Friction welding of selective laser melted Ti6Al4V parts

Friction welding of selective laser melted Ti6Al4V parts

Author’s Accepted Manuscript Friction welding of selective laser melted Ti6Al4V parts K.G. Prashanth, R. Damodaram, T. Maity, P. Wang, J. Eckert www.e...

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Author’s Accepted Manuscript Friction welding of selective laser melted Ti6Al4V parts K.G. Prashanth, R. Damodaram, T. Maity, P. Wang, J. Eckert

PII: DOI: Reference:

S0921-5093(17)31010-9 MSA35352

To appear in: Materials Science & Engineering A Received date: 15 April 2017 Revised date: 23 June 2017 Accepted date: 2 August 2017 Cite this article as: K.G. Prashanth, R. Damodaram, T. Maity, P. Wang and J. Eckert, Friction welding of selective laser melted Ti6Al4V parts, Materials Science & Engineering A, This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Friction welding of selective laser melted Ti6Al4V parts K.G. Prashanth1,2,*, R. Damodaram3, T. Maity4, P. Wang5, J. Eckert2,4


Norwegian University of Science and Technology, Teknologivegen 22, 2815, Gjøvik, Norway


Erich Schmid Institute of Materials Science, Austrian Academy of Sciences, Jahnstraße 12, A-

8700 Leoben, Austria 3


SSN College of Engineering, Chennai, India Department Materials Physics, Montanuniversität Leoben, Jahnstraße 12, A-8700 Leoben,

Austria 5

Institute for Complex Materials, IFW Dresden, Helmholtzstraße 20, D-01069 Dresden,


Ti6Al4V alloy samples fabricated by selective laser melting (SLM) were subjected to solid-state welding friction welding (FW). The welded alloy exhibits a α´-martensitic microstructure in the form of platelets with dimensions in the submicron regime. The base alloy has a relatively coarser microstructure consisting of both α´ and β-phases, as compared to the asprepared SLM microstructure (single-phase α´martensite). Hardness measurements revealed an increase of hardness in the weld zone due to the refined α´platelets. A marginal drop in hardness along the base alloy is observed that may be attributed to the imposed thermal cycle during the FW process. Tensile tests reveal an improved ductility for the FW samples at the expense of a marginal drop in strength, compared to the as-prepared SLM samples. The present work 1

illustrates the ability of solid-state welding processes for successfully joining SLM parts and in improving, the ductility of the SLM parts and offers the opportunity to work with the additive manufacturing processes without size limits.

Keywords: Selective Laser Melting; Ti-based alloys; Friction Welding, Tensile behavior.


Corresponding author. Tel.: +47 97364667 E-mail address: [email protected] (K.G. Prashanth)


1. Introduction Powder bed laser fusion process (PBLP) like selective laser melting (SLM) has gained increased attention because of their ability to produce three-dimensional parts with theoretically any shape from a 3D CAD model [1-4]. SLM offer a novel approach for materials fabrication, where variation of the processing parameters and/or the scanning strategies can tune the mechanical properties of the parts to a significant extent [5-7]. Moreover, SLM allows to process a wide spectrum of materials ranging from Al-based alloys [8-11], Ti-based alloys [12-15], Nibased alloys [16-18], Fe-based alloys [19-21], Co-based alloys [22-24], Cu-based alloys [25] etc., unlike its competitor (EBM), which is limited to selected category of materials (like Ti, Ti6Al4V, CoCrMo, IN718, etc.,). In addition, SLM provides very high cooling rates in the order of ~105-106 K/s [26,27], thus often leading to the formation of metastable microstructures [28]. Extensive research has been carried out to evaluate the different properties of material manufactured by SLM, like mechanical properties, corrosion properties, tribological properties, fatigue, fracture toughness etc. [29-35]. Reports have shown that SLM fabricated components exhibit at least similar but often superior properties compared to cast counterparts and the results are very much reproducible [36]. However, some other reports disagree with the reproducibility concepts [37]. Similarly, researchers now disagree with the concept of energy density / heat input, which is often used as a guideline in optimizing the laser parameters for each alloy [38,39]. Even though there are several contradictions concerning the reproducibility of parts produced by SLM, this technique nowadays is widely used for variety of industrial applications [40]. A major limitation that prevails since a long time is the size limitation of the parts that can be built using the SLM process, which depends greatly on the size of the chamber. SLM


parts are limited in dimensions, since increasing the size of the chamber attracts several problems: (a) use of multiple sources of laser may have the following impact: inaccurate calibration between the laser sources and non-uniform decay process between the laser sources may lead to local differences in the mechanical properties of the parts. (b) Extending the use of laser source for wider area may lead to differences in the laser focus diameter (differences in the offset lengths) at the corners compared to the center of the building chamber, which in turn influences the microstructure and their mechanical properties. (c) De-powdering of largely built parts are difficult, especially it has fine features/thin cooling/internal channels. (d) There is a risk to flush the whole chamber with huge quantities of inert gas (which becomes expensive), in order to avoid the oxidation of parts during the fabrication process. Hence, there is a need to look for alternative secondary processes like additional/conventional metal joining techniques that can join smaller SLM components at a later stage. Even though there are several metal joining processes available, only some of them may be suitable for joining SLM fabricated parts. This is because the SLM fabricated parts have a very fine microstructure [9,32], which is not possible to be produced using conventional manufacturing processes [25]. Hence, if any of the fusion welding processes are employed for joining, the weld zone will be completely re-melted and solidified again at slower cooling rates than the ones realized upon SLM processing [41]. This leads to a relatively coarser microstructure and the advantages of the fine microstructures obtained using the SLM process is lost and so the mechanical properties of the parts are expected to deteriorate [7,9,32]. Moreover, SLM processed materials often contain high internal stresses [42-45] and hence cannot accommodate additional stresses imparted during the conventional fusion welding processes [46-48].


Because of these reasons, the solid-state welding processes the friction welding (FW) or friction stir welding (FSW) processes are preferred in this study, which mostly produce coalescence between parts at temperatures below the melting point of the materials being joined [49]. To the best of our knowledge, the FW welding of SLM processed Al-12Si is the only publication that deals with the joining of additive manufactured components [50]. The most interesting feature of the FW of Al-12Si is that in the weld zone, the hardness of the material deteriorates; thereby increasing the overall ductility of the FW specimens compared to the asprepared SLM Al-12Si specimens [50]. Moreover, the fracture of the welded samples take place in the base alloy and not in the weldment, which is considered as an advantage of using FW process for such SLM parts [50]. The present manuscript deals with the joining of Ti6Al4V parts fabricated by SLM using the FW process. This is because Ti6Al4V is one of the widely fabricated material by SLM and finds its application as both structural and biomaterial. The welded parts were studied for both structural and microstructural characterization techniques and the samples were tested for their mechanical properties using hardness and tensile tests. A detailed fracture analysis was carried out on the fractured surface to evaluate the possible reasons for failure.

2. Experimental details and sample preparation Cylindrical 12 mm diameter samples with a length of 60 mm were fabricated using Ti6Al4V spherical powders, purchased from concept laser. The powder particles had an average diameter of 30±4 µm. A concept laser MLAB device (with Nb-YAG laser) was used to process the Ti6Al4V samples. The fabrication of the samples were carried out at room temperature and a high purity argon gas was flushed inside the processing chamber to avoid any possible oxygen 5

contamination during the process. The following parameters were used to fabricate the Ti6Al4V samples: laser power - 95 W, scanning speed - 650 mm/sec, layer thickness - 30 µm, hatch spacing - 110 µm, hatch overlap - 30% and a straight line hatch with a rotation of 90° was used between successive layers. Adequate support structures were constructed between the rods and the substrate plate (made also of Ti6Al4V) in order to have adequate heat dissipation and to ensure good stability of the SLM parts during the fabrication process. FW of the samples was carried our using a continuous drive friction welding machine of 200 kN capacity (from MakeEta Technology, Bangalore, India). The following parameters were used for the FW process: friction pressure – 275 MPa, upset pressure – 550 MPa, burn off length - 3 mm and a rotational speed of 1500 rpm. The surfaces of the samples were machined before the welding process, in order to have smooth oxide-free surfaces, since oxidized surface regions may hamper the joint strength of the weldment [51]. In addition, the machining of the sample surface also helps in ensuring the perpendicularity of the surfaces, which is crucial for achieving sound welds. Structural characterization of the SLM and FW samples was carried out by X-ray diffraction (XRD) using a D3290 PANalytical X’pert PRO device. The XRD was carried out with Co-Kα radiation and a wavelength of 0.17889 nm using a Bragg-Brentano configuration. Microstructural characterization was done with a scanning electron microscope (SEM) using SEM – LEO 1525, coupled with an energy-dispersive X-ray spectroscopy setup. The macroscopic images were captured using an OLYMPUS SZX16 stereomicroscope, equipped with a CANON EOS 600D camera. Optical microscope images were recorded using an OLYMPUS BX51 microscope. The samples were polished using the following polishing papers: 400, 800, 1200 and 4000 grit. They were then subjected to cloth polishing with 3 µm diamond paste for 2 min and then with 2 µm diamond paste for 1 min followed by 0.25 µm diamond paste


for 1 min. The samples were subsequently cleaned using an ultrasonic bath filled with ethanol and finally etched with Kroll’s reagent (composition: 6% nitric acid, 2% hydrofluoric acid and 92% distilled water) for 5-10 seconds. The dimensions of the platelets were measured from the SEM micrographs using the imageJ software. At least 200 data points were collected to measure each dimension from at least five different SEM images. The Vickers microhardness measurements were performed using a manual Buehler microhardness tester with a pyramid diamond indenter, having a square base and at an angle of 136° between the opposite faces. A load of 0.5 kgf (~5 N) was applied for 10 s during each indentation measurement. Tensile tests were performed using an Instron 8562 facility under quasistatic loading conditions at room temperature with a strain rate of ~1 × 10−4 s−1. The tensile test samples (according to ASTM standard: ASTM:E8/E8M – 13a) were cylindrical tensile bars with 52 mm total length that were machined from the welded samples in such a way that the weldment lies along the center of the tensile bar. The dimensions along the gauge length of the tensile bar were: length - 17.5 mm and diameter - 3.5 mm. The strain measurement during the tensile tests were carried out using a Fiedler laser-extensometer. The mechanical tests were conducted for three samples depending on the scatter. Both the SLM as-prepared as well as the SLM-friction welded sample do not show significant scatter and is reproducible. The fracture surfaces of the samples were investigated using SEM to analyze the causes for the failure.

3. Results and discussion 3.1. Structural analysis The structural characterization of the Ti6Al4V as-prepared SLM samples is shown in Fig. 1. The XRD pattern shows the presence of hexagonally closed packed (hcp) α´martensitic peaks, 7

which is typical for SLM processed Ti6Al4V samples, similar to other published reports [52-54]. The SEM image in Fig. 2 shows the microstructure of the as-prepared and etched SLM samples. Fine acicular martensitic (α´) platelets are observed that are formed due to high cooling rates established during the SLM process [26-27]. Again, these findings are similar as in other published reports [52-54]. The average dimensions of the martensite platelets were found to be ~7±3 µm in length and ~0.77±0.21 µm in width. Some β-phase was also observed in the SEM images, which is consistent with published data [52-54]. However, the peaks of β-phase are not observed in the XRD patterns, which may be due to the low concentrations of this phase (~4-7 vol.%) that may be difficult to be resolved during the XRD measurements. The Vickers microhardness of the as-prepared Ti6Al4V samples is ~422±18 HV0.5. Fig. 3 displays a low magnification stereomicroscope image of the FW Ti6Al4V SLM joint. It shows flash formation, which corresponds to the plastically deformed material [51]. The flash morphology depends greatly on the type of material processed and in case of the Ti6Al4V FW samples, a single-sided flash morphology is observed, which is typical for Ti-based alloys, unlike as for Al-based/Fe-based alloys [55-57]. An appreciable amount of flash formation was observed in the joint, which is a clear indication of an adequate amount of heat generation and subsequent plastic deformation at the weld interface. It is interesting to note that the fusion zone in the weldment is quite narrow around ~150±15 µm and extends to a length of ~1.5±0.1 mm, before the start of the plastically deformed material being pushed towards the edges, leading to the formation of the flash in the FW Ti6Al4V SLM samples. A very thin heat affected zone is observed next to the weldment (10 – 20 µm), where the features from the SLM processing are not prominent. Moreover, it is interesting to note that there are no visible defects observed along the weldment, suggesting a sound quality of the joint.


Further structural characterization of the FW Ti6Al4V SLM samples was carried out using XRD near the weld zone (Fig. 2). The XRD pattern again shows the presence of hexagonally closed packed (hcp) α´martensitic peaks, as it was originally observed for the asprepared SLM material. However, there is a crystallographic texture along the (0002) peak as compared to the (1011) peak, which is absent in the as-prepared SLM samples. In addition, the peaks of the FW Ti6Al4V SLM samples are less intense for the as-prepared SLM samples. This may be due to excessive strain observed along the weld zone, as a results of excessive plastic deformation and flash formation. Moreover, additional peaks of the β-phase are also observed for the FW Ti6Al4V SLM sample. The microstructure of the FW Ti6Al4V SLM samples is shown in Fig. 4. It is interesting to note that the flash part of the sample has rather coarse microstructural features compared to the weldment along the center of the sample (Fig. 4(a-c)). Both the flash and the weld zone contains acicular hexagonally closed packed (hcp) α´martensite; however, their dimensions are different. The martensite platelets observed in the weld zone are ~0.5±0.1 µm in length and ~0.1±0.05 µm in width (Fig. 4 (f-h)), whereas the martensite platelets in the flash region are about ~1.3±0.5 µm long and ~0.5±0.1 µm wide. The laser tracks in the base alloy are still recognizable even after the FW process (Fig. 4(b-d)), suggesting that the base alloy still persists with the SLM features at a macrostructural level along with few changes at the microstructural scale. The length and width of the α´martensite phase is found to be ~12±3 µm and 0.85±0.18 µm respectively (Fig. 4(d,e)), suggesting a growth of the α´martensite from ~7±3 µm µm in length and ~0.77±0.21 µm in width, respectively. This growth is attributed to the heat dissipation process during the FW. The heat generated in the weld zone has to be dissipated to attain thermal equilibrium and the heat is conducted via the base alloy. Since the base alloy acts as medium for


heat dissipation, its microstructure is subjected to certain thermal cycle for considerable period of time, which is sufficient enough to influence the microstructural features in the base alloy. In addition, Fig. 4(f) reveals that there is a small, 10-20 μm wide, heat affected zone between the weld zone and the base alloy. The size of the martensite platelets (~0.5±0.1 µm in length and ~0.1±0.05 µm in width) in the heat affected zone is bigger than in the weld zone (~4±1 µm in length and ~0.8±0.1 µm in width), but smaller than in the base alloy (~12±3 µm in length and 0.85±0.18 µm in width).

3.2. Mechanical properties The Vickers microhardness profile along the weld interface for two FW Ti6Al4V SLM samples are shown in Fig. 5 (a) with the black curve representing sample 1 and the other sample corresponding to the red curve. Both samples exhibit similar hardness profiles, revealing a good consistency of results between samples. The microhardness profiles show a nearly symmetrical behavior with respect to the weld interface in the FW Ti6Al4V SLM samples, which is very similar to the behavior observed for FW Al-12Si SLM samples [50]. The hardness at the weld interface increases to 449±5 HV0.5, which is ~30 HV0.5 higher than the hardness of the original as-prepared SLM Ti6Al4V material. This hardness increase along the weld zone may be attributed to the formation of a refined α´martensitic microstructure and its morphology in the weld zone, compared to the initial as-prepared SLM Ti6Al4V material. On the other hand, the hardness of the base alloy show a drop to about ~395±20 HV0.5, which is ~25 HV0.5 lower than for the initial as-prepared SLM Ti6Al4V samples. The hardness drop in the base alloy for the FW Ti6Al4V SLM sample is due to the growth of the martensite platelets, which is attributed to the heat dissipation effects and the prevailed thermal energy supplied to the base alloy. 10

Characteristic tensile test curves of both as-prepared SLM and FW SLM Ti6Al4V samples at room temperature are shown in Fig. 5(b). The tensile test curves for the different specimens overlap until they reach strength levels of ~1000 MPa. The as-prepared SLM samples have a yield strength of ~1250±15 MPa and an ultimate strength of ~1355±20 MPa with ~7.5±0.5% ductility. These data are similar to results reported by Vrancken et al. [52], Krakhmalev et al. [54] and Wanjara et al. [55]. The FW Ti6Al4V SLM samples exhibit a marginal deterioration in strength, with a yield strength around ~1145±10 MPa and an ultimate strength around ~1300±15 MPa, which is ~105 MPa (yield strength) and ~55 MPa (ultimate strength) lower than for the as-prepared SLM samples. However, the ductility of the FW Ti6Al4V SLM samples reaches ~12±1%, which is ~4.5% higher than for the as-prepared SLM sample. The strength effects observed in for the FW Ti6Al4V SLM samples are similar to what was found for FW Al-12Si SLM samples, where both the yield and ultimate tensile strength of FW samples show a marginal decrease but also becomes more ductile, because of the relatively coarser microstructure observed in the base alloy after the FW welding process [50]. The fracture surfaces of as-prepared SLM and FW samples are shown in Fig. 6. Both the as-prepared SLM and FW samples fracture well away from the center of the gauge length/weld (Fig. 6(a,e)). Even though the hardness of the FW sample increases along the weld zone, the fracture of the sample takes place in the base alloy unlike the conventional FW samples [58,59]. The as-prepared SLM samples show a relatively more brittle fracture compared to the FW samples (Fig. 6(b,f)), where the FW samples show pronounced necking as compared to the asprepared SLM samples. Both types of samples show dimples corresponding to ductile failure mode (Fig. 6(c,d,g,h)) and the morphology of the dimples varies between the two samples depending on the degree of deformation before the failure. However, in the as-prepared SLM


samples, there are some regions showing quasi-cleavage (Fig. 6(c)), which is mainly responsible for the decrease in ductility for the as-prepared SLM samples.

4. Conclusions The effect of FW on the microstructure and mechanical properties of Ti6Al4V parts produced by SLM has been studied in detail. FW was primarily selected to join the Ti6Al4V SLM parts in order to avoid different solidification related problems, namely hot cracking. Phase analysis of the Ti6Al4V FW SLM parts show the presence of both α´- martensitic phase and βphase unlike the parent as-prepared SLM material, where only the α´- martensitic phase was observed from the XRD pattern. Microstructural studies show that the SLM-weld zone exhibits significant changes in the morphology of the α´- martensitic phase. The size of the α´- martensite platelets in the weld zone was found to be ~0.5±0.1 µm in length and ~0.1±0.05 µm in width as opposed to the parent as-prepared SLM material, where the size of the α´- martensite platelets is ~7±3 µm in length and ~0.77±0.21 µm in width. Moreover, the microstructure in the base alloy consists of binary phases: α´- martensite and β-phase. The size of the α´-martensite platelets in the base alloy is larger in the as-prepared SLM samples, which can be attributed to the thermal energy generated during the FW process and is dissipation through the base material. Such microstructural variations influence the mechanical properties of the FW samples, where an increased hardness ~449±5 HV0.5 is observed along the weld zone. On the other hand, the base alloy softens, and the hardness drops to ~395±20 HV0.5 as opposed to a hardness of ~422±18 HV0.5 for the parent material. Tensile tests revealed that the FW samples become softer compared to the as-prepared SLM specimens: the yield strength drops to ~1145±10 MPa (from ~1250±15 MPa for as-prepared SLM samples), the ultimate 12

strength drops to ~1300±15 MPa (from ~1355±20 MPa) and the ductility increases to ~12±1% (from ~7.5±0.5%). Fracture surface analysis revealed that the fracture of the as-prepared SLM samples exhibits both dimple fracture and quasi-cleavage features but the FW samples predominantly fail by dimple fracture, i.e. exhibit better ductility. These findings demonstrate that solid-state processes like friction welding can be successfully used to join SLM fabricated materials and may also be beneficial for improving the ductility at only a marginal decrease in strength. This offers a possible solution for overcoming the problem of limited dimensions achievable with additive manufacturing processes, especially powder bed fusion processes.

Acknowledgement The authors would like to thank Dr. G. D. Janaki Ram, Indian Institute of Technology Madras (IITM), Chennai for his support to carry out the friction welding experiments at IITM, India.

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Figure captions Figure 1. XRD patterns (λ = 0.17889 nm) of the Ti6Al4V samples in the as-prepared SLM and friction welded condition.

Figure 2. SEM micrographs of an Ti6Al4V as-prepared SLM sample at different magnifications after etching with Kroll’s reagent for 5-10 seconds.

Figure 3. Typical image of a friction welded Ti6Al4V SLM joint with single side smooth flash at the joint.

Figure 4. Optical micrographs of a friction welded Ti6Al4V SLM sample after etching with Kroll’s reagent for 5-10 seconds showing (a) flash, (b,c) weld joint and base alloy. SEM micrographs of a friction welded Ti6Al4V sample after etching with Kroll’s reagent for 5-10 seconds showing (d,e) base alloy, (f) base alloy, weld zone and heat affected zone and (g,h) weld zone.

Figure 5. (a) Vickers microhardness profile measured across the weld interface of Ti6Al4V SLM samples and (b) room temperature tensile curves of the SLM specimens tested in the asprepared and welded conditions.

Figure 6. Fracture morphology of the Ti6Al4V samples in the (a-d) as-prepared SLM condition and (e-h) friction welded condition.


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