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Microstructural characterization and wear behavior of (Fe,Ni)–TiC MMC prepared by DMLS
Journal of Alloys and Compounds 421 (2006) 166–171
Microstructural characterization and wear behavior of (Fe,Ni)–TiC MMC prepared by DMLS A. G˚aa˚ rd ∗ , P. Krakhmalev, J. Bergstr¨om Department of Materials Engineering, Karlstad University, SE-651 88 Karlstad, Sweden Received 25 March 2005; accepted 2 September 2005 Available online 15 December 2005
Abstract Direct metal laser sintering (DMLS) is an evolving technique for the production of parts with complex geometry and unique microstructure, implemented in tooling applications by manufacturing of tool inserts. In the present investigation, microstructure, phase composition, mechanical and tribological properties of (Fe,Ni)–TiC composites prepared by DMLS were investigated. Thermal cracking during sintering was observed, which degraded the produced parts, influenced by porosity and particular spherical particles. Dissolution of TiC led to subsequent FCC to BCC matrix phase transformation and to an increase of the average coefficient of thermal expansion. Therefore, a combination of high TiC amount, loss of low-CTE properties of the matrix and matrix phase transformation led to cracking of the composites. Decrease of TiC content to 30 wt.% led to a crack-free, but 10% porous specimen. Bending strength, hardness and wear resistance in sliding wear of the crack free specimen were found in the lower range of conventionally manufactured hard metals due to porosity, low TiC amount and low matrix hardness. © 2005 Elsevier B.V. All rights reserved. Keywords: Laser processing; Composite materials; Hard metals; Invar-TiC; Sliding wear; Selective laser sintering
1. Introduction Hard metals, such as WC–Co and TiC–Ni, find their main use as cutting tools and in wear resistant applications because of their superior hardness and reasonable fracture toughness. Generally, hard metals are produced by conventional sintering methods, which are productive and produce fully dense parts. However, conventional techniques limit complexity in geometry and WC–Co hard metals are quite expensive due to the Co content. Therefore, new sintering techniques and composites with cheaper binder phase are of scientific interest. Direct metal laser sintering (DMLS) is a selective laser sintering (SLS) technique, in use as a complement to conventional sintering methods. The DMLS technology produces parts in a layer-by-layer fashion by direct transfer from a computer-aided design (CAD) model. From the CAD model, thin cross-sections of the part are calculated and represented by layers of powder in ∗
Corresponding author. Tel.: +46 547001262; fax: +46 547001449. E-mail addresses: [email protected] (A. G˚aa˚ rd), [email protected] (P. Krakhmalev), [email protected] (J. Bergstr¨om). 0925-8388/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2005.09.084
the sinterstation. Each layer of powder is fused by the heat of a focused laser beam. After consolidation of one cross-section, a new layer of powder is applied and the process is iterated until a complete 3D part is generated. Laser sintering was initially developed for plastic prototypes, but during the last decade, the technique has evolved to comprise metals and metallic based composites as well [1–5]. The technique is prospective to be implemented in complex tooling applications by manufacturing of tool inserts. DMLS, generally, employs a mixture of high and low melting-point powders, denominated as structural and binder particles, respectively. Laser radiation selectively melts the binder particles and the liquid metal is driven by capillary forces into the voids between the structural, non-melting, particles. Consolidation of the cross-section is attained by cooling and solidification of the liquid metal [6–8]. Hard metals, being a mixture of two phases with different melting points, are potential materials for the DMLS technology. Research on fabricating of hard metals by DMLS has been done, but with limited success in achieving full density. In laser sintering of WC–Co systems, cracking was found to depend on powder particle size and laser speed [9], while [10] used a selective
A. G˚aa˚ rd et al. / Journal of Alloys and Compounds 421 (2006) 166–171
laser sintering process followed by infiltration of the porous body. The mixed powders generally possess different absorption to laser radiation, depending on laser wavelength and the nature of the materials. Therefore, the powder bed is heated heterogeneously. If the structural particles have higher absorptance than the binder particles they will absorb more energy. Depending on particle size, this could cause melting of the structural particles, even though their melting point is higher. To increase the densification, the binder particle size should be smaller than the structural particles. Detailed reviews of the laser and powder interaction have been published in [7,11]. Formation of the final microstructure is preceded by heat conduction, melting, rearrangement, densification and cooling phenomena [6]. In the present research work, a series of (Fe,Ni)–TiC metal matrix composite (MMC) materials were prepared by direct metal laser sintering. The influence of powder mixture on microstructure, porosity and crack formation was studied. Eventually, mechanical and tribological properties of a crack-free specimen were examined. 2. Experimental Invar36, utilized in this investigation, is a Fe–Ni alloy with a low coefficient of thermal expansion (CTE), Table 1. Above 300 ◦ C mechanical properties such as the yield strength decrease rapidly. Hence, during solidification and cooldown less elastic energy will be stored in the matrix, which is presumed to reduce thermal stresses during sintering. The Invar36 powder was produced by gas atomization by Sprey Metal Ltd. with chemical composition (wt.%): 36.0 Ni, 0.27 Mn, 0.21 Si, 0.002 C, balance Fe. The TiC-powder was produced by milling by Atlantic Equipment Engineers with 99.9% TiC pureness. The size and morphology of the powder particles were estimated by scanning electron microscopy. The spherical Invar36 particles diameter was in the range of 1–50 m with mean diameter of 7 m. The angular TiC-powder consisted of a large amount of very small particles, less than 1 m in size. Mean size of the larger TiC-particles was estimated to be about 30–50 m. Specimens were prepared using an EOSINT M 250 extended sinterstation equipped with a CO2 laser and argon atmosphere. The powder mixtures investigated were Invar36 with 30, 60 and 80 wt.% TiC, respectively. All sintered samples were ground and polished to mirror like surface with 1 m diamond paste finishing and examined in scanning electron microscope LEO 1530 and Jeol JSM-820. Porosity was measured by computerized digital analysis on prepared cross-sections with accuracy within the range ±5%. Three-point bending tests were performed on specimens with dimensions of 55 mm × 5 mm × 3 mm in an Instron 8500 plus servo hydraulic machine. Vickers hardness measurements were made in a Leco V-100-C2 within the range ±50 HV. Phase analysis was performed by X-ray diffraction (XRD) in a Seifert XRD 3003 PTS X-ray diffractometer using Cr K␣ X-ray radiation. Tribological tests were performed on a crack free specimen in a block-oncylinder tribometer. A full description of the test equipment could be found in [14]. The Invar36-TiC specimen was employed as a block with dimensions of 5 mm × 20 mm × 8 mm and polished. A 0.12% C steel was used for the cylinder, prepared by turning to Ra = 0.73 µm prior to testing, with a hardness of 200 HV. The specimen was pressed against the cylinder with a normal load of 40 N at
167
constant sliding speed of 0.6 m/s, under room temperature conditions in air. The worn volume was measured by optical profilometry with a Wyko NT3300, and the wear coefficient was calculated as [15]: k=
V Fn s
(1)
where V is the volume loss, Fn is the normal load and s is the sliding distance. Scar morphology was investigated using SEM.
3. Results The microstructure of the sintered composites is illustrated in Fig. 1. It was found that small and large TiC-particles were homogeneously distributed in the Invar36 matrix. Typical for liquid phase sintered microstructures, slightly rounded shapes of large particles were observed. Spherical inclusions having a thin shell-like TiC surface layer and dendritic TiC structure inside, Fig. 1, were observed in some amount in all composites sintered. All as-sintered materials were porous, Table 2, and a maximal porosity of 23 vol.% corresponded to the intermediate Invar36-60 wt.%TiC composition. Thermal stresses provided significant cracking in materials comprising 80 and 60 wt.% TiC, but the Invar36-30 wt.%TiC composite was found crack-free. Crack growth from pore to pore, Fig. 2a, through both the Invar36 matrix and the TiC particles was observed, Fig. 2b. XRD phase composition analysis of the as-sintered Invar36 matrix showed that increasing total TiC amount led to transformation of the FCC phase to BCC, Fig. 3, with the amount of the BCC phase depending on total TiC content, Table 2.
Fig. 1. Sintered microstructure in Invar36-30 wt.% TiC composite. Dark phase corresponds to TiC particles and bright to the Invar36 matrix, SEM image.
Table 2 Properties of the sintered materials at laser speed 50 mm/s
Table 1 Material properties of Invar36 and TiC [12,13] Material
CTE, 10−6 (◦ C)−1
Tm (◦ C)
Absorptance to CO2 radiation
Invar36 TiC
1.60 7.76
1427 3067
0.43 0.46
TiC (wt.%)
Porosity
Thermal cracks
HV
␣/␥ Phase
Ti in matrix (wt.%)
30 60 80
10 23 8
No Yes Yes
380 400 850
0/1 1/3 3/1
3.0 4.4 6.3
168
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4. Discussion
Fig. 2. (a) Thermal crack paths in materials comprising 60 and 80 wt.% TiC. The cracks propagated from pore to pore. (b) Thermal crack originating from sintering. Crack growth through both Invar36 and TiC phases is illustrated.
Additionally, increasing total TiC amount led to an increase of Ti concentration in the matrix as observed by SEM EDX, Table 2. Mean bending rupture strength of the crack free Invar 3630 wt.%TiC material was found to be 912 MPa and hardness was about 380 HV, Table 2. Investigations of the bending fracture surfaces showed a ductile fracture in the Invar36 matrix and brittle cleavage fracture through the TiC particles, Fig. 4. Crack growth around spherical particles, found on the fracture surfaces, was observed, Fig. 4. The morphology of the worn Invar36-30 wt.%TiC composite surface after 6000 m run is illustrated in Fig. 5. A significant amount of material, about 2 mm3 , was removed by wear and a 0.25 mm deep wear scar, with a shape conforming to the cylinder periphery was formed, Fig. 6. Abrasive wear was identified as the dominant wear mechanism and the wear coefficient, at 6000 m sliding length, was 8 × 10−6 mm3 /Nm, Eq. (1). Additionally, an iron oxide film, as confirmed by SEM EDX, was found as scattered local islands on the worn composite surface, Fig. 5a. Fig. 5b illustrates multiple furrows, with an iron oxide particle embedded in the metal matrix at the end of one of the wear scars.
Formation of spherical particles in selective laser sintering is a known phenomena, and has recently been reported of in [8]. Poor wetting and excess amount of liquid phase were proposed contributory factors to the spheres formation mechanism. In the present investigation, formation of the spheres, and particularly the internal dendritic structure, implies a local overheating with melting of both metallic and ceramic phases during sintering. Related regular dendritic structures have been reported in [16], where both TiC and Invar were melted by a YAG laser. In the present work it was assumed that dissolution of TiC in the molten metal and subsequent cooling led to crystallization of the hightemperature TiC phase, forming a shell-like surface layer around an internal dendritic structure, Fig. 1. The wetting properties of the liquid to the structural phase and amount of liquid phase is essential for formation of a dense layer structure [7,8]. These parameters have to be sufficiently high to form a continuous infiltration front to fill the voids between the structural particles. Additionally, for rapid pore evacuation, a high infiltration rate of the liquid into the powder bed is needed. For a three dimensional porous media, the infiltration rate is dependent on the liquid–vapor surface tension and the liquid viscosity [8]. In the present research work, it was assumed that the dissolution of the TiC phase may have influenced the viscosity of the liquid metal and that primary re-crystallization of TiC at the front of the liquid phase decreased the penetration rate of the liquid into the powder bed, leading to porosity. During sintering, stresses develop in already sintered layers if the thermal expansion, induced by the temperature change, is hindered, either by external or internal constraints. In [17], thermal stresses in multiphase materials were discussed as a consequence of internal constraints caused by different coefficients of thermal expansion between the phases. If the thermal expansion is hindered by external constraints or uneven temperature distributions, bending stresses proportional to the average composite CTE, are generated [18]: L/2 6 σb = − 2 Eαt TX dx (2) L −(L/2) where L is the layer thickness, E is the elastic modulus, αt is the average CTE of the composite, T is the temperature difference and X is equal to x − L/2, x is the distance measured from the layer surface. According to the Turner model [19], the average CTE, αt , of a multi-phase composite is derived as follows: αt =
α1 V1 K1 + α2 V2 K2 + α3 V3 K3 · · · , V1 K1 + V 2 K2 + V 3 K3 · · ·
(3)
where V is the volume fraction, K is the bulk modulus and αi is the CTE of each phase. Following Eq. (2) the thermally induced bending stress depends on the average composite CTE, which according to Eq. (3) depends on the particular CTE of the material constituents. For the present material combination, the bending stress increases with increase of TiC content, due to higher CTE
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169
Fig. 3. XRD patterns of composites with 30, 60 and 80 wt.% TiC, respectively. Formation of BCC Fe–Ni–Ti–C phase with increasing TiC amount was observed.
of TiC compared to Invar36. But, as confirmed by SEM EDX, increasing total TiC amount led to dissolution of TiC and subsequent increase of Ti and C concentration in the Invar36 matrix, Table 2, resulting in considerable changes of the material properties. Dissolution of Ti and C in Invar36 significantly increases the alloys CTE [20], and has a solid solution strengthening effect, with possible loss in ductility as a consequence. Furthermore,
Fig. 4. Fracture surface obtained in bending test of Invar36-30 wt.% TiC composite. Ductile fracture of the Invar36 matrix, brittle fracture of TiC phase and delamination around spherical particles are illustrated.
Ti, being a strong BCC stabilizer [21], provided transformation of the matrix FCC phase to BCC phase, observed by XRD, Fig. 3. Matrix strengthening by multiple phases, could intensify the possible loss in ductility. Therefore, several factors participated in affecting the magnitude of the average composite CTE; the powder composition and the variation of the matrix CTE due to FCC to BCC phase transformation and enrichment of Ti and C. At high TiC quantities the influence of the matrix CTE is small, Eq. (3), but at the intermediate Invar36-60 wt.%TiC composition the effect is more significant. Hence, materials comprising 60 and 80 wt.% TiC phase, suffered of severe thermal cracking during sintering due to loss of low-CTE properties of the Invar36 matrix and the influence of TiC on the average composite CTE. Conventionally sintered WC–Co hard metals possess bending rupture strengths from 1700 MPa and hardness values from about 780 HV [22]. In the investigated crack free Invar3630 wt.%TiC material, mean bending rupture strength and hardness was about 900 MPa and 400 HV, respectively. The lower values could be attributed to low TiC amount, low matrix hardness and porosity. Crack propagation, in bending, through both phases, Fig. 2b, implies a strong bonding, pointing at sufficient wetting between the molten Invar36 metal and the TiC during sintering. However, the spherical particles observed on the fracture surfaces,
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Fig. 6. (a) Optical profilometry three-dimensional wear track of the Invar3630 wt.%TiC composite after 6000 m run. (b) Line profile taken along the white line in Fig. 6a.
Fig. 5. (a) Smeared iron oxide film, bright areas, on the worn surface of the Invar36-30 wt.%TiC composite after 6000 m run. (b) Abrasive wear caused by iron oxide particles.
Fig. 2b, indicates delamination and consequently a weak bonding between matrix and spheres. The wear coefficient in sliding was in the order of one magnitude higher than values obtained for TiC–NiMo composites, tested under the same conditions as in the present work, [15]. But on the contrary to [15], where adhesive wear was stated as the ascendant wear mechanism, abrasive wear was the dominating mechanism in the present work. A substantial amount of the TiC particles had a size of less than 1 m, and as illustrated in Fig. 5b, the width of the grooves were approximately 20 m. In multiphase materials, with a relatively low amount of reinforcement particles, the particle size should be of the same order as the abrasive particles to contribute to the total composite wear resistance, [23]. Hence, the wear resistance contribution of the finer TiC particles in the Invar-30 wt.%TiC material was negligible, which facilitated the abrasive wear mechanism. The relatively high binder phase content, led to frequent metal-metal contact, promoting strong adhesion between the mating tribo-surfaces. The smeared film of iron oxide, Fig. 5a, located as islands on the composite surface, is a possible consequence of metal-metal contact. Adhesive junctions may break within one of the mating surfaces, leading to a material transfer. Subsequently, the transferred material oxidizes, which is
promoted by the temperature increase associated with frictional heating during sliding. The transfer layer may come loose and be removed as a debris particle, or it can be reattached to one of the surfaces. In the present work, the latter phenomenon, with reattachment to the steel cylinder and two-body abrasion as a consequence, is a credible cause to the dominant abrasive wear mechanism observed. A second origin to the abrasive particles, are oxide particles originating from the steel cylinder, which was found severely oxidized during the experiment. 5. Conclusions DMLS of Invar36-TiC composites, led to formation of microstructures with homogeneous distributions of TiC particles in the metal matrix. However, melting of both metallic and ceramic phases occurred during sintering, which led to dissolution of Ti and C in the Invar36 matrix. Changes in chemical composition of the liquid Invar36 alloy led to spherical particle formation, changes in phase composition during solidification and loss of its low CTE properties. Consequently, with increasing TiC content, significant cracking of the specimen occurred during sintering, due to increased matrix CTE and transition from FCC to BCC phase. A maintained FCC structure and low amounts of dissolved TiC in the matrix, were realized only in the Invar36-30 wt.%TiC composite, where no cracks were observed. All sintered materials contained pores and the crack free Invar36-30 wt.% TiC composite corresponded to 10 vol.% porosity. Bending rupture strength and hardness of the crack
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free specimen were relatively low and could be attributed to low TiC content, low matrix hardness and porosity. The specific wear rate was 8 × 10−6 mm3 /Nm and SEM investigations concluded abrasive wear as the dominant wear mechanism. Acknowledgement The authors greatly acknowledge Mr. Jorgen Andersson at Uddeholm Tooling AB for support in X-ray investigations. References [1] K. Cooper, Rapid Prototyping Technology: Selection and Application, M. Dekker, New York, 2001. [2] D. Pham, Rapid Manufacturing, Springer, London, 2001. [3] D. King, T. Tansey, J. Mater. Process. Technol. 132 (2003) 42–48. [4] P. Jacobs, Society of Manufacturing Engineers, Dearborn Michigan, 1993. [5] X. Wang, T. Laoui, J. Bonse, J. Kruth, B. Lauwers, L. Froyen, Int. J. Adv. Manuf. Technol. 19 (2003) 351–357. [6] Y. Kathuria, Surf. Coat. Technol. 132 (2003) 42–48. [7] N. Tolochko, S. Mozzharov, T. Laoui, L. Froyen, Rapid Prototyping J. 9 (2003) 68.
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[8] S. Srivatsan, J. Moore (Eds.), Processing and Fabrication of Advanced Materials IV, The Minerals, Metals and Materials Society, Cleveland, OH, 1996, p. 17. [9] T. Laoui, L. Froyen, J.P. Kruth, Met. Powder Rep. 55 (2000) 40. [10] K. Maeda, T. Childs, J. Mater. Process. Technol. 149 (2004) 609– 615. [11] P. Fischer, V. Romano, H. Weber, N. Karapatis, E. Boillat, R. Glardon, Acta Mater. 51 (2003) 1651–1652. [12] N. Tolochko, T. Laoui, Y. Khlopkov, S. Mozzharov, V. Titov, M. Ignatiev, Rapid Prototyping J. 6 (2000) 155–160. [13] G. Song, Y. Wu, Q. Li, J. Eur. Ceram. Soc. 22 (2002) 559–566. [14] T. Bjork, J. Bergstrom, S. Hogmark, Wear 224 (1999) 216–225. [15] J. Pirso, M. Viljus, S. Letunovits, Tribol. Int. 37 (2004) 817–824. [16] X. Li, J. Stampfl, F. Printz, Mater. Sci. Eng., A 282 (2000) 86–90. [17] S. Suresh, Fatigue of Materials, Cambridge University, Cambridge, 1994. [18] A. Weronski, T. Hejwowski, Thermal Fatigue of Metals, Dekker, New York, 1991. [19] A. Fahmy, A. Ragai, J. Appl. Phys. 41 (1970) 5108–5111. [20] ASM Handbook, vol. 2, ASM International, 1990. [21] G. Krauss, Steels Heat Treatment and Processing Principles, ASM International, 1990. [22] G. Upadhyaya, Mater. Des. 22 (2001) 483–489. [23] S. Jacobson, S. Hogmark, Tribology; Friction, Lubrication and Wear, Berlings, 1996.
Microstructural characterization and wear behavior of (Fe,Ni)–TiC MMC prepared by DMLS A. G˚aa˚ rd ∗ , P. Krakhmalev, J. Bergstr¨om Department of Materials Engineering, Karlstad University, SE-651 88 Karlstad, Sweden Received 25 March 2005; accepted 2 September 2005 Available online 15 December 2005
Abstract Direct metal laser sintering (DMLS) is an evolving technique for the production of parts with complex geometry and unique microstructure, implemented in tooling applications by manufacturing of tool inserts. In the present investigation, microstructure, phase composition, mechanical and tribological properties of (Fe,Ni)–TiC composites prepared by DMLS were investigated. Thermal cracking during sintering was observed, which degraded the produced parts, influenced by porosity and particular spherical particles. Dissolution of TiC led to subsequent FCC to BCC matrix phase transformation and to an increase of the average coefficient of thermal expansion. Therefore, a combination of high TiC amount, loss of low-CTE properties of the matrix and matrix phase transformation led to cracking of the composites. Decrease of TiC content to 30 wt.% led to a crack-free, but 10% porous specimen. Bending strength, hardness and wear resistance in sliding wear of the crack free specimen were found in the lower range of conventionally manufactured hard metals due to porosity, low TiC amount and low matrix hardness. © 2005 Elsevier B.V. All rights reserved. Keywords: Laser processing; Composite materials; Hard metals; Invar-TiC; Sliding wear; Selective laser sintering
1. Introduction Hard metals, such as WC–Co and TiC–Ni, find their main use as cutting tools and in wear resistant applications because of their superior hardness and reasonable fracture toughness. Generally, hard metals are produced by conventional sintering methods, which are productive and produce fully dense parts. However, conventional techniques limit complexity in geometry and WC–Co hard metals are quite expensive due to the Co content. Therefore, new sintering techniques and composites with cheaper binder phase are of scientific interest. Direct metal laser sintering (DMLS) is a selective laser sintering (SLS) technique, in use as a complement to conventional sintering methods. The DMLS technology produces parts in a layer-by-layer fashion by direct transfer from a computer-aided design (CAD) model. From the CAD model, thin cross-sections of the part are calculated and represented by layers of powder in ∗
Corresponding author. Tel.: +46 547001262; fax: +46 547001449. E-mail addresses: [email protected] (A. G˚aa˚ rd), [email protected] (P. Krakhmalev), [email protected] (J. Bergstr¨om). 0925-8388/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2005.09.084
the sinterstation. Each layer of powder is fused by the heat of a focused laser beam. After consolidation of one cross-section, a new layer of powder is applied and the process is iterated until a complete 3D part is generated. Laser sintering was initially developed for plastic prototypes, but during the last decade, the technique has evolved to comprise metals and metallic based composites as well [1–5]. The technique is prospective to be implemented in complex tooling applications by manufacturing of tool inserts. DMLS, generally, employs a mixture of high and low melting-point powders, denominated as structural and binder particles, respectively. Laser radiation selectively melts the binder particles and the liquid metal is driven by capillary forces into the voids between the structural, non-melting, particles. Consolidation of the cross-section is attained by cooling and solidification of the liquid metal [6–8]. Hard metals, being a mixture of two phases with different melting points, are potential materials for the DMLS technology. Research on fabricating of hard metals by DMLS has been done, but with limited success in achieving full density. In laser sintering of WC–Co systems, cracking was found to depend on powder particle size and laser speed [9], while [10] used a selective
A. G˚aa˚ rd et al. / Journal of Alloys and Compounds 421 (2006) 166–171
laser sintering process followed by infiltration of the porous body. The mixed powders generally possess different absorption to laser radiation, depending on laser wavelength and the nature of the materials. Therefore, the powder bed is heated heterogeneously. If the structural particles have higher absorptance than the binder particles they will absorb more energy. Depending on particle size, this could cause melting of the structural particles, even though their melting point is higher. To increase the densification, the binder particle size should be smaller than the structural particles. Detailed reviews of the laser and powder interaction have been published in [7,11]. Formation of the final microstructure is preceded by heat conduction, melting, rearrangement, densification and cooling phenomena [6]. In the present research work, a series of (Fe,Ni)–TiC metal matrix composite (MMC) materials were prepared by direct metal laser sintering. The influence of powder mixture on microstructure, porosity and crack formation was studied. Eventually, mechanical and tribological properties of a crack-free specimen were examined. 2. Experimental Invar36, utilized in this investigation, is a Fe–Ni alloy with a low coefficient of thermal expansion (CTE), Table 1. Above 300 ◦ C mechanical properties such as the yield strength decrease rapidly. Hence, during solidification and cooldown less elastic energy will be stored in the matrix, which is presumed to reduce thermal stresses during sintering. The Invar36 powder was produced by gas atomization by Sprey Metal Ltd. with chemical composition (wt.%): 36.0 Ni, 0.27 Mn, 0.21 Si, 0.002 C, balance Fe. The TiC-powder was produced by milling by Atlantic Equipment Engineers with 99.9% TiC pureness. The size and morphology of the powder particles were estimated by scanning electron microscopy. The spherical Invar36 particles diameter was in the range of 1–50 m with mean diameter of 7 m. The angular TiC-powder consisted of a large amount of very small particles, less than 1 m in size. Mean size of the larger TiC-particles was estimated to be about 30–50 m. Specimens were prepared using an EOSINT M 250 extended sinterstation equipped with a CO2 laser and argon atmosphere. The powder mixtures investigated were Invar36 with 30, 60 and 80 wt.% TiC, respectively. All sintered samples were ground and polished to mirror like surface with 1 m diamond paste finishing and examined in scanning electron microscope LEO 1530 and Jeol JSM-820. Porosity was measured by computerized digital analysis on prepared cross-sections with accuracy within the range ±5%. Three-point bending tests were performed on specimens with dimensions of 55 mm × 5 mm × 3 mm in an Instron 8500 plus servo hydraulic machine. Vickers hardness measurements were made in a Leco V-100-C2 within the range ±50 HV. Phase analysis was performed by X-ray diffraction (XRD) in a Seifert XRD 3003 PTS X-ray diffractometer using Cr K␣ X-ray radiation. Tribological tests were performed on a crack free specimen in a block-oncylinder tribometer. A full description of the test equipment could be found in [14]. The Invar36-TiC specimen was employed as a block with dimensions of 5 mm × 20 mm × 8 mm and polished. A 0.12% C steel was used for the cylinder, prepared by turning to Ra = 0.73 µm prior to testing, with a hardness of 200 HV. The specimen was pressed against the cylinder with a normal load of 40 N at
167
constant sliding speed of 0.6 m/s, under room temperature conditions in air. The worn volume was measured by optical profilometry with a Wyko NT3300, and the wear coefficient was calculated as [15]: k=
V Fn s
(1)
where V is the volume loss, Fn is the normal load and s is the sliding distance. Scar morphology was investigated using SEM.
3. Results The microstructure of the sintered composites is illustrated in Fig. 1. It was found that small and large TiC-particles were homogeneously distributed in the Invar36 matrix. Typical for liquid phase sintered microstructures, slightly rounded shapes of large particles were observed. Spherical inclusions having a thin shell-like TiC surface layer and dendritic TiC structure inside, Fig. 1, were observed in some amount in all composites sintered. All as-sintered materials were porous, Table 2, and a maximal porosity of 23 vol.% corresponded to the intermediate Invar36-60 wt.%TiC composition. Thermal stresses provided significant cracking in materials comprising 80 and 60 wt.% TiC, but the Invar36-30 wt.%TiC composite was found crack-free. Crack growth from pore to pore, Fig. 2a, through both the Invar36 matrix and the TiC particles was observed, Fig. 2b. XRD phase composition analysis of the as-sintered Invar36 matrix showed that increasing total TiC amount led to transformation of the FCC phase to BCC, Fig. 3, with the amount of the BCC phase depending on total TiC content, Table 2.
Fig. 1. Sintered microstructure in Invar36-30 wt.% TiC composite. Dark phase corresponds to TiC particles and bright to the Invar36 matrix, SEM image.
Table 2 Properties of the sintered materials at laser speed 50 mm/s
Table 1 Material properties of Invar36 and TiC [12,13] Material
CTE, 10−6 (◦ C)−1
Tm (◦ C)
Absorptance to CO2 radiation
Invar36 TiC
1.60 7.76
1427 3067
0.43 0.46
TiC (wt.%)
Porosity
Thermal cracks
HV
␣/␥ Phase
Ti in matrix (wt.%)
30 60 80
10 23 8
No Yes Yes
380 400 850
0/1 1/3 3/1
3.0 4.4 6.3
168
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4. Discussion
Fig. 2. (a) Thermal crack paths in materials comprising 60 and 80 wt.% TiC. The cracks propagated from pore to pore. (b) Thermal crack originating from sintering. Crack growth through both Invar36 and TiC phases is illustrated.
Additionally, increasing total TiC amount led to an increase of Ti concentration in the matrix as observed by SEM EDX, Table 2. Mean bending rupture strength of the crack free Invar 3630 wt.%TiC material was found to be 912 MPa and hardness was about 380 HV, Table 2. Investigations of the bending fracture surfaces showed a ductile fracture in the Invar36 matrix and brittle cleavage fracture through the TiC particles, Fig. 4. Crack growth around spherical particles, found on the fracture surfaces, was observed, Fig. 4. The morphology of the worn Invar36-30 wt.%TiC composite surface after 6000 m run is illustrated in Fig. 5. A significant amount of material, about 2 mm3 , was removed by wear and a 0.25 mm deep wear scar, with a shape conforming to the cylinder periphery was formed, Fig. 6. Abrasive wear was identified as the dominant wear mechanism and the wear coefficient, at 6000 m sliding length, was 8 × 10−6 mm3 /Nm, Eq. (1). Additionally, an iron oxide film, as confirmed by SEM EDX, was found as scattered local islands on the worn composite surface, Fig. 5a. Fig. 5b illustrates multiple furrows, with an iron oxide particle embedded in the metal matrix at the end of one of the wear scars.
Formation of spherical particles in selective laser sintering is a known phenomena, and has recently been reported of in [8]. Poor wetting and excess amount of liquid phase were proposed contributory factors to the spheres formation mechanism. In the present investigation, formation of the spheres, and particularly the internal dendritic structure, implies a local overheating with melting of both metallic and ceramic phases during sintering. Related regular dendritic structures have been reported in [16], where both TiC and Invar were melted by a YAG laser. In the present work it was assumed that dissolution of TiC in the molten metal and subsequent cooling led to crystallization of the hightemperature TiC phase, forming a shell-like surface layer around an internal dendritic structure, Fig. 1. The wetting properties of the liquid to the structural phase and amount of liquid phase is essential for formation of a dense layer structure [7,8]. These parameters have to be sufficiently high to form a continuous infiltration front to fill the voids between the structural particles. Additionally, for rapid pore evacuation, a high infiltration rate of the liquid into the powder bed is needed. For a three dimensional porous media, the infiltration rate is dependent on the liquid–vapor surface tension and the liquid viscosity [8]. In the present research work, it was assumed that the dissolution of the TiC phase may have influenced the viscosity of the liquid metal and that primary re-crystallization of TiC at the front of the liquid phase decreased the penetration rate of the liquid into the powder bed, leading to porosity. During sintering, stresses develop in already sintered layers if the thermal expansion, induced by the temperature change, is hindered, either by external or internal constraints. In [17], thermal stresses in multiphase materials were discussed as a consequence of internal constraints caused by different coefficients of thermal expansion between the phases. If the thermal expansion is hindered by external constraints or uneven temperature distributions, bending stresses proportional to the average composite CTE, are generated [18]: L/2 6 σb = − 2 Eαt TX dx (2) L −(L/2) where L is the layer thickness, E is the elastic modulus, αt is the average CTE of the composite, T is the temperature difference and X is equal to x − L/2, x is the distance measured from the layer surface. According to the Turner model [19], the average CTE, αt , of a multi-phase composite is derived as follows: αt =
α1 V1 K1 + α2 V2 K2 + α3 V3 K3 · · · , V1 K1 + V 2 K2 + V 3 K3 · · ·
(3)
where V is the volume fraction, K is the bulk modulus and αi is the CTE of each phase. Following Eq. (2) the thermally induced bending stress depends on the average composite CTE, which according to Eq. (3) depends on the particular CTE of the material constituents. For the present material combination, the bending stress increases with increase of TiC content, due to higher CTE
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Fig. 3. XRD patterns of composites with 30, 60 and 80 wt.% TiC, respectively. Formation of BCC Fe–Ni–Ti–C phase with increasing TiC amount was observed.
of TiC compared to Invar36. But, as confirmed by SEM EDX, increasing total TiC amount led to dissolution of TiC and subsequent increase of Ti and C concentration in the Invar36 matrix, Table 2, resulting in considerable changes of the material properties. Dissolution of Ti and C in Invar36 significantly increases the alloys CTE [20], and has a solid solution strengthening effect, with possible loss in ductility as a consequence. Furthermore,
Fig. 4. Fracture surface obtained in bending test of Invar36-30 wt.% TiC composite. Ductile fracture of the Invar36 matrix, brittle fracture of TiC phase and delamination around spherical particles are illustrated.
Ti, being a strong BCC stabilizer [21], provided transformation of the matrix FCC phase to BCC phase, observed by XRD, Fig. 3. Matrix strengthening by multiple phases, could intensify the possible loss in ductility. Therefore, several factors participated in affecting the magnitude of the average composite CTE; the powder composition and the variation of the matrix CTE due to FCC to BCC phase transformation and enrichment of Ti and C. At high TiC quantities the influence of the matrix CTE is small, Eq. (3), but at the intermediate Invar36-60 wt.%TiC composition the effect is more significant. Hence, materials comprising 60 and 80 wt.% TiC phase, suffered of severe thermal cracking during sintering due to loss of low-CTE properties of the Invar36 matrix and the influence of TiC on the average composite CTE. Conventionally sintered WC–Co hard metals possess bending rupture strengths from 1700 MPa and hardness values from about 780 HV [22]. In the investigated crack free Invar3630 wt.%TiC material, mean bending rupture strength and hardness was about 900 MPa and 400 HV, respectively. The lower values could be attributed to low TiC amount, low matrix hardness and porosity. Crack propagation, in bending, through both phases, Fig. 2b, implies a strong bonding, pointing at sufficient wetting between the molten Invar36 metal and the TiC during sintering. However, the spherical particles observed on the fracture surfaces,
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Fig. 6. (a) Optical profilometry three-dimensional wear track of the Invar3630 wt.%TiC composite after 6000 m run. (b) Line profile taken along the white line in Fig. 6a.
Fig. 5. (a) Smeared iron oxide film, bright areas, on the worn surface of the Invar36-30 wt.%TiC composite after 6000 m run. (b) Abrasive wear caused by iron oxide particles.
Fig. 2b, indicates delamination and consequently a weak bonding between matrix and spheres. The wear coefficient in sliding was in the order of one magnitude higher than values obtained for TiC–NiMo composites, tested under the same conditions as in the present work, [15]. But on the contrary to [15], where adhesive wear was stated as the ascendant wear mechanism, abrasive wear was the dominating mechanism in the present work. A substantial amount of the TiC particles had a size of less than 1 m, and as illustrated in Fig. 5b, the width of the grooves were approximately 20 m. In multiphase materials, with a relatively low amount of reinforcement particles, the particle size should be of the same order as the abrasive particles to contribute to the total composite wear resistance, [23]. Hence, the wear resistance contribution of the finer TiC particles in the Invar-30 wt.%TiC material was negligible, which facilitated the abrasive wear mechanism. The relatively high binder phase content, led to frequent metal-metal contact, promoting strong adhesion between the mating tribo-surfaces. The smeared film of iron oxide, Fig. 5a, located as islands on the composite surface, is a possible consequence of metal-metal contact. Adhesive junctions may break within one of the mating surfaces, leading to a material transfer. Subsequently, the transferred material oxidizes, which is
promoted by the temperature increase associated with frictional heating during sliding. The transfer layer may come loose and be removed as a debris particle, or it can be reattached to one of the surfaces. In the present work, the latter phenomenon, with reattachment to the steel cylinder and two-body abrasion as a consequence, is a credible cause to the dominant abrasive wear mechanism observed. A second origin to the abrasive particles, are oxide particles originating from the steel cylinder, which was found severely oxidized during the experiment. 5. Conclusions DMLS of Invar36-TiC composites, led to formation of microstructures with homogeneous distributions of TiC particles in the metal matrix. However, melting of both metallic and ceramic phases occurred during sintering, which led to dissolution of Ti and C in the Invar36 matrix. Changes in chemical composition of the liquid Invar36 alloy led to spherical particle formation, changes in phase composition during solidification and loss of its low CTE properties. Consequently, with increasing TiC content, significant cracking of the specimen occurred during sintering, due to increased matrix CTE and transition from FCC to BCC phase. A maintained FCC structure and low amounts of dissolved TiC in the matrix, were realized only in the Invar36-30 wt.%TiC composite, where no cracks were observed. All sintered materials contained pores and the crack free Invar36-30 wt.% TiC composite corresponded to 10 vol.% porosity. Bending rupture strength and hardness of the crack
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free specimen were relatively low and could be attributed to low TiC content, low matrix hardness and porosity. The specific wear rate was 8 × 10−6 mm3 /Nm and SEM investigations concluded abrasive wear as the dominant wear mechanism. Acknowledgement The authors greatly acknowledge Mr. Jorgen Andersson at Uddeholm Tooling AB for support in X-ray investigations. References [1] K. Cooper, Rapid Prototyping Technology: Selection and Application, M. Dekker, New York, 2001. [2] D. Pham, Rapid Manufacturing, Springer, London, 2001. [3] D. King, T. Tansey, J. Mater. Process. Technol. 132 (2003) 42–48. [4] P. Jacobs, Society of Manufacturing Engineers, Dearborn Michigan, 1993. [5] X. Wang, T. Laoui, J. Bonse, J. Kruth, B. Lauwers, L. Froyen, Int. J. Adv. Manuf. Technol. 19 (2003) 351–357. [6] Y. Kathuria, Surf. Coat. Technol. 132 (2003) 42–48. [7] N. Tolochko, S. Mozzharov, T. Laoui, L. Froyen, Rapid Prototyping J. 9 (2003) 68.
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