Wire based additive layer manufacturing: Comparison of microstructure and mechanical properties of Ti–6Al–4V components fabricated by laser-beam deposition and shaped metal deposition

Journal of Materials Processing Technology 211 (2011) 1146–1158

Contents lists available at ScienceDirect

Journal of Materials Processing Technology journal homepage: www.elsevier.com/locate/jmatprotec

Wire based additive layer manufacturing: Comparison of microstructure and mechanical properties of Ti–6Al–4V components fabricated by laser-beam deposition and shaped metal deposition Bernd Baufeld a,∗ , Erhard Brandl b , Omer van der Biest a a b

Department of Metallurgy and Materials Engineering, Katholieke Universiteit Leuven, Belgium EADS Innovation Works, Metallic Technologies & Surface Engineering, Munich, Germany

a r t i c l e

i n f o

Article history: Received 9 August 2010 Received in revised form 5 January 2011 Accepted 24 January 2011 Available online 1 February 2011 Keywords: Additive layer manufacturing Shaped metal deposition (SMD) Laser beam melting Ti–6Al–4V Ultimate tensile strength High cycle fatigue

a b s t r a c t The microstructure and the mechanical properties of Ti–6Al–4V components, fabricated by two different wire based additive layer manufacturing techniques, namely laser-beam deposition and shaped metal deposition, are presented. Both techniques resulted in dense components with lamellar ˛/ˇ microstructure. Large ultimate tensile strength values between 900 and 1000 MPa were observed. The strain at failure strongly depends on the orientation, where highest values up to 19% were obtained in direction of the building direction. Heat treatment increased the highest strain at failure up to 22%. The fatigue limit was observed to be higher than 770 MPa. © 2011 Elsevier B.V. All rights reserved.

1. Introduction Manufacturing components in a layer-by-layer fashion offers a high geometrical flexibility and great potential of time and cost savings in comparison to conventional manufacturing technologies. The technology is known under many names, such as rapid prototyping, rapid manufacturing, free form fabrication, or additive layer manufacturing (ALM). Nowadays, not only prototypes are demanded from these techniques, but also serial production parts are envisaged. Ti–6Al–4V is the most widely used titanium alloy and one of the most common aerospace alloys (Peters and Leyens, 2003). The additive manufacture of small and medium-sized Ti–6Al–4V parts represents an interesting business case for the aerospace industry. However, additive processes are currently not used for the manufacture of serial production parts for aerospace applications. This is primarily due to outstanding issues regarding material properties and repeatability. Most additive manufactured Ti–6Al–4V parts are built up from powderized feedstock in either a powder-bed or a powder-feed process. Generally, powder has a big surface-to-

weight ratio and therefore a high contamination risk. Especially for Ti–6Al–4V, contaminations have a strong impact on the mechanical properties. Wire-feed deposition is discussed as a promising technology in this area (Brandl et al., 2008). The generally lower contamination of using wire than powder is an advantage regarding material quality (Mok et al., 2008a). Technologies applying wire in contrast to powder seem to offer higher repeatability levels, due to its simpler process setup and operation, and higher deposition rates. However, experience and development activities regarding wire based manufacturing of Ti–6Al–4V components are relatively limited to date. This paper aims to investigate and evaluate the material properties derived from two different wire-feed based systems: in one system Ti–6Al–4V is deposited by a Nd:YAG laser beam, in the other by a tungsten inert gas torch, a process also known as shaped metal deposition (SMD). Microstructure and mechanical properties of the deposits are presented and evaluated from an aerospace point of view. The mechanical tests including static tension and high cycle fatigue were performed in as-built, stress-relieved and annealed conditions. 2. Experimental

∗ Corresponding author. Present address: NAMRC, University of Sheffield, Brunel Way, Catcliffe, Rotherham S60 5WG, UK. Tel.: +44 1142229919. E-mail addresses: b.baufeld@sheffield.ac.uk (B. Baufeld), [email protected] (E. Brandl), [email protected] (O. van der Biest). 0924-0136/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2011.01.018

2.1. Wire based additive layer manufacturing The experimental setup of the laser beam ALM essentially comprises a Trumpf HLD 3504 Nd:YAG rod laser (diode pumped) with

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158 Table 3 Mechanical testing module and number of tests.

Table 1 Chemical composition of wires used and final ALM components.

ELI wire (laser ALM) Conventional grade wire (SMD) Final component (laser ALM) Final component (SMD)

O [wt.%]

N [wt.%]

0.045 0.145 0.073 0.150

0.010 0.008 0.105 0.002

a maximum power of 3.5 kW, a Weldaix wire feeder and a Kuka KR 100 HA (high accuracy) 6-axis robot. In order to prevent oxidation during deposition the ALM was performed in an open box, which is permanently flooded by argon (99.9996% purity) from its base (design of the box based on (Bergmann, 2004)). The setup was developed at EADS Innovation Works in Ottobrunn/Munich (Germany) (Brandl et al., 2008, 2009a,b). The SMD cell consists of a tungsten inert gas welding torch attached to a 6-axis Kuka robot which is linked to a 2-axis table. The whole setup is enclosed in an airtight chamber filled with argon (99.999% purity) and the moisture is controlled. This allows obtaining a final product with low oxygen and nitrogen contamination. More details about the setup can be found in (Baufeld and van der Biest, 2009; Baufeld et al., 2009, 2010a,b). The controlling of the robot is performed by a CAD/CAM process where-by an off-line program provides the robot with the necessary weld path information. For both systems, Ti–6Al–4V welding wire is deposited layer by layer onto a Ti–6Al–4V substrate. For the laser beam ALM extra low interstitial (ELI) material and for SMD conventional grade was applied. In Table 1 the amount of oxygen and nitrogen, which have a particularly high influence on the properties, are given. The components were built by single beads. In the case of laser beam ALM several walls adopting a single deposition direction and in the case of SMD one tubular component with a square base (275 mm × 275 mm) were produced. The chosen coordinate system is the direction of the welding speed x, the direction of the wall width y, and the direction of the wall height z. The laser beam ALM was paused between each layer until the temperature of the previous layer fell below 300 ◦ C, while SMD was performed continuously. In the case of laser beam ALM the welding speed derives only from the movement from the laser beam, while in the case of SMD the welding speed is a combination of rotation and tilting of the table and the movement of the torch. The process parameters for both techniques were derived in individual optimization procedures and are summarized in Table 2, indicating that power, welding speed, wire feed speed and height increment are comparable. The difference in wall widths (also given in Table 2) may result from a larger weld pool, a slightly higher wire feed speed, and a slower welding speed of the SMD. 2.2. Mechanical testing The static tensile tests were performed on a Z250 (Zwick) universal testing machine at room temperature according to EN 10002 Table 2 Process parameters of laser ALM and SMD.

Power (kW) Welding speed (mm/s) Wire-feed speed (mm/s) Wire diameter (mm) Height increment (mm) Deposition rate (kg/h) Wall thickness (mm)

1147

Laser ALM

SMD

3.5 10 40 1.2 1.0 0.7 4–5

2.2 5.0 33 1.2 1.0 0.6 9.1

Laserbeam x Tensile testing As fabricated 3 4 600 ◦ C/4 h 2 843 ◦ C/2 h High cycle fatigue (HCF) As fabricated 0 13 843 ◦ C/2 h

Laserbeam z

SMD x

SMD z

4 5 4

4 4 4

4 4 4

0 17

8 8

9 9

(DeutscheNorm, 2001) with a test velocity 0.4 mm/min. The cylindrical, dog-boned specimens had a gauge length of 12 mm and a gauge cross section of 2 mm. The high cycle fatigue (HCF) specimens were tested according to DIN 50100 (DeutscheNorm, 1978) at room temperature on a Microtron 654 (Rumul) resonance tester. A test frequency of approximately 100 Hz and a load ratio of R = 0.1 were used. The 26 mm long, hour glass shaped specimens had a minimal cross section of 1.5 mm. The stress concentration factor is approximately ˛K = 1. Tests were terminated at 4 × 107 cycles. Besides as-fabricated specimens, heat treated specimens were investigated. One heat treatment was the so-called stress-relieving, commonly applied in aerospace industry for Ti–6Al–4V parts, which is meant to reduce residual stresses without substantial change of the microstructure (SAE Aerospace, 2003). The other treatment was annealing at 834 ◦ C to eliminate possible precipitation hardening or to develop a stable initial state (DeutscheNorm, 1990). Stress relieving was performed at 600 ◦ C for 4 h and annealing at 843 ◦ C for 2 h, both in vacuum in the range of 10−4 to 10−5 mbar followed by furnace cooling with roughly 6 ◦ C/min at high temperatures. All samples were electrochemically polished (∼60 ␮m diameter reduction) to remove potential residual stresses, hardening effects and contaminations at the surface, which may derive from machining/milling and, where applicable, from heat treatments and which would cause erroneous cycle fatigue results. For both mechanical testing types two differently oriented specimens were fabricated, i.e. with mechanical loading direction parallel to the direction of the welding speed x, respectively parallel to the direction of the wall height z. These specimens will be called x-, respectively z-specimens in this paper. The number, type, orientation, and thermal history of the mechanically tested specimens are summarized in Table 3. As will be shown in the results, the very top of the walls (∼10 mm) exhibits a different microstructure than the bottom. Due to the larger wall height of the SMD component, it was possible to extract specimens where the tested parts always were outside of this very top region. In the case of laser beam ALM components, however, some of the x-specimens were derived from this top region. 2.3. Microstructure and hardness The microstructure was investigated using optical and scanning electron microscopy (SEM, FEI XL30FEG) of polished or etched cross sections. The etching agent was 6 vol.% nitrid acid and 3 vol.% hydrofluoric acid in water. Back scattered electron (BSE) imaging with very high contrast allows discerning the ˛ and ˇ phases without etching and therefore avoiding possible artifacts from the etching. The Vickers microhardness tests were performed on cross sections of as fabricated and heat treated specimens by a Leitz/Durimet 2 microhardness tester using a weight of 100 g (HV0.1). Gen-

1148

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

Fig. 1. Surface morphology of the SMD component (a, growth direction of the columnar prior ˇ grains highlighted by line) and etched cross sections of laser beam (b) and SMD (c) ALM components.

erally, for each sample the average of 20 measurements and a confidence interval of 95%, based on the variance  n−1 , were determined. In addition, in order to investigate possible influence of the hardness on the location within the components, indentations were performed on y–z planes from top to bottom in 0.5 mm intervals.

3. Results 3.1. Morphology and microstructure The surface of the laser beam ALM components is dull and gray, while the one of the SMD component is shiny and has variation in

Fig. 2. Microstructure of a laser beam ALM component, as-fabricated, near the top (a, b), and in the central region (c, d).

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

1149

Fig. 4. Light microscopical image of an etched cross section exhibiting the morphology, including a grain boundary, from the top region of a laser beam ALM component. The arrow indicates a lamella outgrowing other members of the same grain boundary colony forming a grid with similar more successful lamellae.

Fig. 3. Top region of a laser beam ALM component. (a) Etched cross section by light microscopy. (b) Polished cross section of detail indicated by box in (a).

color of brown and blue. The surfaces of components built by both techniques show, viewed from the side, layered bulges parallel to the deposition plane and large, elongated columnar grains (Fig. 1a). These elongated grains are prior ˇ grains which grew epitaxially across the welding layers (Baufeld et al., 2009). They are inclined relative to the z direction, following the largest temperature gradient resulting from the moving laser beam, respectively welding torch. The layered bulges indicate the sequential deposition layers. The top of the components is round due to the surface tension of the melt. The etched cross sections (Fig. 1b and c) exhibit aslant cut prior ˇ grains (the one of laser ALM are smaller than the one of SMD). Furthermore, two different regions can be discerned with a top region and a bottom region. The bottom region is characterized by bands parallel to the base plate, while the top region lacks these parallel bands (Fig. 1b and c). Examples of the microstructures observed in components in the as-fabricated state are given in Figs. 2–4 for laser beam ALM and in Fig. 5 for SMD. Generally, the microstructure consists for both ALM techniques of ˛ phase lamellae (gray shades in the SEM micrographs) in a ˇ matrix (white contrast in the SEM micrographs). The different gray shades derive from orientation contrast and not from material contrast (Baufeld et al., 2009). The morphology can

be differentiated between basket weave structures (for example Figs. 2c and 5c), possibly the main type in the bottom region, and colony structures of parallel bundles (Figs. 2a, b and 5a). Typically, colony structures can be observed at wall surfaces, grain boundaries of prior ˇ grains (Fig. 3a) and along the ˇ transus lines of the ALM steps (Baufeld et al., 2009, 2010a). The top region of the laser beam ALM is dominated by colonies of lamellae starting at grain boundaries. Occasionally, for both types of ALM a rectangular grid can be observed by light optical means of etched cross sections at relatively low magnification (Figs. 3a and 4; Fig. 4 showing only lamellae forming one direction of the grid). These examples are from the top region of a laser beam ALM component, but similar grids can be found also in SMD components and are not restricted to the top region (Baufeld et al., 2009). SEM investigations reveal that the grids consist of sets of very long lamellae perpendicular to each other (Fig. 3b). Such lamellae are surrounded by the ˇ phase (white contrast) and the space between the grid is filled with smaller lamellae. In the case of the top region of a laser beam ALM, this space is filled by colonies (Fig. 3b). In the case of SMD from the bottom region basket weave microstructure is reported to fill this space. (Baufeld et al., 2009). Detailed investigation of the lamellae in Fig. 3b by EDS revealed, that both types of lamellae contain a similar composition of 6 wt.% Al and 2 wt.% V. The white phase exhibit a lower amount of Al and a higher amount of V, but the dimensions of the white phase are too small to determine reliable values. Fig. 4 suggests that the lamellae forming the grid are similar to the lamellae forming the grain boundary colonies, but by chance were growing just more successful than the other lamellae of their colony. Regions with ˇ phase developed during first cooling after deposition may experience diffusional partitioning by being subjected to temperatures within the ˛/ˇ phase field during subsequent ALM steps. This may lead to the formation of secondary ˛ lamellae. An example for this phenomenon is shown in Fig. 6. The lamellae widths in dependence of the location are given in Fig. 7, each data point derives from 50 measurements of a micrograph with 2000× magnification. The error bars describe the standard variations. Both components exhibit very thin lamellae with a width of about 0.6 ␮m in the top region and broader lamellae in the region below. Yet, even in the bottom regions areas can be found with smaller lamellae. In addition, the component from laser beam ALM exhibits a decrease of lamellae width near the base plate

1150

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

Fig. 5. Microstructure of the SMD component, as-fabricated, near the top (a, b) and in the central region (c, d).

which coincides with increased occurrence of smaller ˇ grains and grid structures commonly attributed to martensites. Figs. 8 and 9 show the microstructure of laser beam ALM and SMD after heat treatment at 600◦ , respectively at 843 ◦ C. Apparently, these heat treatments do not change significantly the lamellar structure. The measured lamellae width all are in the range of 1 ␮m indicating that no apparent coarsening can be reported.

Fig. 10 exhibits the stress–strain curves of laser beam ALM (a–c) and SMD specimens (d–f) for the different heat treatments and ori-

entations. Obviously, the deformation behavior depends strongly on the type, orientation and heat treatment. Furthermore, within one set some variation can be observed. The specimens show some amount of plastic deformation with limited work hardening and in some cases apparent work softening attributed to extensive necking. The stress–strain curves of the x-specimens from the laser beam ALM components can be discerned into to different sets, one set with significantly higher strength than the other (Fig. 10a: the stresses of specimen 6-5 larger than of specimens 66 and 6-7, Fig. 10b: the stresses of specimens 5-1 and 5-2 larger than of specimens 5-3 and 5-4). The x-specimens with the higher strength (6-5, 5-1, 5-2) all were extracted from the top region of the

Fig. 6. Secondary ˛ lamellae in the central region (highlighted by inset) of an asfabricated laser beam ALM component due to temperature treatment within the ˛/ˇ phase field developed during deposition.

Fig. 7. Lamellae width, including standard deviation, derived from SEM micrographs of a laser beam and a SMD ALM component in dependence on the distance from the top. The vertical line indicates the level of the plate in the case of laser beam ALM.

3.2. Tensile testing

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

1151

Fig. 8. Microstructure of a laser-beam ALM component from the central region, after heat treatment at 600 ◦ C (a, b) and at 843 ◦ C (c, d).

components, while all other x-specimens derive from the bottom region. Fig. 11 gives a compilation of the results in form of UTS versus elongation at failure plots. It can be concluded that generally under similar circumstances (i.e. similar orientation and heat treatment)

the UTS of specimens from laser beam ALM components is higher than the UTS of specimens from SMD components. The UTS of x-specimens from laser beam ALM are larger than the one of zspecimens. This tendency is not so clear for SMD specimens and especially the SMD specimens after heat treatment at 600 ◦ C have

Fig. 9. Microstructure of the SMD component, from the central region, after heat treatment at 600 ◦ C (a, b) and at 843 ◦ C (c, d).

1152

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

Fig. 10. Stress–strain curves of laser beam (a–c) and SMD ALM specimens (d–f) in the as-fabricated state (a, d), after heat treatment at 600 ◦ C (b, e) and 843 ◦ C (c, f).

similar UTS in x and z direction. x-Specimens clearly exhibit a much smaller elongation at failure than z-specimens. Furthermore, the elongation at failure generally is smaller for laser beam ALM than for SMD components. Within the significant scatter of UTS and elongation at failure it is difficult to discern a reliable effect of heat treatment. Generally, the UTS seems hardly to be affected by the heat treatment. The elongation at failure, however, is increased by the heat treatment at 843 ◦ C.

A comparison with the minimum requirements for wrought (AMS 4928), respectively for cast (ASTM F1108) material indicates that the demanded ultimate tensile strength usually is fulfilled by the current ALM material. Much more critical, however, is the elongation at failure. Only SMD z-specimens in all states of heat treatment, and SMD x-specimens after heat treatment at 843 ◦ C fulfill the fierce requirements put on wrought material.

Fig. 11. Plot of ultimate tensile strength (UTS) versus elongation at failure of tensile tests on specimens from laser beam (a) and SMD ALM components (b) in the as-fabricated state, and after heat treatment at 600 ◦ C and 843 ◦ C. The minimum requirements for cast and wrought material are indicated by lines. The three specimens derived from the top region of laser ALM components indicated by ellipse (a).

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

1153

Fig. 12. High cycle fatigue properties of laser beam and SMD ALM specimens, tested along (x) and across (z) deposition direction; one data point represents one tested specimen. The line represents the upper fatigue limit required for wrought annealed material.

3.3. Fatigue testing The range of maximum stress applicable in fatigue testing is very limited, since the upper limit is given by the yield strength, as plastic deformation has to be avoided in HCF tests (DeutscheNorm, 1978), and the lower limit by the run out within reasonable testing time. Accordingly, the applied maximum stress values were between 750 and 913 MPa. The result of the high cycle fatigue tests is presented in Fig. 12. The fatigue limit is difficult to determine, since fatigue experiments suffer intrinsically under the large fluctuation of life to failure. Especially the as-fabricated SMD specimens exhibit a large fluctuation in fatigue life with no clear indication of a SN curve. The heat treated SMD specimens display much less fluctuations. Nevertheless, it may be concluded, that for both orientations, and independent on the heat-treatment, similar results can be obtained with a fatigue limit probably higher than 770 MPa. The heat treated laser beam ALM x-specimens have a significant higher fatigue limit, probably larger than 840 MPa. The laser beam ALM z-specimens have lower fatigue limit than the x-specimens and exhibit a number of very early failures. The fracture surfaces usually are very rough; sometimes transgranular fracture along the colonies can be recognized (Fig. 13a). In one case, the failure can be attributed to a pore. This SMD xspecimen failed already after 1.7 × 105 cycles with a maximum load of 800 MPa and it is noteworthy that it failed not at the thinnest cross section like all the other specimens. Fig. 13b shows the whole fracture surface and the pore as the origin. The fracture crack preferential runs along the ˛/ˇ interface omitting to cut the ˛ lamellae (Fig. 14). 3.4. Hardness The microhardness HV0.1 of the laser beam ALM and of the SMD component in dependence on the location is shown in Fig. 15. For the laser beam ALM components high values are recorded near the top, and decreasing values towards a constant in the bottom. For the SMD component, within the large scatter, the values seem not to depend on the location. The effect of heat treatment is documented in Table 4 for the bottom regions. Due to the hardness gradient in the laser beam ALM components the top and the center area has to be discerned. Apparently, the different heat treatments do not have an influence on the hardness in the central regions. Possibly, the hardness of the top regions of laser beam ALM specimens is decreased by the heat treatments.

Fig. 13. Fracture surface of a SMD x-specimen failed after 1.7 × 105 cycles with a maximum load of 800 MPa.

4. Discussion 4.1. Microstructure Titanium alloys are very well known for their polymorphism which depends critically on the thermal history. Globular microstructures, as resulting for example from hot working (Boyer et al., 1994), naturally does not occur in the case of ALM. Moreover, different types of micro-structures, such as colony and basket weave Widmanstätten, massive ˛ or martensitic ˛ phases, depending on the cooling rates, may develop (Ahmed and Rack, 1998; Boyer et al., 1994). This is due to the fact, that ALM subjects the components to sequential heating and cooling with rates which may vary locally to a large amount. Cooling from the liq-

Table 4 Microhardness of laser deposited and SMD component (below the top region) in the as fabricated state, and after heat treatment at 600 ◦ C, respectively 843 ◦ C. Laser ALM Vickers hardness [HV0.1] As fabricated 392 ± 14 (top 10 mm) 332 ± 7 (center 10 mm) 358 ± 16 (top 10 mm) 600 ◦ C/4 h 326 ± 4 (center 10 mm) ◦ 391 ± 11 (top) 843 C/2 h 330 ± 4 (center)

SMD component 337 ± 7 335 ± 7 336 ± 9

1154

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

Fig. 14. Fatigue crack of a laser beam ALM component in overview (a). The box indicates the magnified area of (b).

uid state, Ti–6Al–4V first traverses the bcc ˇ phase arriving below 1000 ◦ C the extensive ˛ + ˇ phase field (Boyer et al., 1994). Slow cooling leads to colony structure, intermediate cooling to basket weave structure, and fast cooling to diffusion less transformation into martensite ˛ (Boyer et al., 1994; Lütjering and Williams, 2007). It was pointed out, that some authors call the colony and the basket weave together Widmanstätten, while others reserve this expression only to the basket weave structure (Vanderesse et al., 2008). Vanderesse et al. subscribe to the first definition and claim that the average number of lamellae is responsible whether in cross section a microstructure appear as colony or

Fig. 15. Microhardness in dependence on the distance from the top (y–z plane) of laser beam and SMD ALM components.

basket weave, since a low average number goes along with a fine entanglement of the colonies, resulting in a basket weave microstructure and a large average number appear as colony structure. In addition, at cooling rates higher than 20 ◦ C/s so-called massive transformation into ˛m phase is reported (Ahmed and Rack, 1998). These authors observe colonies of lamellae, starting at grain boundaries, which they explain with trans-interphase boundary diffusion based massive transformation. According to this work, at cooling rates between 20 and 410 ◦ C/s the massive transformation occurs in addition to martensitic transformation, decreasing in amount with increasing cooling rates. No significant difference in composition between the martensite and the massive ˛m phase was observed. This is in contrast to colonies forming at much lower cooling rates and to basket weave lamellae forming at intermediate cooling rates, which exhibit due to extensive diffusion increased Al and decreased V concentrations (Boyer et al., 1994). The ˇ phase is stable at room temperature only if it is enriched with more than 15 wt.% V (Boyer et al., 1994). Since ˛, massive ˛m , and martensitic ˛ have all a hexagonal closed packed crystal structure and are difficult to discern by XRD. However, the description and the proof of massive transformation (Ahmed and Rack, 1998) is not conclusive since it is not clearly explained how the authors discern massive ˛m from normal ˛ or martensitic ˛ . According to this work the main characteristic for massive ˛ is the morphology of colonies of lamellae growing from grain boundaries, which have the same composition as the ˇ phase at high temperatures, respectively the ˛ martensites. However, another explanation for ˛ lamellae having similar composition as ˛ martensites may be given by the highly dynamic processing scheme of ALM. It must be kept in mind that most experimental data are based on equilibrium heat treatment followed by different cooling rates. Most of the published phase diagrams are based on the equilibrium condition. In the case of ALM however, a dynamic heat treatment takes place. Possibly, the well known phase diagram (Boyer et al., 1994) does not apply in such a case, and a smaller ˛ + ˇ phase field during rapid cooling must be considered (Banerjee et al., 2003). This would implement a cooling rate dependence of the ˛ phase composition. High cooling rates probably lead to a composition of the ˛ phase close to Ti–4Al–6V (and martensite), while with decreasing cooling rates the Al concentration would increase

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

and the V concentration would decrease towards the equilibrium condition. Considering these different processes, it is obvious that ALM with its sequential heating and cooling steps results in very specific microstructures. The two different regions, visible for etched cross sections (Fig. 1b and c), with a bottom region sporting parallel bands and a top region with out these bands are a consequence of these repeated steps. These parallel bands are not related with the liquidus lines, as suggested by (Mok et al., 2008b), but with the ˇ transus lines of the subsequent ALM steps (Baufeld et al., 2009, 2010a; Kelly and Kampe, 2004). The top region represents the area heated above the ˇ transus during the last ALM step (Baufeld et al., 2009, 2010a; Kelly and Kampe, 2004). The parallel bands derive from the ˇ transus lines of two subsequent ALM steps. This implies that the top region has experienced only one cooling step through the ˛ + ˇ phase field. As a consequence the microstructure in the top region is very fine (Figs. 2a, b and 5a, b). The top region of laser beam ALM components is dominated by colonies starting at grain boundaries, whereas the top region of SMD component features more a mixture of basket weave and colony microstructure. Based on these different morphologies one can attribute the microstructure in the top region of laser beam ALM to relatively high cooling rates compared to SMD, since colonies decorating grain boundaries are reported to be typical for cooling rates between 20 and 410 ◦ C/s (Ahmed and Rack, 1998). Further investigations are necessary to attribute these colonies to massive transformation or to cooling rate dependent compositions. The SMD component exhibits a mixture of basket weave and colonies indicating lower cooling rates close to thermal equilibrium. In the bottom region, the consecutive ALM steps lead to periodic heat treatment in the ˛ + ˇ phase field allowing further diffusion and elemental segregation of Al into the ˛ lamellae and V into the ˇ phase. This generally leads to coarsening of the lamellae (Figs. 2c, d and 5c, d) in comparison to the top region which has experienced only one cooling down through the ˛ + ˇ phase field. The sequential heating and cooling may also lead to the formation of secondary alpha lamellae (Fig. 6). This can be explained by the following: during heating and cooling, the phase fraction as well as the vanadium (V) and aluminum (Al) concentration of the phases change. With increasing temperature the ˛-fraction decreases and the ˇ-fractions increases (Boyer et al., 1994). The Alfraction in ˇ-phase increases (Al is ˛-stabilizer) and the V-fraction (V is ˇ-stabilizer) considerably decreases with increasing temperature (Kellerer, 1970). The relationship between ˛ and ˇ-fraction, Al and V concentration, microstructure and temperature are complex. The ˛ phase that forms during cooling through ˇ transus is called primary ˛ (˛P ) and the ˛ phase that forms after heating and cooling within the ˛ + ˇ phase field is called secondary ˛ (˛S ) (Kelly, 2004). Fig. 16 is based on data from Kellerer (1970) and shows exemplarily a sketched equilibrium Widmannstätten microstructure and the related ˛ and ˇ fraction, and the Al and V compositions. At 600 ◦ C the ˛ phase covering 94.8% of the material is assumed to be primary ˛ which represents mainly the microstructure in the last ALM layer deposited (Fig. 16a). Upon heating in the ˛–ˇ field e.g. due to a next layer deposited, some primary ˛ dissolves into ˇ (Fig. 16b) resulting in a mixture of primary ˛ and ˇ (Fig. 16c). As the ˛ phase dissolves, the total amount of ˇ increases and ˇ is enriched of Al (cAl,ˇ (600 ◦ C + T)) and depleted of V (cV,ˇ (600 ◦ C + T)). The right side of Fig. 16b shows a schematic illustration of the Al and V composition during the dissolution of an ˛ lamella assuming local equilibrium at the ˛/ˇ interface which, however, might not be the case at ALM processes. At 900 ◦ C (Fig. 16c), the (primary) ˛ phase covers only 60% of the material (˛P,900 ◦ C ) containing less V (cV,˛ (900 ◦ C)) and more Al (cAl,˛ (900 ◦ C)) than at 600 ◦ C. During cooling from 900 ◦ C (Fig. 16d)), some of the ˇ decomposes into sec-

1155

ondary ˛ (˛S, 900 ◦ C−T ). At 600 ◦ C (Fig. 16e), the microstructure consists again of 5.2% ˇ and 94.8% ˛, which is now a mixture of secondary ˛ (34.8%) and primary ˛ (60%). Morphologies with grid structures like in Fig. 3a are taken by many authors as evidence for martensites (Boyer et al., 1994; Gil et al., 2001; Lütjering and Williams, 2007). Some authors argue also, that the cooling rates for laser beam ALM are so high that martensitic transformation has to be expected (Brandl et al., 2008, 2009a,b, 2010; Thijs et al., 2010). But due to the great difficulty to discern the ˛ from the ˛ phases, usually the experimental evidence is not conclusive. In the present case, the proof of martensite is also not possible. The grid morphology and high cooling rates would support the existence of martensites. The observation, however, that every ˛ lamella is surrounded by ˇ phase, which is evidently richer in V, suggests diffusion which would oppose the assumption of diffusionless martensitic transformation. Furthermore, the different types of lamellae, i.e. the long lamellae forming the grid and the smaller lamellae in between, have a similar composition, which also seems to contradict the possibility of martensites. A possible solution of this dilemma between morphology and composition could be related with the fact that the cooling rate is not constant but decreasing. After the formation of high cooling rate related phases possibly even for the very top area some diffusion is imaginable resulting in a transformation of the martensite or massive ˛ phase into a ˛/ˇ structure. It was observed, that the laser beam ALM components exhibit near the base plate a decrease in lamellae width (Fig. 7) and an increased amount of smaller prior ˇ grains and martensitic grids. This must be related to the interaction of the plate acting as a cooling reservoir. During ALM the main heat flow goes through the components into the basis plate, and heat flow through the component surface via radiation or convection is much smaller. In other words, the components are cooled by the basis plate. 4.2. Mechanical properties 4.2.1. Static tensile properties The material manufactured by laser beam ALM and SMD can attain static tensile properties of the material specifications AMS 4928 of wrought material and/or ASTM F1108 of cast material. Additive layer manufacturing is, in principle, a mini-casting process characterized by heterogeneous nucleation, and directional and rapid solidification. Hence, its morphology and microstructure is related to castings. Due to the fast solidification, the microstructure of ALM material, however, is generally finer than of castings (Jovanovic et al., 2006). This may be the reason for the relatively high strength of ALM material, reaching values of wrought material. Both different ALM methods exhibit a larger elongation at failure for testing in z- than in x-direction. Similar observations are reported by an other group for laser beam ALM (Mok et al., 2008b) and by the present authors for other deposition parameters for laser beam ALM (Brandl et al., 2009a,b, 2010) and SMD (Baufeld and van der Biest, 2009; Baufeld et al., 2010a,b). Earlier laser beam ALM work claims the opposite behavior (Kobryn and Semiatin, 2001) and lower elongation values, but for these specimens the inferior behavior must be attributed to lack-of-fusion porosity weakening the x–y planes. This orientation dependence of the elongation at failure mainly results from the anisotropic property of the deposition (Mok et al., 2008b). Testing in x-direction means that specimens oriented almost perpendicular to the elongated prior ˇ grains. This leads to a higher amount of grain boundaries for these specimens than for specimens tested in z-direction. It was suggested, that this larger amount of grain boundaries, acting as a potential sources of failure, is the reason of the anisotropic behavior (Baufeld and van der Biest, 2009). This may also be the reason for lower elongation at

1156

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

Fig. 16. Schematic illustration (left) of the diffusion controlled ˇ- and ˛-growth during heating (˛P ↓, ˇ ↑) and cooling (˛p + ˇ → ˛p + ˛S + ˇ) in the ˛–ˇ-field. The graphs on the right side, based on data from (Kellerer, 1970), show the concentration of V (broken line) and Al (full line) along the trajects indicated on the left.

failure values for the laser beam ALM than for the SMD specimens, since laser beam ALM components exhibit smaller prior ˇ grains and therefore more grain boundaries. The three tensile laser beam ALM specimens from the top region have failed very early. Similar inferior elongation at failure values are reported for tensile SMD specimens from the top region (Baufeld and van der Biest, 2009).

The heat treatment at 600 ◦ C does not change the elongation at failure, while the heat treatment at 843 ◦ C increases the strain at failure. Similarly, the UTS and hardness was not influenced significantly by both of these heat treatments. This agrees with observations on SMD components with different deposition parameters than in the present work, also heat treated at 600 ◦ C (Baufeld et al., 2010b). In the case of laser beam ALM for some deposition

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

parameters an increase in strain at failure due to heat treatment at 600 ◦ C, and for other parameters no influence is reported (Brandl et al., 2009a, 2010). Heat treatment at 843 ◦ C resulted into significantly increased strain at failure for testing in x-, but not into an effect for testing in z-direction. Grain growth, stress relief and precipitation hardening by Ti3 Al were suggested for the influence of heat treatment on the mechanical properties (Brandl et al., 2009a, 2010). Future work has to clarify the active processes and explain why they do affect the properties of some of the ALM components and of others not. The UTS of laser beam ALM and SMD specimens from regions outside of the top regions are very similar. The same holds for the Vickers hardness. However, the top region of the laser beam ALM exhibit higher hardness and specimens from this region higher UTS. It is interesting to note, that in the case of SMD, despite different lamellae widths in the top and bottom region, the hardness is not affected. Apparently, the lamellae width is not responsible for the different mechanical performance. Most probably, the higher cooling rates in the laser beam ALM components have lead to microstructures which have higher strength and hardness than in the case of lower cooling rates. As discussed, possible high cooling rate microstructures can be martensites, massive ˛ lamellae, and ˛ lamellae with nonequilibrium composition. These microstructures may be harder and have higher strength due to solid solution hardening or dislocation hardening. 4.2.2. High cycle fatigue properties Despite the large scatter of the fatigue data one can conclude that the fatigue properties are much better than required for cast parts and similar or better than required for wrought parts. The observation of very shallow SN curves for Ti–6Al–4V is in accordance with literature. Marmi et al. (2009), for example, has modeled for R = 0.1 a SN curve between 1000 and 500 MPa for cycles to failure between1 104 and 107 , with experimental data exhibiting a wide scatter within this range, supporting only roughly the model prediction. In the previous chapter it was discussed that the finer microstructure is responsible for the improved strength and ductility of ALM components. All microstructural parameters that increase yield strength and/or reduce slip length should also improve HCF strength (Welsch et al., 1998) depending primarily on resistance to dislocation motion (Lütjering and Williams, 2003; Williams and Lütjering, 1981). According to (Williams and Lütjering, 1981), the most critical parameter in increasing HCF strength is the reduction of the maximum dislocation slip length in the microstructure. In (Nalla et al., 2002), lamellar microstructures showed superior smooth-bar S–N cycle fatigue properties, from ∼105 to >108 cycles at load ratios of R = 0.1 and R = 0.5, than bimodal microstructures. This is explained by its approximately 10% higher tensile strength and coarser microstructural dimensions. The coarse lamellar microstructure limits the number of colonies to be sampled for weak orientations (Nalla et al., 2002). Overall, small slip length is considered a main reason for why the HCF strength of the laser beam ALM and SMD material is comparable to that of wrought material, with high fatigue limits above 700 MPa. Furthermore, the small size of the fatigue samples may contribute to their higher fatigue limit (size effect) (Dowling, 2007). One should, however, also be aware of the extensive smooth-bar fatigue data for different alloy conditions reported in (Sparks and Long, 1974), which are summarized in (Williams and Starke, 1984): The correlation between fatigue limit and tensile strength is smaller than 0.1, indicating that essentially no correlation exists. This points out that the effects of microstructure and strength can be offsetting factors meaning that no change in fatigue performance may

1157

be observed even when strength is increased (Williams and Starke, 1984). 5. Conclusions Wire based laser beam ALM and SMD both successfully fabricate large, near net shape dense components with large ultimate tensile strength values between 900 and 1000 MPa. Both techniques can be applied for rapid prototyping as well as for small lot size fabrication. The big advantage of wire based techniques compared to powder based techniques are the high deposition rates of currently up to 0.7 kg/h, the large component sizes, less contamination and lower material costs. Both ALM techniques show elongated prior ˇ-grains, and layered surfaces, and banded meso-structures. This is in agreement with the results from other ALM groups and techniques. The top region exhibits finer, lamellar ˛/ˇ structures than the bottom region. In the case of laser beam ALM, a hardness gradients within the top region are observed, which is not the case for SMD. This correlates with a higher tensile strength for laser beam ALM in this top region. The strain at failure is in both cases higher for the z than for x direction, and the strain at failure is higher for SMD than for laser beam ALM. The influence of orientation, and in the case of laser beam ALM of location, on the strain of failure must be considered when designing and building for example aerospace components. The direction, where the most ductility is necessary, should be preferentially built with oriented in z direction, and the top region, usually about 10 mm, should be removed to have similar properties along the component. In the present investigation the heat treatment at 600 ◦ C does not have a significant influence on the mechanical properties. The heat treatment at 843 ◦ C, however, increases the strain at failure significantly. The mechanical properties are competitive to the ones of cast material, after the heat treatment at 843 ◦ C the SMD components even may compete with wrought material. ALM components fabricated by both methods exhibit a high fatigue limit of more 770 MPa. These excellent mechanical properties, concurrent with the economical and ecological benefits of ALM, certainly will raise the interest of industry for these techniques. Acknowledgments The SMD research was performed within the RAPOLAC STREP project under contract number 030953 of the 6th Framework Programme of the European Commission (www.RAPOLAC.eu), which is gratefully acknowledged. The support of Dr. Rosemary Gault and her team at AMRC, Sheffield, United Kingdom, where the components have been built is highly appreciated. The activities at EADS Innovation Works were especially supported by Frank Palm, Achim Schoberth and Claudio Dalle Donne as well as by Christoph Leyens (Technical University of Dresden). References Ahmed, T., Rack, H.J., 1998. Phase transformations during cooling in ˛ + ˇ titanium alloys. Mater. Sci. Eng. A 243, 206–211. Banerjee, R., Collins, P.C., Bhattacharyya, D., Banerjee, S., Fraser, H.L., 2003. Microstructural evolution in laser deposited compositionally graded [alpha]/[beta] titanium–vanadium alloys. Acta Mater. 51, 3277–3292. Baufeld, B., van der Biest, O., 2009. Mechanical properties of Ti–6Al–4V specimens produced by Shaped Metal Deposition. Sci. Tech. Adv. Mater. 10, 10. Baufeld, B., Van der Biest, O., Gault, R., 2009. Microstructure of Ti–6Al–4V specimens produced by Shaped Metal Deposition. Int. J. Mater. Res. 100, 1536–1542. Baufeld, B., van der Biest, O., Gault, R., 2010a. Additive manufacturing of Ti–6Al–4V components by Shaped Metal Deposition: microstructure and mechanical properties. Mater. Des. 31, S106–S111. Baufeld, B., van der Biest, O., Gault, R., Ridgway, K., 2010b. Manufacturing Ti–6Al–4V Components by Shaped Metal Deposition: Microstructure and Mechanical Prop-

1158

B. Baufeld et al. / Journal of Materials Processing Technology 211 (2011) 1146–1158

erties, TRAM 2009. IOP Conference Series: Materials Science and Engineering Sheffield, UK. Bergmann, J.P., 2004. Laserstrahlschweißen von Titanwerkstoffen unter Berücksichtigung des Einflusses des Sauerstoffes. Materialwissenschaft und Werkstofftechnik 35, 543–556. Boyer, R., Welsch, G., Collings, E.W., 1994. Materials Properties Handbook: Titanium Alloys. The Materials Information Society. Brandl, E., Baufeld, B., Leyens, C., Gault, R., 2010. Additive manufactured Ti–6Al–4V using welding wire: comparison of laser and arc beam deposition and evaluation with respect to aerospace material specifications. Phys. Procedia 5, 595–606. Brandl, E., Leyens, C., Dalle-Donne, C., Holzinger, V., 2009a. Deposition of Ti–6Al–4V using Nd:YAG laser & wire: microstructure and mechanical properties. In: NATO AVT-163 Specialists Meeting on Additive Technology for Repair of Military Hardware , NATO/PFP, Bonn, pp. 1–13. Brandl, E., Leyens, C., Palm, F., 2009b. Mechanical Properties of Additive Manufactured Ti–6Al–4V Using Wire and Powder based Processes. Trends in Aerospace Manufacturing (TRAM), Sheffield, UK. Brandl, E., Leyens, C., Palm, F., Schoberth, A., Onteniente, P., 2008. Wire Inste Ad of Powder? Properties of Additive Manufactured Ti–6Al–4V for Aerospace Applications. Euro-uRapid, Fraunhofer-Allianz, Berlin, Germany. DeutscheNorm, 1978. Werkstoffprüfung; Dauerschwingversuch, Begriffe, Zeichen, Durchführung, Auswertung. Deutsches Institut für Normung DIN 50100. DeutscheNorm, 1990. Wämebehandlung von Titan-Knetlegierungen. Deutsches Institut für Normung DIN 65084. DeutscheNorm, 2001. Metallische Werkstoffe – Zugversuch – Teil 1: Prüfverfahren bei Raumtemperatur, DIN EN 10002-1. Deutsches Institut für Normung. Dowling, N.E., 2007. Mechanical Behavior of Materials, third ed. Prentice Hall, Blacksburg Virginia. Gil, F.J., Ginebra, M.P., Manero, J.M., Planell, J.A., 2001. Formation of [alpha]Widmanstatten structure: effects of grain size and cooling rate on the Widmanstatten morphologies and on the mechanical properties in Ti6Al4V alloy. J. Alloys Compd. 329, 142–152. Jovanovic, M.T., Tadic, S., Zec, S., Miskovic, Z., Bobic, I., 2006. The effect of annealing temperatures and cooling rates on microstructure and mechanical properties of investment cast Ti–6Al–4V alloy. Mater. Des. 27, 192–199. Kellerer, 1970. Übersicht über die Wärmebehandlung von TiAl6V4. Härterei-Techn. Mitt. 25, 242–253. Kelly, S.M., 2004. Thermal and Microstructure Modeling of Metal Deposition Processes with Application to Ti–6Al–4V. Virginia Polytechnic Institute and State University, Blacksburg, p. 297.

Kelly, S.M., Kampe, S.L., 2004. Microstructural evolution in laser-deposited multilayer Ti–6Al–4V builds. Part II. Thermal modeling. Metall. Mater. Trans. A 35A, 1869–1879. Kobryn, P.A., Semiatin, S.L., 2001. Mechanical properties of laser-deposited Ti–6Al–4V. In: Bourell, D.L., Beaman, J.J., Crawford, R.H., Marcus, H.L., Wood, L., Barlow, J.W. (Eds.), Solid Freeform Fabrication. The University of Texas at Austin, Austin, TX, USA. Lütjering, G., Williams, J.C., 2003. Titanium. Springer, Berlin, Heidelberg, New York. Lütjering, G., Williams, J.C., 2007. Titanium, 2nd ed. Springer, Berlin. Marmi, A.K., Habraken, A.M., Duchene, L., 2009. Multiaxial fatigue damage modelling at macro scale of Ti–6Al–4V alloy. Int. J. Fatigue 31, 2031– 2040. Mok, S.H., Bi, G., Folkes, J., Pashby, I., 2008a. Deposition of Ti–6Al–4V using a high power diode laser and wire. Part I. Investigation on the process characteristics. Surf. Coat. Technol. 202, 3933–3939. Mok, S.H., Bi, G., Folkes, J., Pashby, I., Segal, J., 2008b. Deposition of Ti–6Al–4V using a high power diode laser and wire. Part II. Investigation on the mechanical properties. Surf. Coat. Technol. 202, 4613–4619. Nalla, R.K., Boyce, B.L., Campbell, J.P., Peters, J.O., Ritchie, R.O., 2002. Influence of microstructure on high-cycle fatigue of Ti–6Al–4V: bimodal vs. lamellar structures. Metall. Mater. Trans., 899–918. Peters, M., Leyens, C., 2003. Titanium and Titanium Alloys. Wiley-VCH, Weinheim. SAE Aerospace, S., 2003. Heat treatment of titanium and titanium alloys. Aerospace Material Specification, AMS-H-81200A. Sparks, R.B., Long, J.R., 1974. Improved manufacturing methods for producing high integrity more reliable titanium forgings. AFML-TR-73301. Thijs, L., Verhaeghe, F., Craeghs, T., Humbeeck, J.V., Kruth, J.-P., 2010. A study of the microstructural evolution during selective laser melting of Ti–6Al–4V. Acta Mater. 58, 3303–3312. Vanderesse, N., Maire, E., Darrieulat, M., Montheillet, F., Moreaud, M., Jeulin, D., 2008. Three-dimensional microtomographic study of Widmanstatten microstructures in an alpha/beta titanium alloy. Scripta Mater. 58, 512–515. Welsch, G., Boyer, R., Collings, E.W., 1998. Materials Properties Handbook: Titanium Alloys, 2nd ed. ASM International, Materials Park, Ohio. Williams, J.C., Starke jr., E.A., 1984. Deformation, Processing, and Structure. American Society of Metals, Metals Park, OH. Williams, J.C., Lütjering, G., 1981. Titanium’80 Science and Technology TMS-AIME.