- Email: [email protected]
M) AZ31B magnesium alloy
Materials Science and Engineering A 529 (2011) 143–150
Contents lists available at SciVerse ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
A fundamental study of laser welding of hot extruded powder metallurgy (P/M) AZ31B magnesium alloy M. Wahba a,b,∗ , Y. Kawahito c , K. Kondoh c , S. Katayama c a b c
Central Metallurgical Research & Development Institute, Helwan 11421, Egypt Graduate School of Engineering, Osaka University, 1-1 Yamadaoka, Suita, Osaka 565-0871, Japan Joining and Welding Research Institute, Osaka University, 11-1 Mihogaoka, Ibaraki, Osaka 567-0047, Japan
a r t i c l e
i n f o
Article history: Received 16 March 2011 Received in revised form 7 July 2011 Accepted 6 September 2011 Available online 12 September 2011 Keywords: Magnesium alloys Powder metallurgy Laser welding Porosity
a b s t r a c t Laser welding characteristics of a hot extruded AZ31B magnesium alloy produced by rapid solidification powder metallurgy (P/M) were investigated. Extremely porous weld metals were formed under all the employed welding conditions although the base material possessed a fully dense microstructure without pores. Features of porosity morphology in the weld metals as well as X-ray transmission real-time observation results of the molten pool during welding implied that pores originated from the base material. Investigation of the chemical compositions of the base material revealed the presence of a small amount of oxygen and nitrogen. However, the high pressure applied in the initial production processes of the base material caused these gases to expand to large volumes once released when the material was melted during welding, resulting in the formation of large pores. The use of insert-layer of conventionally extruded AZ31B alloy did not improve the situation because a large amount of porosity was formed even when melting of the P/M alloy was minimized. © 2011 Elsevier B.V. All rights reserved.
1. Introduction A growing interest is being given to magnesium alloys due to their extremely light weight. Taking into account other unique characteristics of magnesium alloys such as high specific strength, excellent machinability, good electromagnetic shielding capability and recyclability [1–3], these alloys are very promising to replace steel and aluminum parts in several industrial applications especially where light weight is a critical requirement. Nevertheless, the poor room temperature formability of magnesium alloys is still limiting their widespread applications [4]. Therefore, various techniques including shear rolling [5], rare earths addition [6], equal-channel angular processing [7] and rapid solidification [8] have been investigated to overcome this problem. Based on these investigation results, it was pointed out that high performance magnesium alloys could be produced. It goes without saying that wider application of components made of the newly developed magnesium alloys could be realized by the development of adequate joining technology to integrate these components into different engineering structures and facil-
∗ Corresponding author at: Joining and Welding Research Institute, 11-1 Mihogaoka, Ibaraki, Osaka 567-0047, Japan. Tel.: +81 80 3843 0055; fax: +81 66879 8689. E-mail addresses: [email protected], ma [email protected] (M. Wahba). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.09.010
itate the manufacturing of products with complex geometrical design. The objective of the present work, therefore, is to investigate the laser weldability of rapidly solidified powder metallurgy (P/M) magnesium alloys as no work has been reported regarding this issue. The initial powders were prepared by spinning water atomization process (SWAP) [9]; a process capable of achieving very high cooling rates in the range of 106 K/s by quenching the gas-atomized droplets into a stream of spinning water. Such high cooling rates produce powders with very fine grained microstructure that can improve the mechanical properties of the consolidated material [10]. Previous studies reported on joining P/M parts [11] indicate that porous or low density parts could be joined by solid state joining processes such as friction welding [12] and diffusion bonding [13]. Meanwhile, parts with higher densities or minimal porosity are easily joined by fusion welding processes such as gas tungsten arc welding [14] and laser welding [15]. The material under consideration features a fully dense microstructure, suggesting that fusion welding is applicable. On the other hand, the grain fineness of the microstructure necessitates minimal welding heat input to be involved in order not to deteriorate the mechanical properties upon welding. Owing to its high power density, laser welding can produce deep penetration welds at very high welding speeds [16]. The achievable high welding speeds in laser welding process can minimize the heat input and secure cooling rates fast enough to avoid grain coarsening in the welded material. This advantage has already been utilized in welding ultra-fine grained steel where joints with
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Fig. 2. Schematic illustration of microfocused X-ray transmission real-time imaging system.
2.3. X-ray transmission real-time observation of porosity formation during welding
Fig. 1. Microstructure of base material.
mechanical performance higher than that of the base material could be obtained, as reported in [17]. Moreover, the development of the new generation of high-power and high beam-quality laser sources such as fiber and disc lasers has further increased the capability of the laser welding process [18]. In the present study, a high-brightness disc laser source was used to investigate the weldability of AZ31B magnesium alloy produced by hot extrusion of green compacted SWAP powders. 2. Experimental material and procedures 2.1. Material The base material investigated in this study was 3 mm thick plates of hot extruded P/M AZ31B magnesium alloy (3.14 wt.% Al, 1.08 wt.% Zn, 0.39 wt.% Mn and balance Mg) produced by SWAP. A scanning electron microscope (SEM) photo of a cross-section of the base material is shown in Fig. 1. It is observed that the material exhibits a fully dense and very fine microstructure. Prior to welding, specimens were ground to remove the oxide layer and then cleaned with acetone. 2.2. Laser welding parameters The laser source employed was 16 kW continuous-wave disc laser (Trumpf-TruDisk 16002) producing a beam with 1.03 m wavelength and 8 mm mrad beam parameter product (BPP). A laser beam was transmitted through an optical fiber of 200 m diameter and focused on the workpiece surface by a lens of 280 mm focal length. Two shielding configurations were employed. One includes shielding of the top and bottom surfaces of the specimens by Ar gas flowing at 30 L/min through a 16 mm diameter nozzle and 20 L/min through the holding fixture, respectively. In the second configuration shielding was provided through a gas chamber filled with Ar flowing at 20 L/min. For the sake of simplicity, bead-on-plate welding was first performed parallel to the longitudinal direction in the (XY) plane to investigate the influence of different processing parameters. The applied laser power was varied from 1 to 4 kW with an increment of 1 kW, and the welding speed was varied from 2 to 10 m/min with an increment of 2 m/min. The defocused distance was set at −7 mm (7 mm below the surface).
The keyhole behavior and the situation inside the molten pool during laser welding were observed through a high-speed X-ray transmission real-time imaging system, developed by the authors [19], at a framing rate of 250 f/s. A schematic representation of the system arrangement and specimen sitting with respect to laser beam are displayed in Figs. 2 and 3, respectively. This system consists of a microfocused X-ray tube (160 kV, 1 mA), two image intensifiers and a high speed video camera. The first image intensifier converts the X-ray transmission image into a visible image that is amplified by the second image intensifier. Observations were conducted during bead-on-plate welding parallel to the longitudinal direction in the (XZ) plane. 2.4. Metallography Macrostructures of the weld fusion zones were examined with optical microscope. Metallographic samples were cut across the weld line, mounted, polished and etched with an aqueous solution of 25 mL acetic acid, 25 mL water, 10 g picric acid and 175 mL ethanol. The microstructures and the chemical compositions were investigated using a SEM coupled with energy dispersive X-ray spectroscopy (EDS) and an electron probe micro-analyzer (EPMA). 2.5. Analysis of gas compositions Oxygen and nitrogen contents in the base material were measured by Horiba Oxygen/Nitrogen Analyzer EMGA-520. Gas compositions inside the porosity were analyzed using quadrupole mass analyzer (ANELVA AGS-7000, modified type). The technique involves drilling a hole of about 2 mm diameter inside the weld metal under high vacuum (4 × 10−6 Pa). The mass spectra of the released gases are then measured by a Q-mass spectrometer.
Fig. 3. Schematic illustration of specimen geometry and setting during real-time X-ray observations.
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Fig. 4. Cross-sectional macrostructure of weld beads produced at laser power and welding speed of (a) 4 kW and 4 m/min, (b) 3 kW and 6 m/min, (c) 2 kW and 8 m/min and (d) 1 kW and 10 m/min, respectively.
2.6. Welding experiments for porosity reduction In the final set of welding experiments, butt joints with an insert-layer of conventionally extruded-cast AZ31B magnesium alloy between specimens of extruded-P/M AZ31B magnesium alloy were produced. The insert-layer was 1.3 mm thick. Laser power was set at 3 kW and welding speed was varied from 4 to 8 m/min.
3. Results and discussion 3.1. Bead-on-plate welding Contrary to expected, an extremely porous fusion zone was formed in all welded specimens. Fig. 4 shows some typical exam-
ples of fusion zone macrostructure of weld beads produced at different laser powers and welding speeds. Variation of laser welding parameters influenced only the fusion zone size and did not affect the porosity formation propensity. A large number of micro and macropores with spherical or irregular shapes were formed under all the welding conditions. This abnormal amount of porosity displaced the weld metal resulting in large excess weld metal and excessive penetration, as indicated in Fig. 4(b) and (c). Moreover, the welded joints were so weak that they could be easily broken by hand. It should be noted that the shielding configuration did not have any influence as the same results were obtained in the case of shielding nozzle as well as in the case of gas chamber. Investigation of the fusion zone microstructure revealed that many micropores were present near and along the fusion boundary, as shown in Fig. 5. Also, it was observed that some pores near
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Fig. 5. Pores and their morphology near fusion boundary.
the fusion boundary were elongated and featuring a growth direction towards the inside of the fusion zone starting from the fusion boundary. Two examples of such pores are marked with white arrows in Fig. 5. In general, the mechanism of porosity formation during laser welding of magnesium alloys should be considered in relation with the production method of the base material. The authors have previously investigated the laser weldability of extruded AZ31B and AZ61A alloys that did not contain any pre-existing porosity and diecast AZ91D alloy that had 2% pre-existing porosity content. Porosity formation in the case of wrought alloys is not a serious problem, and joints with tensile strength similar to that of the base material could be obtained by stabilizing the keyhole [20]. On the other hand, gases entrapped in the pre-existing pores in the base material resulted in the formation of porous weld metal in the case of the die-cast alloy [21]. Similar conclusions were reported regarding wrought [22] as well as die-cast [23] alloys. The pore morphology displayed in Fig. 5 suggests similar behavior to that of the die-cast alloy, i.e. the porosity was formed by infusion of gases from the base material. However, it has previously been indicated that the base material microstructure did not contain any porosity (Fig. 1).
Fig. 7. Cross-sectional macrostructure of welded specimen subjected to X-ray transmission real-time observation.
Also, the observed pores and weld metal displacement are much larger than those in the case of the die-cast alloy. Therefore, more in-depth investigations were carried out to reveal the mechanism of this abnormal pore formation and consequently to find a suitable measure. 3.2. X-ray transmission real-time observation of keyhole and molten pool During laser beam welding, the keyhole behavior and porosity evolution inside the weld pool were observed using a high-speed real-time X-ray imaging system. It was observed that once a laser beam was irradiated and a keyhole was formed, many bubbles started to emerge from the weld pool. Bubbles could be identified only at the upper part of the weld pool. No melt flows were observed to drift the bubbles backwards in the weld pool instead; bubbles displaced the weld metal on their way out from the weld pool. The molten metal was displaced to a height of approximately
Fig. 6. Bubbles generation and molten metal displacement.
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Fig. 8. Microstructure and EDS analysis of the base material: (a) SEM photo of the microstructure indicating the locations for the analysis, (b) EDS spectrum for the matrix (spectrum 1 in (a)) and (c) EDS spectrum for the white particles (spectrum 2 in (a)).
4 mm over the specimen surface while the keyhole penetrated to a depth of approximately 6 mm. Some bubbles could escape out from the weld metal but many others were captured by the rapid solidification resulting in porosity. Snapshots from the real-time X-ray transmission observation results during laser beam welding, taken at 0.048 s, 0.104 s and 0.664 s after the laser beam was switched on, are shown in Fig. 6 illustrating the process of bubbles generation and weld metal displacement. At time t = 0.048 s the height of the displaced molten metal was approximately 2 mm. This height was increased to 4 mm at time t = 0.104 s. As time progressed, the situation remained unchanged; generation of many bubbles and displacement of the molten metal, as shown at time t = 0.664 s (see the Supplementary video).
Table 1 Gas compositions in porosity, %. Gas component a
H2 (m/z 2) CHGb (m/z 15) N2 (m/z 28) Ar (m/z 40) CHG (m/z 43) CO2 (m/z 44) CHG (m/z 57) a b
m/z refers to mass-to-charge ratio. CHG refers to hydrocarbon gas.
Point 1 – – 100 <0.1 – – –
Point 2 – – 99.7 0.3 – – –
Also it was observed that neither the rear wall of the keyhole nor bubbles coming from the bottom part of the weld pool could be identified. This is believed to be due to the presence of large volume of gases and a small amount of molten metal in this part of the weld pool. Examination of the weld bead cross-section confirmed the presence of very little weld metal in the weld fusion zone part below the excess weld metal, as displayed in Fig. 7. 3.3. Analysis of gas compositions inside the porosity Gases retained inside the weld metal porosity were analyzed at two different locations referred to as point 1 and point 2. The data are given in Table 1. The analytical results reveal that the porosity contained only nitrogen gas and a very small amount of argon. Recalling that welding inside a gas chamber did not prevent the porosity formation, this nitrogen is believed to come from the base material. It is also notable that hydrogen was not detected in the porosity. Hydrogen rejection at the solid–liquid interface can contribute to porosity formation at the last solidification stages of Table 2 Nitrogen and oxygen contents in the base material, ppm. Specimen
Oxygen Mean
Extruded-P/M AZ31B Extruded-cast AZ31B
14.95 6.2
Nitrogen STD 1.8 0.3
Mean 24.93 16.4
STD 0.5 1.3
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Fig. 9. EPMA analytical results of base material, showing backscattered electron image (a), and distribution maps of magnesium (b), oxygen (c), aluminum (d), zinc (f), and manganese (e).
Fig. 10. Fractured surfaces and EDS analysis of a welded joint: (a) SEM photo of the fractured surface, (b) magnified view of the black boxed region in (a), (c) magnified view of the white boxed region in (a), (d) EDS spectrum for the white particles (spectrum 1 in (b) and (c)), and (e) EDS spectrum for the fracture surface (spectrum 2 in (b) and (c)).
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Fig. 11. Macrostructural weld fusion zones of butt joints with insert-layer in extruded-P/M AZ31B magnesium alloy produced at 3 kW laser power and (a) 4 m/min, (b) 6 m/min and (c) 8 m/min welding speed.
magnesium alloys and some hydrogen had already been detected inside the porosity of the die-cast AZ91D alloy welds [21]. However, hydrogen rejection was reported to occur at the Mg17 Al12 intermetallic compound in the alloy AZ91 (containing approx. 9 wt.% Al) [23]. Besides the difference in hydrogen solubility between the liquid and solid phases in magnesium is less than that in aluminum [24] which means that hydrogen rejection is less in the case of magnesium. Therefore, the absence of Mg17 Al12 phase in the rapidly solidified P/M alloy microstructure (as will be indicated later) might be the reason why hydrogen was not detected in the porosity. 3.4. Chemical compositions analysis of base material Fig. 8 shows a SEM image with some EDS spectrum analyses of the base material. In general, the microstructure of a conventionally extruded AZ31B magnesium alloy consists of ␣-magnesium matrix and intermetallic compound -Mg17 Al12 [1]. However, the very rapid solidification rates associated with the SWAP retard the precipitation of the intermetallic compounds and produces super saturated powders, as reported in [8]. Fig. 8(a) indicates that the microstructure comprises a fine grained matrix and white particles aligned in the extrusion direction. As can be seen from Fig. 8(b) and (c), nitrogen was not detected in neither the matrix nor the white particles and these particles contained mainly magnesium and oxygen. This was confirmed by EPMA mapping photos displayed in Fig. 9. Accordingly, the base material microstructure could be considered to consist of a super saturated ␣-magnesium matrix containing particles of magnesium oxide and no nitrogencontaining compounds could be detected. To clarify the source of the nitrogen gas detected in the weld metal porosity, a specimen was prepared from the base material and tested in a gas analyzer where it is melted and the evolved gases are analyzed. The results are given in Table 2. A specimen of conventionally extruded AZ31B alloy was also examined for comparison. It is seen that the extruded-P/M alloy contained slightly higher amounts of oxygen and nitrogen compared with the extrudedcast alloy. It is believed that these gases were contained in the base material in the initial production processes since the green compacted powder had a relative density of 92%. Besides, SWAP powders were reported to contain some oxides [9]. Due to its higher affinity to magnesium, oxygen reacted with magnesium forming oxides while nitrogen remained in an uncombined state inside the base material since it was not detected by EDS or EPMA
analysis. This would also explain the reason why only nitrogen gas was detected in the weld metal porosity. Once the base material is melted during welding, nitrogen and some of the oxygen are released forming bubbles in the molten metal. Oxygen reacts with magnesium whereas nitrogen is retained inside the weld metal porosity. Fig. 10 shows SEM photos of the fractured surfaces of a welded joint and the corresponding EDS spectrum. Small white particles were observed inside and outside the pores, as seen in Fig. 10(b) and (c). The EDS spectrum in Fig. 10(d) indicates that these white particles are magnesium oxide. These particles, most probably, were formed by the reaction of oxygen with magnesium from the metal vapor. Oxygen was also detected on the fractured surface inside and outside the pores, as shown in Fig. 10(e). Based on these results it could be concluded that when the material was melted during welding, nitrogen that was initially entrapped in the base material during the production processes was released and led to the formation of large pores in the weld metal. However, such small amount of nitrogen detected in the base material cannot explain alone the abnormally large pores observed in the weld metal particularly porosity was hardly formed in the extruded-cast alloy welds [20] although the difference in gas contents between the two materials is not so high. The applied pressure during powder compaction and extrusion processes has to be taken into consideration to adequately analyze the abnormal pore formation upon welding. The initial SWAP powder was compacted under a pressure of more than 3800 atm. and then hot extruded with an extrusion ratio of 42. This means that the pressure of gases inside the material would be very much higher than the atmospheric pressure. The liberated gas during welding expands to release the pressure forming large pores and displacing the molten metal out from the weld pool. Another factor that might have contributed to the abnormal pore formation is that nitrogen does not dissolve in liquid magnesium as I.J. Polmear indicated that hydrogen is the only gas that dissolves in molten magnesium [24]. 3.5. Porosity reduction approach The above mentioned results indicate that the weld metal porosity originated from the pre-existing gases in the base material. Therefore, minimizing the portion of the base material to be melted during welding might be a good measure to reduce the pore formation. Accordingly, butt joints of extruded-P/M AZ31B magnesium
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alloy were designed with an insert-layer of extruded-cast AZ31B magnesium alloy comprising the same chemical compositions. Despite there is no big difference in the gas contents between the two materials, this approach might help reduce the porosity because gases inside the extruded-cast alloy are not retained under high pressure. Besides, the approach has already been applied in the case of die-cast alloy and the porosity was significantly reduced [21]. Fig. 11 displays some typical examples of fusion zone macrostructure of butt-welded joints with the insert-layer. In general, a smaller number of pores were formed compared with autogenous welds. However, the pore proportion is still high as well as the excess weld metal and the excessive penetration, as seen in Fig. 11(a). A reduction in the energy input by increasing the welding speed led to a decrease in the melted part of P/M alloy and consequently a reduction in the weld metal porosity, as observed in Fig. 11(b) and (c). Nevertheless, the energy input was not sufficient to achieve adequate fusion. Moreover, even in this minimized melted part of the P/M alloy significant porosity was formed particularly near and along the fusion boundary. The difference in the effectiveness of the porosity reduction approach between this case and the die-cast alloy case [21] might be attributed to the higher pressure levels of the entrapped gases of the P/M alloy. These results confirm the proposed mechanism of porosity evolution on one hand, but on the other hand, indicate that it is very difficult to produce sound welded joints in the material under consideration using fusion welding. Solid-state welding might give better results. 4. Conclusions Laser beam weldability of hot extruded P/M AZ31B magnesium alloy was investigated. The welding process was characterized by abnormal pore formation and ejection of the weld metal although no porosity was observed in the base material microstructure. This abnormal pore formation was attributed to the expansion of gases that were entrapped in the base material under high pressure during the initial production processes and later on released during welding. Welding with an insert-layer of conventionally extruded alloy was not effective because large proportion of porosity was
still produced even by minimizing the melted part of the P/M alloy, revealing that fusion welding of this alloy is very difficult. Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at doi:10.1016/j.msea.2011.09.010. References [1] M. Avedesian, H. Baker, ASM Specialty Handbook: Magnesium and Magnesium Alloys, ASM International, Ohio, 1999. [2] K.U. Kainer, F. von Buch, in: K.U. Kainer (Ed.), Magnesium Alloys and Technology, Wiley-VCH, Weinheim, 2003, pp. 1–22. [3] B.L. Mordike, T. Ebert, Mater. Sci. Eng. A 302 (2001) 37–45. [4] S.R. Agnew, M.H. Yoo, C.N. Tome, Acta Mater. 49 (2001) 4277–4289. [5] X. Huang, K. Suzuki, A. Watazu, I. Shigematsu, N. Saito, J. Alloys Compd. 457 (2008) 408–412. [6] Y. Chino, M. Kado, M. Mabuchi, Mater. Sci. Eng. A 494 (2008) 343–349. [7] S. Suwas, G. Gottstein, R. Kumar, Mater. Sci. Eng. A 471 (2007) 1–14. [8] E. Ayman, U. Junko, K. Katsuyoshi, Acta Mater. 59 (2011) 273–282. [9] I. Endo, I. Otsuka, R. Okuno, A. Shintami, M. Yoshino, M. Yagi, IEEE Trans. Magn. 35 (1999) 3385–3387. [10] I. Hisashi, K. Masashi, K. Katsuyoshi, O. Isamu, I. Hiroshi, Trans. JWRI 36 (2007) 33–38. [11] C. Selcuk, S. Bond, P. Woollin, Powder Metall. 53 (2010) 7–11. [12] K. Jayabharath, M. Ashfaq, P. Venugopal, D.R.G. Achar, Mater. Sci. Eng. A 454–455 (2007) 114–123. [13] M.G. Fillabi, A. Simchi, A.H. Kokabi, Mater. Des. 29 (2008) 411–417. [14] M.V. Suresh, B. Vamsi Krishna, P. Venugopal, K. Prasad Rao, Sci. Technol. Weld. Join. 9 (2004) 362–368. [15] J.A. Hamill, Jr., Peter Wirth, SAE Tech. Paper 940355, SAE International Congress, Detroit, MI, February 1994. [16] W.M. Steen, Laser Material Processing, 3rd ed., Springer, London, 2003. [17] M. Gao, X. Zeng, Q. Hu, J. Yan, J. Mater. Process. Technol. 209 (2009) 785–791. [18] H. Hügel, Opt. Laser Eng. 34 (2000) 213–229. [19] N. Seto, S. Katayama, A. Matsunawa, J. Laser Appl. 12 (2000) 245–250. [20] M. Wahba, M. Mizutani, Y. Kawahito, S. Katayama, Sci. Technol. Weld. Join. 15 (2010) 559–566. [21] M. Wahba, M. Mizutani, Y. Kawahito, S. Katayama, Mater. Des. (2011), doi:10.1016/j.matdes.2011.05.016. [22] R.S. Coelho, A. Kostka, H. Pinto, S. Reikeher, M. Kocak, A.R. Pyzalla, Mater. Sci. Eng. A 485 (2008) 20–30. [23] H. Zhao, T. DebRoy, Weld. J. 80 (2001) 204–210. [24] I.J. Polmear, Light Alloys from Traditional Alloys to Nanocrystals, 4th ed., Butterworth-Heinemann, Oxford, 2006.
Contents lists available at SciVerse ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
A fundamental study of laser welding of hot extruded powder metallurgy (P/M) AZ31B magnesium alloy M. Wahba a,b,∗ , Y. Kawahito c , K. Kondoh c , S. Katayama c a b c
Central Metallurgical Research & Development Institute, Helwan 11421, Egypt Graduate School of Engineering, Osaka University, 1-1 Yamadaoka, Suita, Osaka 565-0871, Japan Joining and Welding Research Institute, Osaka University, 11-1 Mihogaoka, Ibaraki, Osaka 567-0047, Japan
a r t i c l e
i n f o
Article history: Received 16 March 2011 Received in revised form 7 July 2011 Accepted 6 September 2011 Available online 12 September 2011 Keywords: Magnesium alloys Powder metallurgy Laser welding Porosity
a b s t r a c t Laser welding characteristics of a hot extruded AZ31B magnesium alloy produced by rapid solidification powder metallurgy (P/M) were investigated. Extremely porous weld metals were formed under all the employed welding conditions although the base material possessed a fully dense microstructure without pores. Features of porosity morphology in the weld metals as well as X-ray transmission real-time observation results of the molten pool during welding implied that pores originated from the base material. Investigation of the chemical compositions of the base material revealed the presence of a small amount of oxygen and nitrogen. However, the high pressure applied in the initial production processes of the base material caused these gases to expand to large volumes once released when the material was melted during welding, resulting in the formation of large pores. The use of insert-layer of conventionally extruded AZ31B alloy did not improve the situation because a large amount of porosity was formed even when melting of the P/M alloy was minimized. © 2011 Elsevier B.V. All rights reserved.
1. Introduction A growing interest is being given to magnesium alloys due to their extremely light weight. Taking into account other unique characteristics of magnesium alloys such as high specific strength, excellent machinability, good electromagnetic shielding capability and recyclability [1–3], these alloys are very promising to replace steel and aluminum parts in several industrial applications especially where light weight is a critical requirement. Nevertheless, the poor room temperature formability of magnesium alloys is still limiting their widespread applications [4]. Therefore, various techniques including shear rolling [5], rare earths addition [6], equal-channel angular processing [7] and rapid solidification [8] have been investigated to overcome this problem. Based on these investigation results, it was pointed out that high performance magnesium alloys could be produced. It goes without saying that wider application of components made of the newly developed magnesium alloys could be realized by the development of adequate joining technology to integrate these components into different engineering structures and facil-
∗ Corresponding author at: Joining and Welding Research Institute, 11-1 Mihogaoka, Ibaraki, Osaka 567-0047, Japan. Tel.: +81 80 3843 0055; fax: +81 66879 8689. E-mail addresses: [email protected], ma [email protected] (M. Wahba). 0921-5093/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2011.09.010
itate the manufacturing of products with complex geometrical design. The objective of the present work, therefore, is to investigate the laser weldability of rapidly solidified powder metallurgy (P/M) magnesium alloys as no work has been reported regarding this issue. The initial powders were prepared by spinning water atomization process (SWAP) [9]; a process capable of achieving very high cooling rates in the range of 106 K/s by quenching the gas-atomized droplets into a stream of spinning water. Such high cooling rates produce powders with very fine grained microstructure that can improve the mechanical properties of the consolidated material [10]. Previous studies reported on joining P/M parts [11] indicate that porous or low density parts could be joined by solid state joining processes such as friction welding [12] and diffusion bonding [13]. Meanwhile, parts with higher densities or minimal porosity are easily joined by fusion welding processes such as gas tungsten arc welding [14] and laser welding [15]. The material under consideration features a fully dense microstructure, suggesting that fusion welding is applicable. On the other hand, the grain fineness of the microstructure necessitates minimal welding heat input to be involved in order not to deteriorate the mechanical properties upon welding. Owing to its high power density, laser welding can produce deep penetration welds at very high welding speeds [16]. The achievable high welding speeds in laser welding process can minimize the heat input and secure cooling rates fast enough to avoid grain coarsening in the welded material. This advantage has already been utilized in welding ultra-fine grained steel where joints with
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Fig. 2. Schematic illustration of microfocused X-ray transmission real-time imaging system.
2.3. X-ray transmission real-time observation of porosity formation during welding
Fig. 1. Microstructure of base material.
mechanical performance higher than that of the base material could be obtained, as reported in [17]. Moreover, the development of the new generation of high-power and high beam-quality laser sources such as fiber and disc lasers has further increased the capability of the laser welding process [18]. In the present study, a high-brightness disc laser source was used to investigate the weldability of AZ31B magnesium alloy produced by hot extrusion of green compacted SWAP powders. 2. Experimental material and procedures 2.1. Material The base material investigated in this study was 3 mm thick plates of hot extruded P/M AZ31B magnesium alloy (3.14 wt.% Al, 1.08 wt.% Zn, 0.39 wt.% Mn and balance Mg) produced by SWAP. A scanning electron microscope (SEM) photo of a cross-section of the base material is shown in Fig. 1. It is observed that the material exhibits a fully dense and very fine microstructure. Prior to welding, specimens were ground to remove the oxide layer and then cleaned with acetone. 2.2. Laser welding parameters The laser source employed was 16 kW continuous-wave disc laser (Trumpf-TruDisk 16002) producing a beam with 1.03 m wavelength and 8 mm mrad beam parameter product (BPP). A laser beam was transmitted through an optical fiber of 200 m diameter and focused on the workpiece surface by a lens of 280 mm focal length. Two shielding configurations were employed. One includes shielding of the top and bottom surfaces of the specimens by Ar gas flowing at 30 L/min through a 16 mm diameter nozzle and 20 L/min through the holding fixture, respectively. In the second configuration shielding was provided through a gas chamber filled with Ar flowing at 20 L/min. For the sake of simplicity, bead-on-plate welding was first performed parallel to the longitudinal direction in the (XY) plane to investigate the influence of different processing parameters. The applied laser power was varied from 1 to 4 kW with an increment of 1 kW, and the welding speed was varied from 2 to 10 m/min with an increment of 2 m/min. The defocused distance was set at −7 mm (7 mm below the surface).
The keyhole behavior and the situation inside the molten pool during laser welding were observed through a high-speed X-ray transmission real-time imaging system, developed by the authors [19], at a framing rate of 250 f/s. A schematic representation of the system arrangement and specimen sitting with respect to laser beam are displayed in Figs. 2 and 3, respectively. This system consists of a microfocused X-ray tube (160 kV, 1 mA), two image intensifiers and a high speed video camera. The first image intensifier converts the X-ray transmission image into a visible image that is amplified by the second image intensifier. Observations were conducted during bead-on-plate welding parallel to the longitudinal direction in the (XZ) plane. 2.4. Metallography Macrostructures of the weld fusion zones were examined with optical microscope. Metallographic samples were cut across the weld line, mounted, polished and etched with an aqueous solution of 25 mL acetic acid, 25 mL water, 10 g picric acid and 175 mL ethanol. The microstructures and the chemical compositions were investigated using a SEM coupled with energy dispersive X-ray spectroscopy (EDS) and an electron probe micro-analyzer (EPMA). 2.5. Analysis of gas compositions Oxygen and nitrogen contents in the base material were measured by Horiba Oxygen/Nitrogen Analyzer EMGA-520. Gas compositions inside the porosity were analyzed using quadrupole mass analyzer (ANELVA AGS-7000, modified type). The technique involves drilling a hole of about 2 mm diameter inside the weld metal under high vacuum (4 × 10−6 Pa). The mass spectra of the released gases are then measured by a Q-mass spectrometer.
Fig. 3. Schematic illustration of specimen geometry and setting during real-time X-ray observations.
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Fig. 4. Cross-sectional macrostructure of weld beads produced at laser power and welding speed of (a) 4 kW and 4 m/min, (b) 3 kW and 6 m/min, (c) 2 kW and 8 m/min and (d) 1 kW and 10 m/min, respectively.
2.6. Welding experiments for porosity reduction In the final set of welding experiments, butt joints with an insert-layer of conventionally extruded-cast AZ31B magnesium alloy between specimens of extruded-P/M AZ31B magnesium alloy were produced. The insert-layer was 1.3 mm thick. Laser power was set at 3 kW and welding speed was varied from 4 to 8 m/min.
3. Results and discussion 3.1. Bead-on-plate welding Contrary to expected, an extremely porous fusion zone was formed in all welded specimens. Fig. 4 shows some typical exam-
ples of fusion zone macrostructure of weld beads produced at different laser powers and welding speeds. Variation of laser welding parameters influenced only the fusion zone size and did not affect the porosity formation propensity. A large number of micro and macropores with spherical or irregular shapes were formed under all the welding conditions. This abnormal amount of porosity displaced the weld metal resulting in large excess weld metal and excessive penetration, as indicated in Fig. 4(b) and (c). Moreover, the welded joints were so weak that they could be easily broken by hand. It should be noted that the shielding configuration did not have any influence as the same results were obtained in the case of shielding nozzle as well as in the case of gas chamber. Investigation of the fusion zone microstructure revealed that many micropores were present near and along the fusion boundary, as shown in Fig. 5. Also, it was observed that some pores near
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Fig. 5. Pores and their morphology near fusion boundary.
the fusion boundary were elongated and featuring a growth direction towards the inside of the fusion zone starting from the fusion boundary. Two examples of such pores are marked with white arrows in Fig. 5. In general, the mechanism of porosity formation during laser welding of magnesium alloys should be considered in relation with the production method of the base material. The authors have previously investigated the laser weldability of extruded AZ31B and AZ61A alloys that did not contain any pre-existing porosity and diecast AZ91D alloy that had 2% pre-existing porosity content. Porosity formation in the case of wrought alloys is not a serious problem, and joints with tensile strength similar to that of the base material could be obtained by stabilizing the keyhole [20]. On the other hand, gases entrapped in the pre-existing pores in the base material resulted in the formation of porous weld metal in the case of the die-cast alloy [21]. Similar conclusions were reported regarding wrought [22] as well as die-cast [23] alloys. The pore morphology displayed in Fig. 5 suggests similar behavior to that of the die-cast alloy, i.e. the porosity was formed by infusion of gases from the base material. However, it has previously been indicated that the base material microstructure did not contain any porosity (Fig. 1).
Fig. 7. Cross-sectional macrostructure of welded specimen subjected to X-ray transmission real-time observation.
Also, the observed pores and weld metal displacement are much larger than those in the case of the die-cast alloy. Therefore, more in-depth investigations were carried out to reveal the mechanism of this abnormal pore formation and consequently to find a suitable measure. 3.2. X-ray transmission real-time observation of keyhole and molten pool During laser beam welding, the keyhole behavior and porosity evolution inside the weld pool were observed using a high-speed real-time X-ray imaging system. It was observed that once a laser beam was irradiated and a keyhole was formed, many bubbles started to emerge from the weld pool. Bubbles could be identified only at the upper part of the weld pool. No melt flows were observed to drift the bubbles backwards in the weld pool instead; bubbles displaced the weld metal on their way out from the weld pool. The molten metal was displaced to a height of approximately
Fig. 6. Bubbles generation and molten metal displacement.
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Fig. 8. Microstructure and EDS analysis of the base material: (a) SEM photo of the microstructure indicating the locations for the analysis, (b) EDS spectrum for the matrix (spectrum 1 in (a)) and (c) EDS spectrum for the white particles (spectrum 2 in (a)).
4 mm over the specimen surface while the keyhole penetrated to a depth of approximately 6 mm. Some bubbles could escape out from the weld metal but many others were captured by the rapid solidification resulting in porosity. Snapshots from the real-time X-ray transmission observation results during laser beam welding, taken at 0.048 s, 0.104 s and 0.664 s after the laser beam was switched on, are shown in Fig. 6 illustrating the process of bubbles generation and weld metal displacement. At time t = 0.048 s the height of the displaced molten metal was approximately 2 mm. This height was increased to 4 mm at time t = 0.104 s. As time progressed, the situation remained unchanged; generation of many bubbles and displacement of the molten metal, as shown at time t = 0.664 s (see the Supplementary video).
Table 1 Gas compositions in porosity, %. Gas component a
H2 (m/z 2) CHGb (m/z 15) N2 (m/z 28) Ar (m/z 40) CHG (m/z 43) CO2 (m/z 44) CHG (m/z 57) a b
m/z refers to mass-to-charge ratio. CHG refers to hydrocarbon gas.
Point 1 – – 100 <0.1 – – –
Point 2 – – 99.7 0.3 – – –
Also it was observed that neither the rear wall of the keyhole nor bubbles coming from the bottom part of the weld pool could be identified. This is believed to be due to the presence of large volume of gases and a small amount of molten metal in this part of the weld pool. Examination of the weld bead cross-section confirmed the presence of very little weld metal in the weld fusion zone part below the excess weld metal, as displayed in Fig. 7. 3.3. Analysis of gas compositions inside the porosity Gases retained inside the weld metal porosity were analyzed at two different locations referred to as point 1 and point 2. The data are given in Table 1. The analytical results reveal that the porosity contained only nitrogen gas and a very small amount of argon. Recalling that welding inside a gas chamber did not prevent the porosity formation, this nitrogen is believed to come from the base material. It is also notable that hydrogen was not detected in the porosity. Hydrogen rejection at the solid–liquid interface can contribute to porosity formation at the last solidification stages of Table 2 Nitrogen and oxygen contents in the base material, ppm. Specimen
Oxygen Mean
Extruded-P/M AZ31B Extruded-cast AZ31B
14.95 6.2
Nitrogen STD 1.8 0.3
Mean 24.93 16.4
STD 0.5 1.3
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Fig. 9. EPMA analytical results of base material, showing backscattered electron image (a), and distribution maps of magnesium (b), oxygen (c), aluminum (d), zinc (f), and manganese (e).
Fig. 10. Fractured surfaces and EDS analysis of a welded joint: (a) SEM photo of the fractured surface, (b) magnified view of the black boxed region in (a), (c) magnified view of the white boxed region in (a), (d) EDS spectrum for the white particles (spectrum 1 in (b) and (c)), and (e) EDS spectrum for the fracture surface (spectrum 2 in (b) and (c)).
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Fig. 11. Macrostructural weld fusion zones of butt joints with insert-layer in extruded-P/M AZ31B magnesium alloy produced at 3 kW laser power and (a) 4 m/min, (b) 6 m/min and (c) 8 m/min welding speed.
magnesium alloys and some hydrogen had already been detected inside the porosity of the die-cast AZ91D alloy welds [21]. However, hydrogen rejection was reported to occur at the Mg17 Al12 intermetallic compound in the alloy AZ91 (containing approx. 9 wt.% Al) [23]. Besides the difference in hydrogen solubility between the liquid and solid phases in magnesium is less than that in aluminum [24] which means that hydrogen rejection is less in the case of magnesium. Therefore, the absence of Mg17 Al12 phase in the rapidly solidified P/M alloy microstructure (as will be indicated later) might be the reason why hydrogen was not detected in the porosity. 3.4. Chemical compositions analysis of base material Fig. 8 shows a SEM image with some EDS spectrum analyses of the base material. In general, the microstructure of a conventionally extruded AZ31B magnesium alloy consists of ␣-magnesium matrix and intermetallic compound -Mg17 Al12 [1]. However, the very rapid solidification rates associated with the SWAP retard the precipitation of the intermetallic compounds and produces super saturated powders, as reported in [8]. Fig. 8(a) indicates that the microstructure comprises a fine grained matrix and white particles aligned in the extrusion direction. As can be seen from Fig. 8(b) and (c), nitrogen was not detected in neither the matrix nor the white particles and these particles contained mainly magnesium and oxygen. This was confirmed by EPMA mapping photos displayed in Fig. 9. Accordingly, the base material microstructure could be considered to consist of a super saturated ␣-magnesium matrix containing particles of magnesium oxide and no nitrogencontaining compounds could be detected. To clarify the source of the nitrogen gas detected in the weld metal porosity, a specimen was prepared from the base material and tested in a gas analyzer where it is melted and the evolved gases are analyzed. The results are given in Table 2. A specimen of conventionally extruded AZ31B alloy was also examined for comparison. It is seen that the extruded-P/M alloy contained slightly higher amounts of oxygen and nitrogen compared with the extrudedcast alloy. It is believed that these gases were contained in the base material in the initial production processes since the green compacted powder had a relative density of 92%. Besides, SWAP powders were reported to contain some oxides [9]. Due to its higher affinity to magnesium, oxygen reacted with magnesium forming oxides while nitrogen remained in an uncombined state inside the base material since it was not detected by EDS or EPMA
analysis. This would also explain the reason why only nitrogen gas was detected in the weld metal porosity. Once the base material is melted during welding, nitrogen and some of the oxygen are released forming bubbles in the molten metal. Oxygen reacts with magnesium whereas nitrogen is retained inside the weld metal porosity. Fig. 10 shows SEM photos of the fractured surfaces of a welded joint and the corresponding EDS spectrum. Small white particles were observed inside and outside the pores, as seen in Fig. 10(b) and (c). The EDS spectrum in Fig. 10(d) indicates that these white particles are magnesium oxide. These particles, most probably, were formed by the reaction of oxygen with magnesium from the metal vapor. Oxygen was also detected on the fractured surface inside and outside the pores, as shown in Fig. 10(e). Based on these results it could be concluded that when the material was melted during welding, nitrogen that was initially entrapped in the base material during the production processes was released and led to the formation of large pores in the weld metal. However, such small amount of nitrogen detected in the base material cannot explain alone the abnormally large pores observed in the weld metal particularly porosity was hardly formed in the extruded-cast alloy welds [20] although the difference in gas contents between the two materials is not so high. The applied pressure during powder compaction and extrusion processes has to be taken into consideration to adequately analyze the abnormal pore formation upon welding. The initial SWAP powder was compacted under a pressure of more than 3800 atm. and then hot extruded with an extrusion ratio of 42. This means that the pressure of gases inside the material would be very much higher than the atmospheric pressure. The liberated gas during welding expands to release the pressure forming large pores and displacing the molten metal out from the weld pool. Another factor that might have contributed to the abnormal pore formation is that nitrogen does not dissolve in liquid magnesium as I.J. Polmear indicated that hydrogen is the only gas that dissolves in molten magnesium [24]. 3.5. Porosity reduction approach The above mentioned results indicate that the weld metal porosity originated from the pre-existing gases in the base material. Therefore, minimizing the portion of the base material to be melted during welding might be a good measure to reduce the pore formation. Accordingly, butt joints of extruded-P/M AZ31B magnesium
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alloy were designed with an insert-layer of extruded-cast AZ31B magnesium alloy comprising the same chemical compositions. Despite there is no big difference in the gas contents between the two materials, this approach might help reduce the porosity because gases inside the extruded-cast alloy are not retained under high pressure. Besides, the approach has already been applied in the case of die-cast alloy and the porosity was significantly reduced [21]. Fig. 11 displays some typical examples of fusion zone macrostructure of butt-welded joints with the insert-layer. In general, a smaller number of pores were formed compared with autogenous welds. However, the pore proportion is still high as well as the excess weld metal and the excessive penetration, as seen in Fig. 11(a). A reduction in the energy input by increasing the welding speed led to a decrease in the melted part of P/M alloy and consequently a reduction in the weld metal porosity, as observed in Fig. 11(b) and (c). Nevertheless, the energy input was not sufficient to achieve adequate fusion. Moreover, even in this minimized melted part of the P/M alloy significant porosity was formed particularly near and along the fusion boundary. The difference in the effectiveness of the porosity reduction approach between this case and the die-cast alloy case [21] might be attributed to the higher pressure levels of the entrapped gases of the P/M alloy. These results confirm the proposed mechanism of porosity evolution on one hand, but on the other hand, indicate that it is very difficult to produce sound welded joints in the material under consideration using fusion welding. Solid-state welding might give better results. 4. Conclusions Laser beam weldability of hot extruded P/M AZ31B magnesium alloy was investigated. The welding process was characterized by abnormal pore formation and ejection of the weld metal although no porosity was observed in the base material microstructure. This abnormal pore formation was attributed to the expansion of gases that were entrapped in the base material under high pressure during the initial production processes and later on released during welding. Welding with an insert-layer of conventionally extruded alloy was not effective because large proportion of porosity was
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