (2006) 250–254. [21] S.H.C. Park, Y.S. Sato, H. Kokawa, K. Okamoto, S. Hirano, M. Inagaki, Metall. Mater. Trans. A 40A (2009) 625–636. [22] U.F.H.R. Suhuddin, S. Mironov, Y.S. Sato, H. Kokawa, Mater. Sci. Eng. A 527 (2010) 1962–1969. [23] J.Q. Su, T.W. Nelson, R. Mishra, M. Mahoney, Acta Mater. 51 (2003) 713–729. [24] R.E. Schramm, R.P. Reed, Metall. Trans. A 7A (1976) 359–363. [25] A. Vinogradov, S. Hashimoto, V.I. Kopylov, Mater. Sci. Eng. A 355 (2003) 277–285.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and properties of friction stir welded high strength Fe–36 wt%Ni alloy Yue Zhao a,b,c,∗ , Yutaka S. Sato c , Hiroyuki Kokawa c , Aiping Wu a,b a b c
Department of Mechanical Engineering, Tsinghua University, PR China Key Laboratory for Advanced Materials Processing Technology, Ministry of Education, PR China Department of Materials Processing, Graduate School of Engineering, Tohoku University, 6-6-02 Aramaki-aza-Aoba, Aoba-ku, Sendai 980-8579, Japan
a r t i c l e
i n f o
Article history: Received 31 August 2010 Received in revised form 27 June 2011 Accepted 28 June 2011 Available online 5 July 2011 Keywords: Friction stir welding Fe–36 wt%Ni alloy Microstructure Mechanical properties Thermal expansion coefficient Electron back-scatter diffraction
a b s t r a c t High strength Fe–36 wt%Ni alloy sheets with 3 mm thickness were successfully friction stir welded and defect-free welds were acquired at travelling speed of 2 mm/s and rotational speed between 600 and 1000 rpm. Essentially, the friction stir welding process did not change the excellent thermal expansion properties of Fe–36 wt%Ni alloy. Friction stir welding resulted in a uniform coarser-grained microstructure in the stir zone with comparable fractions of low angle boundaries and 3 twin boundaries to the base material. However, approximate 10% reductions of hardness and tensile strength were found in the joint. The relationship between grain size and mechanical properties of friction stir welds was discussed. The base material has higher mechanical properties, not only because of the finer average grain size, but also attribute to its inhomogeneous grain structure. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Fe–36 wt%Ni alloy, also named Invar 36 alloy, consists of an austenite-single phase structure and is notable for its uniquely low thermal expansion coefficient (TEC) under its Curie temperature (about 230 ◦ C). Fe–36 wt%Ni alloy has been widely used in temperature-independent instruments, such as precision measuring devices which required excellent dimensional stability in a large temperature range. Recently, Fe–36 wt%Ni alloy have been increasingly used as an important, high-reliability structural material for cryogenic liquid storage and transport [1,2]. In general, conventional fusion welding processes, such as gas tungsten arc welding, are used for Fe–36 wt%Ni alloy joint. However, they tend to lead hot cracking, resulting from solidification and high thermal effect. The addition of Ti, Mn and Mo elements in the filler metal would be conductive to reduce the hot cracking sensibility of the weld [3–5], but obviously increase its TEC of the welds [4]. Friction stir welding (FSW), invented at The Welding Institute of UK in 1991, is a solid-state joining process in which a rotating tool
∗ Corresponding author at: Department of Mechanical Engineering, Tsinghua University, PR China. Fax: +86 10 62773859. E-mail address: [email protected] (Y. Zhao). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.06.082
consisting of a shoulder and a pin inserts the plates to be welded and moves along the welding direction [6,7]. FSW was initially applied to aluminum and some other low melting-temperature materials. Nowadays, with the development of welding tools, it has been expanded to higher melting-temperature materials, such as carbon steel [8–10], stainless steel [11–15] and titanium [16,17]. Compared with the fusion welding processes, FSW has much lower heat input and will reduce the sensitivity to hot creaking. Additionally, FSW would produce defect-free weld with homogeneous microstructure more easily than conventional fusion processes, resulting in good mechanical properties of the weld. Therefore, FSW might be an alternative method for Fe–36 wt%Ni alloy fusion welding processes. However, there is only one report concerning the FSW application for Fe–36 wt%Ni. Jasthi et al. [18] achieved to produce defect-free friction stir welds of normal strength (ultimate strength of 461 MPa) Fe–36 wt%Ni alloy and the results showed that its TEC, tensile strength and micro hardness were essentially matching the base material. However, the microstructural characteristics of the welds were not examined and microstructural factors governing mechanical properties were not systematically studied. Moreover, with the development of the manufacturing technology, the strength of Fe–36 wt%Ni alloy has improved from about 450 MPa to over 500 MPa. Will the FSW be applicative for the high
Y. Zhao et al. / Materials Science and Engineering A 528 (2011) 7768–7773
strength Fe–36 wt%Ni alloy and produce defect-free welds with comparable properties as the base materials has not been reported yet. The present study is first aim to test the feasibility of FSW on high strength Fe–36 wt%Ni alloy (ultimate strength of 516 MPa). Then mainly focus on the investigation of microstructure characteristics of FS welds, and discussion of the relationship between microstructure and properties.
2. Experimental procedures 3-mm thick high strength Fe–36 wt%Ni alloy sheet was used as the base material in this study. Its chemical composition (wt%) was Fe–36.0Ni–0.3Mn–0.19Si–0.03C–0.001S–0.002P. A PCBN tool and a stainless steel backing plate were used to perform the bead-onplate FSW. The PCBN tool had a convex shoulder with step-spiral pattern having a diameter of 15.3 mm and a 2 mm length tapered pin. The pin tapered from 6.9 mm at the shoulder to 4.1 mm at the pin tip. The load to the axial direction of tool was controlled to obtain a plunge depth of about 0.4 mm. The travelling speed of the welding tool changed at 2 mm/s, 3 mm/s and 5 mm/s. The rotational speed was varied among 600 rpm, 800 rpm and 1000 rpm. The defect of welds was detected by X-ray inspection and crosssectional observation. Microstructure was characterized by optical microscopy. Electron backscatter diffraction (EBSD), equipped on the Hitachi S-4300SE scanning electron microscope (SEM), was also employed to examine the orientation characteristics. Inclusions in welds were identified by energy-dispersive X-ray spectroscopy (EDS) analysis system on a SEM. Specimens for optical microscopy were cut perpendicular to the welding direction, mechanically polished with 3 m and 1 m diamond paste, and then chemically etched in a 60 vol% nitric acid + 40 vol% acetic acid solution. Specimens for EBSD were cut and mechanically polished, and finally electro-polished in a solution of 75 vol% acetic acid + 5 vol% perchloric acid + 20 vol% ethanol at 33 V for 150 s. The scanning step of EBSD data was 1 m. The average confidential indexes for each EBSD map were from 0.7 to 0.9. The experiments on f.c.c. materials have shown that when the confidential indexes are greater than 0.1, the fraction of correctly indexed pattern is over 95% [19]. Consequently, the confidence level of the present EBSD maps was high enough. The non-indexed data points and points with low confidential index were usually associated with grain boundary regions. In order to obtain more reliable map, all small grains including less than 5 pixels and data point with confidential index less than 0.1 were nullified (set to zero solutions). To eliminate spurious boundaries caused by orientation noise, a lower limit boundary misorientation was cut-off of 2◦ . The grain size was specified by diameter, which was calculated by determining the area of a grain and then assuming the grain as a circle. The average grain size was calculated by weighted average of the area fraction of grains. Linear TEC was measured by a thermal dilatometer. The dimensional change in specimen was measured during the heating process, from 50 ◦ C to 500 ◦ C, at the heating rate of 5 ◦ C/min. The cylindrical specimens for the TEC test were cut from the base material and center of the welds, with diameter of 3 mm and length of 10 mm. Vickers hardness test and transverse tensile test were employed to examine mechanical properties of the welds at room temperature. Vickers hardness profile was measured with a load of 9.8 N and a dwell time of 10 s along a 1 mm location under the top surface on the cross section of the welds. The dog-bone-like transverse tensile specimens were cut perpendicular to the welding direction with a gauge length of 30 mm. Weld was positioned in the center of gauge length.
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3. Results and discussion Defect-free welds were only acquired at the travelling speed of 2 mm/s by detection of X-ray inspection. Cross-sectional overview of defect-free and defect welds were shown in Fig. 1. Cavities were found in the advancing side (AS) of all 3 mm/s and 5 mm/s welds. These cavities penetrating through the welding direction are the typical tunnel defect which was reported due to the insufficient heat input during FSW [20]. In the present study, travelling speed was the dominate factor to make defect-free weld. So, besides the heat input which was mainly governed by rotational speed, material flow also plays an important role to the formation of FS welds. In relatively higher-travelling-speed welding process (3 mm/s, 5 mm/s), material was stirred from advancing side (AS) to retreating side (RS), but did not have enough capability return to the AS, finally leaded to cavity defect in the bottom of AS. With the decrease of welding speed, the cavities in cross-section trends to disappear, as shown in Fig. 1a, d, and e. The following microstructure and properties investigations will focus on the defect-free welds.
3.1. Microstructure of FS welded Fe–36 wt%Ni alloy As shown in Fig. 1, The stir zones of defect-free FS welds have homogeneous structure in general, except some white zones (WZ) in the top and bottom of AS. High density of inclusions and pits were observed in the WZs, as shown in Fig. 2a. Some WZs also have banded structure in the bottom AS, as shown in Fig. 2b. EDS analysis was used on as-polished WZ to clarify the composition of inclusion. Flaky-shaped inclusions were observed, as shown in Fig. 2c. The EDS spectrums reveal that WZ consists of Fe, Ni peaks as well as B, N peaks. Therefore, the inclusions could be inferred as the CBN debris arising from wearing of the PCBN tool during FSW process. The wear of PCBN tool was also observed in FSW applications to some other materials, such as stainless steel [21]. Optical microstructure in the stir zone center of defect-free FS welds and the base material were shown in Fig. 3. All of them have single austenitic equiaxed microstructure with annealing twins. However, the stir zones are relatively coarser than the base material. The grain boundary maps of base material and the stir zone center of FS welds were shown in Fig. 4. Low angle boundaries (LABs) having misorientation angles of 2–15◦ were marked by thinner pink lines; random high angle boundaries (HABs) having misorientation angles higher than 15◦ were marked by thicker black lines; and 3 twin boundaries, identified with the misorientation angles around 60◦ and misorientation axis around 1 1 1 direction, were marked by thicker gray lines. Factions fraction of LABs and 3 twin boundaries and average grain size in etch grain boundary map were summarized in Table 1. The grain structure, grain boundary characteristics and the recrystallization of FS welded Fe–36 wt%Ni alloy were studied base on the grain boundary maps as follows. Firstly, the grain size as well as grain structure of Fe–36 wt%Ni alloy changed after FSW. The grain size of stir zones is relatively bigger and its distribution is more homogeneous than that of the base material. For the base material, its average grain size is 15 m. It contains some 25–30 m big grains and many 5–10 m fine grains, the big grains are surrounded by fine grains. Accordingly, for the 600 rpm–2 mm/s FS weld, the average grain size is 20 m. Compared with the base material, it is slightly bigger, but the amount of fine grains is obvious reduced, i.e., the grain structure of FS weld becomes more homogeneous. With the increase of rotational speed, the grain structure becomes more homogeneous and coarse. The grain size is 31 m and 34 m in average for 800 rpm and 1000 rpm–2 mm/s welds, respectively.
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Fig. 1. Cross-sectional overviews of defect-free and defect FS welds.
Fig. 2. Optical microstructure (a, b) and EDS analysis of white zone in stir zone (c, d).
Table 1 Grain boundary characteristics and average grain size of the base material and stir zone center of FS welds.
LABs fraction 3 fraction Grain size
Base material
600 rpm, 2 mm/s weld
800 rpm, 2 mm/s weld
1000 rpm, 2 mm/s weld
23% 10% 15 m
25% 9% 20 m
22% 12% 31 m
19% 12% 34 m
Secondly, the LABs fraction and 3 twin boundary fraction of FS welds are comparable to those of the base material, and trends to decrease and increase with increasing of the rotational speed, respectively. The LABs fraction of base material is 23%, while LABs fractions for 600 rpm, 800 rpm and 1000 rpm–2 mm/s welds are 25%, 22% and 19%, respectively. The LABs fractions of FS welds trend to decrease with the increasing of rotational speed in general. The 3 twin boundary fraction of base material is 10%, the 3 twin boundary for 600 rpm, 800 rpm and 1000 rpm–2 mm/s welds are 9%, 12% and 12%, respectively. The 3 twin boundary of FS welds
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Fig. 3. Optical microstructure of base material and the stir zone center of FS welds.
trend to increase with the increasing of the rotational speed. As well known, LABs are always considered as forming during recovery in the deformed microstructure; 3 twin boundaries, which also named annealing twin, are believed forming with grain growth after recrystallization. So, these tendencies of LABs and 3 twin boundary fraction are reasonable, since the increasing of rotational speed results in the enhancement of heat input during welding. The greater heat input leads to more completed recovery, recrystallization even grain growth (i.e., to lower the system energy), produce a microstructure with less LABs and more 3 twin boundaries. Finally, evidence of continuous dynamic recrystallization during FSW of Fe–36 wt%Ni alloy is found in the present grain boundary map. A recovered grain “A” containing of many subgrains was observed, as shown in Fig. 4b. And some subgrain boundaries accumulate a misorientation exceeding 15◦ to high angle boundaries, which was also marked by arrowhead in the figure. This evidence is most likely associated with grain subdivision driven continuous dynamic recrystallization. Some researchers [22,23] also reported the continuous dynamic recrystallization of f.c.c. materials. Moreover, as a medium stacking-fault-energy material (163 mJ/m2 [24]), discontinuous dynamic recrystallization also probably occurred during welding process. Since it is difficult to find the transitional region between the stir zone and base material, i.e., thermal mechanical affected zone, the proof for discontinues dynamic recrystallization during FSW still requires further investigation. 3.2. TEC and mechanical properties The TEC – temperature curves of base material and FS welds were shown in Fig. 5. The curves are significantly increased
around 230 ◦ C, which related to the magnetic transformation of Fe–36 wt%Ni alloy. All welds had comparable TEC to the base material. The small amount of CBN particles did not change the good thermal expansion properties of the base material. Vickers hardness profiles of the FS welds were shown in Fig. 6. The average hardness of 600 rpm, 800 rpm and 1000 rpm–2 mm/s welds are 144 Hv , 140 Hv and 140 Hv , respectively, which is a little lower than the base material (157 Hv ). Besides, with the increasing of rotational speed, the hardness of stir zones trends to decrease. But this tendency is not significant. Transverse tensile properties of the base material and the welds were summarized in Table 2. All welds were broken in the stir zone during tensile test, indicating the stir zone had the lowest hardness in the weld. Similar to the hardness results, the welds exhibit lower ultimate tensile strength and yield strength than the base material. The elongation of FS welds is also decreased, compared with the base material. The reduction of elongation is because of the preferentially yielded stir zone, in other words, the inhomogeneity of deformation of the joint leads to the lower elongation. The strength and elongation of FS welds also have faintly trends to decrease with the increasing of rotational speed. Although the mechanical properties of FS welds are reduced compared with the base material, they are still superior to those of the conventional arc welding and previous FSW of Fe–36 wt%Ni alloy. Gottlieb reported ultimate strength of 407 MPa for gas tungsten arc welded Invar alloy with similar-to-Invar filler [3]. Jasthi reported ultimate strength of 461 MPa and yield strength of 296 MPa for FS welded Fe–36 wt%Ni alloy with PCBN tool [18].
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Fig. 4. Grain boundary maps of base material and the stir zone center of FS welds.
3.3. Relationship of microstructure characteristic and mechanical properties According to the grain boundary strengthening theory, the decrease of grain size will improve the mechanical properties. The relationship between yield strength and grain size of most polycrystalline metals could be described by Hall–Petch relationship, i.e., y = 0 + kd−1/2 , where y and d are yield strength and grain size, respectively; 0 and k are constants independent of grain size.
Table 2 Transverse tensile properties of base material and FS welds.
Fig. 5. Thermal expansion coefficient of base material and stir zones.
Ultimate strength, MPa Yield strength, MPa Elongation, %
Base material
600 rpm, 2 mm/s weld
800 rpm, 2 mm/s weld
1000 rpm, 2 mm/s weld
516 ± 5
485 ± 1
472 ± 15
475 ± 1
393 ± 17 33 ± 3
336 ± 5 19 ± 0
321 ± 26 18 ± 1
327 ± 4 19 ± 1
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strengthen of base material is probably attributed to the difference of grain structure in base material and FS welds. Base material has an inhomogeneous grain size distribution with some big grains surrounded by many fine grains. These fine grains should be helpful to strengthen the base material. 4. Conclusions Friction stir welding was successfully applied to high strength Fe–36 wt%Ni alloy and defect-free welds were acquired at travelling speed of 2 mm/s and rotational speed of 600, 800 and 1000 rpm with PCBN tool. The welds have homogeneous coarse austenite microstructure in the stir zone with comparable LABs and 3 twin boundary fraction to the base material. The outstanding TEC remains for the FS welds, while their mechanical properties are slightly reduced. The reduction of mechanical properties is probably caused by the bigger average grain size and uniform grain structure compared with the base material. Fig. 6. Microhardness of base material and FS welds.
Acknowledgements The authors are grateful to Mr. A. Honda, Prof. G. Miyamoto, Mr. Z.D. Li, Dr. S. Mironov and Dr. K. Kobayashi for technical assistance, and thank Prof. K. Maruyama and Prof. T. Furuhara for their helpful discussions. One of the authors (Y. Zhao) thanks Joint Education Program between Tohoku University and Tsinghua University. We also acknowledge financial support from the Japanese Ministry of Education, Science, Sports and Culture with a Grant-in-Aid from the Global COE program in Materials Integration International Center of Education and Research at Tohoku University. References
Fig. 7. The yield strength and grain size relationship of base material and FS welds.
Since the Hall–Petch relationship was always linear fitted by large grain size range to ensure its accuracy, it is difficult to calculate accurate Hall–Petch relationship of FS welded Fe–36 wt%Ni alloy with the existing data, for the present grain size range was limited from 20 to 34 m. Vinogradov and his colleagues [25] reported the yield strength and grain size (0.18–100 m) of equalchannel-angular-pressed (ECAP-ed) Fe–36 wt%Ni alloy and their results were linear fitted to Hall–Petch relationship. The plot of yield strength against inverse square root of grain size with base material and FS welds was shown and compared with the reported Hall-patch relationship of ECAP-ed Fe–36 wt%Ni alloy in Fig. 7. The FS welds have an about 50 MPa higher strength than the reported Hall–Petch relationship, while the deviation of the base material is about 100 MPa. All the present base material and FS welds deviated from the reported Hall–Petch relationship. It might due to the different preparation methods of base materials and specializations between FSW and ECAP process. However, it could be noticed that the base material has a much bigger deviation from the reported Hall–Petch relationship compared with the FS welds. As shown in Fig. 4, the base material has comparable fraction of 3 annealing twin boundaries and LABs of FS welds. So the
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