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Solid state joining of nickel based alloy, Haynes 230
Journal of Materials Processing Technology 225 (2015) 492–499
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Journal of Materials Processing Technology journal homepage: www.elsevier.com/locate/jmatprotec
Solid state joining of nickel based alloy, Haynes 230 J.A. Schneider a,∗ , D. Williston a , T.L. Murphy a , C. Varner a , J. Hawkins a , B. Walker b a b
Department of Mechanical Engineering, Mississippi State University, Starkville, MS 39762, United States Keystone Synergistic Enterprises, Inc., Palm City, FL 34990, United States
a r t i c l e
i n f o
Article history: Received 14 November 2014 Received in revised form 14 April 2015 Accepted 30 April 2015 Available online 12 May 2015 Keywords: Solid state joining Friction stir welding Thermal stir welding Haynes 230
a b s t r a c t Haynes 230 is a nickel based, solid-solution strengthened alloy that is used for high-temperature applications in the aero-engine and power generation industries. Addition of carbon promotes the formation of carbides to enhance creep resistance and control grain size. Although the alloy has been formulated for high temperature stability, cracking issues are still observed in fusion welding applications. To avoid formation of common fusion weld defects, this study considers the use of a solid state technique variation of friction stir welding (FSW) called thermal stir welding (TSW). An 85% reduction in grain size was observed in the TSW region along with a 30–60% reduction in the size of the W6 C carbide phase. Negligible change in the area fraction of the tungsten rich carbides indicated the W6 C carbides fragmented rather than dissolved. A slight decrease of 12% from the parent material properties was observed in the tensile strength and elongation to failure of the TSW specimens. No cracks were observed in the TSW region during bend testing. Based on the TSW region morphology and lack of cracking, use of solid state joining provides an alternative to fusion welding and its associated defects. © 2015 Elsevier B.V. All rights reserved.
1. Introduction Haynes 230 alloy was developed in the 1980s by Haynes International Inc. to provide higher temperature stability within the materials developed for high temperature corrosion and creep resistance. Initial development of the alloy is discussed by Tawancy et al. (1984) with further discussion of the design rationale by Klarstrom (2001). The matrix is primarily a solid solution strengthened matrix based on nickel (Ni)-Chromium (Cr)-Tungsten (W)-Molybdenum (Mo). Carbon (C) is added within the solid solution strengthened matrix to form 2 major carbides with the excess W and Cr. Primary W6 C carbides form during solidification and help to control grain size. Secondary Cr23 C6 carbide precipitates have been described by Klarstrom (2001) as fine particles that tend to preferentially precipitate along stacking faults and twin boundaries to promote creep resistance. In fusion welding of nickel based superalloys, Dupont et al. (2009) has summarized various defects which can occur including: solidification cracking, liquation cracking, and subsolidus intergrannular cracking. Both solidification and liquation cracking are driven by the stresses that develop between the liquid and solid phases during solidification and are characterized by a second
∗ Corresponding author. Tel.: +1 662 325 9154; fax: +1 662 325 7223. E-mail address: [email protected] (J.A. Schneider). http://dx.doi.org/10.1016/j.jmatprotec.2015.04.034 0924-0136/© 2015 Elsevier B.V. All rights reserved.
phase film which is left at the grain boundary. Solidification cracking occurs as the shrinkage across partially solidified boundaries forms tensile stresses and is most noticeable along the last to solidify regions of dendritic structures located within the weld zone. Liquation cracking generally occurs along grain boundaries in the partially melted zone. A liquid film forms as a result of a eutectic reaction between a secondary phase and the matrix due to non-equilibrium heating resulting in constitutive liquation. Tensile stresses caused by differences in coefficient of thermal expansion drive the crack through the liquid film at the grain boundary. In Haynes 230, liquation cracking has been reported by Ernst (1994) to occur along grain boundary films formed by the constitutional liquation of the W6 C primary carbides with the matrix. Cracks can also occur at subsolidus temperatures without the formation of a second phase film at the grain boundaries. Subsolidus intergranular cracking, also called ductility dip cracking (DDC), occurs at intermediate temperatures associated with the reduced ductility phenomenon in austenitic steels and nickel alloys. At intermediate temperatures, the ductility dip is associated with a change in deformation accommodation mechanisms from grain boundary sliding to grain boundary migration. At the fusion welding temperatures the Cr23 C6 carbides are driven into solution and during cooling reprecipitate along grain boundaries as the metal solidifies. If the carbides are present in sufficient quantities, the resulting thermal coefficient mismatch between the second phases
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metal in front of the weld tool, there is less need for the frictional heating from the shoulder. Thus as shown schematically in Fig. 1b, the shoulder diameter can be reduced as compared to standard FSW process. It is hypothesized for this study, that the use of solid state joining techniques may mitigate various defects that Noecker and Du-Pont (2009) have reported to occur in fusion welds. Since the material remains in the solid state, solidification cracking will not occur. If the solid state joining time is short enough or the temperatures are low enough, the non-equilibrium formation of localized liquidus phase at the interface between the primary W6 C carbide and the matrix will not occur. Similar to the grain size reduction summarized by Reynolds (2007) in aluminum alloys, a similar grain reduction is expected in the Haynes 230. The refined grain size also results in torturous grain boundaries which Dupont et al. (2009) has correlated with minimizing the DDC phenomena. 2. Materials and methods
Fig. 1. Schematic of FSW process (a) compared to TSW process (b).
and the matrix can result in stresses which drive cracking along the grain boundaries at temperatures where the ductility is already reduced. Friction stir welding (FSW), shown schematically in Fig. 1, is a solid state joining process in which shoulder driven frictional and pin driven deformational heating combine to heat the metal to a plasticized state. Unlike other solid state joining processes, FSW is similar to fusion welding in forming butt welds along any length. Due to the solid state nature, differences in thermal expansion are minimized thus reducing stresses within the weldment that can lead to cracking. Schneider (2007) has summarized the temperatures reported within the stir zone to be within the homologous temperature range of 0.7 to 0.9. For the solidus temperature of Haynes 230, this would correspond to a FSW temperature range of 820 to 1134 ◦ C. The upper end of this range is within the recommended hot working temperature for Haynes 230 of 980 to 1205 ◦ C. Within this temperature range, the movement of the rotating tool stirs the plasticized metal together to produce a solid state joint. Various pin geometries have been used for the FSW process such as the threaded pin, shown schematically in Fig. 1a, or a smooth and tapered pin. While the FSW process has been successful in joining many aluminum alloys, the higher melting temperature alloys challenge tool survival as noted by Farias et al. (2013). Ding (2011) has patented a method to decouple the heating and stirring processes of FSW in a process called TSW. By incorporation of an induction coil in advance of the rotating tool, as shown in Fig. 1b, the workpiece can be independently preheated rather than relying on the heat generation from the interaction between the weld tool and the workpiece. As the induction coil preheats the
Hot rolled (HR), solution annealed (SA) Haynes 230 panels of nominal 10 cm wide × 46 cm long by 0.64 cm thick were used in this study. All TSWs were butt welded at 300 RPM and 25 mm/m in position control with a 1◦ back tilt. Fig. 2 shows the tool which is made from a 2.5 cm diameter extruded W—25% Rhenium (Re) with 4% Hafnium Carbide (HfC) rod. A truncated taper forms the pin with an included angle of 60◦ . The flat shoulder has a 2.5 cm diameter and is smooth. To prevent galling of the surface of the panel during the TSW, the shoulder has a generous radius. Thermocouples were positioned in grooves cut into the back, or root side, of the instrumented panel with the thermocouple tip placed along the faying surface at a distance of 0.03 cm from the crown surface. The 1st thermocouple was placed 6.5 cm from the TSW start, for a total of 5 thermocouples spaced 6.5 cm apart along the 26 cm long TSW joint. To provide the preheating, an 18 cm long linear induction coil from Fluxtrol, Inc. was positioned 0.64 cm above the crown surface, approximately 2.5 cm in advance of the TSW tool. The coil was operated at average conditions of 390 V, 2.6 kW, and 1.3 kHz. At the start of each TSW panel, the induction coil dwelled for 7 min to bring the local workpiece temperature to a target temperature of 873 K (600 ◦ C). After reaching the target temperature, the tool was inserted, rotation begun, and the TSW initiated with tool travel. After the TSW, 2 sets of panels were sectioned into strips as illustrated in Fig. 3. To evaluate uniformity along the length of the TSW, the specimens were removed from the beginning, middle, and end of each panel. Each panel produced a total of 9 specimens for tensile testing, 9 specimens for bend testing, and 3 specimens for metallographic characterization. All cuts were made using a water jet to minimize thermal effects. Nominal dimensions of the tensile and bend specimens were 0.64 cm wide × 0.64 cm wide × 20 cm
Fig. 2. W—25% Re—4%HfC TSW tool geometry.
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Tensile Specimens (3 locations)
5 cm
15 cm Weld
20 cm End
Begin
Bend Specimens (3 locations) Fig. 3. Representative cut plan of TSWed Haynes 230 panel.
long. To maximize the number of test specimens per panel, flat strips were used for the tensile specimens. Because of the nonstandard specimen geometry, additional testing including hardness and bending were used to correlate with tensile properties and elongation. Tensile testing of the TSWs and parent material was conducted in an Instron Model 5869 load frame equipped with 50 kN load cell operated at a constant crosshead velocity. A constant crosshead rate of 0.05 mm/m was used to pre load the specimen to 0.5 kN, after which the cross head rate was increased to 1.3 mm/min for the duration of the test. Due to the non-homogenous nature of the weld specimens, bend tests were also conducted to verify the ability to accommodate strain within the TSW nugget. Bend tests were conducted using a mandrel diameter of 1.9 cm to apply a strain across the weld region of 25%. Hardness testing was conducted on a Wilson Instruments model Rockwell 574 using a Rockwell B scale. Fig. 4 shows the location of the indents made across the transverse section of the thermomechanically affected zone (TMAZ)/heat affected zone (HAZ) and stir zone (SZ) or nugget. Parent material measurements were made outside of the TSW zone, and are located outside the region shown in Fig. 4. As welds typically have a hardness gradient across the regions, these locations were selected for evaluation. Metallographic specimens were prepared from sections oriented transverse to the welding direction of the TSW process. Specimens were mounted in a phenolic and ground using a sequence of varying grit silicon carbide papers and polished using a series of alumina powders. A mixed acid (1/3 acetic acid, 1/3 HNO3 ,
1/3 HCL) etchant was used to reveal the grain structure. Macro images were taken on a Canon EOS Rebel T1i equipped with a Canon MP-E 65 mm lens and micro images were captured with a Leica DMI500 M. Scanning electron microscopy (SEM) imaging and energydispersive X-ray spectroscopy (EDS) were performed on unetched samples in a JEOL field emission (FE) SEM equipped with an Oxford “X-Max 50” EDS detector. SEM secondary electron images (SEI) and backscattered images (BSI) were used for image analysis to determine the grain size of the matrix and W6 C particles, respectively. Grain and particle size were determined on the basis of the feret diameter using public domain software Image J, Version 1.45S from National Institute of Health (2006). Multiple images were used at the same magnification within different regions of the TSW macrostructure to obtain an average over 100 grains or particles. The location of the hardness indentations was used to visually identify the regions for SEM imaging. X-ray diffraction analysis was performed on the metallographically mounted specimens in a Rigaku Ultima III X-Ray Diffraction System (XRD) in an attempt to identify the carbides present. The samples were scanned in the 2 range of 35◦ to 50◦ at a scanning speed of 0.05◦ min−1 . A JEOL 2100 transmission electron microscope (TEM) was used for imaging of the smaller Cr carbides. The TEM was equipped with an Oxford “X-Max-80T EDS. One foil was extracted from the TSW nugget using standard metallographic procedures. A 3 mm diameter disk was punched from a thin slice, and then dimpled on both sides. The disk was ion milled in a Fischione Model 1010 at a milling angle of 9◦ and voltage of 4.5 keV until electron transparency was obtained. 3. Results As measured by the Type K thermocouple placed along the faying surface just in advance of the TSW start position, a 7 min dwell brought the local workpiece temperature to 835 K (562 ◦ C). As the tool was inserted into a predrilled hole, the rotation of the tool began along with the tool travel. As the TSW tool passed through the preheated material, the additional contributions from the tool/workpiece interaction raised the nugget zone to average temperature of 1291 + 311 K (1018◦ + 38 ◦ C) along the length of the TSW. This corresponds to a homologous temperature of 0.74 to 0.78, typical of the range reported for FSWs as summarized by Schneider (2007). The macrostructure of the initial Haynes 230 parent material is shown in Fig. 5. The W rich phases are labeled along with assumed Cr rich phases which are located along twin boundaries as reported by Tawancy et al. (1984). Table 1 compares the tensile results for the TSW with published data for the parent material by Haynes International (2007). A slight
Fig. 4. Location of hardness indents across section of the TSW oriented transverse to the welding direction.
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Fig. 5. Haynes 230 parent material macro-structure.
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Fig. 7. Comparison of hardness data of parent material to the TSW nugget.
Table 1 Tension test data summary.
Parent material from Haynes Intl (2007) TSW
UTS (MPa)
εf (%)
840 737.8 ± 6.9
47.7 41.8 ± 1.6
12.1% reduction in ultimate tensile strength (UTS) and 12.4% reduction elongation to failure (εf ) were observed. Fig. 6 shows the bend tests specimens with no observed cracks within the weld region. Rockwell hardness B values, shown in Fig. 7, were recorded across a section oriented transverse to the welding direction of the TSW process and compared with the parent material values. A 4% increase in hardness is observed in the TSW nugget which decreases to match parent material values on either side. XRD analysis was performed on bulk samples to identify the carbides present. Although the diffraction peaks are of low intensity in Fig. 8, the W6 C phase is observed which agrees with the SEM/EDS analysis of elemental content and observed morphology. Intensity peak overlaps between the Ni and the Cr23 C6 phase obscure verification of the Cr23 C6 phase with XRD. Figs. 9 and 10 present SEM images of the parent material vs. that within the TSW nugget, respectively. The secondary electron images (SEI) show the grain size difference and the backscattered electron images (BSI) help to highlight the W6 C particle size and distribution. Table 2 summarizes the grain size for the parent material and TSW nugget based on an average of 100 grains. The 85% reduction in the grain size of the TSW nugget would be expected to result in the measured increased hardness due to Hall-Petch strengthening.
Fig. 8. XRD analysis of the TSW nugget.
Using the SEM images, the area fraction and particle size of approximately 100 particles of W6 C are summarized in Table 3. Along with the range of reduction in W6 C particle size from 30 to 60%, there is a corresponding difference in area fraction depending on the location within the TSW zone. The sum of the area fractions across the TSW zone is noted to be equivalent to that measured in the parent material. A decrease in the aspect ratio of the W6 C particles is also observed from the parent material to that within the TSW region. While the SEM images of the TSW region show relatively clean grain boundaries, tortuous grain boundaries are observed as illus-
Fig. 6. No cracks were observed in bend test specimens from either the (a) crown side or (b) root side.
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Fig. 9. (a) SEI of the parent material showing initial grain size and (b) BSI highlighting W6 C particle size and distribution in the parent material.
Fig. 10. (a) SEI of the TSW nugget showing refined grain size and (b) BSI highlighting W6 C particle size and distribution in the TSW nugget. Table 2 Summary of grain sizes in the parent material vs the TSW nugget. Average grain size (m) 75.5 ± 3.4 11.2 ± 2.7
Parent material TSW
Table 3 Area fraction of W6 C phase in parent material and TSW nugget.
Parent material TSW RS TSW center SZ TSW AS
To obtain further information on the presence of Cr23 C6 carbides along the grain boundaries in the TSW nugget, TEM imaging was used. Fig. 12 shows carbides present along a grain boundary. EDS analysis confirmed these particles were Cr rich and were similar in size to Tawancy et al. (1984) published range of 500 to 1000 nm for the Cr23 C6 carbide. 4. Discussion
Diameter of W6 C (m)
Area fraction of W6 C (%)
Maximum/ minimum aspect ratio
5.1 ± 3.5 2.0 ± 1.6 3.5 ± 4.5 3.6 ± 3.3
1.62 0.04 0.41 1.29
6.2/1.1 3.0/1.0 3.9/1.0 3.0/1.0
trated in Fig. 11. This indicates that the Cr23 C6 phases formerly present on stacking faults and twin boundaries serve to pin the migrating grain boundaries during the deformation and recrystallization process as discussed by Humphreys and Hatherly (2004).
Heating during a conventional FSW process is based on a combination of frictional heating between the shoulder and the workpiece, and deformational heating in the shearing of material in the through material thickness by the pin. The TSW process decouples these heating modes during the solid state joining process by use of an induction coil to generate preheat internally in the workpiece. By passing a high-frequency alternating current through an electromagnetic coil, a rapidly alternating magnetic field is applied to penetrate the workpiece. This results in rapid internal heating of the workpiece as a result of the internal magnetic hysteresis losses. To optimize the coil design and parameters, the current frequency is tuned to the material thickness, material type, and the pene-
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Fig. 11. Carbide pinning of grain boundaries in the TSW nugget.
Fig. 12. TEM image of a Cr-rich phase along a grain boundary.
tration depth. The current coil configuration and parameters were able to raise the workpiece temperature to 835 K (562 ◦ C) within 7 min, slightly less than the target temperature of 873 K (600 ◦ C). As the workpiece was heated internally, longer holding times were ineffective in raising the temperature to the target value due to conduction losses through the backing anvil. Although the target temperature was not reached in this study, the initial data can be used for optimizing the induction coil design in subsequent studies. By preheating the workpiece to 43% of the solid state joining temperature, less preheating is required from the frictional heating of the shoulder on the workpiece. This allows a reduction in the shoulder diameter on the TSW tool as compared to conventional FSW process. As the TSW tool and deformation path advanced at a velocity of 25 mm/m, the TSW metal experienced a maximum temperature of 1291 K (1018 ◦ C) as the pin stirred the metal together. Thus based on the thermocouple measurements, 729 K (456 ◦ C) of the heat input can be attributed to the interaction of the tool and workpiece. Assuming the only heating was from the deformation as the pin sheared the workpiece, a comparison can be made with an estimation of temperature increase due to adiabatic heating. Using an expression in Eq. (1) developed by Farren and Taylor (1925) and Taylor and Quinney (1933), the temperature increase (T) due to conversion of deformation energy into heat can be obtained: T =
ˇ ×c
eff dE
(1)
where eff is the effective shear stress, eff is the effective shear strain, is the density; and c is the specific heat capacity for Haynes 230 at 835 K (562 ◦ C). Based on the early experimental work of Taylor and Quinney (1933), the efficiency of deformation heat conversion, ˇ, has been assumed to be a constant ranging from 0.85 to 0.95. Using a ˇ value of 0.9, an estimated 344 K (71 ◦ C) of adiabatic heat would be expected to be generated from the deformation
in the Haynes 230 TSW nugget. The slightly higher measured value indicates there was some frictional heating provided by the reduced diameter shoulder. The melting temperature for Haynes 230 as published by Haynes International (2007) is in the range of 1574 to 1644 K (1301 to 1371 ◦ C). As the TSW temperature was in the homologous range of 0.74 to 0.78, no melting is expected eliminating the potential for solidification cracking. Based on SEM images and EDS analysis, no evidence of a second phase grain boundary film was found in the SEM images which also confirm that no localized liquidus or eutectic formation occurred between the carbide particles and the nickel matrix. Thus both the W rich and Cr rich carbides appear to be thermally stable at the short times at the TSW temperature. Subsolidus intergrannular cracking, also known as DDC, has been reported in wrought nickel based superalloys at intermediate temperatures where ductility is reduced. Studies on an austenite AISI 316 stainless steel by Mannan et al. (1985) correlated the occurrence with the grain size. With smaller grains, the DDC temperature is reduced as the ductility minimum corresponds with the maximum amount of grain boundary sliding. As the temperature increases, the increased grain boundary diffusion associated with the refined grains becomes rate controlling. While the DDC phenomenon occurs in wrought material, it is also problematic in fusion welds due to the elongated grain structure. A series of fusion welding articles by Collins and Lippold (2003), Collins et al. (2003, 2004) and Ramirez et al. (2006) focus primarily on selection of filler materials to help alleviate DDC in fusion weld regions by forming of tortuous grain boundaries to impede crack propagation. Analyses of the SEM images taken of the TSW nugget show the grains were refined from a starting 75.5 m to 11.2 m. This is similar to the degree of grain refinement in aluminum alloys as summarized by Reynolds (2007) in FSWs of aluminum alloys. The grain refinement would result in Hall-Petch grain size strengthening as evidenced by the increased hardness in the nugget. In addition, the equiaxed refined grains present a tortuous grain boundary path which has been identified by Noecker and Du-Pont (2009) to resist the occurrence of DDC. Within the W6 C TSW region, the carbides appear to have fragmented rather than diffused into the matrix as shown in Fig. 13. While there was a distribution of sizes across the TSW region, the total area fraction of the fragmented W6 C carbides remains similar to the parent material. No change is observed in the Cr23 C6 carbides which remain submicron in size and are located at grain boundaries in the post weld region. Similar particle breakage has been observed in FSW studies of metal matrix composites. Storjohann et al. (2005) and Marzoli et al.
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Fig. 13. SEM image of W6 C carbides fragmenting within the TSW HAZ region.
(2006) both investigated the FSW of AA6061 with 20% Al2 O3 particles of average 20 m diameter. Although there were variations in their weld parameters, Storjohann et al. (2005) used one set of fixed parameters whereas Marzoli et al. (2006) used a range of tool rotations and travel speeds, similar results were found in the refinement of the particle size but not area fraction within the FSW nugget attributed to particle breaking. Uzun (2007) also reported refinement of SiC reinforcing particles in a composite plate of AA2124. In these studies, no details were provided of the tool geometry and the possible influence of features such as threads, on the resulting distribution of particles observed. The fragmentation of the W6 C carbides in this study occurs as the particles enter the rotation plug of metal around the weld tool. This can be due to either forces resulting from the severe shear zone or the temperature gradients within this region. Using a material independent, kinematic model developed by Nunes (2003), the severe shear zone is estimated to impart an instantaneously high shear strain rate in the range of 103 s−1 while the metal undergoes shear strains on the order of 50 to 100. During the solid state joining process, the flow of metal around the weld tool can be thought of as following slip lines as illustrated in Fig. 14. As the tool rotates, a non-symmetric flow field is generated. The retreating side (RS) of the weld is where the rotation vector and travel vector are in the opposite direction and the advancing side (AS) of the weld is where they are in the same direction. As the tool rotates past the center line of the weld, metal is picked up as each slip line is passed. This leads to a thicker region of rotating metal on the RS of the weld as described by Nunes (2003). As the tool rotation passes the center line in the wake of the weld, metal is deposited in the order accumulated. Thus larger particles would tend to be deposited in the nugget center and AS TMAZ as the slip lines converge. Metal that moves multiple times around the tool would capture the smaller particles which are eventually deposited on the RS as reported in studies by London et al. (2003). Fig. 15 illustrates the slip line entering the severe shear zone which demarks the rotating plug of metal from the workpiece
Fig. 15. Plan view illustration of shear zone surrounding the weld tool.
(Nunes, 2003). The shear zone is a region of finite thickness (ı) located a finite distance from the weld tool. Brittle particles of sufficient length to diameter ratio (L/D) would break due to the shearing forces.
5. Conclusions (1) Haynes 230 was shown to be successfully joined using the TSW process with the weld region exhibiting equiaxed grains. The grains were reduced by 85% to 11.2 m equiaxed diameter and the size of the W6 C phase was reduced over a range of 30 to 60%. (2) As compared to the wrought properties, the TSW displayed a slight decrease of 12% in the tensile strength and elongation to failure. (3) The TSW nugget displayed a higher hardness than the parent material which is attributed to Hall-Petch grain size strengthening. (4) The area fraction of the W6 C within the TSW region was similar to that of the starting parent material indicating the stability of the W6 C phase during the solid state TSW process. (5) The wiping motion of the smooth tool in the TSW zone results in an angular reduction of the shear zone around the tool letting smaller particles pass multiple times around the tool before being deposited on the retreating side. This results in changes in the W6 C size distribution across the TSW section oriented transverse to the welding direction.
Acknowledgements
Fig. 14. Illustration of metal flow following slip lines around the weld tool.
Funding was provided by the NASA-STTR Phase II grant no. KSE 11008 10070691 in collaboration with Keystone Synergistic Enterprises, Inc. All TSWs were conducted at the NASA-Marshall Space Flight Center under a Keystone Space Act Agreement in collaboration with Jeff Ding.
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Characterization equipment used on this project was purchased under NSF-IMR grant nos. DMR0216703 and 02070615 for the scanning electron microscope and detectors, NSF-MRI grant no. DMR0619773 for the X-ray diffractometer, and NSF DBI grant no. 1126743 for the transmission electron microscope and EDS detector. References Collins, M.G., Lippold, J.C., 2003. An investigation of ductility dip cracking in nickelbased filler materials—Part I. Weld. J. 82 (10), 288s–295s. Collins, M.G., Ramirez, A.J., Lippold, J.C., 2003. An investigation of ductility dip cracking in nickel-based filler materials—Part II. Weld. J. 82 (12), 348s–354s. Collins, M.G., Ramirez, A.J., Lippold, J.C., 2004. An investigation of ductility dip cracking in nickel-based filler materials—Part III. Weld. J. 83 (2), 39s–49s. Ding, R.J., 2011. Thermal Stir Welding Apparatus (US7980449 B2, July 19, 2011). Dupont, J., Lippold, J., Kiser, S., 2009. Welding Metallurgy and Weldability of NickelBased Alloys, first ed. Wiley and Sons Inc. Pub, Hoboken, NJ, pp. 131–143. Ernst, S.C., 1994. Weldability studies of Haynes 230 alloy. Weld. J. 73 (4), 80s–89s. Farias, A., Batalha, G.F., Prados, E.F., 2013. Tool wear evaluations in friction stir processing of commercial titanium Ti–6Al–4V. Wear 302 (1–2), 1327–1333. Farren, W.S., Taylor, G.I., 1925. The heat developed during plastic extension of metals. Proc. R. Soc., A 107, 422–451. Haynes International, 2007. Haynes 230 alloy. In: Report #H-3000H, http://www. haynesintl.com/pdf/h3000.pdf (accessed 04/06/2015). Humphreys, F.J., Hatherly, M., 2004. Recrystallization and Related Annealing Phenomena, second ed. Elsevier Pub, Oxford, UK, pp. 430–431. Klarstrom, D.L., 2001. The development of HAYNES 230. In: Zhao, J.C., Fahrmann, M., Pollock, T.M. (Eds.), Proceedings of Mat’ls Design Approaches & Experiences. TMS Pub, Warrendale, PA, pp. 297–307.
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Journal of Materials Processing Technology journal homepage: www.elsevier.com/locate/jmatprotec
Solid state joining of nickel based alloy, Haynes 230 J.A. Schneider a,∗ , D. Williston a , T.L. Murphy a , C. Varner a , J. Hawkins a , B. Walker b a b
Department of Mechanical Engineering, Mississippi State University, Starkville, MS 39762, United States Keystone Synergistic Enterprises, Inc., Palm City, FL 34990, United States
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Article history: Received 14 November 2014 Received in revised form 14 April 2015 Accepted 30 April 2015 Available online 12 May 2015 Keywords: Solid state joining Friction stir welding Thermal stir welding Haynes 230
a b s t r a c t Haynes 230 is a nickel based, solid-solution strengthened alloy that is used for high-temperature applications in the aero-engine and power generation industries. Addition of carbon promotes the formation of carbides to enhance creep resistance and control grain size. Although the alloy has been formulated for high temperature stability, cracking issues are still observed in fusion welding applications. To avoid formation of common fusion weld defects, this study considers the use of a solid state technique variation of friction stir welding (FSW) called thermal stir welding (TSW). An 85% reduction in grain size was observed in the TSW region along with a 30–60% reduction in the size of the W6 C carbide phase. Negligible change in the area fraction of the tungsten rich carbides indicated the W6 C carbides fragmented rather than dissolved. A slight decrease of 12% from the parent material properties was observed in the tensile strength and elongation to failure of the TSW specimens. No cracks were observed in the TSW region during bend testing. Based on the TSW region morphology and lack of cracking, use of solid state joining provides an alternative to fusion welding and its associated defects. © 2015 Elsevier B.V. All rights reserved.
1. Introduction Haynes 230 alloy was developed in the 1980s by Haynes International Inc. to provide higher temperature stability within the materials developed for high temperature corrosion and creep resistance. Initial development of the alloy is discussed by Tawancy et al. (1984) with further discussion of the design rationale by Klarstrom (2001). The matrix is primarily a solid solution strengthened matrix based on nickel (Ni)-Chromium (Cr)-Tungsten (W)-Molybdenum (Mo). Carbon (C) is added within the solid solution strengthened matrix to form 2 major carbides with the excess W and Cr. Primary W6 C carbides form during solidification and help to control grain size. Secondary Cr23 C6 carbide precipitates have been described by Klarstrom (2001) as fine particles that tend to preferentially precipitate along stacking faults and twin boundaries to promote creep resistance. In fusion welding of nickel based superalloys, Dupont et al. (2009) has summarized various defects which can occur including: solidification cracking, liquation cracking, and subsolidus intergrannular cracking. Both solidification and liquation cracking are driven by the stresses that develop between the liquid and solid phases during solidification and are characterized by a second
∗ Corresponding author. Tel.: +1 662 325 9154; fax: +1 662 325 7223. E-mail address: [email protected] (J.A. Schneider). http://dx.doi.org/10.1016/j.jmatprotec.2015.04.034 0924-0136/© 2015 Elsevier B.V. All rights reserved.
phase film which is left at the grain boundary. Solidification cracking occurs as the shrinkage across partially solidified boundaries forms tensile stresses and is most noticeable along the last to solidify regions of dendritic structures located within the weld zone. Liquation cracking generally occurs along grain boundaries in the partially melted zone. A liquid film forms as a result of a eutectic reaction between a secondary phase and the matrix due to non-equilibrium heating resulting in constitutive liquation. Tensile stresses caused by differences in coefficient of thermal expansion drive the crack through the liquid film at the grain boundary. In Haynes 230, liquation cracking has been reported by Ernst (1994) to occur along grain boundary films formed by the constitutional liquation of the W6 C primary carbides with the matrix. Cracks can also occur at subsolidus temperatures without the formation of a second phase film at the grain boundaries. Subsolidus intergranular cracking, also called ductility dip cracking (DDC), occurs at intermediate temperatures associated with the reduced ductility phenomenon in austenitic steels and nickel alloys. At intermediate temperatures, the ductility dip is associated with a change in deformation accommodation mechanisms from grain boundary sliding to grain boundary migration. At the fusion welding temperatures the Cr23 C6 carbides are driven into solution and during cooling reprecipitate along grain boundaries as the metal solidifies. If the carbides are present in sufficient quantities, the resulting thermal coefficient mismatch between the second phases
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metal in front of the weld tool, there is less need for the frictional heating from the shoulder. Thus as shown schematically in Fig. 1b, the shoulder diameter can be reduced as compared to standard FSW process. It is hypothesized for this study, that the use of solid state joining techniques may mitigate various defects that Noecker and Du-Pont (2009) have reported to occur in fusion welds. Since the material remains in the solid state, solidification cracking will not occur. If the solid state joining time is short enough or the temperatures are low enough, the non-equilibrium formation of localized liquidus phase at the interface between the primary W6 C carbide and the matrix will not occur. Similar to the grain size reduction summarized by Reynolds (2007) in aluminum alloys, a similar grain reduction is expected in the Haynes 230. The refined grain size also results in torturous grain boundaries which Dupont et al. (2009) has correlated with minimizing the DDC phenomena. 2. Materials and methods
Fig. 1. Schematic of FSW process (a) compared to TSW process (b).
and the matrix can result in stresses which drive cracking along the grain boundaries at temperatures where the ductility is already reduced. Friction stir welding (FSW), shown schematically in Fig. 1, is a solid state joining process in which shoulder driven frictional and pin driven deformational heating combine to heat the metal to a plasticized state. Unlike other solid state joining processes, FSW is similar to fusion welding in forming butt welds along any length. Due to the solid state nature, differences in thermal expansion are minimized thus reducing stresses within the weldment that can lead to cracking. Schneider (2007) has summarized the temperatures reported within the stir zone to be within the homologous temperature range of 0.7 to 0.9. For the solidus temperature of Haynes 230, this would correspond to a FSW temperature range of 820 to 1134 ◦ C. The upper end of this range is within the recommended hot working temperature for Haynes 230 of 980 to 1205 ◦ C. Within this temperature range, the movement of the rotating tool stirs the plasticized metal together to produce a solid state joint. Various pin geometries have been used for the FSW process such as the threaded pin, shown schematically in Fig. 1a, or a smooth and tapered pin. While the FSW process has been successful in joining many aluminum alloys, the higher melting temperature alloys challenge tool survival as noted by Farias et al. (2013). Ding (2011) has patented a method to decouple the heating and stirring processes of FSW in a process called TSW. By incorporation of an induction coil in advance of the rotating tool, as shown in Fig. 1b, the workpiece can be independently preheated rather than relying on the heat generation from the interaction between the weld tool and the workpiece. As the induction coil preheats the
Hot rolled (HR), solution annealed (SA) Haynes 230 panels of nominal 10 cm wide × 46 cm long by 0.64 cm thick were used in this study. All TSWs were butt welded at 300 RPM and 25 mm/m in position control with a 1◦ back tilt. Fig. 2 shows the tool which is made from a 2.5 cm diameter extruded W—25% Rhenium (Re) with 4% Hafnium Carbide (HfC) rod. A truncated taper forms the pin with an included angle of 60◦ . The flat shoulder has a 2.5 cm diameter and is smooth. To prevent galling of the surface of the panel during the TSW, the shoulder has a generous radius. Thermocouples were positioned in grooves cut into the back, or root side, of the instrumented panel with the thermocouple tip placed along the faying surface at a distance of 0.03 cm from the crown surface. The 1st thermocouple was placed 6.5 cm from the TSW start, for a total of 5 thermocouples spaced 6.5 cm apart along the 26 cm long TSW joint. To provide the preheating, an 18 cm long linear induction coil from Fluxtrol, Inc. was positioned 0.64 cm above the crown surface, approximately 2.5 cm in advance of the TSW tool. The coil was operated at average conditions of 390 V, 2.6 kW, and 1.3 kHz. At the start of each TSW panel, the induction coil dwelled for 7 min to bring the local workpiece temperature to a target temperature of 873 K (600 ◦ C). After reaching the target temperature, the tool was inserted, rotation begun, and the TSW initiated with tool travel. After the TSW, 2 sets of panels were sectioned into strips as illustrated in Fig. 3. To evaluate uniformity along the length of the TSW, the specimens were removed from the beginning, middle, and end of each panel. Each panel produced a total of 9 specimens for tensile testing, 9 specimens for bend testing, and 3 specimens for metallographic characterization. All cuts were made using a water jet to minimize thermal effects. Nominal dimensions of the tensile and bend specimens were 0.64 cm wide × 0.64 cm wide × 20 cm
Fig. 2. W—25% Re—4%HfC TSW tool geometry.
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Tensile Specimens (3 locations)
5 cm
15 cm Weld
20 cm End
Begin
Bend Specimens (3 locations) Fig. 3. Representative cut plan of TSWed Haynes 230 panel.
long. To maximize the number of test specimens per panel, flat strips were used for the tensile specimens. Because of the nonstandard specimen geometry, additional testing including hardness and bending were used to correlate with tensile properties and elongation. Tensile testing of the TSWs and parent material was conducted in an Instron Model 5869 load frame equipped with 50 kN load cell operated at a constant crosshead velocity. A constant crosshead rate of 0.05 mm/m was used to pre load the specimen to 0.5 kN, after which the cross head rate was increased to 1.3 mm/min for the duration of the test. Due to the non-homogenous nature of the weld specimens, bend tests were also conducted to verify the ability to accommodate strain within the TSW nugget. Bend tests were conducted using a mandrel diameter of 1.9 cm to apply a strain across the weld region of 25%. Hardness testing was conducted on a Wilson Instruments model Rockwell 574 using a Rockwell B scale. Fig. 4 shows the location of the indents made across the transverse section of the thermomechanically affected zone (TMAZ)/heat affected zone (HAZ) and stir zone (SZ) or nugget. Parent material measurements were made outside of the TSW zone, and are located outside the region shown in Fig. 4. As welds typically have a hardness gradient across the regions, these locations were selected for evaluation. Metallographic specimens were prepared from sections oriented transverse to the welding direction of the TSW process. Specimens were mounted in a phenolic and ground using a sequence of varying grit silicon carbide papers and polished using a series of alumina powders. A mixed acid (1/3 acetic acid, 1/3 HNO3 ,
1/3 HCL) etchant was used to reveal the grain structure. Macro images were taken on a Canon EOS Rebel T1i equipped with a Canon MP-E 65 mm lens and micro images were captured with a Leica DMI500 M. Scanning electron microscopy (SEM) imaging and energydispersive X-ray spectroscopy (EDS) were performed on unetched samples in a JEOL field emission (FE) SEM equipped with an Oxford “X-Max 50” EDS detector. SEM secondary electron images (SEI) and backscattered images (BSI) were used for image analysis to determine the grain size of the matrix and W6 C particles, respectively. Grain and particle size were determined on the basis of the feret diameter using public domain software Image J, Version 1.45S from National Institute of Health (2006). Multiple images were used at the same magnification within different regions of the TSW macrostructure to obtain an average over 100 grains or particles. The location of the hardness indentations was used to visually identify the regions for SEM imaging. X-ray diffraction analysis was performed on the metallographically mounted specimens in a Rigaku Ultima III X-Ray Diffraction System (XRD) in an attempt to identify the carbides present. The samples were scanned in the 2 range of 35◦ to 50◦ at a scanning speed of 0.05◦ min−1 . A JEOL 2100 transmission electron microscope (TEM) was used for imaging of the smaller Cr carbides. The TEM was equipped with an Oxford “X-Max-80T EDS. One foil was extracted from the TSW nugget using standard metallographic procedures. A 3 mm diameter disk was punched from a thin slice, and then dimpled on both sides. The disk was ion milled in a Fischione Model 1010 at a milling angle of 9◦ and voltage of 4.5 keV until electron transparency was obtained. 3. Results As measured by the Type K thermocouple placed along the faying surface just in advance of the TSW start position, a 7 min dwell brought the local workpiece temperature to 835 K (562 ◦ C). As the tool was inserted into a predrilled hole, the rotation of the tool began along with the tool travel. As the TSW tool passed through the preheated material, the additional contributions from the tool/workpiece interaction raised the nugget zone to average temperature of 1291 + 311 K (1018◦ + 38 ◦ C) along the length of the TSW. This corresponds to a homologous temperature of 0.74 to 0.78, typical of the range reported for FSWs as summarized by Schneider (2007). The macrostructure of the initial Haynes 230 parent material is shown in Fig. 5. The W rich phases are labeled along with assumed Cr rich phases which are located along twin boundaries as reported by Tawancy et al. (1984). Table 1 compares the tensile results for the TSW with published data for the parent material by Haynes International (2007). A slight
Fig. 4. Location of hardness indents across section of the TSW oriented transverse to the welding direction.
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Fig. 5. Haynes 230 parent material macro-structure.
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Fig. 7. Comparison of hardness data of parent material to the TSW nugget.
Table 1 Tension test data summary.
Parent material from Haynes Intl (2007) TSW
UTS (MPa)
εf (%)
840 737.8 ± 6.9
47.7 41.8 ± 1.6
12.1% reduction in ultimate tensile strength (UTS) and 12.4% reduction elongation to failure (εf ) were observed. Fig. 6 shows the bend tests specimens with no observed cracks within the weld region. Rockwell hardness B values, shown in Fig. 7, were recorded across a section oriented transverse to the welding direction of the TSW process and compared with the parent material values. A 4% increase in hardness is observed in the TSW nugget which decreases to match parent material values on either side. XRD analysis was performed on bulk samples to identify the carbides present. Although the diffraction peaks are of low intensity in Fig. 8, the W6 C phase is observed which agrees with the SEM/EDS analysis of elemental content and observed morphology. Intensity peak overlaps between the Ni and the Cr23 C6 phase obscure verification of the Cr23 C6 phase with XRD. Figs. 9 and 10 present SEM images of the parent material vs. that within the TSW nugget, respectively. The secondary electron images (SEI) show the grain size difference and the backscattered electron images (BSI) help to highlight the W6 C particle size and distribution. Table 2 summarizes the grain size for the parent material and TSW nugget based on an average of 100 grains. The 85% reduction in the grain size of the TSW nugget would be expected to result in the measured increased hardness due to Hall-Petch strengthening.
Fig. 8. XRD analysis of the TSW nugget.
Using the SEM images, the area fraction and particle size of approximately 100 particles of W6 C are summarized in Table 3. Along with the range of reduction in W6 C particle size from 30 to 60%, there is a corresponding difference in area fraction depending on the location within the TSW zone. The sum of the area fractions across the TSW zone is noted to be equivalent to that measured in the parent material. A decrease in the aspect ratio of the W6 C particles is also observed from the parent material to that within the TSW region. While the SEM images of the TSW region show relatively clean grain boundaries, tortuous grain boundaries are observed as illus-
Fig. 6. No cracks were observed in bend test specimens from either the (a) crown side or (b) root side.
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Fig. 9. (a) SEI of the parent material showing initial grain size and (b) BSI highlighting W6 C particle size and distribution in the parent material.
Fig. 10. (a) SEI of the TSW nugget showing refined grain size and (b) BSI highlighting W6 C particle size and distribution in the TSW nugget. Table 2 Summary of grain sizes in the parent material vs the TSW nugget. Average grain size (m) 75.5 ± 3.4 11.2 ± 2.7
Parent material TSW
Table 3 Area fraction of W6 C phase in parent material and TSW nugget.
Parent material TSW RS TSW center SZ TSW AS
To obtain further information on the presence of Cr23 C6 carbides along the grain boundaries in the TSW nugget, TEM imaging was used. Fig. 12 shows carbides present along a grain boundary. EDS analysis confirmed these particles were Cr rich and were similar in size to Tawancy et al. (1984) published range of 500 to 1000 nm for the Cr23 C6 carbide. 4. Discussion
Diameter of W6 C (m)
Area fraction of W6 C (%)
Maximum/ minimum aspect ratio
5.1 ± 3.5 2.0 ± 1.6 3.5 ± 4.5 3.6 ± 3.3
1.62 0.04 0.41 1.29
6.2/1.1 3.0/1.0 3.9/1.0 3.0/1.0
trated in Fig. 11. This indicates that the Cr23 C6 phases formerly present on stacking faults and twin boundaries serve to pin the migrating grain boundaries during the deformation and recrystallization process as discussed by Humphreys and Hatherly (2004).
Heating during a conventional FSW process is based on a combination of frictional heating between the shoulder and the workpiece, and deformational heating in the shearing of material in the through material thickness by the pin. The TSW process decouples these heating modes during the solid state joining process by use of an induction coil to generate preheat internally in the workpiece. By passing a high-frequency alternating current through an electromagnetic coil, a rapidly alternating magnetic field is applied to penetrate the workpiece. This results in rapid internal heating of the workpiece as a result of the internal magnetic hysteresis losses. To optimize the coil design and parameters, the current frequency is tuned to the material thickness, material type, and the pene-
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Fig. 11. Carbide pinning of grain boundaries in the TSW nugget.
Fig. 12. TEM image of a Cr-rich phase along a grain boundary.
tration depth. The current coil configuration and parameters were able to raise the workpiece temperature to 835 K (562 ◦ C) within 7 min, slightly less than the target temperature of 873 K (600 ◦ C). As the workpiece was heated internally, longer holding times were ineffective in raising the temperature to the target value due to conduction losses through the backing anvil. Although the target temperature was not reached in this study, the initial data can be used for optimizing the induction coil design in subsequent studies. By preheating the workpiece to 43% of the solid state joining temperature, less preheating is required from the frictional heating of the shoulder on the workpiece. This allows a reduction in the shoulder diameter on the TSW tool as compared to conventional FSW process. As the TSW tool and deformation path advanced at a velocity of 25 mm/m, the TSW metal experienced a maximum temperature of 1291 K (1018 ◦ C) as the pin stirred the metal together. Thus based on the thermocouple measurements, 729 K (456 ◦ C) of the heat input can be attributed to the interaction of the tool and workpiece. Assuming the only heating was from the deformation as the pin sheared the workpiece, a comparison can be made with an estimation of temperature increase due to adiabatic heating. Using an expression in Eq. (1) developed by Farren and Taylor (1925) and Taylor and Quinney (1933), the temperature increase (T) due to conversion of deformation energy into heat can be obtained: T =
ˇ ×c
eff dE
(1)
where eff is the effective shear stress, eff is the effective shear strain, is the density; and c is the specific heat capacity for Haynes 230 at 835 K (562 ◦ C). Based on the early experimental work of Taylor and Quinney (1933), the efficiency of deformation heat conversion, ˇ, has been assumed to be a constant ranging from 0.85 to 0.95. Using a ˇ value of 0.9, an estimated 344 K (71 ◦ C) of adiabatic heat would be expected to be generated from the deformation
in the Haynes 230 TSW nugget. The slightly higher measured value indicates there was some frictional heating provided by the reduced diameter shoulder. The melting temperature for Haynes 230 as published by Haynes International (2007) is in the range of 1574 to 1644 K (1301 to 1371 ◦ C). As the TSW temperature was in the homologous range of 0.74 to 0.78, no melting is expected eliminating the potential for solidification cracking. Based on SEM images and EDS analysis, no evidence of a second phase grain boundary film was found in the SEM images which also confirm that no localized liquidus or eutectic formation occurred between the carbide particles and the nickel matrix. Thus both the W rich and Cr rich carbides appear to be thermally stable at the short times at the TSW temperature. Subsolidus intergrannular cracking, also known as DDC, has been reported in wrought nickel based superalloys at intermediate temperatures where ductility is reduced. Studies on an austenite AISI 316 stainless steel by Mannan et al. (1985) correlated the occurrence with the grain size. With smaller grains, the DDC temperature is reduced as the ductility minimum corresponds with the maximum amount of grain boundary sliding. As the temperature increases, the increased grain boundary diffusion associated with the refined grains becomes rate controlling. While the DDC phenomenon occurs in wrought material, it is also problematic in fusion welds due to the elongated grain structure. A series of fusion welding articles by Collins and Lippold (2003), Collins et al. (2003, 2004) and Ramirez et al. (2006) focus primarily on selection of filler materials to help alleviate DDC in fusion weld regions by forming of tortuous grain boundaries to impede crack propagation. Analyses of the SEM images taken of the TSW nugget show the grains were refined from a starting 75.5 m to 11.2 m. This is similar to the degree of grain refinement in aluminum alloys as summarized by Reynolds (2007) in FSWs of aluminum alloys. The grain refinement would result in Hall-Petch grain size strengthening as evidenced by the increased hardness in the nugget. In addition, the equiaxed refined grains present a tortuous grain boundary path which has been identified by Noecker and Du-Pont (2009) to resist the occurrence of DDC. Within the W6 C TSW region, the carbides appear to have fragmented rather than diffused into the matrix as shown in Fig. 13. While there was a distribution of sizes across the TSW region, the total area fraction of the fragmented W6 C carbides remains similar to the parent material. No change is observed in the Cr23 C6 carbides which remain submicron in size and are located at grain boundaries in the post weld region. Similar particle breakage has been observed in FSW studies of metal matrix composites. Storjohann et al. (2005) and Marzoli et al.
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Fig. 13. SEM image of W6 C carbides fragmenting within the TSW HAZ region.
(2006) both investigated the FSW of AA6061 with 20% Al2 O3 particles of average 20 m diameter. Although there were variations in their weld parameters, Storjohann et al. (2005) used one set of fixed parameters whereas Marzoli et al. (2006) used a range of tool rotations and travel speeds, similar results were found in the refinement of the particle size but not area fraction within the FSW nugget attributed to particle breaking. Uzun (2007) also reported refinement of SiC reinforcing particles in a composite plate of AA2124. In these studies, no details were provided of the tool geometry and the possible influence of features such as threads, on the resulting distribution of particles observed. The fragmentation of the W6 C carbides in this study occurs as the particles enter the rotation plug of metal around the weld tool. This can be due to either forces resulting from the severe shear zone or the temperature gradients within this region. Using a material independent, kinematic model developed by Nunes (2003), the severe shear zone is estimated to impart an instantaneously high shear strain rate in the range of 103 s−1 while the metal undergoes shear strains on the order of 50 to 100. During the solid state joining process, the flow of metal around the weld tool can be thought of as following slip lines as illustrated in Fig. 14. As the tool rotates, a non-symmetric flow field is generated. The retreating side (RS) of the weld is where the rotation vector and travel vector are in the opposite direction and the advancing side (AS) of the weld is where they are in the same direction. As the tool rotates past the center line of the weld, metal is picked up as each slip line is passed. This leads to a thicker region of rotating metal on the RS of the weld as described by Nunes (2003). As the tool rotation passes the center line in the wake of the weld, metal is deposited in the order accumulated. Thus larger particles would tend to be deposited in the nugget center and AS TMAZ as the slip lines converge. Metal that moves multiple times around the tool would capture the smaller particles which are eventually deposited on the RS as reported in studies by London et al. (2003). Fig. 15 illustrates the slip line entering the severe shear zone which demarks the rotating plug of metal from the workpiece
Fig. 15. Plan view illustration of shear zone surrounding the weld tool.
(Nunes, 2003). The shear zone is a region of finite thickness (ı) located a finite distance from the weld tool. Brittle particles of sufficient length to diameter ratio (L/D) would break due to the shearing forces.
5. Conclusions (1) Haynes 230 was shown to be successfully joined using the TSW process with the weld region exhibiting equiaxed grains. The grains were reduced by 85% to 11.2 m equiaxed diameter and the size of the W6 C phase was reduced over a range of 30 to 60%. (2) As compared to the wrought properties, the TSW displayed a slight decrease of 12% in the tensile strength and elongation to failure. (3) The TSW nugget displayed a higher hardness than the parent material which is attributed to Hall-Petch grain size strengthening. (4) The area fraction of the W6 C within the TSW region was similar to that of the starting parent material indicating the stability of the W6 C phase during the solid state TSW process. (5) The wiping motion of the smooth tool in the TSW zone results in an angular reduction of the shear zone around the tool letting smaller particles pass multiple times around the tool before being deposited on the retreating side. This results in changes in the W6 C size distribution across the TSW section oriented transverse to the welding direction.
Acknowledgements
Fig. 14. Illustration of metal flow following slip lines around the weld tool.
Funding was provided by the NASA-STTR Phase II grant no. KSE 11008 10070691 in collaboration with Keystone Synergistic Enterprises, Inc. All TSWs were conducted at the NASA-Marshall Space Flight Center under a Keystone Space Act Agreement in collaboration with Jeff Ding.
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Characterization equipment used on this project was purchased under NSF-IMR grant nos. DMR0216703 and 02070615 for the scanning electron microscope and detectors, NSF-MRI grant no. DMR0619773 for the X-ray diffractometer, and NSF DBI grant no. 1126743 for the transmission electron microscope and EDS detector. References Collins, M.G., Lippold, J.C., 2003. An investigation of ductility dip cracking in nickelbased filler materials—Part I. Weld. J. 82 (10), 288s–295s. Collins, M.G., Ramirez, A.J., Lippold, J.C., 2003. An investigation of ductility dip cracking in nickel-based filler materials—Part II. Weld. J. 82 (12), 348s–354s. Collins, M.G., Ramirez, A.J., Lippold, J.C., 2004. An investigation of ductility dip cracking in nickel-based filler materials—Part III. Weld. J. 83 (2), 39s–49s. Ding, R.J., 2011. Thermal Stir Welding Apparatus (US7980449 B2, July 19, 2011). Dupont, J., Lippold, J., Kiser, S., 2009. Welding Metallurgy and Weldability of NickelBased Alloys, first ed. Wiley and Sons Inc. Pub, Hoboken, NJ, pp. 131–143. Ernst, S.C., 1994. Weldability studies of Haynes 230 alloy. Weld. J. 73 (4), 80s–89s. Farias, A., Batalha, G.F., Prados, E.F., 2013. Tool wear evaluations in friction stir processing of commercial titanium Ti–6Al–4V. Wear 302 (1–2), 1327–1333. Farren, W.S., Taylor, G.I., 1925. The heat developed during plastic extension of metals. Proc. R. Soc., A 107, 422–451. Haynes International, 2007. Haynes 230 alloy. In: Report #H-3000H, http://www. haynesintl.com/pdf/h3000.pdf (accessed 04/06/2015). Humphreys, F.J., Hatherly, M., 2004. Recrystallization and Related Annealing Phenomena, second ed. Elsevier Pub, Oxford, UK, pp. 430–431. Klarstrom, D.L., 2001. The development of HAYNES 230. In: Zhao, J.C., Fahrmann, M., Pollock, T.M. (Eds.), Proceedings of Mat’ls Design Approaches & Experiences. TMS Pub, Warrendale, PA, pp. 297–307.
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