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Microstructure and properties of Ti-6Al-4V fabricated by low-power pulsed laser directed energy deposition
Journal of Materials Science & Technology 35 (2019) 2027–2037
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Research Article
Microstructure and properties of Ti-6Al-4V fabricated by low-power pulsed laser directed energy deposition Hua Tan a,b,∗ , Mengle Guo a , Adam T. Clare c,d , Xin Lin a,b,∗ , Jing Chen a,b , Weidong Huang a a
State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an, 710072, China Key Laboratory of Metal High Performance Additive Manufacturing and Innovative Design, MIIT China, Northwestern Polytechnical University, Xi’an, 710072, China c Institute for Advanced Manufacturing, Faculty of Engineering, University Of Nottingham, NG7 2RD, UK d Department of Mechanical, Materials and Manufacturing Engineering, Faculty of Science and Engineering, University of Nottingham China, 199 Taikang East Road, University Park, Ningbo, 315100, China b
a r t i c l e
i n f o
Article history: Received 13 January 2019 Received in revised form 7 March 2019 Accepted 25 April 2019 Available online 9 May 2019 Keywords: Directed energy deposition Ti-6Al-4V Microstructure Thermal cycle history Mechanical properties
a b s t r a c t Thin-wall structures of Ti-6Al-4V were fabricated by low-power pulsed laser directed energy deposition. During deposition, consistent with prior reports, columnar grains were observed which grew from the bottom toward the top of melt pool tail. This resulted in a microstructure mainly composed of long and thin prior epitaxial  columnar grains (average width ≈200 m). A periodic pattern in epitaxial growth of grains was observed, which was shown to depend upon laser traverse direction. Utilizing this, a novel means was proposed to determine accurately the fusion boundary of each deposited layer by inspection of the periodic wave patterns. As a result it was applied to investigate the influence of thermal cycling on microstructure evolution. Results showed that acicular martensite, ␣’ phase, and a small amount of Widmanstätten, ␣ laths, gradually converted to elongated acicular ␣ and a large fraction of Widmanstätten ␣ laths under layer-wise thermal cycling. Tensile tests showed that the yield strength, ultimate tensile strength and elongation of Ti-6Al-4V thin wall in the build direction were 9.1%, 17.3% and 42% higher respectively than those typically observed in forged solids of the same alloy. It also showed ˜ the yield strength and ultimate tensile strength of the transverse tensile samples both were 13.3% higher than those from the build direction due to the strengthening effect of a large number of vertical  grain boundaries, but the elongation was 69.7% lower than that of the build direction due to the uneven grain deformation of  grains. © 2019 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
1. Introduction Directed energy deposition (DED) is an Additive Manufacturing (AM) technology for the fabrication and repair of high value components [1,2]. It has been shown to significantly reduce the lead time between initial concept and finished part by eliminating several intermediate steps [3]. Moreover, it allows the production of graded materials and complex components [4,5]. Titanium alloys, especially Ti-6Al-4V, are widely used as structural components for aircraft applications due to their high specific strength and excellent corrosion resistance. DED for this material has focused on process control, microstructure and properties of fabricated titanium alloy structures [6–9]. Understanding the depo-
∗ Corresponding authors at: State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an, 710072, China. E-mail addresses: [email protected] (H. Tan), [email protected] (X. Lin).
sition mechanism of titanium alloys in thin walls is an important technique which simulates the production of high aspect structures often used in aerospace applications such as blanks for brackets. Contributions to the literature in the field of Ti-6Al-4V by DED to date are broad. Gharbi et al. [10] investigated the influence of the process parameters on the surface finish of Ti-6Al-4V multi-layer structure, and proposed solutions for improving surface finish. Kelly and Kampel [11,12] deposited thin wall structures using an 11 kW CO2 laser and investigated the microstructural evolution in multilayer Ti-6Al-4V deposited structures. Further, the mechanism of microstructure formation based on a thermal modeling analysis is reported here. Qian et al. [13] built a finite element model to predict the thermal history of direct laser fabricated Ti6Al-4V thin wall samples. This was correlated with microstructural evaluation in-process. These works have all contributed to understanding thermal history and the resulting microstructure in DED titanium alloys particularly in thin wall structures. In these studies, the energy source used was usually a continuous wave type laser
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and the laser power was usually over 500 W, which resulted in a relatively high thermal accumulation through build duration. Some researchers have studied the DED of pulsed lasers. Pinkerton et al. [14] deposited 316 L stainless steel part with a pulsed laser beam and found that the porosity and surface roughness reduced when using a pulsed beam rather than using a continuous beam, and the hardness of the deposited part increased. Banday et al. [15] prepared Ti/MoS2 coating on Al-Si alloy by pulsed laser deposition to invesigate the nanoscratch resistance and nanotribological performance of the coating. Gharbi et al. [16] demonstrated that using quasi-continuous instead of fully continuous laser irradiations could add further improvement to the surface finish of Ti-6Al-4V parts produced by DED. Ahsan et al. [17] fabricated Ti6Al-4V porous structures by pulsed laser metal deposition with high power, and focused on the formation of pores. In fact, lowpower pulsed laser DED can be proposed as a method to precisely fabricate and repair thin wall with its much lower average energy input. In this study, a low-power pulsed laser is used to deposit Ti6Al-4V alloy and can provide high cooling rate (<103 K/s [18]) while causing sufficient melting for a high integrity deposit to form. The resulting microstructure and mechanical properties are worthy of further investigation. Therefore, it is necessary to comprehensively study the microstructural evolution and mechanical properties in DED utilizing a low-power pulsed laser source.
mens were prepared by standard mechanical polishing method, and etched with a solution of HF: HNO3 : H2 O with a ratio of 1: 3: 20. Macro- and Micro-structure of the alloys were examined by an OLYMPUS optical microscope (OM) and a Hitachi S-4800 scanning electron microscope (SEM). Quantitative measurements of microstructure involved in this study were conducted using ImagePro Plus software.
2. Experimental procedure
3.1. Morphology of prior ˇ grains
2.1. Low-power pulse laser DED
Fig. 3(a)-(c) shows the typical microstructures of the bottom, middle and top of Ti-6Al-4V thin wall respectively. These were extracted from blue dashed boxes in Fig. 2(a). It can be seen that the microstructure of the thin-wall specimen is composed of elongated epitaxial  columnar grains, which are almost perpendicular to the substrate surface and clearly show the growth is driven by conductive heat flux toward the substrate. From Fig. 3(a), at the bottom of thin-wall sample, it can be seen that the width of prior  columnar grains is smaller at proximity to the substrate, and as the Z distance increases the width of the  columnar grains increases. When Z ≈ 1 mm, the width of the  columnar grains reaches a steady state with an average value of about 200 m. From Fig. 3(b) and (c), it can also be seen that the average width of the  columnar grains almost remains constant at the middle and top of thin-wall sample. Fig. 3(d) shows a local enlargement of the area marked by the red box in Fig. 3(c). The grain boundaries of the epitaxial  grains exhibit clear periodic wave patterns, and this is a notable feature in this figure. The periodic wavelength is about 0.2 mm, and just is the thickness of two layers. The morphology of prior  grains is very different from that in continuous laser DED. Usually, the microstructure of Ti-6Al-4V fabricated by conventional continuous laser DED is composed of coarse epitaxial  columnar grains or mixed structure of columnar grains and equiaxed grains, and the columnar grains grow straight along built direction with less obvious grain steering [7,22].
Thin wall (single track thickness) deposition of Ti-6Al-4V was performed using a low-power pulsed laser DED system as illustrated in Fig. 1(a). This system consists of a 300 W HANS Nd: YAG pulsed laser, a three-axis CNC unit to control the movement of X-Y table and control the Z-axis (build direction), a glove box which contained an argon atmosphere (oxygen content below 50 ppm), a powder feeder and a coaxial nozzle conveying powder to the deposition point by a pressurized carrier gas (argon). The powder material for the experiment was Ti-6Al-4V with particle size ranging from 200 to 300 mesh powder particle and chemical composition (wt%) as follows: 6.17 Al, 4.14 V, 0.15 O, 0.02 H, 0.005 N, 0.016 Si, 0.006 C, 0.032 Fe and Ti balance. The powder was supplied by YuGuang Phelly Ltd (China). Ti-6Al-4V plate (140 mm × 50 mm × 6 mm) was used as a substrate, which was burnished and cleaned before deposition. The powder was dried in a vacuum oven at 120 ◦ C for 2 h to remove moisture. Thin-walls were deposited with the following process parameters: laser pulse width 6 ms, pulse frequency 20 Hz, current intensity 40 A, laser spot diameter 1 mm, laser scanning speed 4 mm/s, powder feeding rate 1 g/min, carrier gas flow rate 200 L/h, and layer thickness 0.1 mm. The process parameters were chosen based on previous researches to ensure the metallurgical quality. Liu [19] and Hou [20] in our research group had used these process parameters for pulsed laser DED experiments to fabricate Ti-6Al-4V parts and got good results. In this way a thin-wall sample was built using single-pass and multi-layer laser deposition using a ‘back and forth’ strategy. Thus deposition directions of the adjacent tracks are opposing, as shown in Fig. 1(b). Finally, a thin wall with a geometric size of 75 mm × 35 mm × 1.2 mm was fabricated, as shown in Fig. 1(c), and was directly aged at 560 ◦ C for 2 h in order to relieve residual stresses. 2.2. Microstructural characterization The bottom, middle and top zones in cross-section XZ (which are marked respectively by blue dashed boxes in Fig. 2(a)), were extracted for microstructure observation. Metallographic speci-
2.3. Mechanical testing Hardness was measured by a Vickers tester (Struers DuraminA300) using a load of 0.3 kg and loading time 15 s. Hardness values were averaged over three measurements. In this study, nonstandard tensile specimens were used to investigate the tensile properties in the deposition height direction and the laser traverse direction. The extraction of specimens is illustrated in Fig. 2(a). The dimensions of tensile specimens and actual specimens are shown in Fig. 2(b) and (c). Tensile tests were carried out using an Instron 3382 universal material testing machine, and stress-strain data were recorded in real time using a Digital Image Correlation System (DIC) (XJTUDIC). Tensile properties were calculated from the results of DIC based on a method reported elsewhere [21]. 3. Results and discussion
3.2. Formation of ˇ grains During the initial deposition to the substrate, the  grains epitaxially grow from the fine equiaxed grains of the substrate, which is shown in Fig. 3(a). Due to proximity to the substrate and high conductivity the cooling rate is also greater, so the width of prior  columnar grains is relatively small. Upon multi-layer deposition toward a thin-wall structure, the temperature of the succeeding deposited layer would increase due to reduced heat sinkage, resulting in a significant decrease of the temperature gradient, thus a rapid increase of grain size. However, the energy delivered by the low power pulsed laser deposition is significantly less than the reported laser deposition process (usually a laser power
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Fig. 1. Illustration of pulsed laser deposition experiment, (a) schematic diagram of system, (b) schematic diagram of thin-wall laser deposition process, (c) Ti-6Al-4V thin-wall sample deposited.
Fig. 2. Preparation of the samples for metallographic observation and tensile test, (a) schematic illustration of specimen extractions, (b) dimensions of the test specimen, (c) tensile test samples arranged as built.
of 500–8000 W [7,11,12,17,21–23]), so the process would rapidly reach a thermal steady-state, which means that the temperature gradient of the melt pool zone would remain constant. Therefore, the average width of the  columnar grains will remain constant. During the solidification process, the high temperature gradient, G, and low solidification velocity, Vs , will tend to result in colum-
nar grain growth, while low G and high Vs will trend to result in equiaxed growth [24]. Fig. 4(a) shows the schematic relationship between scanning velocity, Vb , solidification velocity, Vs , and the melt pool boundary. It can be seen that Vs is normal to the melt pool boundary, and the relationship between Vb and Vs is defined by Vs = Vb ·cos [25], where is the angle between the vectors of Vb and
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Fig. 3. Morphologies of prior  grains of Ti-6Al-4V thin wall, (a) the bottom, (b) the middle, (c) the top and (d) local enlargement of typical morphology.
Fig. 4. Solidification analysis of the melt pool, (a) schematic relationship among scanning velocity Vb , solidification velocity Vs and melt pool boundary, (b) the calculated CET curve (calculated from Lin’s model [28]) and schematic solidification condition at the melt pool tail.
Vs . Therefore, from the bottom to top of the melt pool, Vs gradually increases. According to literature [26], G always decreases from the bottom to the top of the melt pool. Fig. 4(b) shows the critical condition curve of columnar to equiaxed transition (CET) of Ti-6Al-4V. This is calculated from Lin’s model [27]. The red dashed line indicates the schematic solidification condition at the melt pool tail. In this case, the pulsed laser results in a high temperature gradient, which would promote epitaxial columnar grain growth. It can be seen that the solidification condition is in the columnar grain growth zone at all times, which indicates that CET would not occur at the solid-liquid interface during the solidification of the molten pool. It is different from the usual C/W mode laser DED, and the  grains usually present mixed growth of columnar and equiaxed forms, as shown in previous studies [9,28]. Further, it can also be seen in Fig. 3(c) that there are no equiaxed grains at the top of the sample, indicating that solidification process is dominated by the columnar grain growth from the bottom to the top of the molten pool tail. Therefore, the solidification microstructure of the thinwall sample presents prior  columnar grains with highly epitaxial growth.
The formation mechanism of the periodic wave boundary of epitaxial  grains is also worthy of further analyses. According to Han et al. [29], in moving the melt pool, the columnar grains grow from the fusion line with the planar solid-liquid interface, and then curve gradually to follow the direction of the highest temperature gradient (normal to the local tailing pool boundary). Fig. 5(a) shows simulated results of typical thermal behavior in laser deposition of thin wall using a finite element model developed by Bontha et al. [30]. It can be seen that the direction of local highest temperature gradient Gm is always normal to local tailing pool boundary, and always diverts to the laser scanning direction. Therefore, as shown in Fig. 5(b), when the layer i is deposited, the melt pool moves to the right and Gm at the melt pool tail leans to the right, which causes the columnar grains to curve gradually to the right. This is clearly also evident in the experimental results presented here. Also, when the layer i+1 is deposited, the melt pool moves to the left, and Gm at the melt pool tail leans to the left, causing the columnar grains to curve gradually to the left. When layer i+2 is deposited, it mimics the situation reported in the layer i. Therefore, with the laser scanning back and forth, the direction of Gm at the melt pool tail
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Fig. 5. Mechanism of the wavy epitaxial growth of  grain, (a) temperature fields simulation by FEM with laser scanning back and forth, (b) schematic of the wavy growth.
laser deposition. In addition, from Fig. 3(c), a microstructure transitional zone of 500 m (the thickness of 5 layers of deposition) can be observed at the top of thin-wall sample, which exhibits a different microstructure from that of a stable typical zone.
Fig. 6. Schematics of melt pools with different depth-to-width ratios.
periodically changes, resulting in the periodic deflection of epitaxial grain growth. However, for conventional continuous laser DED, the deflections of epitaxial grains are not as obvious. This is because the peak energy density of pulsed laser may be higher than that of the continuous laser, resulting in a greater depth-to-width ratio of the melt pool, indicating that the tail boundary of melt pool caused by pulsed laser is steeper, as observed by Assuncao et al. [31], who reported this phenomenon when comparing the melt pools caused by continuous and pulsed lasers. Fig. 6 shows the schematics of the melt pools with different depth-to-width ratios. It can be found that the angle that Gm leans toward the laser scanning direction increases significantly at the tail of melt pool with greater depth-towidth ratio, resulting in an increase of the deflection of the epitaxial grain. As a result of this, the unusual and interesting periodic wave epitaxial grains (shown in Fig. 3(d)) were formed during the thinwall deposition with pulsed laser scanning back and forth. 3.3. Microstructure in prior ˇ grains 3.3.1. Microstructural overview From Fig. 3(d),  grains in different gray levels can be observed in the OM photograph. Fig. 7(a) shows the typical microstructure in “bright”  grain (marked in Fig. 3(d)). It can be observed in Fig. 7(a) that the typical microstructure in the region is mainly composed of many ␣ laths and a small amount of acicular ␣ phase, with a width of 1˜ m and a length between 10–30 m. Fig. 7(b) shows the typical microstructure in “dark”  grain (marked in Fig. 3(d)). It can be seen that the typical microstructure in the region contains a large number of elongated acicular ␣ phases, with a length of over 100 m and a width of 1˜ m, and a large amount of lath ␣ phase with a width of only 0.3 m and a length of 5˜ m. The difference of phase between “bright” and “dark” grains may be caused due to different crystal orientations of prior  grain. The “bright” and “dark”  grains both contain a large number of acicular ␣ phases, indicating that the cooling rate is high during low power pulsed
3.3.2. Microstructural evolution It is well known that this microstructure transitional region at the top of a build is formed from the last several thermal cycles, and can be used to analyze the evolution of the microstructure. As shown in Fig. 8(a), during thin wall deposition, the last layer n undergoes a thermal cycle, the layer n-1 undergoes two thermal cycles, and the layer n-2 undergoes three thermal cycles, and so on. Therefore, if each deposited layer at the top of the sample can be clearly distinguished, the influence of thermal cycles on microstructure evolution can be directly observed. However, it is usually difficult to distinguish the fusion line. Fig. 8(b) shows the local amplification of the blue dotted frame area at the top zone in Fig. 3(c). It can be also seen that it is difficult to discern the melting boundary of each layer in the top transitional region. However, based on previous analyses, we can see that each deflection of the periodic wave  grain represents a change of laser scanning direction, and the periodic wavelength is approximately 0.2 mm (just is the thickness of two layers deposition). In other words, the periodic wave pattern of one  grain can be used to accurately determine the fusion boundary of each layer (that is to say the peak and trough of the wave are exactly the fusion line’s positions at each layer). Therefore, the influence of thermal cycling on the layer-by-layer microstructure evolution can be investigated based on the fusion boundary of each deposited layer by evaluation of the periodic wave pattern of the epitaxial  grain. In this study, in order to determine the correct area and then observe the microstructure in each layer under a high magnification scanning electron microscope, some stamps were marked by the Vickers tester at the peak and trough of wave of  grain using a load of 0.1 kg and loading time 11 s, which allows us to determine the final several layers at the top of the sample, as shown in Fig. 8(b). Here, we defined the last layer as Layer n, and those layers before the last layer were defined as follows: layer n-1, n-2, n-3, n-4, and so on. From the marks, the thickness of each layer can be measured, ˜ mm (it agrees with the set value of layer thickness), which and is 0.1 indicates that the deflection analysis of the wave epitaxial grain is accurate and reliable. Then, by microstructure observation in the last several layers, the layer-by-layer influence of thermal cycles on the solid phase transformation can be obtained. Therefore, this study will present a novel and reliable method that can be used
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Fig. 7. Microstructures in  columnar grain, (a) “bright” grain, (b) “dark” grain.
Fig. 8. Analysis of the top of the sample, (a) schematic of deposited layer formation at the top of the sample, (b) microstructure observed at the top of the sample and diagram of deposition layer marking.
to investigate the microstructure evolution under thermal cycling conditions of DED. In order to disclose the influence of thermal cycle on the microstructure evolution, in this study, the red line frames in Fig. 8(b) were selected as the exemplar region in each layer. Fig. 9 showed the typical SEM microstructure images in each layer of the final five layers. As shown in Fig. 9(a) and (b), the last layer zone exhibits elongated and parallel acicular martensitic ␣’ phase with widths of 1–2 m, and a small amount of ␣ phase (lathes). These
with an average width of 0.3 m can be observed especially around the acicular martensitic ␣’ phase. At layer n-1, the acicular martensitic ␣’ phase is still present, but the amount of lath ␣ is observed to increase and enlarge slightly, with an average width of 0.5 m, as shown in Fig. 9(c) and (d). From Fig. 9(e) and (f), it can be seen that the microstructure features change significantly at layer n-2. The acicular ␣’ phase disappears, while many long and thin ␣ laths exhibited in the microstructure, and are approximately perpendicular to parallel acicular martensitic ␣’ phase in Fig. 9(a)-(d), the
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Fig. 9. Microstructures in the same  grain and final five layers of the sample, (a), (b) the microstructures of the last layer, (c), (d) the microstructures of the last second layer, (e), (f) the microstructures of the last third layer, (g), (h) the microstructures of the last fourth layer, and (i), (j) the microstructures of the last fifth layer.
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Fig. 10. Schematic diagram of the thermal cycle history and microstructure formation.
amount of lath ␣ continuing to increase as well. From Fig. 9(g)(j), it can be seen that long and thin ␣ phase becomes gradually shorter and less prevalent, while the size and amount of lath ␣ gradually increase. Overall, it can be clearly seen that from the final layer to the layer n-4 (as thermal cycle increases), the lath ␣ phase uniformly increases in size. Moreover, the microstructure in the  grain at layer n-4 (suffering five thermal cycles) is basically similar to typical microstructures (see Fig. 7(b)) in the same  grain, indicating that the microstructure of the deposition layer tends to be stable after five thermal cycles. 3.3.3. Microstructure evolution mechanism The microstructure evolution in dark  grains is determined via thermal cycle history conditions. In order to make the discussion representative, thermal cycle histories of each layer during steady deposition were used in this study. For ease of explanation, three important temperatures of Ti-6Al-4V were noted. The TL and T are liquidus and  transus lines, respectively. Tdiss stands for ␣-dissolution temperature, where ␣ phase starts to dissolve into  under equilibrium heating conditions. When the temperature is below Tdiss , the microstructure in  grain will not change. Fig. 10 shows the schematic diagram of the thermal cycle history and microstructure formation. The thermal cycle time that the layer n (final layer) underwent is tn , the thermal cycle time that the layer n-1 underwent is tn-1 , and the thermal cycle time that the layer n-2 underwent is tn-2 , and so on. It is important to recall that cooling rate of the thermal cycle gradually decreases in a laser multi-layer deposition due to reducing proximity to the heat source. During the deposition of layer n, its temperature sharply decreases, from above TL to T , due to the very high cooling rate caused by the pulsed laser, the martensitic ␣’ phase would be formed at T , and with decreasing temperature, small amounts of Widmanstätten ␣ lath are formed (see the observation results of layer n in Fig. 9(a) and (b)). Layer n-1 underwent two thermal cycles within tn-1 . The first thermal cycle is the same as that of layer n, and with the second thermal cycle, when the temperature rapidly rises to above T , the acicular ␣’ phase and ␣ lath formed under the first thermal cycle would be completely dissolved. When the temperature decreases rapidly to T , some new martensitic ␣’ phase will be formed again due to the high cooling rate, and then, with a decrease of temperature at a relatively lower cooling rate than that of layer n, some Widmanstätten ␣ laths (more than those of layer n) will be formed as well (see the observation results of layer n-1 in Fig. 9(c) and (d)). Layer n-2 underwent three thermal cycles within tn-2 , with the third thermal cycle, when temperature increases to above T , the acicular ␣’ phase and ␣ lath formed under the second thermal cycle will be completely dissolved again. When the temperature decreases to T , the new martensitic ␣’ phase was not formed again, which is because the cooling rate is lower than that of ␣’ formation.
However, the cooling rate remains relatively high, so some elongated acicular ␣ phase would be formed, and there would be more Widmanstätten ␣ laths than those of layer n-1 formed as well, as shown in Fig. 9(e) and (f). Ahmed et al. [32] made clear that for ␣ +  titanium alloys, when the cooling rate is lower than that of martensitic ␣’ formation but remains relatively high, the elongated acicular ␣ phase will be formed on the grain boundary, and rapidly grow towards the grain interior. So, in this study, a large number of elongated acicular ␣ particle will rapidly grow from the boundary of thin grain and distribute on the whole grain interior. For the layer n-3, similarly, with the fourth thermal cycle, when the temperature rises to between T and Tdiss , the acicular ␣ phase and ␣ lath formed in third thermal cycle would be partially dissolved, so the acicular ␣ phase would get shorter and thinner. As the temperature continues to decline at a lower cooling rate, more Widmanstätten ␣ lath structures will be formed and grow bigger, as shown in Fig. 9(g) and (h). For layer n-4, the acicular ␣ phase would get much shorter and thinner, and ␣ lath becomes more prevalent and enlarged, as shown in Fig. 9(i) and (j). Compared to Fig. 7(b), the microstructure in layer n-4 has been basically similar to the typical microstructures in the same  grain. Furthermore, for layer n-5, due to the peak temperature of the sixth thermal cycle below Tdiss , the microstructure remains unchanged. Therefore, a steady microstructure would be reached in layer n-4 undergoing five thermal cycles under these experimental conditions. In summary, using a pulsed laser, the solidification velocity Vs and the temperature gradient G are larger, so primary dendritic spacing 1 is smaller according to 1 ∝Vs –a G−b (a, b are constants) and fine epitaxial growth columnar grains are formed. Meanwhile, ␣ phase nucleates at the grain boundary and grows through the whole fine columnar grain, and with the thermal cycles the acicular ␣ and Widmanstätten ␣ lath are finally formed, while using conventional continuous laser there exhibits basket weave ␣ lath structures in  grain interior [13,33]. 3.4. Mechanical properties 3.4.1. Hardness Microstructures in  grains exhibit apparent differences. Thus the microhardness of different  grains was further tested every 0.5 mm within the range of 4 mm along the deposition height direction, as shown in Fig. 11(a). The hardness results were averaged over three measurements at the same height, and the results are shown in Fig. 11(b). It can be seen that the “bright” and “dark”  grains ˜ have high hardness values (375-389 HV in the “bright”  grain, and ˜ 398-410 HV in the “dark”  grain). This is because they both contain a large number of elongated acicular ␣ phase due to high cooling rate caused by the low power pulsed laser. Microhardness distribution in the “dark”  grain is higher than that in the “bright”  grain due to greater strengthening by a large number of dense parallel acicular ␣ particles (see Fig. 7(b)). 3.4.2. Tensile testing Fig. 12 shows the tensile properties of thin wall samples. It can be seen that the average yield strength ( y ) and ultimate tensile strength ( b ) of transverse tensile samples reached 1020 MPa and 1190 MPa, respectively, but the average elongation was only 4.3%. The average yield strength and ultimate tensile strength of height direction tensile samples was 900 MPa and 1050 MPa, respectively, and the average elongation was 14.2%, indicating an excellent combination of strength and elongation of height direction. Obviously the strength of the Ti-6Al-4V thin wall samples deposited are higher than that of the forged standard (yield strength 828 MPa and ultimate tensile strength 895 MPa) [34] due to the cooling rates. Surprisingly, the strength of the transverse tensile samples is significantly higher (13.3%) than those of the build direction tensile
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Fig. 11. Micro-hardness test in different  grains, (a) the schematic of test position, (b) hardness results.
Fig. 12. Tensile properties of specimens taken from the thin wall.
samples, but the elongation of the transverse tensile samples is significantly lower (69.7%) than those of the height direction tensile samples. This is because in the case of the transverse tensile test, there is a large amount of  grain boundaries being perpendicular to the tensile direction (see Fig. 13(a)). From the previous research, we know that the average value of the width of the  columnar grains is about 200 m and the effective length of the tensile test specimen is 10 mm, so we can quantify that there are five  grain boundaries per mm. According to plastic deformation theory [35], the grain boundary density serves to impede dislocation movement, resulting in a notable strengthening effect. So the higher strength in the transverse tensile test was exhibited due to the strengthening effect of many vertical grain boundaries. In
addition, during the tensile process, the sample will deform into bamboo joint pattern due to the strengthening of these vertical grains boundaries, as shown in Fig. 13(a). Due to the anisotropy of the grains’ property (see the hardness test results in Fig. 11), the  grains in the test sample could not uniformly deform during transverse tension. So the weakest  grain will first reach the ultimate deformation to break, resulting in the serious decrease of elongation to failure. However, in the case of the build direction tensile test, the amount of  grain boundaries is reduced, and these boundaries lie largely in the tensile direction (as shown in Fig. 13(b)), so the strengthening effect is not as significant, and is lower than that of the transverse direction. During the tensile testing, the thin and long epitaxial grains in the test samples would be able to deform in unison, resulting in excellent elongation to failure. Fig. 14 shows the SEM micrographs of the typical fracture surfaces of the two direction tensile samples. The contraction of the fracture cross-section is obvious (see Fig. 14(a)), and the fracture morphologies of the height direction sample show the typical high fracture toughness characteristics associated with ductile dimples (see Fig. 14(b)), which shows that the fracture mechanism is microvoid accumulation fracture. The tensile fracture surface in the height direction sample is dominated by large and round or equiaxed dimples, with an average diameter of 4˜ m, indicating excellent ductility, which agrees with the elongation measured by the tensile test. From Fig. 14(c), the shrink of the fracture crosssection is not obvious, and there are some shallow and elongated dimples on the fracture surface of the transverse tensile sample, as shown in Fig. 14(d). Furthermore, river patterns and cleavage terraces can also be observed, which shows the fracture morphologies are of cleavage character. In other words, the fracture mechanism of the transverse direction sample is ductile/brittle mixed fracture
Fig. 13. Schematic of tensile process analysis, (a) transverse directoin, (b) build direction.
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Fig. 14. SEM micrographs of the typical fracture surfaces of two direction tensile samples, (a), (b) height direction, (c), (d) transverse direction.
which indicates poor plasticity, agreeing with the elongation measured from the transverse sample. 4. Conclusions A Ti-6Al-4V thin wall was fabricated using low-power pulsed laser DED to elucidate information relating to microstructure formation phenomena. Solidification from the bottom to the top of the melt pool tail is dominated by the columnar grain growth at all times, so the solidification microstructure of the thin-wall sample is composed of long and thin prior  columnar grains, which are almost perpendicular to the laser scanning direction. The boundaries of the epitaxial  grains exhibited clear periodic wave patterns, which can be used to identify the fusion boundary of each layer (the periodicity matches the fusion line positions of each layer). Based on the determination of the fusion line of each layer by the periodic wave pattern of epitaxial  grain at the top of the thin wall, the influence of the thermal cycle on microstructure layer-bylayer evolution was investigated. It was observed that the acicular martensitic ␣’ phase and a small amount of Widmanstätten ␣ lath converted to long and thin acicular ␣ phase and larger amounts of Widmanstätten ␣ lath under thermal cycle conditioning. Elongated acicular ␣ phase was formed on the  grain boundary, and rapidly extended into the grain interior. Therefore, the thin acicular ␣ phase easily distributes in the grain interior. The microstructure difference between  grains will cause the property difference. The strength of Ti-6Al-4V thin wall samples are higher than that of the forged standard due to high cooling rates, and the elongation of build direction is much higher than that of the forging standard, indicating an excellent combination of strength and elongation in the build direction. The strength of the transverse tensile samples is 13.3% higher than that of the build direction samples due to the significant strengthening effect of a large amount of vertical  grain boundaries. However, the elongation of transverse tensile samples is 69.7% lower than those of build direction samples due to the uneven deformation of  grains.
Acknowledgements This work was supported by the National Key Research and Development Program of China (No. 2018YFB1106302), the National Natural Science Foundation of China (Grant No. 51475380) and the Aeronautical Science Foundation of China (Grant No. 2016ZE53). References [1] L.J. Kumar, C.G.K. Nair, Mater. Today: Pro. 4 (2017) 11068–11077. [2] N.K. Dey, F.W. Liou, C. Nedic, International SFF Symposium, 2013, pp. 853–858. [3] J. Mazumder, D. Dutta, N. Kikuchi, A. Ghosh, Opt. Lasers Eng. 34 (2000) 397–414. [4] S. Kamran, H.I. Ul, K. Ashfaq, S.S. Ali, K. Mushtaq, J.P. Andrew, Mater. Des. 54 (2014) 531–538. [5] C.L. Qiu, G.A. Ravi, C. Dance, A. Ranson, S. Dilworth, M.M. Attallah, J. Alloys. Compd. 629 (2015) 351–361. [6] A.R. Nassar, J.S. Keist, E.W. Reutzel, T.J. Spurgeon, Addit. Manuf. 6 (2015) 39–52. [7] Y.M. Ren, X. Lin, X. Fu, H. Tan, J. Chen, W.D. Huang, Acta Mater. 132 (2017) 82–95. [8] C. Qiu, G.A. Ravi, M.M. Attallah, J. Mater. Des. 81 (2015) 21–30. [9] B.E. Carroll, A. Palmer, A.M. Beese, Acta Mater. 87 (2015) 309–320. [10] M. Gharbi, P. Peyrea, C. Gorny, M. Carin, S. Morville, P.L. Masson, D. Carron, R. Fabbro, J. Mater. Process. Technol. 213 (2013) 791–800. [11] S.M. Kelly, S.L. Kampe, Metall. Mater. Trans. A 35 (2004) 1861–1867. [12] S.M. Kelly, S.L. Kampe, Metall. Mater. Trans. A 35 (2004) 1869–1879. [13] L. Qian, J. Mei, J. Liang, X. Wu, Mater. Sci. Technol. 21 (2005) 597–605. [14] A.J. Pinkerton, L. Li, Appl. Surf. Sci. 208-209 (2003) 405–410. [15] S. Banday, M.F. Wani,J. Tribol. 141 (2018), 022003. [16] M. Gharbi, P. Peyre, C. Gorny, M. Carin, S. Morville, P.L. Masson, D. Carron, R. Fabbro, J. Mater. Process. Technol. 214 (2014) 485–495. [17] M.N. Ahsan, C.P. Paul, L.M. Kukreja, A.J. Pinkerton, J. Mater. Process. Technol. 211 (2011) 602–609. [18] D. Herzog, V. Seyda, E. Wycisk, C. Emmelmann, Acta Mater. 117 (2016) 371–392. [19] H. Liu, Electrochemical Behavior and Smoothing Law During Surface Finishing of Laser Solid Formed Ti-6Al-4V Parts, Thesis, Northwestern Polytechnical University, 2016 (in Chinese). [20] W. Hou, The Deformation Behavior of Laser Solid Formed Ti-6Al-4V Thin-wall Components, Thesis, Northwestern Polytechnical University, 2017 (in Chinese). [21] Y.H. Wang, J.H. Jiang, C. Wanintrudal, C. Du, D. Zhou, L.M. Smith, L.X. Yang, Exp. Tech. Urol. Nephrol. 34 (2010) 54–59.
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Research Article
Microstructure and properties of Ti-6Al-4V fabricated by low-power pulsed laser directed energy deposition Hua Tan a,b,∗ , Mengle Guo a , Adam T. Clare c,d , Xin Lin a,b,∗ , Jing Chen a,b , Weidong Huang a a
State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an, 710072, China Key Laboratory of Metal High Performance Additive Manufacturing and Innovative Design, MIIT China, Northwestern Polytechnical University, Xi’an, 710072, China c Institute for Advanced Manufacturing, Faculty of Engineering, University Of Nottingham, NG7 2RD, UK d Department of Mechanical, Materials and Manufacturing Engineering, Faculty of Science and Engineering, University of Nottingham China, 199 Taikang East Road, University Park, Ningbo, 315100, China b
a r t i c l e
i n f o
Article history: Received 13 January 2019 Received in revised form 7 March 2019 Accepted 25 April 2019 Available online 9 May 2019 Keywords: Directed energy deposition Ti-6Al-4V Microstructure Thermal cycle history Mechanical properties
a b s t r a c t Thin-wall structures of Ti-6Al-4V were fabricated by low-power pulsed laser directed energy deposition. During deposition, consistent with prior reports, columnar grains were observed which grew from the bottom toward the top of melt pool tail. This resulted in a microstructure mainly composed of long and thin prior epitaxial  columnar grains (average width ≈200 m). A periodic pattern in epitaxial growth of grains was observed, which was shown to depend upon laser traverse direction. Utilizing this, a novel means was proposed to determine accurately the fusion boundary of each deposited layer by inspection of the periodic wave patterns. As a result it was applied to investigate the influence of thermal cycling on microstructure evolution. Results showed that acicular martensite, ␣’ phase, and a small amount of Widmanstätten, ␣ laths, gradually converted to elongated acicular ␣ and a large fraction of Widmanstätten ␣ laths under layer-wise thermal cycling. Tensile tests showed that the yield strength, ultimate tensile strength and elongation of Ti-6Al-4V thin wall in the build direction were 9.1%, 17.3% and 42% higher respectively than those typically observed in forged solids of the same alloy. It also showed ˜ the yield strength and ultimate tensile strength of the transverse tensile samples both were 13.3% higher than those from the build direction due to the strengthening effect of a large number of vertical  grain boundaries, but the elongation was 69.7% lower than that of the build direction due to the uneven grain deformation of  grains. © 2019 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
1. Introduction Directed energy deposition (DED) is an Additive Manufacturing (AM) technology for the fabrication and repair of high value components [1,2]. It has been shown to significantly reduce the lead time between initial concept and finished part by eliminating several intermediate steps [3]. Moreover, it allows the production of graded materials and complex components [4,5]. Titanium alloys, especially Ti-6Al-4V, are widely used as structural components for aircraft applications due to their high specific strength and excellent corrosion resistance. DED for this material has focused on process control, microstructure and properties of fabricated titanium alloy structures [6–9]. Understanding the depo-
∗ Corresponding authors at: State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an, 710072, China. E-mail addresses: [email protected] (H. Tan), [email protected] (X. Lin).
sition mechanism of titanium alloys in thin walls is an important technique which simulates the production of high aspect structures often used in aerospace applications such as blanks for brackets. Contributions to the literature in the field of Ti-6Al-4V by DED to date are broad. Gharbi et al. [10] investigated the influence of the process parameters on the surface finish of Ti-6Al-4V multi-layer structure, and proposed solutions for improving surface finish. Kelly and Kampel [11,12] deposited thin wall structures using an 11 kW CO2 laser and investigated the microstructural evolution in multilayer Ti-6Al-4V deposited structures. Further, the mechanism of microstructure formation based on a thermal modeling analysis is reported here. Qian et al. [13] built a finite element model to predict the thermal history of direct laser fabricated Ti6Al-4V thin wall samples. This was correlated with microstructural evaluation in-process. These works have all contributed to understanding thermal history and the resulting microstructure in DED titanium alloys particularly in thin wall structures. In these studies, the energy source used was usually a continuous wave type laser
https://doi.org/10.1016/j.jmst.2019.05.008 1005-0302/© 2019 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
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and the laser power was usually over 500 W, which resulted in a relatively high thermal accumulation through build duration. Some researchers have studied the DED of pulsed lasers. Pinkerton et al. [14] deposited 316 L stainless steel part with a pulsed laser beam and found that the porosity and surface roughness reduced when using a pulsed beam rather than using a continuous beam, and the hardness of the deposited part increased. Banday et al. [15] prepared Ti/MoS2 coating on Al-Si alloy by pulsed laser deposition to invesigate the nanoscratch resistance and nanotribological performance of the coating. Gharbi et al. [16] demonstrated that using quasi-continuous instead of fully continuous laser irradiations could add further improvement to the surface finish of Ti-6Al-4V parts produced by DED. Ahsan et al. [17] fabricated Ti6Al-4V porous structures by pulsed laser metal deposition with high power, and focused on the formation of pores. In fact, lowpower pulsed laser DED can be proposed as a method to precisely fabricate and repair thin wall with its much lower average energy input. In this study, a low-power pulsed laser is used to deposit Ti6Al-4V alloy and can provide high cooling rate (<103 K/s [18]) while causing sufficient melting for a high integrity deposit to form. The resulting microstructure and mechanical properties are worthy of further investigation. Therefore, it is necessary to comprehensively study the microstructural evolution and mechanical properties in DED utilizing a low-power pulsed laser source.
mens were prepared by standard mechanical polishing method, and etched with a solution of HF: HNO3 : H2 O with a ratio of 1: 3: 20. Macro- and Micro-structure of the alloys were examined by an OLYMPUS optical microscope (OM) and a Hitachi S-4800 scanning electron microscope (SEM). Quantitative measurements of microstructure involved in this study were conducted using ImagePro Plus software.
2. Experimental procedure
3.1. Morphology of prior ˇ grains
2.1. Low-power pulse laser DED
Fig. 3(a)-(c) shows the typical microstructures of the bottom, middle and top of Ti-6Al-4V thin wall respectively. These were extracted from blue dashed boxes in Fig. 2(a). It can be seen that the microstructure of the thin-wall specimen is composed of elongated epitaxial  columnar grains, which are almost perpendicular to the substrate surface and clearly show the growth is driven by conductive heat flux toward the substrate. From Fig. 3(a), at the bottom of thin-wall sample, it can be seen that the width of prior  columnar grains is smaller at proximity to the substrate, and as the Z distance increases the width of the  columnar grains increases. When Z ≈ 1 mm, the width of the  columnar grains reaches a steady state with an average value of about 200 m. From Fig. 3(b) and (c), it can also be seen that the average width of the  columnar grains almost remains constant at the middle and top of thin-wall sample. Fig. 3(d) shows a local enlargement of the area marked by the red box in Fig. 3(c). The grain boundaries of the epitaxial  grains exhibit clear periodic wave patterns, and this is a notable feature in this figure. The periodic wavelength is about 0.2 mm, and just is the thickness of two layers. The morphology of prior  grains is very different from that in continuous laser DED. Usually, the microstructure of Ti-6Al-4V fabricated by conventional continuous laser DED is composed of coarse epitaxial  columnar grains or mixed structure of columnar grains and equiaxed grains, and the columnar grains grow straight along built direction with less obvious grain steering [7,22].
Thin wall (single track thickness) deposition of Ti-6Al-4V was performed using a low-power pulsed laser DED system as illustrated in Fig. 1(a). This system consists of a 300 W HANS Nd: YAG pulsed laser, a three-axis CNC unit to control the movement of X-Y table and control the Z-axis (build direction), a glove box which contained an argon atmosphere (oxygen content below 50 ppm), a powder feeder and a coaxial nozzle conveying powder to the deposition point by a pressurized carrier gas (argon). The powder material for the experiment was Ti-6Al-4V with particle size ranging from 200 to 300 mesh powder particle and chemical composition (wt%) as follows: 6.17 Al, 4.14 V, 0.15 O, 0.02 H, 0.005 N, 0.016 Si, 0.006 C, 0.032 Fe and Ti balance. The powder was supplied by YuGuang Phelly Ltd (China). Ti-6Al-4V plate (140 mm × 50 mm × 6 mm) was used as a substrate, which was burnished and cleaned before deposition. The powder was dried in a vacuum oven at 120 ◦ C for 2 h to remove moisture. Thin-walls were deposited with the following process parameters: laser pulse width 6 ms, pulse frequency 20 Hz, current intensity 40 A, laser spot diameter 1 mm, laser scanning speed 4 mm/s, powder feeding rate 1 g/min, carrier gas flow rate 200 L/h, and layer thickness 0.1 mm. The process parameters were chosen based on previous researches to ensure the metallurgical quality. Liu [19] and Hou [20] in our research group had used these process parameters for pulsed laser DED experiments to fabricate Ti-6Al-4V parts and got good results. In this way a thin-wall sample was built using single-pass and multi-layer laser deposition using a ‘back and forth’ strategy. Thus deposition directions of the adjacent tracks are opposing, as shown in Fig. 1(b). Finally, a thin wall with a geometric size of 75 mm × 35 mm × 1.2 mm was fabricated, as shown in Fig. 1(c), and was directly aged at 560 ◦ C for 2 h in order to relieve residual stresses. 2.2. Microstructural characterization The bottom, middle and top zones in cross-section XZ (which are marked respectively by blue dashed boxes in Fig. 2(a)), were extracted for microstructure observation. Metallographic speci-
2.3. Mechanical testing Hardness was measured by a Vickers tester (Struers DuraminA300) using a load of 0.3 kg and loading time 15 s. Hardness values were averaged over three measurements. In this study, nonstandard tensile specimens were used to investigate the tensile properties in the deposition height direction and the laser traverse direction. The extraction of specimens is illustrated in Fig. 2(a). The dimensions of tensile specimens and actual specimens are shown in Fig. 2(b) and (c). Tensile tests were carried out using an Instron 3382 universal material testing machine, and stress-strain data were recorded in real time using a Digital Image Correlation System (DIC) (XJTUDIC). Tensile properties were calculated from the results of DIC based on a method reported elsewhere [21]. 3. Results and discussion
3.2. Formation of ˇ grains During the initial deposition to the substrate, the  grains epitaxially grow from the fine equiaxed grains of the substrate, which is shown in Fig. 3(a). Due to proximity to the substrate and high conductivity the cooling rate is also greater, so the width of prior  columnar grains is relatively small. Upon multi-layer deposition toward a thin-wall structure, the temperature of the succeeding deposited layer would increase due to reduced heat sinkage, resulting in a significant decrease of the temperature gradient, thus a rapid increase of grain size. However, the energy delivered by the low power pulsed laser deposition is significantly less than the reported laser deposition process (usually a laser power
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Fig. 1. Illustration of pulsed laser deposition experiment, (a) schematic diagram of system, (b) schematic diagram of thin-wall laser deposition process, (c) Ti-6Al-4V thin-wall sample deposited.
Fig. 2. Preparation of the samples for metallographic observation and tensile test, (a) schematic illustration of specimen extractions, (b) dimensions of the test specimen, (c) tensile test samples arranged as built.
of 500–8000 W [7,11,12,17,21–23]), so the process would rapidly reach a thermal steady-state, which means that the temperature gradient of the melt pool zone would remain constant. Therefore, the average width of the  columnar grains will remain constant. During the solidification process, the high temperature gradient, G, and low solidification velocity, Vs , will tend to result in colum-
nar grain growth, while low G and high Vs will trend to result in equiaxed growth [24]. Fig. 4(a) shows the schematic relationship between scanning velocity, Vb , solidification velocity, Vs , and the melt pool boundary. It can be seen that Vs is normal to the melt pool boundary, and the relationship between Vb and Vs is defined by Vs = Vb ·cos [25], where is the angle between the vectors of Vb and
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Fig. 3. Morphologies of prior  grains of Ti-6Al-4V thin wall, (a) the bottom, (b) the middle, (c) the top and (d) local enlargement of typical morphology.
Fig. 4. Solidification analysis of the melt pool, (a) schematic relationship among scanning velocity Vb , solidification velocity Vs and melt pool boundary, (b) the calculated CET curve (calculated from Lin’s model [28]) and schematic solidification condition at the melt pool tail.
Vs . Therefore, from the bottom to top of the melt pool, Vs gradually increases. According to literature [26], G always decreases from the bottom to the top of the melt pool. Fig. 4(b) shows the critical condition curve of columnar to equiaxed transition (CET) of Ti-6Al-4V. This is calculated from Lin’s model [27]. The red dashed line indicates the schematic solidification condition at the melt pool tail. In this case, the pulsed laser results in a high temperature gradient, which would promote epitaxial columnar grain growth. It can be seen that the solidification condition is in the columnar grain growth zone at all times, which indicates that CET would not occur at the solid-liquid interface during the solidification of the molten pool. It is different from the usual C/W mode laser DED, and the  grains usually present mixed growth of columnar and equiaxed forms, as shown in previous studies [9,28]. Further, it can also be seen in Fig. 3(c) that there are no equiaxed grains at the top of the sample, indicating that solidification process is dominated by the columnar grain growth from the bottom to the top of the molten pool tail. Therefore, the solidification microstructure of the thinwall sample presents prior  columnar grains with highly epitaxial growth.
The formation mechanism of the periodic wave boundary of epitaxial  grains is also worthy of further analyses. According to Han et al. [29], in moving the melt pool, the columnar grains grow from the fusion line with the planar solid-liquid interface, and then curve gradually to follow the direction of the highest temperature gradient (normal to the local tailing pool boundary). Fig. 5(a) shows simulated results of typical thermal behavior in laser deposition of thin wall using a finite element model developed by Bontha et al. [30]. It can be seen that the direction of local highest temperature gradient Gm is always normal to local tailing pool boundary, and always diverts to the laser scanning direction. Therefore, as shown in Fig. 5(b), when the layer i is deposited, the melt pool moves to the right and Gm at the melt pool tail leans to the right, which causes the columnar grains to curve gradually to the right. This is clearly also evident in the experimental results presented here. Also, when the layer i+1 is deposited, the melt pool moves to the left, and Gm at the melt pool tail leans to the left, causing the columnar grains to curve gradually to the left. When layer i+2 is deposited, it mimics the situation reported in the layer i. Therefore, with the laser scanning back and forth, the direction of Gm at the melt pool tail
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Fig. 5. Mechanism of the wavy epitaxial growth of  grain, (a) temperature fields simulation by FEM with laser scanning back and forth, (b) schematic of the wavy growth.
laser deposition. In addition, from Fig. 3(c), a microstructure transitional zone of 500 m (the thickness of 5 layers of deposition) can be observed at the top of thin-wall sample, which exhibits a different microstructure from that of a stable typical zone.
Fig. 6. Schematics of melt pools with different depth-to-width ratios.
periodically changes, resulting in the periodic deflection of epitaxial grain growth. However, for conventional continuous laser DED, the deflections of epitaxial grains are not as obvious. This is because the peak energy density of pulsed laser may be higher than that of the continuous laser, resulting in a greater depth-to-width ratio of the melt pool, indicating that the tail boundary of melt pool caused by pulsed laser is steeper, as observed by Assuncao et al. [31], who reported this phenomenon when comparing the melt pools caused by continuous and pulsed lasers. Fig. 6 shows the schematics of the melt pools with different depth-to-width ratios. It can be found that the angle that Gm leans toward the laser scanning direction increases significantly at the tail of melt pool with greater depth-towidth ratio, resulting in an increase of the deflection of the epitaxial grain. As a result of this, the unusual and interesting periodic wave epitaxial grains (shown in Fig. 3(d)) were formed during the thinwall deposition with pulsed laser scanning back and forth. 3.3. Microstructure in prior ˇ grains 3.3.1. Microstructural overview From Fig. 3(d),  grains in different gray levels can be observed in the OM photograph. Fig. 7(a) shows the typical microstructure in “bright”  grain (marked in Fig. 3(d)). It can be observed in Fig. 7(a) that the typical microstructure in the region is mainly composed of many ␣ laths and a small amount of acicular ␣ phase, with a width of 1˜ m and a length between 10–30 m. Fig. 7(b) shows the typical microstructure in “dark”  grain (marked in Fig. 3(d)). It can be seen that the typical microstructure in the region contains a large number of elongated acicular ␣ phases, with a length of over 100 m and a width of 1˜ m, and a large amount of lath ␣ phase with a width of only 0.3 m and a length of 5˜ m. The difference of phase between “bright” and “dark” grains may be caused due to different crystal orientations of prior  grain. The “bright” and “dark”  grains both contain a large number of acicular ␣ phases, indicating that the cooling rate is high during low power pulsed
3.3.2. Microstructural evolution It is well known that this microstructure transitional region at the top of a build is formed from the last several thermal cycles, and can be used to analyze the evolution of the microstructure. As shown in Fig. 8(a), during thin wall deposition, the last layer n undergoes a thermal cycle, the layer n-1 undergoes two thermal cycles, and the layer n-2 undergoes three thermal cycles, and so on. Therefore, if each deposited layer at the top of the sample can be clearly distinguished, the influence of thermal cycles on microstructure evolution can be directly observed. However, it is usually difficult to distinguish the fusion line. Fig. 8(b) shows the local amplification of the blue dotted frame area at the top zone in Fig. 3(c). It can be also seen that it is difficult to discern the melting boundary of each layer in the top transitional region. However, based on previous analyses, we can see that each deflection of the periodic wave  grain represents a change of laser scanning direction, and the periodic wavelength is approximately 0.2 mm (just is the thickness of two layers deposition). In other words, the periodic wave pattern of one  grain can be used to accurately determine the fusion boundary of each layer (that is to say the peak and trough of the wave are exactly the fusion line’s positions at each layer). Therefore, the influence of thermal cycling on the layer-by-layer microstructure evolution can be investigated based on the fusion boundary of each deposited layer by evaluation of the periodic wave pattern of the epitaxial  grain. In this study, in order to determine the correct area and then observe the microstructure in each layer under a high magnification scanning electron microscope, some stamps were marked by the Vickers tester at the peak and trough of wave of  grain using a load of 0.1 kg and loading time 11 s, which allows us to determine the final several layers at the top of the sample, as shown in Fig. 8(b). Here, we defined the last layer as Layer n, and those layers before the last layer were defined as follows: layer n-1, n-2, n-3, n-4, and so on. From the marks, the thickness of each layer can be measured, ˜ mm (it agrees with the set value of layer thickness), which and is 0.1 indicates that the deflection analysis of the wave epitaxial grain is accurate and reliable. Then, by microstructure observation in the last several layers, the layer-by-layer influence of thermal cycles on the solid phase transformation can be obtained. Therefore, this study will present a novel and reliable method that can be used
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Fig. 7. Microstructures in  columnar grain, (a) “bright” grain, (b) “dark” grain.
Fig. 8. Analysis of the top of the sample, (a) schematic of deposited layer formation at the top of the sample, (b) microstructure observed at the top of the sample and diagram of deposition layer marking.
to investigate the microstructure evolution under thermal cycling conditions of DED. In order to disclose the influence of thermal cycle on the microstructure evolution, in this study, the red line frames in Fig. 8(b) were selected as the exemplar region in each layer. Fig. 9 showed the typical SEM microstructure images in each layer of the final five layers. As shown in Fig. 9(a) and (b), the last layer zone exhibits elongated and parallel acicular martensitic ␣’ phase with widths of 1–2 m, and a small amount of ␣ phase (lathes). These
with an average width of 0.3 m can be observed especially around the acicular martensitic ␣’ phase. At layer n-1, the acicular martensitic ␣’ phase is still present, but the amount of lath ␣ is observed to increase and enlarge slightly, with an average width of 0.5 m, as shown in Fig. 9(c) and (d). From Fig. 9(e) and (f), it can be seen that the microstructure features change significantly at layer n-2. The acicular ␣’ phase disappears, while many long and thin ␣ laths exhibited in the microstructure, and are approximately perpendicular to parallel acicular martensitic ␣’ phase in Fig. 9(a)-(d), the
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Fig. 9. Microstructures in the same  grain and final five layers of the sample, (a), (b) the microstructures of the last layer, (c), (d) the microstructures of the last second layer, (e), (f) the microstructures of the last third layer, (g), (h) the microstructures of the last fourth layer, and (i), (j) the microstructures of the last fifth layer.
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Fig. 10. Schematic diagram of the thermal cycle history and microstructure formation.
amount of lath ␣ continuing to increase as well. From Fig. 9(g)(j), it can be seen that long and thin ␣ phase becomes gradually shorter and less prevalent, while the size and amount of lath ␣ gradually increase. Overall, it can be clearly seen that from the final layer to the layer n-4 (as thermal cycle increases), the lath ␣ phase uniformly increases in size. Moreover, the microstructure in the  grain at layer n-4 (suffering five thermal cycles) is basically similar to typical microstructures (see Fig. 7(b)) in the same  grain, indicating that the microstructure of the deposition layer tends to be stable after five thermal cycles. 3.3.3. Microstructure evolution mechanism The microstructure evolution in dark  grains is determined via thermal cycle history conditions. In order to make the discussion representative, thermal cycle histories of each layer during steady deposition were used in this study. For ease of explanation, three important temperatures of Ti-6Al-4V were noted. The TL and T are liquidus and  transus lines, respectively. Tdiss stands for ␣-dissolution temperature, where ␣ phase starts to dissolve into  under equilibrium heating conditions. When the temperature is below Tdiss , the microstructure in  grain will not change. Fig. 10 shows the schematic diagram of the thermal cycle history and microstructure formation. The thermal cycle time that the layer n (final layer) underwent is tn , the thermal cycle time that the layer n-1 underwent is tn-1 , and the thermal cycle time that the layer n-2 underwent is tn-2 , and so on. It is important to recall that cooling rate of the thermal cycle gradually decreases in a laser multi-layer deposition due to reducing proximity to the heat source. During the deposition of layer n, its temperature sharply decreases, from above TL to T , due to the very high cooling rate caused by the pulsed laser, the martensitic ␣’ phase would be formed at T , and with decreasing temperature, small amounts of Widmanstätten ␣ lath are formed (see the observation results of layer n in Fig. 9(a) and (b)). Layer n-1 underwent two thermal cycles within tn-1 . The first thermal cycle is the same as that of layer n, and with the second thermal cycle, when the temperature rapidly rises to above T , the acicular ␣’ phase and ␣ lath formed under the first thermal cycle would be completely dissolved. When the temperature decreases rapidly to T , some new martensitic ␣’ phase will be formed again due to the high cooling rate, and then, with a decrease of temperature at a relatively lower cooling rate than that of layer n, some Widmanstätten ␣ laths (more than those of layer n) will be formed as well (see the observation results of layer n-1 in Fig. 9(c) and (d)). Layer n-2 underwent three thermal cycles within tn-2 , with the third thermal cycle, when temperature increases to above T , the acicular ␣’ phase and ␣ lath formed under the second thermal cycle will be completely dissolved again. When the temperature decreases to T , the new martensitic ␣’ phase was not formed again, which is because the cooling rate is lower than that of ␣’ formation.
However, the cooling rate remains relatively high, so some elongated acicular ␣ phase would be formed, and there would be more Widmanstätten ␣ laths than those of layer n-1 formed as well, as shown in Fig. 9(e) and (f). Ahmed et al. [32] made clear that for ␣ +  titanium alloys, when the cooling rate is lower than that of martensitic ␣’ formation but remains relatively high, the elongated acicular ␣ phase will be formed on the grain boundary, and rapidly grow towards the grain interior. So, in this study, a large number of elongated acicular ␣ particle will rapidly grow from the boundary of thin grain and distribute on the whole grain interior. For the layer n-3, similarly, with the fourth thermal cycle, when the temperature rises to between T and Tdiss , the acicular ␣ phase and ␣ lath formed in third thermal cycle would be partially dissolved, so the acicular ␣ phase would get shorter and thinner. As the temperature continues to decline at a lower cooling rate, more Widmanstätten ␣ lath structures will be formed and grow bigger, as shown in Fig. 9(g) and (h). For layer n-4, the acicular ␣ phase would get much shorter and thinner, and ␣ lath becomes more prevalent and enlarged, as shown in Fig. 9(i) and (j). Compared to Fig. 7(b), the microstructure in layer n-4 has been basically similar to the typical microstructures in the same  grain. Furthermore, for layer n-5, due to the peak temperature of the sixth thermal cycle below Tdiss , the microstructure remains unchanged. Therefore, a steady microstructure would be reached in layer n-4 undergoing five thermal cycles under these experimental conditions. In summary, using a pulsed laser, the solidification velocity Vs and the temperature gradient G are larger, so primary dendritic spacing 1 is smaller according to 1 ∝Vs –a G−b (a, b are constants) and fine epitaxial growth columnar grains are formed. Meanwhile, ␣ phase nucleates at the grain boundary and grows through the whole fine columnar grain, and with the thermal cycles the acicular ␣ and Widmanstätten ␣ lath are finally formed, while using conventional continuous laser there exhibits basket weave ␣ lath structures in  grain interior [13,33]. 3.4. Mechanical properties 3.4.1. Hardness Microstructures in  grains exhibit apparent differences. Thus the microhardness of different  grains was further tested every 0.5 mm within the range of 4 mm along the deposition height direction, as shown in Fig. 11(a). The hardness results were averaged over three measurements at the same height, and the results are shown in Fig. 11(b). It can be seen that the “bright” and “dark”  grains ˜ have high hardness values (375-389 HV in the “bright”  grain, and ˜ 398-410 HV in the “dark”  grain). This is because they both contain a large number of elongated acicular ␣ phase due to high cooling rate caused by the low power pulsed laser. Microhardness distribution in the “dark”  grain is higher than that in the “bright”  grain due to greater strengthening by a large number of dense parallel acicular ␣ particles (see Fig. 7(b)). 3.4.2. Tensile testing Fig. 12 shows the tensile properties of thin wall samples. It can be seen that the average yield strength ( y ) and ultimate tensile strength ( b ) of transverse tensile samples reached 1020 MPa and 1190 MPa, respectively, but the average elongation was only 4.3%. The average yield strength and ultimate tensile strength of height direction tensile samples was 900 MPa and 1050 MPa, respectively, and the average elongation was 14.2%, indicating an excellent combination of strength and elongation of height direction. Obviously the strength of the Ti-6Al-4V thin wall samples deposited are higher than that of the forged standard (yield strength 828 MPa and ultimate tensile strength 895 MPa) [34] due to the cooling rates. Surprisingly, the strength of the transverse tensile samples is significantly higher (13.3%) than those of the build direction tensile
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Fig. 11. Micro-hardness test in different  grains, (a) the schematic of test position, (b) hardness results.
Fig. 12. Tensile properties of specimens taken from the thin wall.
samples, but the elongation of the transverse tensile samples is significantly lower (69.7%) than those of the height direction tensile samples. This is because in the case of the transverse tensile test, there is a large amount of  grain boundaries being perpendicular to the tensile direction (see Fig. 13(a)). From the previous research, we know that the average value of the width of the  columnar grains is about 200 m and the effective length of the tensile test specimen is 10 mm, so we can quantify that there are five  grain boundaries per mm. According to plastic deformation theory [35], the grain boundary density serves to impede dislocation movement, resulting in a notable strengthening effect. So the higher strength in the transverse tensile test was exhibited due to the strengthening effect of many vertical grain boundaries. In
addition, during the tensile process, the sample will deform into bamboo joint pattern due to the strengthening of these vertical grains boundaries, as shown in Fig. 13(a). Due to the anisotropy of the grains’ property (see the hardness test results in Fig. 11), the  grains in the test sample could not uniformly deform during transverse tension. So the weakest  grain will first reach the ultimate deformation to break, resulting in the serious decrease of elongation to failure. However, in the case of the build direction tensile test, the amount of  grain boundaries is reduced, and these boundaries lie largely in the tensile direction (as shown in Fig. 13(b)), so the strengthening effect is not as significant, and is lower than that of the transverse direction. During the tensile testing, the thin and long epitaxial grains in the test samples would be able to deform in unison, resulting in excellent elongation to failure. Fig. 14 shows the SEM micrographs of the typical fracture surfaces of the two direction tensile samples. The contraction of the fracture cross-section is obvious (see Fig. 14(a)), and the fracture morphologies of the height direction sample show the typical high fracture toughness characteristics associated with ductile dimples (see Fig. 14(b)), which shows that the fracture mechanism is microvoid accumulation fracture. The tensile fracture surface in the height direction sample is dominated by large and round or equiaxed dimples, with an average diameter of 4˜ m, indicating excellent ductility, which agrees with the elongation measured by the tensile test. From Fig. 14(c), the shrink of the fracture crosssection is not obvious, and there are some shallow and elongated dimples on the fracture surface of the transverse tensile sample, as shown in Fig. 14(d). Furthermore, river patterns and cleavage terraces can also be observed, which shows the fracture morphologies are of cleavage character. In other words, the fracture mechanism of the transverse direction sample is ductile/brittle mixed fracture
Fig. 13. Schematic of tensile process analysis, (a) transverse directoin, (b) build direction.
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Fig. 14. SEM micrographs of the typical fracture surfaces of two direction tensile samples, (a), (b) height direction, (c), (d) transverse direction.
which indicates poor plasticity, agreeing with the elongation measured from the transverse sample. 4. Conclusions A Ti-6Al-4V thin wall was fabricated using low-power pulsed laser DED to elucidate information relating to microstructure formation phenomena. Solidification from the bottom to the top of the melt pool tail is dominated by the columnar grain growth at all times, so the solidification microstructure of the thin-wall sample is composed of long and thin prior  columnar grains, which are almost perpendicular to the laser scanning direction. The boundaries of the epitaxial  grains exhibited clear periodic wave patterns, which can be used to identify the fusion boundary of each layer (the periodicity matches the fusion line positions of each layer). Based on the determination of the fusion line of each layer by the periodic wave pattern of epitaxial  grain at the top of the thin wall, the influence of the thermal cycle on microstructure layer-bylayer evolution was investigated. It was observed that the acicular martensitic ␣’ phase and a small amount of Widmanstätten ␣ lath converted to long and thin acicular ␣ phase and larger amounts of Widmanstätten ␣ lath under thermal cycle conditioning. Elongated acicular ␣ phase was formed on the  grain boundary, and rapidly extended into the grain interior. Therefore, the thin acicular ␣ phase easily distributes in the grain interior. The microstructure difference between  grains will cause the property difference. The strength of Ti-6Al-4V thin wall samples are higher than that of the forged standard due to high cooling rates, and the elongation of build direction is much higher than that of the forging standard, indicating an excellent combination of strength and elongation in the build direction. The strength of the transverse tensile samples is 13.3% higher than that of the build direction samples due to the significant strengthening effect of a large amount of vertical  grain boundaries. However, the elongation of transverse tensile samples is 69.7% lower than those of build direction samples due to the uneven deformation of  grains.
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