March 16, Berlin / GER The Novel Scanning Strategy For Fabrication Metallic Glasses By Selective Laser Melting Łukasz Żrodowski* 1, Bartłomiej Wysocki 1, Rafał Wróblewski 1, Krzysztof Jan Kurzydłowski 1, Wojciech Święszkowski 1 1 Warsaw University of Technology, Faculty of Materials Science and Engineering, Woloska 141 St., 02-507 Warsaw, Poland * Corresponding Author: lukasz.zrodowski@gmail.com +48 692 366 939 Abstract Metallic Glasses (MGs) can be described as a stable in the room temperature metallic materials with a disordered liquid-like structure produced during rapid cooling of a molten alloy. Due to the limited Glass Forming Ability (GFA), most of the MGs are produced as a thin ribbons through melt-spinning or as a fine powder during atomization processes. Certain multicomponent alloys with exceptionally high GFA, known as Bulk Metallic Glasses (BMGs), can retain disordered structure during copper-mold casting. The size and the complexity of objects produced in a such way is limited by the critical cooling rate. Additive Manufacturing (AM) methods like Selective Laser Melting (SLM) have been shown in a few studies as very perspective for producing BMGs without such limitations. In this work Realizer SLM-50, a desktop SLM machine, was used to selectively melt Kuamet52 Fe-Si-B-Cr-C metallic glass powder with various laser parameters. Material was rescanned with high power density and a novel Pulse-Random (P-R) strategy to achieve high content of glassy phase, despite low GFA of the alloy. Optical Microscopy (OM) was used to preliminary determine material structure. Amorphous structure was confirmed by the Differential Scanning Calorimetry (DSC) while residual crystalline phases were identified by the X-Ray Diffraction (XRD). Microstructural observations revealed, after the first melting in conduction mode, mostly crystalline phases made in samples. After the second melting in key-hole mode, which was performed with P-R strategy, there was observed great increase of amorphous phase content. Samples fabricated with new scanning strategies had amorphisation degree exceeding 60%, with retained α-fe 3 Si and FeB 2 crystalline phases. Model of the laser melted MGs crystallization, focused on devitrification during laser heating in Heat Affected Zone (HAZ), has been proposed to explain observed phenomena. 1. Introduction Due to exceptional mechanical and chemical properties Metallic Glasses (MGs) are promising structural [1] and bio materials [2,3], but the high cost of advanced molding systems and technical difficulties still limit their applications. The solution for enhancing the applicability of metallic glasses is processing them by Additive Manufacturing (AM) methods. Selective Laser Melting (SLM) drew attention as a suitable method for processing amorphous alloys due to the cooling rates higher than these needed for amorphisation of various MGs. Recently, a few examples where amorphous structure has been achieved via this technique were published [4,5]. Still alloys used in those experiments have exceptionally high Glass Forming Ability (GFA) and they can be produced in bulk form and casted into copper mold. Another example showed that MGs with lower GFA induces severe crystallization during the process [6]. Especially Heat Affected Zone (HAZ) between layers or melt pools are sensitive to crystallization caused by accumulation of structural relaxation [7]. In this work, we have sucesfully selectively laser melted Fe-based metallic glass with low GFA using various laser parameters. The introduced Pulse-Random (P-R) remelting strategy reduced influence of the heat accumulation and increased Amorphisation Degree (AD). The new crystallization model for SLM processing of MGs, based on Ozawa equation [8], has been proposed. 2. Materials and methods The commercial amorphous powder Kuamet 52 (Atmix-Epson, Japan) with the nominal composition: Fe 71 Si 10 B 11 C 6 Cr 2 and mean diameter 24.5 µm has been used in the experiment. Cylindrical samples of the 5 mm diameter and 5 mm height were produced with constant layer thickness of 50 µm using SLM-50 (Realizer GmbH) desktop machine for manufacturing metal components. Each layer during the samples fabrication was scanned two times. During the first melting there was used checkerboard strategy with 1 mm edge length and the laser power W. The second melting was made with the P-R strategy and maximum laser power W. The detailed set of the parameters is summarized in Table 1. Minimum distance between subsequent points during the P-R strategy was set to 1 mm. Scheme of the P-R strategy is shown in Fig. 1. Two sets of samples, with and without the remelting, ISBN 978-3-8396-1-5 Fraunhofer / DDMC 16 1
March 16, Berlin / GER were created to show the influence of remelting strategy. The Scanning Electron Microscope (SEM) observations of the powder were performed using HITACHI S3500 microscope. In order to prepare the metallographic sections, samples were cut in the middle perpendicularly to the base plate using wire Electro Discharge Machining (w-edm). After moulding in epoxy resin, samples were prepared by metallographic standard methods like grinding and polishing. Etching was performed with a solution composed of ml 98% ethanol, 10 ml 65% nitric acid, and ml of distilled water. The metallographic observations were performed using NIKON EPIPHOT 0 light microscope. The DSC analysis was performed for both initial powder and every sample ( mg in graphite crucible) using Perkin Perkin Elmer DSC 8000 with constant heating rate 40 K/min. The AD % was calculated as the crystallization enthalpy of the sample divided by the crystallization enthalpy of the powder. In order to determine the activation energy of the first crystallization step there was used Ozawa method [7]. For this purpose, initial powder peak crystallization temperature was measured at the heating rate β =, 40, 80, 160, and 3 K/min. The X-ray diffraction (XRD) was performed using Rigaku MiniFlex II diffractometer with Cu Kα X-Ray source on the initial powder, sample with the highest AD, and on the annealed powder (1123 K 1 h). The goniometer step was set to 0,01 with measurement time 5 s. The diffraction patterns were compared to standard PDF4+ database. 50 µm Figure 2: SEM image of Kuamet 52 powder a) Table 1: Laser parameters used during the first melting Sample A B C Exposure time [µs] 500 500 0 P-R Point distance [µm] 65 Hatch distance [µm] 65 X Laser power [W] Energy density [J/mm3] 25 59 50 5 b) Fig. 1. Fabrication scheme of one layer of a metallic glass with the novel scanning strategy; a) loose powder; b) first melting of loose power with checkerboard strategy; c) second melting of previously melted layer by random pulses numbers represent melting order; d) fully remelted layer after P-R scanning strategy ISBN 978-3-8396-1-5 c) Fig. 3. Etched cross sections of samples without remelting; a) Sample A, AD:13,2%; b) Sample B, AD:5,7%; c) Sample C, AD:3,6. Amorphous phases marked by arrows, crystalline phases marked as stars, pores marked as squares Fraunhofer / DDMC 16 2
March 16, Berlin / GER 3. Results and discussion The SEM images of the initial metallic glass powder are shown in Fig. 2. Most of the powder particles had a spherical or oval shape, together with a smooth and featureless surface, which explains a good powder processability. The microstructures of the single melted samples together with the DSC-measured AD values are shown in Fig. 3. The values of used laser parameters had influence on both porosity and AD of the fabricated samples. Increased exposure time or reduced point/hatch distance effected in the lower porosity but also in the lower AD. This effect was well visible between microstructures showed in Fig. 3a-c. The lower porosity of the samples B and C than A can be attributed to the higher energy density delivered to the material during the melting. The lower AD content in the samples B and C than A can be explained by increased heat shock effect in HAZ [7]. The metallographic cross-sections of samples, which layers were melted during the fabrication just once with checkboard strategy, consist mostly regions strongly etched by the chemical reagent (marked as stars). These dark etched regions were identified as the crystalline phase. In the contrast to the crystalline dark regions metallic glass could be observed as the bright non etched dots (marked by arrows). It should be denoted that, regardless of the used parameters, amorphous phase formed only separated melt pools around the pores and these regions were not connected through the layers. On the other hand the pores tended to protect amorphous melt pools from devitrification during subsequent layer melting. The different AD values and the melt pools morphology were observed on the samples fabricated when the second melting of each layer with the P-R strategy was performed (Fig. 4a-c). The AD values for all samples after remelting have increased approximately four times. The minor reduction of the porosity in comparison to the samples fabricated without remelting was also observed. Similarly to the samples fabricated without P-R remelting the sample A exhibited the highest AD, while the samples B and C had lower and almost equal AD. In the sample A the most amorphous melt pools were connected without crystallization in HAZ and their length exceed thickness of a few layers. a) b) c) Fig. 4. Etched cross sections of samples made with P-R remelting; a) Sample A, AD:62,3%; b) Sample B, AD:,7%; c) Sample C, AD:,6%. µm Amorphous phases marked by arrows, crystalline phases marked as stars, pores marked as squares Fig. 5. Microstructure of sample B after remelting; Melt pool after first scanning is marked with a rhombus; Melt pool after P-R is marked with a square ISBN 978-3-8396-1-5 Fraunhofer / DDMC 16 3
March 16, Berlin / GER The higher magnification of the remelted B sample is shown in Fig. 5. There can be observed two different types of melt pools: the crystalline and the amorphous melt pools with the different aspect ratio. The crystalline melt pools after the first melting were 30-60 µm deep and 80-1 µm wide. The amorphous melt pools after the P-R strategy remelting were - 0 µm deep and 30-60 µm wide. There was a little or no crystallization in the HAZ of the P-R melt pools. In the field of a laser processing change of a melt pool aspect ratio is well known as a result of the transition between the conduction melting and the key-hole melting modes. During a key-hole mode melting power density is high enough to vaporize a material and allow a laser beam to penetrate deeper inside it. In this experiment, key-hole mode melt pools were responsible for high AD values. On the contrary to the key-hole mode, just after the conduction mode the microstructure of samples did not consist of connected regions of Heat flow [a.u.] endo T r T x T p1 T p2 300 400 500 600 Temperature [ C] amorphous phase. Fig. 6. The XRD patterns for the annealed powder, the remelted A sample and the initial powder. Peaks marked with circle were identifed as α-fe 3 (Si) while with square as Fe 2 B The XRD patterns for the annealed powder, the remelted Sample A and the initial powder are shown in Fig. 6. There were no peaks on the XRD pattern of the initial powder, but only a wide amorphous hump, therefore initial powder was considered to be fully amorphous. The DSC calculations of the initial powder AD confirmed that result. The annealed powder had equilibibirum phase composition which was identified as α-fe 3 (Si) with addition of Fe 2 B. Due to the multicomponent composition of the alloy there were also other minor peaks of an unidentified phase or phases. The remelted Sample A was a halfway between the initial and the annealed powder. On the amorphous hump there were sharp peaks of α-fe 3 (Si) and barely distinguishable from background Fe 2 B peaks. No peaks shift was observed for annealed powder and the remelted Sample A, which excluded occurrence of stress concentration in the crystalline phase. Fig. 7. The DSC curve for the Sample A with remelting (upper line) and the initial powder (bottom line); T r - relaxation start; T x - crystalization temperature onset; T p1 - first crystallization stage peak temperature; T p2 - Second crystallization stage peak temperature The AD values for the all fabricated samples was calculated as the crystallization enthalpy of the sample divided by the crystallization enthalpy of the initial powder. The DSC curves for the amorphous powder and the remelted sample A are shown in Fig. 7. Both curves had the same character with weak exothermal effect visible at 351 C followed by two overlapping crystallization effects. Glass transition temperature is not visible for this alloy. For the initial powder and the remelted Sample A characteristic temperatures were: T x = 572 C, T p1 = 579 C, T p2 = 595 C and T x = 573 C, T p1 = 582 C, T p2 = 600 C respectively. Minor characteristic temperatures shift and the similar devitrification effect prove that initial glassy structure was kept. The crystallization enthalpy for the powder was 76.8 J/g and 47.6 J/g for the remelted Sample A, thus AD of the sample A with P-R remelting was 62.3 %. This value was the highest for all fabricated samples. Both Sample A and the initial powder are considered to be nonrelaxed due to long exothermal effect from lasting from T r to T x. As the relaxation usually causes embrittlement of MGs [9] it is important not only to retain an amorphous structure, but also prevent a relaxation. To investigate the effect of the pulse duration on the crystallization in HAZ, the activation energy (E a ) for the ISBN 978-3-8396-1-5 Fraunhofer / DDMC 16 4
March 16, Berlin / GER first stage of the crystallization, was calculated by the Ozawa equation (1). For this purpose linear function was fitted to ln(β) vs 0/T p plot, where β was heating rate and T p was temperature peak. llllllll ββββ = EEEE + CCCC (1) β - constant heating rate[k/s], T p [K] crystallization peak temperature, E a - crystallization activation energy [J/mol], R- gas constant [J/Kmol], C- constant. The calculated activation energy equals 260.6 kj/mol for our alloy. The Ozawa model makes an assumption that during constant heating 63% of the volume fraction of the glass crystalizes at the T p. Due to devitrification kinetics model there is no defined β at which the crystallization completely stops. Thus to simplify the model we have assumed that a glass heated during a laser pulse to the maximum temperature (T max ) devitrifies if heating rate is lower than critical. Critical heating rate (β c ) is defined by equation (2). ln(β) 25 15 10 5 1. Melting before devitri ication 4. Melting after devitri ication ββββ = e ( ) (2) Taking melting point T m = 1250 K and taking that it was independent from β, we have calculated critical heating rate β c = 8.77 * 10 6 K/s at which no glass devitrifies until material melts. T m Measured Tp value 500 µs pulse µs pulse β for µs pulse β for 500 µs pulse 0 0,7 0,8 0,9 1 1,1 1,2 0/T Fig. 8. The thermal model conditions for laser heated glass. ln(β) vs 0/T p plot with linear fit for Ozawa equation; β - constant heating rate[k/s]; T - temperature [K];T m - melting point of the alloy O l Plotting two intersecting lines (T m, and Ozawa linear fit) we have proposed four different thermal conditions for the laser heated glasses (Fig. 8). The first condition above T m and Ozawa line (O l ) corresponds to metallic glass melted without devitrification (1. Melting before devitrification). The second one below T m and above O l represents glass in HAZ which avoided devitrification due to high heating rate (2. Amorphous HAZ). The third region which is below T m and O l, where β was not sufficient, thus glass devitrified (3. Crystalline HAZ). The fourth, above T m and below O l represent glass melted after devitrification (4. Melting after devitrification). Assuming that the real heating rate in the melt pool and the HAZ were close to the T m divided by the total exposure time, two points were marked on the plot. The estimated heating rate for first melting (pulse duration 500 µs) was β 500 = 2,5 * 10 6 K/s. The estimated heating rate for second melting (P-R remelting with pulse duration µs) was β = 6.25 * 10 7 K/s. The proposed model explained different microstructures for the µs and 500 µs pulse melt pools. The µs pulse provided heating rate well above critical at melting point. It was fully located within first and second part of the plot, thus no crystallization occurred. On the other side the 500 µs pulse provided heating rate below critical (β c) at melting point. This pulse crossed both third and fourth part of the plot, thus it could be affected by the crystallization in HAZ and melting after devitrification. The model was focused on the constant heating condition but it could be used to the better understanding of crystallization during the cooling step. If the heating conditions are in the fourth part of the plot, the melt pool appears and its border contains crystalline phases. Thus during a melt pool cooling, crystallites might be uplifted by Marangoni convection and promote crystallization acting as homogenous nuclei. According to the experimental results and known asymmetry of the critical cooling and heating rates [10] we have stated that producing fully amorphous MGs part by SLM is limited not only by the cooling rate, but foremost by the heating rate required to counteract devitrification in HAZ. 4. Conclusions The Point-Random (P-R) rescanning strategy has restored original amorphous structure of the initial powder despite its low GFA and crystallization during the first melting. After the P-R processing amorphous phase retained connected between layers instead of forming the separate areas as observed in the samples made just with the single melting. The high aspect ratio of the amorphous areas fabricated in P-R mode was attributed to transition between conduction and key-hole mode melting. The difference between crystallization in the HAZ during the long and short pulses irradiation ISBN 978-3-8396-1-5 Fraunhofer / DDMC 16 5
March 16, Berlin / GER was attributed to the different heating rate and it was an effect of the devitrification kinetics. As a final concussion we have found that remelting with the key-hole mode is advantageous for the SLM processing of metallic glasses with low GFA. Furthermore amorphisation of the MGs during SLM is limited by the heating rate. Acknowledgements The authors would like to thank the NCBiR (National Center for Research and Development) for providing financial support to project LasIMP (Grant No. PBS3/A5/53/15). P-R scanning strategy is a part of patent pending method of producing MGs parts via SLM. Patent application priority date is 14/1/16. Literature [1] M. Ashby and A.Greer, Metallic glasses as structural materials.scripta Materialia, vol 54, no. 3, 06, pp. 321-326 [2] X. Gu,Y. Zheng,S. Zhong,T, Xi, J. Wang, and Wang, W., Corrosion of, and cellular responses to MgZnCa bulk metallic glasses, Biomaterials vol. 31, 10, pp. 1093 1103. [3] L. Huang, D. Qiao, B. Green, P. Liaw, J. Wang, S. Pang, Bio-corrosion study on zirconium-based bulk-metallic glasses Intermetallics vol.17, 09, pp. 195-199 [4] S. Pauly, L. Löber; R. Petters, M. Stoica, S. Scudino, U. Kühn, J. Eckert, Processing metallic glasses by selective laser melting. Mater Today, vol.16, no.1/2,13, pp 37-41 [5] H. Jung, S. Choi, K. Prashanth, M. Stoica, S. Scudino, S. Yib, U. Kühn, D. H. Kim, K. B. Kim, J. Eckert, Fabrication of Fe-based bulk metallic glass by selective laser melting: A parameter study:. Materials & Design, vol. 86, 15, pp 703-708. [6] X. Li,C. Kang, H. Kuang, T. Sercombe, The role of a low-energy density re-scan in fabricating crack-free Al85Ni5Y6Co2Fe2 bulk metallic glass composites via selective laser melting, Materials and Desing, vol. 46, 14, pp. 408-411 [7] G. Yang, X. Lin, F. Liu, Q. Hu, L. Ma, J. Li, W. Huang, Laser solid forming Zr-based bulk metallic glass, Intermetallics, vol. 22, 12, pp. 110-115 [8] T. Ozawa, Kinetic analysis of derivative curves in thermal analysis, Journal of Thermal Analysis, vol.2, 1970, pp. 301 324 [9] P. Murali, U. Ramamurty, Embrittlement of a bulk metallic glass due to sub-tg annealing, Acta Materialia, vol 53, 05, pp. 1467-1478. [10] J. Schroers, A. Masuhr, W. Johnson, Pronounced asymmetry in the crystallization behavior during constant heating and cooling of a bulk metallic glass-forming liquid, Physical Review B, vol. 60, no. 17, 1999 ISBN 978-3-8396-1-5 Fraunhofer / DDMC 16 6