In-situ Formation of Titanium Boride and Titanium ...

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Physics Procedia 39 (2012) 432 – 438

LANE 2012

In-situ Formation of Titanium Boride and Titanium Carbide by Selective Laser Melting Timo Kühnle , Knut Partes BIAS - Bremer Institut für angewandte Strahltechnik GmbH, Klagenfurter Str. 2, 28359 Bremen, Germany

Abstract In the present study the Selective Laser Melting of a high-energy ball milled Ti - B 4 C powder blend was performed. The mixing ratio of the powder blend was settled to fulfill the reaction 3 Ti + B 4 C TiC + 2 TiB 2 . The in-situ synthesis of intermediate phases was performed by single line welding experiments on a pure titanium substrate. Phase identification by X-ray diffraction gave no indication for the formation of new phases due to the mixing of the powder in the ball mill. However, within the weld seam the XRD analysis proved the existence of Ti 2 C and TiB. © 2012 2011 Published Publishedby by Elsevier ElsevierB.V. Ltd. Selection © Selection and/or and/or review review under under responsibility responsibility of of Bayerisches Bayerisches Laserzentrum Laserzentrum GmbH GmbH Keywords: Selective Laser Melting; In-situ phase formation; particle reinforced material; Titanium Diboride; Titanium Carbide

1. Introduction Selective Laser Melting (SLM) is an emerging technology for the additive manufacturing of metal parts. A typical build rate in additive manufacturing for metallic powders is 2 mm³/s. Recent developments show that it is possible to increase the build rate up to 10 mm³/s with a high power laser and a skin-core scanning strategy [1]. The demands on the material properties of the SLM products are increasing: especially for rapid tooling applications a high relative density, wear resistance, strength, hardness with a fine, homogeneous microstructure are required. A state of the art material for rapid tooling is for example the stainless steel 18Ni-300. With an appropriate heat treatment its possible to increase the material hardness up to 58 HRC [2]. In order to increase the hardness even further a higher carbon content is necessary. However, a drawback of high carbon steels is the poor weldability of the material due to the phenomenon of cold-cracking. According to [3] the maximum hardness of a full martensite microstructure is limited to ~65 HRC. To achieve further improvement of wear resistance a

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1875-3892 © 2012 Published by Elsevier B.V. Selection and/or review under responsibility of Bayerisches Laserzentrum GmbH doi:10.1016/j.phpro.2012.10.058

Timo Kühnle and Knut Partes / Physics Procedia 39 (2012) 432 – 438

reinforcement of a metal base matrix with ceramic particles or intermetallic compounds is eligible. An efficient way to modify the surface properties of tools is provided by laser material processing. For example, laser hardening of tool surfaces or laser cladding of wear resistance coatings are well known methods for industrial applications [4]. Laser beam cladding of gradient coatings with hard reinforcement particles, e.g. Cr 3 C 2 , NiBSi or WC is also a promising method to improve the wear resistance [5]. In that sense, the present study deals with a strategy to create particle reinforced materials with an SLMmachine. The in-situ synthesis of hard intermetallic constituents is a promising method to create particle reinforced metal matrix composites. The target result of in-situ phase formation after solidification is a fine and homogeneous microstructure, which is thermally stable. Furthermore, the intermetallic constituents have good compatibility with the matrix and therefore the created composite has elevated mechanical properties [6]. In the present study in-situ formation of new phases means, that the desired chemical reaction is induced by the laser beam and occurs in the molten pool of a premixed powder blend during SLM. There are several examples for the in-situ synthesis of hard particles during SLM. For instance, it is possible to reinforce a nickel base alloy with tungsten carbide by SLM using a mixture of nickel, tungsten and carbon powder. In addition to tungsten carbide the M 6 C-type carbide Ni 2 W 4 C is formed [7]. An other example for in-situ synthesis during SLM is the reinforcement of titanium with the reaction 8 Ti + 3 SiC Ti 5 Si 3 + 3 TiC using a powder blend of titanium and silica-carbide as raw material [8]. The possibility for in-situ generation of titanium carbide and titanium diboride by argon arc cladding is given by [9]. The synthesis of titanium carbide and titanium diboride using SLM has not yet been investigated. TiB 2 -TiC composites exhibit superior properties including a high hardness up to -1/2 and a bending strength in the range of 400-600 MPa 25 GPa, high fracture toughness up to 12 [10]. 2. Experimental The aim of this study is the in-situ synthesis of TiB2 and TiC hard particles starting with a Ti + B4C powder. The mixing ratio of the powder blend was settled to a mass ratio of Ti:B4C = 72:28 to fulfill the reaction 3 Ti + B4C TiC + 2 TiB2. Milling was performed with a Retsch PM400 planetary ball mill in argon atmosphere to homogenize the powder blend. The total milling time was 15 hours at 200 rpm in intervals of 15 minutes with pauses of 5 minutes. For in-situ phase synthesis single line welding experiments were conducted in a thin powder layer with a thickness of 100 μm on a pure titanium substrate. All welding experiments were carried out on a Realizer SLM-250 machine with an IPG YLR-200-SM ytterbium fiber laser. The single mode laser was operating in continuous wave mode. The corresponding laser power was set to 100 W with a scan velocity of 20 mm/s and the beam diameter was 150 μm. The powder morphology before and after milling was examined via SEM microscopy. The particle size distribution was examined with a Beckman Coulter LS laser diffraction particle size analyzer. XRD analysis was performed to proof the in-situ synthesis of new phases after selective laser melting of the thin powder layer. The microstructure in the weld bead was analyzed with optical microscopy.

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3. Results and Discussion The morphology of the basic material is show in Figure 1. The size specification of the B4C particles was -60 +30 μm and the irregular shape with sharp edges and fractured surfaces indicates that particles were crushed. The Ti 99.9 powder was specified with -60 μm and Figure 1b shows many small particles with a size of ~10 μm and smaller. Figure 1c shows the premixed powder blend after 15 h high energy ball milling. Surprisingly, the morphology of the constituents did not change significantly. However, through ball milling it was possible to guarantee a homogeneous mixed powder blend. A disadvantage of the crushed particle morphology is the lack of flowability of the loose powder blend. Because of the irregular geometry the particles have the tendency stick to each other and form agglomerates. Therefore it was not possible to create a powder layer with a thickness of less than 100 μm with the standard wiper mechanism for powder deposition in the SLM machine.

Fig. 1. Powder morphology: a) B 4 C powder; b) Ti 99.9 powder; c) powder blend Ti + B 4 C after 15h ball milling

Figure 2 shows the particle size distribution of the two powders and the powder blend. The percentage values on the ordinate are related to the total powder volume. This result confirms the visual impression of the SEM Image in Figure 1c: although the powder blend was milled for a time period of 15 hours, it was not possible to reduce the size of the comparatively large B 4 C particles. The maximum of the size distribution curve shifts from 14.8 vol.% at 55 μm for the B 4 C powder to 7.6 vol.% at 50 μm for the powder blend. The measured peak volume percentage of the powder blend is in good agreement with theoretical value, if we assume a linear mixing rule of the initial constituents.

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Fig. 2. Particle size analysis of the powder materials (Ti and B 4 C) and the powder blend after 15h ball milling

In addition, Table 1 provides a numerical comparison of the particle size properties. One example for Ti 99.9 to clarify the notation in the table: d 25,3 = 13.7 μm means that 25 vol.% of the powder particles have a diameter which is smaller than 13.7 μm. In context with the size distribution in Figure 2, we conclude that particles are not crushed during the milling process. Presumably, the energy input during the milling process was insufficient to break the particles. Table 1. Classification of the powder particle size with respect to the percentage of the total powder volume

powder type

particle size / μm d 10,3

d 25,3

d 50,3

d 75,3

d 90,3

Ti 99.9 powder

7.7

13.7

22.6

33.6

44.5

B 4 C powder

36.7

44.2

53.0

62.4

71.6

ball milled powder blend Ti + B 4 C

10.6

19.6

36.3

51.2

62.7

Figure 3a shows the result of the single layer - single line selective laser melting experiments. The cross-section of the single track reveals the weld seam geometry. Remarkable is the high degree of dilution between the base material and the powder. However, there is no indication for un- or partially melted particles in the weld seam. The magnified detail D-1 in Figure 3b shows a dentritic microstructure. Although the laser beam diameter is set to 150 μm, the width of the weld seam is 300 μm. A possible reason for this discrepancy is the fluid-dynamics of the liquid material in the melt pool, for instance the margangoni convection.

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Fig. 3. a) Cross section of the single track – single layer SLM-experiment; b) Detail D-1: dentritic microstructure after rapid solidification in the centre of the weld bead. Etching: 100ml H 2 O + 2g NH 4 HF 2 . Process parameters: laser power 100 W; -1 ; beam diameter 150 μm; powder layer thickness 100 μm scan velocity 20

The result of the XRD analysis is shown in Figure 4. The lower graph represents the properties of the milled powder blend prior laser melting. It is possible to identify the two phases of the base material: B 4 C + Ti. However, the important information is that there is no clear indication for titanium borides or titanium carbides. Therefore, no additional phases were synthesized due to high energy ball milling. The graph on top is representative for melted and solidified material in the weld bead. The weld bead was irritated from above to measure the XRD spectrum, which means the upper graph represents the properties at the weld bead surface. The most important information of the XRD graph is the fact that after laser melting the B 4 C-peaks disappear. Instead of B 4 C it is possible to identify TiB and Ti 2 C. We conclude, that boron carbide reacted in the melt pool with titanium to form titanium boride and titanium carbide. However, the initially intended synthesis of TiB 2 and TiC was not achieved. All synthesized compounds had a higher titanium content than the targeted phases. The cross section in Figure 3a shows a weld seam with a penetration depth of 70 μm, which implies that a large amount of substrate material was molten during the welding process. This molten substrate material increases the titanium content in the weld pool. The strong dilution within the weld seam indicates an excess of energy input. On the one hand it is desirable to ensure a low dilution between substrate material and powder material to guarantee the correct mixing ratio of the constituents. On the other hand it is necessary to provide enough energy input to ensure complete melting of the B 4 C-particels to start the in-situ reaction. Further examinations are necessary to find the balance between low dilution and complete melting of the powder blend.


BIAS ID 120619

Fig. 4. XRD-analysis of the Ti+B4C powder blend prior and after laser melting upon a titanium substrate

4. Conclusions No additional phases were synthesized due to high energy ball milling of a Ti + B 4 C powder blend. The melting of a premixed powder blend with a laser beam is a successful method for in-situ synthesis of titan borides and titanium carbides. The original phase B 4 C disappeared after laser melting of the powder blend and the study provides evidence for the existence of the phases TiB and Ti 2 C in the weld bead. The intended in-situ reaction 3 Ti + B 4 C TiC + 2 TiB 2 did not occur due to the strong dilution of the molten powder blend with the molten substrate material.

Acknowledgement The authors gratefully acknowledge the financial support by DFG (German Research Foundation) for subproject C1 “Werkzeugwerkstoffe” within the SFB747 (Collaborative Research Center) “Mikrokaltumformen – Prozesse, Charakterisierung, Optimierung”. Our special thanks go to Karoline Pardun from the “advanced ceramics group” for the support with high energy ball milling and to Alwin Schulz from “Stiftung Institut für Werkstofftechnik” for the assistance and fruitful discussion in context with the XRD analysis.

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