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Nanocrystalline TiC-reinforced H13 steel matrix nanocomposites fabricated by selective laser melting
Materials and Design 96 (2016) 150–161
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Nanocrystalline TiC-reinforced H13 steel matrix nanocomposites fabricated by selective laser melting☆ Bandar AlMangour a,⁎, Dariusz Grzesiak b, Jenn-Ming Yang a a b
Department of Materials Science and Engineering, University of California Los Angeles, Los Angeles, CA 90095, USA Department of Mechanical Engineering and Mechatronics, West Pomeranian University of Technology, Szczecin, Poland
a r t i c l e
i n f o
Article history: Received 15 December 2015 Received in revised form 4 February 2016 Accepted 5 February 2016 Available online 7 February 2016 Keywords: Nanocomposite High-energy ball milling Selective laser melting Hardness Friction/wear
a b s t r a c t Additive manufacturing (AM) has a strong potential for the formation of a new class of multifunctional nanocomposites. In this study, nanocomposite feedstock powders were prepared by a mechanical alloying method based on high-energy ball milling. The evolution of constitutional phases and microstructural features of the milled powders was investigated as a function of milling time. The results showed that the milled powder particles experienced significant cold-welding during the entire milling time, with a wide size distribution. Selective laser melting (SLM), a promising AM fabrication technique, was applied to produce nanoscale 15% (by volume) TiC-reinforced H13 steel matrix nanocomposites. After SLM, uniformly dispersed nanoscale TiC particles with a mean particle size of 50 nm were obtained and a fine heterogeneous structure was observed. Relative to the unreinforced H13 steel part, the TiC/H13 steel nanocomposite parts exhibited higher hardness and elastic modulus, along with lower friction and a lower wear rate; these improvements are attributed to the combined effects of grain refinement and grain boundary strengthening. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction H13 steel type is a medium-carbon, hot work steel used to make tools for cutting, forming, or shaping materials [1]. Sufficient elastic strength, wear resistance, and high temperature stability are all necessary properties when selecting a tool steel [2, 3]. Many tools are subjected to extremely high loads that are applied rapidly and under high temperature gradients. Tools must withstand these loads repeatedly without breaking and without undergoing excessive wear or deformation. One way to overcome such problems is by incorporating hard second-phase carbides into the steel matrix. Among metal matrix composites, particulate-reinforced metal matrix nanocomposites (MMNCs) have received considerable attention because of their improved wear resistance, reduced cost, isotropic properties, and low density [4, 5]. Stable carbides coarsen more slowly than cementite and are therefore much more effective than cementite at higher temperatures. Along with the hard martensitic matrix, the addition of TiC, which is thermodynamically stable when in contact with the steel matrix [6], has been observed to greatly increase stiffness, hardness, and wear resistance [7, 8]. However, most conventional tool steels are wrought products (e.g., produced by casting) that often contain coarse, nonuniform microstructure (e.g., carbide segregation), accompanied by problems with hardness uniformity and poor transverse properties [9, 10]. It is also difficult to achieve a uniform distribution of nanoscale ☆ Part of this paper was presented at the EUROMAT 2015 congress in Warsaw, Poland. ⁎ Corresponding author.
http://dx.doi.org/10.1016/j.matdes.2016.02.022 0264-1275/© 2016 Elsevier Ltd. All rights reserved.
reinforcement through the steel matrix because of the large van der Waals attractive force between neighboring nanoparticles, which ultimately leads to agglomeration into coarse clusters. A recent novel processing method to fabricate tool steels involves powder metallurgy and laser-based additive manufacturing (AM) technology. AM refers to a class of technologies that can automatically construct physical parts directly from computer aided design (CAD) data [11, 12]. AM is potentially a powerful technology for customer-driven product development by reducing the product fabrication cycle time and the cost of tooling [13]; however, only a few materials have currently been developed for producing metal prototypes and tooling. Selective laser melting (SLM), as one of the AM processes, is a promising candidate to produce metal prototypes and tooling using metal-based powders [14]. The rapid solidification rate of the atomized powders eliminates phase segregation (i.e., promotes better chemical homogeneity) and produces a very fine microstructure with uniform size distribution of carbides and nonmetallic inclusions [15, 16]. This would improve the high-edge toughness of the tools, resulting in better cutting performance under difficult conditions such as interrupted cutting, where micro-chipping of the cutting edge can occur. Other advantages of AM tool steels compared to their wrought counterparts include superior machinability and dimensional control during heat treatment [17]. In addition, the alloying flexibility of the AM process may allow the production of new tool steels with special compositions that cannot be produced by conventional casting processes because of segregation-related hot-workability problems [18]. The higher hardness attainable with AM tool steels, along with their greater amount of
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carbides, constitutes a significant advantage over wrought tool steels. Wear resistance is largely determined by the volume content and distribution of the reinforcement carbides and is promoted by vanadium, molybdenum, and cobalt. These elements can be used in larger quantities in AM tool steels than in wrought steels without degrading the properties. A number of studies have been performed to evaluate the use of AM for fabricating H13 steel. Mazumder et al. [14] and Choi et al. [19] showed the feasibility of fabricating H13 steel by a direct metal deposition technique, and determined its characteristics, while Cormier et al. [20] performed microstructural analyses on H13 steel processed via electron beam melting. Childs et al. [21, 22] and Badrossamay et al. [23], used experimental and modeling approaches to study the influence of the processing parameters on the microstructural evolution of H13 steel fabricated by SLM. Sander et al. in their recent work [24] developed high-strength and dense tool steel processed by SLM. Kumar et al. [25] summarized the recent works of composite materials produced by different AM technologies and the methodologies of formation. Nevertheless, to the best of the authors' knowledge, there have been no comprehensive studies of the laser melting of TiC/H13 nanocomposites. The development of MMNCs by AM technology is still in its infancy; currently, no company or academic institution is able to process complex near-net shaped MMNC's by AM techniques to meet the requirements of user industries [26]. The physical and mechanical properties that can be obtained with TiC/H13 nanocomposites via SLM, discussed above, will make them very attractive candidate materials for future tooling applications. Therefore, in the light of the above-mentioned advantages, this work focuses on the processing of TiC/H13 nanocomposites by the SLM technique. Comprehensive microstructural observations were performed through scanning electron microscopy (SEM), electron backscatter diffraction (EBSD), and transmission electron microscopy (TEM), while the mechanical properties were assessed by hardness, friction and wear testing. 2. Experimental procedures 2.1. Powder preparation H13 steel powder was used as the metal matrix in this study. The H13 powder had a spherical shape and particle size distribution: d10 = 23.84 μm, d50 = 45.29 μm, d90 = 142 μm (Fig. 1A). The reinforcement powder was TiC nanopowder (99 +% purity) with a nearspherical shape and a mean particle size of 50 nm (Fig. 1B). The nanocomposite powders were prepared by a mechanical alloying method using high-energy ball milling (Pulverisette 4 vario-planetary ball mill) of H13 with 15% vol of TiC powder added. The ball-to-powder weight ratio was kept at 5:1, with a main disc rotation speed of
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200 rpm for milling times of 2, 4, 6, and 8 h. The balls and the inner surface of the mill were made of a stainless steel, and the milling process was carried out under a protective atmosphere of Ar. In order to avoid excessive temperature rise within the grinding bowl, there was one 15-min pause in the middle of each 2-h cycle. Clusters of nanoparticles resulting from insufficient mixing or electrostatic phenomena degrade the final microstructure and properties; therefore, it is expected that ball milling should enhance the distribution of the added nanoparticles, resulting in lower microporosity and fewer cracks. 2.2. SLM process The SLM parameters were fixed as follows: laser power, 100 W; spot size, 0.18 mm; scan speed, 250 mm/s; and hatch spacing, 0.12 mm. The powder layer thickness was only 0.05 mm, since higher layer thicknesses result in insufficient melting of the particles and lead to poor bonding between layers. The atmosphere inside the chamber was high purity argon to avoid any oxidation. Afterwards, nanocomposite powder that had been ball milled for 8 h was deposited on the substrate by a layering mechanism. The laser beam scanned the powder bed surface to form a layered profile according to the CAD data for the part, using a “cross-hatching” scanning method. The process was repeated in a layer-by-layer manner until multi-layer cylindrical parts with dimensions of 8 mm × 6 mm were constructed. 2.3. Microstructural observation Phase identification of the ball-milled powders and SLM-processed specimens was performed by X-ray diffraction (XRD) using a PANalytical X'Pert PRO X-ray powder diffractometer with Cu Kα radiation at 45 kV and 40 mA, using a continuous scan mode at 5°/min. Transverse and longitudinal sections of the fabricated samples were removed with a water-cooled cutoff wheel and then mounted on bakelite. The surfaces to be viewed were ground with progressively finer grits of silicon carbide paper and then polished with 1 and 0.3 μ alumina suspensions. Finally, the samples were etched with a 3% Nital solution for 50 s. Microstructures were observed using the Nova 230 SEM. The microstructure was further characterized by EBSD to study the crystal orientations, texture, and grain boundary misorientations for the nanocomposite sample from both the top and side views. A beam-scanning step-size of 40 nm was used and the EBSD data were collected using an EDAX system. The EBSD scans were performed at low magnification to improve the statistical reliability of the data. A clean-up procedure was applied to the EBSD images to remove rogue points. A Philips 420 TEM was also used at 120 kV to examine the internal microstructure of the nanocomposite samples that had been prepared using a focused-ion beam.
Fig. 1. Microstructures of the starting material. (A) H13 powder (the inset in the upper-right corner shows equiaxed grains of a single particle); (B) TiC powder.
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2.4. Hardness and wear behavior The Vickers hardness was measured on the polished sections of SLM parts using a microhardness tester (Leco, LM800AT) at a load of 200 g and an indentation time of 10 s for a minimum of 15 indentations. Also, nanoindentation tests were performed at room temperature, using an MST nanoindenter to measure the nanohardness and elastic modulus. The wear behavior of the fabricated samples was estimated by dry rotative wear tests in a T50 ball-on-disk tribometer at room temperature following the ASTM G99 standard. The surfaces of the specimens were ground and polished to a uniform standard surface finish prior to wear tests. The counter material was a 52100 bearing-steel ball with a diameter of 3 mm. The friction unit was rotated at a speed of 840 rpm for 20 min with a testing load of 3 N and a rotation radius of 2 mm. The coefficient of friction (COF) of the specimens was recorded during the wear tests. The volume of materials lost (V) was measured with the ST400 white light profilometer, which has a vertical resolution and accuracy of 8 and 80 nm respectively, and a lateral resolution of 2 μm. The wear rate (w) was calculated by: w ¼ FVL (mm3/Nm), where F is the contact load, and L is the total sliding distance. The morphologies of the worn surfaces were observed by SEM.
3. Results and discussion 3.1. Characterization of MA processed powders Fig. 2. XRD spectra of the powders. (a) Pure H13 before milling; (b), (c), (d), and (e) the mixture of H13 and 15% TiC after milling for 2, 4, 6, and 8 h, respectively.
3.1.1. Phase evolution Fig. 2 illustrates the XRD spectra of the H13 powder by itself, and also after mixing with 15% TiC for four different milling times. Before and
Fig. 3. SEM images showing characteristics of the milled 15% TiC/H13 nanocomposite powders after various milling times: (A) 2 h, (B) 4 h, (C) 6 h, (D) 8 h. Higher-magnification insets are included in the upper-left corners of (C) and (D) to show additional detail.
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Fig. 4. Size distribution of milled TiC/H13 nanocomposite powders after different milling times: (A) 2 h, (B) 4 h, (C) 6 h, (D) 8 h.
after milling, the strong diffraction peaks for α-Fe corresponding to H13 steel were clearly detected. After 2–8 h of milling, the TiC peaks were observed. Broadening of the peaks increased gradually with increased milling times because of the impact forces imposed on the powder particles by the grinding medium, which led to severe plastic deformation of the particles. The peak broadening indicates the formation of considerably smaller crystallites in the milled powders (Scherrer's equation). In addition to showing peak broadening, the α-Fe peaks generally shifted to lower angles with increased milling times, suggesting the generation of tensile stresses.
3.1.2. Microstructural development Fig. 3 shows the influence of the applied milling time on the morphology evolution of the TiC/H13 milled powders. Fig. 4 depicts the particle size distribution of the milled powders as a function of the milling time, revealing that the microstructural characteristics of the milled powders (e.g., particle shape, particle size and its distribution, and reinforcement dispersion state) were significantly influenced by the milling times. At the early stages of milling, i.e., the first two to four hours, the particles almost retained their spherical shape. As the milling time increased, part of the powder volume fractured intensively to form fine particles, resulting in a broad range of particle sizes. This is because a portion of the milled powders becomes work-hardened and has a limited ability to accept further plastic deformation, resulting in
Fig. 5. XRD spectra of SLM-processed materials. (a) Pure H13; and (b) 15% vol TiC/H13 nanocomposite.
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the absence of strong agglomerating forces [27]. The H13 particles tend to undergo plastic deformation, while the brittle TiC particles tend not to deform plastically. Overall, because of the severe plastic deformation experienced by powder particles during milling, particles tend to become cold-welded to each other, resulting in an increase in the average particle size. A comparative study of Fig. 3A and C reveals that when the milling time was increased to 6 h, a change occurred from a loose particle surface to a flattened and irregular morphology because of the competition between fracturing and cold-welding processes. Traces of fine TiC reinforcements could be easily seen on the rough surface of H13 particles, suggesting enhancement of the atomic diffusion effect and coldwelding mechanism. However, the high magnification image (corner of Fig. 3D) suggests that a longer milling time of 8 h might be optimal to ensure uniform dispersion of the TiC powder around the matrix. When the ball milling was prolonged to 8 h, a dynamic balance was eventually achieved between cold-welding, which tends to increase the particle size, and fracturing, which tends to decrease the mean
composite particle size. It was noted that the average particle size tends to increase with increasing milling time, but we predict that milling for more than 8 h may lead to further fracturing of the particles and a sharp decrease in the particle size [28]. Therefore, the above ball-milling times were carefully limited to 8 h or less in order to maintain a reasonable approximation of the original spherical structure of the H13 powder and thus obtain high flowability of the mixed powder during the SLM processing. 3.2. Characterization of SLM processed parts 3.2.1. Phases Fig. 5 presents the XRD patterns of SLM-processed pure H13 and of TiC/H13 nanocomposites. As expected, the fast cooling leads to a strong and complete martensite corresponding to an α-Fe phase. Similar observations were reported in some previous work [20], but other work [29] reported formation of retained austenite in addition to the martensitic phase. The difference could be due to the processing
Fig. 6. SEM images showing the characteristic morphologies of etched SLM-processed samples. (A) The top view of a pure H13 sample showing three distinct regions, denoted by “A,” “B,” and “C.” (B) A high-magnification top view of region “A”, revealing homogenous and fine equiaxed grains due to the rapid solidification. (C) A side view of pure H13 revealing elongated columnar grains toward the building direction (BD). (D) TiC/H13 nanocomposite illustrating the limited densification level. (E) TiC/H13 nanocomposite viewed from the top. (F) TiC/H13 nanocomposite viewed from the side.
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parameters—mainly the scanning speed that controls the solidification rate. Also, no peaks involving carbides were detected; this could be attributed to the relatively small amount of carbon (e.g., compared to M2 steel) [30]. After the addition of 15% TiC, diffraction peaks appear for the TiC-containing phase. This confirms the formation of TiCreinforced H13 based nanocomposites after SLM. The low intensity of the TiC peaks can be associated with the small initial size of the nanocrystallites and their further refinement after milling, according to Scherrer's formula [31]. 3.2.2. Microstructural observation by SEM The SEM observation of the pure H13 sample from the top reveals three distinct regions suggesting that the final microstructure was a thermally modified solidified structure (Fig. 6A). A continuous grain boundary with fine carbide networks is shown in region A, while region B is composed of cellular dendrite grain structures that grow along the direction of higher temperature gradient. Region C is similar to region
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B, but with a finer structure. The fine secondary dendrite cell spacing is indicative of high strength and hardness [32]. The difference in dendrite growth between regions “B” and “C” is due to the cyclic heat dissipation. The secondary-cell spacing of the dendrite depends on the laser scanning speed, preheating temperature, and the location of the region with respect to the scanning tracks. The cooling rates of the melting pool in the SLM process are generally dependent on the thermal gradients experienced in the build direction and the laser scan direction [33, 34]. The high solidification rate limits the complete growth of the dendrite structure and the formation of lath martensite [35]. The chemical concentration or temperature gradients in the molten pool may generate a surface tension gradient and resultant Marangoni convection, making the solidification a non-steady-state process. These result in a variety of crystal orientations with localized regularity. Fig. 6B is a high magnification image of region “A” in Fig. 6A, showing uniform and fine equiaxed grains with an average size of ~1.5 μm. It is expected that the temperature of the melting pool is high enough to
Fig. 7. (A,B) Unique grain color map of the 15% TiC/H13 nanocomposite; (C,D) grain boundary misorientation maps superimposed on image quality maps of the H13 matrix; (E,F) Kernel average misorientation (KAM) maps. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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dissolve chromium carbides phases (M7C3/M23C6), which are known to be easily soluble in tool steels at temperatures of ~ 900 °C [36]. In the side view (Fig. 6C), the additive manufacturing nature of SLM can be seen through the elongated columnar grains that grow along the building direction (BD). This epitaxial growth is based on re-melted grains on the previously solidified material, which is influenced by the cooling rate, thermal gradient, and growth velocity. With the addition of TiC particles, significant differences occur in the wettability and flowability of the liquid present in the melting pool, leading to some porosity and cracks (Fig. 6D). During SLM, the TiC/ H13 nanocomposite powder is melted line-by-line via the laser beam, forming a continuous and mobile molten pool. The amount of the liquid formed in the molten pool influences thermo-capillary characteristics such as the viscosity, wettability, and liquid-solid rheological properties. The presence of TiC nanoparticles in the H13 matrix liquid tends to increase the viscosity of the melt significantly, limiting the flow and decreasing the overall rheological performance of the composite melt. It has also been reported that the relatively high C content of the H13 steel (compared to 316 L stainless steel, for example) leads to severe balling of the molten pool [37], which interferes with the processing. The high melting viscosity and limited wetting characteristics caused by insufficient laser energy are the key factors in producing the balling effect and interlayer pores, thereby decreasing the densification level of SLM-processed TiC/H13 nanocomposite parts. The average relative
density for the nanocomposite is ~91.1% of bulk in comparison to that of the pure H13 with ~95.46% of bulk. Further optimization of the SLM processing parameters is necessary in order to reduce the porosity level as was shown by the study of Read et al. [38]. Fig. 6E and f are high-magnification images of the TiC/H13 nanocomposite, viewed from the top and side, respectively. A series of TiC nanoparticles formed uniformly in the H13 matrix. Their homogenous and interconnected distribution must be attributed to the strong convection caused by the rapid heating and solidification [39]. Interestingly, the side view (Fig. 6F) did not reveal a clear layered microstructure but showed a microstructure that was similar to the top view: heterogeneous and continuous grain boundaries rich with TiC reinforcement. Yuan et al. [40] found out that increasing the energy density up to a certain limit enhances the distribution of the TiC reinforcements within the matrix. 3.2.3. Microstructural observation by EBSD Fig. 7A and B show the unique grain color maps, obtained by EBSD, from the top and side views, respectively, of the TiC/H13 nanocomposite. The top view shows a structure of fine equiaxed grains oriented parallel to the scanning direction (SD), with an average size ~ 1 μm as a result of the rapid solidification that was discussed earlier. The microstructure when viewed from the side, shows epitaxial columnar grains oriented along the SLM building direction. These grains formed during the solidification of the previous layer elongated toward the building
Fig. 8. (A,B) Crystal orientation maps of the 15% TiC/H13 nanocomposite; (C,D) Inverse pole figures for the α-Fe phase (top view) indicating random texture; (E,F) Pole figures for the α-Fe phase (top view).
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Fig. 9. TEM images for the 15% TiC/H13 nanocomposite at two different regions revealing (A) spherical TiC precipitates at the interface; and (B) an increased dislocation density in the H13 steel matrix. The single arrow in (A) indicates a grain boundary; the triple arrows indicate TiC precipitates near a grain boundary. The arrows in (B) indicate the TiC particle size and shape; (C) selected area diffraction pattern (SADP) showing the existence of TiC and BCC-Fe constitution phases.
direction of heat conduction through the laser scan with the same orientation. Rotation of the scanning direction by 90° usually breaks up the defined epitaxial columnar structure [33]. It was also noticed that the presence of unmelted particles, porosity, or TiC nanoparticles results in a breakup of the epitaxial growth.
The distribution of grain boundary misorientation angles are presented in Fig. 7C and D. Clearly, the grains from the top and side views were separated by a majority of high-angle grain boundaries distribution ([N 15°] [blue color]). This could be due to the high scanning speed implemented in this study, which leads to elongation of the
Table 1 Vickers and nanohardness of SLM-processed samples.
Vicker's and nanohardness (GPa) Elastic modulus (GPa) Anisotropy (%)
Pure H13 (top view)
Pure H13 (side view)
Nanocomposite (top view)
Nanocomposite (side view)
748.04 ± 27.61 (8.36 ± 0.17) 230.71 ± 8.41 ~8.35
810.51 ± 40.11 (8.56 ± 0.11) 236 ± 5.5
811.23 ± 44.18 (9.41 ± 0.30) 255.83 ± 8.14 ~5.82
858.45 ± 36.56 (9.71 ± 0.24) 259 ± 6.47
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crystallographic orientation may be deviated because of repeated solidification and annealing in the narrow-scanned builds.
Table 2 Coefficient of friction (COF) and wear rates of SLM-processed samples.
Average COF Max COF Wear rate (mm3/Nm)
Pure H13
Nanocomposite
0.550 0.989 4.906 × 10−6
0.526 0.822 3.623 × 10−6
molten pool [41]. The small zones containing low angle boundaries (2–5°) that are visible at the borders of the tracks are probably not grain boundaries, but rather sub-grain boundaries formed by dislocation rearrangement. The Kernel average misorientation (KAM), shown in Fig. 7E and F, calculates the average misorientation between a data point and all of its neighbors. No significant differences were noted between the side view and the top view. The higher KAM portion in the grains (i.e., green color) indicates a higher degree of misorientation, which implies high stored energy in the grain. Generally, KAM is high in deformed grains because of higher dislocation density. The crystal orientation maps shown in Fig. 8A and B, demonstrate clear random grain orientations. This was expected because the “cross-hatching” scanning strategy (i.e., the scanning direction is altered by 90° between consecutive layers) results in a reduction of texture along the build and scanning directions, which minimizes the anisotropic properties of such structures. This agrees with the previous observation reported by Boegelein and co-workers [42]. However, looking closely at the inverse pole figures, shown in Fig. 8C and D, a preferential orientation of grain growth in the 〈100〉 direction is observed, which is favored for cubic crystals [43]. When comparing the side view of the SLM build with the top view from the pole figure maps, shown in Fig. 8E and F, one can see some differences in the preferential alignment. This was expected because the same area was melted several times as a result of the narrow scan spacing that was used. Therefore, the
3.2.4. Microstructural observation by TEM Fig. 9 shows the bright-field TEM images of the SLM-processed nanocomposite with selected area diffraction pattern (SADP). Fig. 9A presents the H13 matrix grain boundary with a “metal-ceramic” interface in which the H13 matrix has wet the TiC particles, and small TiC spherical precipitates have formed in the side of the interface (indicated by the triple arrows). Bolton et al. [44] reported minor diffusion of iron from the steel matrix into the TiC particles (limited solubility), which may have caused the formation of Fe-rich precipitates at the interface. It is generally believed that a very thin precipitate-rich interface layer (b200 nm) can be beneficial for the strength of the interface. The reinforcements were not just located at the grain boundaries, which hinders the grain growth at elevated temperature, but also within the grain (indicated by the single arrow). More importantly, TiC nanoparticles with a particle size ~ 50 nm were uniformly dispersed in the grain boundaries of the H13 matrix retaining its size and initial spherical shape (Fig. 9B). The high solidification rate ~106–108 K/s [45], inhibits the grain growth of the TiC nanoparticles because of insufficient time for grain coarsening, hence remaining as the favorable nanostructured TiC reinforcement after SLM. The sample exhibited a high level of dislocations (Fig. 9B), as a result of the fast solidification rate and also as a result of the differences in coefficient of thermal expansion between the matrix and the TiC reinforcement. The TiC nanoparticles act as obstacles to dislocation glide and grain growth. The SADP taken at the grain boundaries of the nanocomposite sample revealed the existence of TiC and BCC-Fe constitution phases (Fig. 9C). No evidence of any amorphous regions was found anywhere within the sample. To confirm that, we tilted the regions that we
Fig. 10. SEM images showing morphologies of the worn surfaces of the nanocomposites. (A,C) Pure H13; (BD) 15% TiC/H13.
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believed might be amorphous, but the tilting of the regions showed a contrast change, suggesting full crystallinity.
3.3. Hardness and wear behavior The effects of TiC reinforcement on the micro- and nano-hardness of SLM-processed H13 steel and nanocomposite are shown in Table 1. The SLM-processed samples demonstrate superior hardness compared to conventionally processed H13 (e.g., casting), with a typical hardness of ~750 HV0.2. This high initial hardness is attributed to the significant grain refinement effect resulting from laser rapid solidification and the presence of a completely martensitic phase that favors high hardness. With the addition of 15% vol TiC, the hardness and elastic modulus were considerably improved by ~ 15% (rule of mixtures), because of the homogenous incorporation of nanoscale TiC reinforcing particles throughout the matrix. The SLM-processed parts are relatively isotropic, as would be expected from the EBSD results, indicating that the orientations of the grains were random and not associated with any preferred directions. Similar results were obtained by Walker et al. [46] who showed that a “cross-hatching” scanning strategy minimizes the anisotropic properties.
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The variation of friction coefficients and wear rates of the pure H13 and TiC/H13 are given in Table 2. Compared to the unreinforced H13 steel, the coefficient of friction of TiC/H13 nanocomposite decreased, and the resultant wear rate was lowered by ~ 26%. Fig. 10 shows the morphology of the worn surfaces of the pure H13 and the nanocomposites. The material loss is believed to result from the removal of material chips from the matrix phase because of micro-cutting by the abrasive particles [7]. The presence of irregularly shaped fragments reveals that local deformation and delamination of the worn surface occurred. With the addition of TiC nanoparticles, the mechanism of material removal during sliding was changed from partial abrasion or a localized tribolayer to a further strain-hardened tribolayer (i.e., shallower grooves and less abrasive fragments). Clearly, the strain-induced tribolayer portion was increased after the incorporation of TiC nanoparticles (Fig. 10C, D). Such a transition is expected to reduce the wear volume after the wear tests, although there was still some minor microplowing. No particle cracking or separation of particles from the matrix can be observed. The improvement in wear performance can thus be attributed to the uniformly dispersed TiC nanoparticle reinforcement and low interfacial stress with the matrix, leading to less breakage of the nanocomposites during sliding. Moreover, the improvement in
Fig. 11. SEM micrographs of the worn surface of the nanocomposites, and the corresponding elemental maps. (a) Pure H13; (b) 15% TiC/H13.
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hardness (Table 1) helped to increase the area of the strain-hardened tribolayer, which prevented further wearing. Ghosh et al. [47] have shown that when the size of reinforcement decreases the wear resistance of the composite improves, with no further enhancement in wear resistance after 20 vol%. Elemental maps of the worn surfaces of the pure steel and the nanocomposites are shown in Fig. 11a and b, respectively. The elemental mapping of the pure steel shows that the worn surface was rich with iron, carbon, and chromium. For the worn surface of the nanocomposite, the map indicates the presence of more iron, but less carbon, suggesting that a small amount of TiC was lost. The TiC particles were present on the worn surface; thus, they protected the matrix during the wear test. The only damage on the surface seems to be some pits from spalled M7C3 carbides and the corresponding mini-grooves produced by them. No presence of oxides after sliding was seen, suggesting that there was no significant temperature rise at the sample surface. Some fine reinforcing particles were observed at the edges of the grooves on the worn surface. 4. Conclusion 1. The mechanical intermixing caused by the high-energy ball milling produces H13 steel reinforced with 15% vol TiC nanocomposite powder. With increased milling time, the powder, which initially had a spherical shape, underwent successive changes in its morphology and particle size distribution. This microstructural evolution is determined by the competition between cold-welding and fracturing of the particles. 2. The microstructure of the SLM-processed nanocomposite viewed from the top reveals equiaxed fine grains as a result of rapid solidification, while the side view clearly shows elongated columnar grains caused by the additive nature of the SLM process. 3. Overall, it was shown that the “cross-hatching” scanning strategy (i.e., the scanning direction is altered 90° between consecutive layers) results in a reduction of texture along the build and scanning directions, which consequently minimizes the anisotropic properties of such structures. 4. In comparison to the unreinforced H13 steel, the average hardness and elastic modulus of the nanocomposite part increased, resulting in enhanced wear resistance. 5. The benefits of AM can be explained by the changes in the microstructure and mechanical properties when compared to cast materials. With a better understanding of these differences, the technology can be further developed and AM methods will be more commonly used. Acknowledgment One of the authors, Bandar AlMangour, gratefully appreciates the financial support from the Saudi Arabia Basic Industries Corporation (SABIC). References [1] J.F. Shackelford, W. Alexander, CRC Materials Science and Engineering Handbook, CRC Press, 2015. [2] G. Krauss, Steels: Processing, Structure, and Performance, ASM International, 2015. [3] S. Hashmi, Comprehensive Materials Processing, Elsevier, 2014. [4] I. Ibrahim, F. Mohamed, E. Lavernia, Particulate reinforced metal matrix composites—a review, J. Mater. Sci. 26 (1991) 1137–1156. [5] S.C. Tjong, Novel nanoparticle-reinforced metal matrix composites with enhanced mechanical properties, Adv. Eng. Mater. 9 (2007) 639–652. [6] E. Pagounis, V. Lindroos, M. Talvitie, Influence of reinforcement volume fraction and size on the microstructure and abrasion wear resistance of hot isostatic pressed white iron matrix composites, Metall. Mater. Trans. A 27 (1996) 4171–4181. [7] F. Akhtar, Microstructure evolution and wear properties of in situ synthesized TiB2 and TiC reinforced steel matrix composites, J. Alloys Compd. 459 (2008) 491–497. [8] W. Jiang, P. Molian, Nanocrystalline TiC powder alloying and glazing of H13 steel using a CO2 laser for improved life of die-casting dies, Surf. Coat. Technol. 135 (2001) 139–149. [9] J. Campbell, Complete Casting Handbook: Metal Casting Processes, Metallurgy, Techniques and Design, Butterworth-Heinemann, 2015.
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Nanocrystalline TiC-reinforced H13 steel matrix nanocomposites fabricated by selective laser melting☆ Bandar AlMangour a,⁎, Dariusz Grzesiak b, Jenn-Ming Yang a a b
Department of Materials Science and Engineering, University of California Los Angeles, Los Angeles, CA 90095, USA Department of Mechanical Engineering and Mechatronics, West Pomeranian University of Technology, Szczecin, Poland
a r t i c l e
i n f o
Article history: Received 15 December 2015 Received in revised form 4 February 2016 Accepted 5 February 2016 Available online 7 February 2016 Keywords: Nanocomposite High-energy ball milling Selective laser melting Hardness Friction/wear
a b s t r a c t Additive manufacturing (AM) has a strong potential for the formation of a new class of multifunctional nanocomposites. In this study, nanocomposite feedstock powders were prepared by a mechanical alloying method based on high-energy ball milling. The evolution of constitutional phases and microstructural features of the milled powders was investigated as a function of milling time. The results showed that the milled powder particles experienced significant cold-welding during the entire milling time, with a wide size distribution. Selective laser melting (SLM), a promising AM fabrication technique, was applied to produce nanoscale 15% (by volume) TiC-reinforced H13 steel matrix nanocomposites. After SLM, uniformly dispersed nanoscale TiC particles with a mean particle size of 50 nm were obtained and a fine heterogeneous structure was observed. Relative to the unreinforced H13 steel part, the TiC/H13 steel nanocomposite parts exhibited higher hardness and elastic modulus, along with lower friction and a lower wear rate; these improvements are attributed to the combined effects of grain refinement and grain boundary strengthening. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction H13 steel type is a medium-carbon, hot work steel used to make tools for cutting, forming, or shaping materials [1]. Sufficient elastic strength, wear resistance, and high temperature stability are all necessary properties when selecting a tool steel [2, 3]. Many tools are subjected to extremely high loads that are applied rapidly and under high temperature gradients. Tools must withstand these loads repeatedly without breaking and without undergoing excessive wear or deformation. One way to overcome such problems is by incorporating hard second-phase carbides into the steel matrix. Among metal matrix composites, particulate-reinforced metal matrix nanocomposites (MMNCs) have received considerable attention because of their improved wear resistance, reduced cost, isotropic properties, and low density [4, 5]. Stable carbides coarsen more slowly than cementite and are therefore much more effective than cementite at higher temperatures. Along with the hard martensitic matrix, the addition of TiC, which is thermodynamically stable when in contact with the steel matrix [6], has been observed to greatly increase stiffness, hardness, and wear resistance [7, 8]. However, most conventional tool steels are wrought products (e.g., produced by casting) that often contain coarse, nonuniform microstructure (e.g., carbide segregation), accompanied by problems with hardness uniformity and poor transverse properties [9, 10]. It is also difficult to achieve a uniform distribution of nanoscale ☆ Part of this paper was presented at the EUROMAT 2015 congress in Warsaw, Poland. ⁎ Corresponding author.
http://dx.doi.org/10.1016/j.matdes.2016.02.022 0264-1275/© 2016 Elsevier Ltd. All rights reserved.
reinforcement through the steel matrix because of the large van der Waals attractive force between neighboring nanoparticles, which ultimately leads to agglomeration into coarse clusters. A recent novel processing method to fabricate tool steels involves powder metallurgy and laser-based additive manufacturing (AM) technology. AM refers to a class of technologies that can automatically construct physical parts directly from computer aided design (CAD) data [11, 12]. AM is potentially a powerful technology for customer-driven product development by reducing the product fabrication cycle time and the cost of tooling [13]; however, only a few materials have currently been developed for producing metal prototypes and tooling. Selective laser melting (SLM), as one of the AM processes, is a promising candidate to produce metal prototypes and tooling using metal-based powders [14]. The rapid solidification rate of the atomized powders eliminates phase segregation (i.e., promotes better chemical homogeneity) and produces a very fine microstructure with uniform size distribution of carbides and nonmetallic inclusions [15, 16]. This would improve the high-edge toughness of the tools, resulting in better cutting performance under difficult conditions such as interrupted cutting, where micro-chipping of the cutting edge can occur. Other advantages of AM tool steels compared to their wrought counterparts include superior machinability and dimensional control during heat treatment [17]. In addition, the alloying flexibility of the AM process may allow the production of new tool steels with special compositions that cannot be produced by conventional casting processes because of segregation-related hot-workability problems [18]. The higher hardness attainable with AM tool steels, along with their greater amount of
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carbides, constitutes a significant advantage over wrought tool steels. Wear resistance is largely determined by the volume content and distribution of the reinforcement carbides and is promoted by vanadium, molybdenum, and cobalt. These elements can be used in larger quantities in AM tool steels than in wrought steels without degrading the properties. A number of studies have been performed to evaluate the use of AM for fabricating H13 steel. Mazumder et al. [14] and Choi et al. [19] showed the feasibility of fabricating H13 steel by a direct metal deposition technique, and determined its characteristics, while Cormier et al. [20] performed microstructural analyses on H13 steel processed via electron beam melting. Childs et al. [21, 22] and Badrossamay et al. [23], used experimental and modeling approaches to study the influence of the processing parameters on the microstructural evolution of H13 steel fabricated by SLM. Sander et al. in their recent work [24] developed high-strength and dense tool steel processed by SLM. Kumar et al. [25] summarized the recent works of composite materials produced by different AM technologies and the methodologies of formation. Nevertheless, to the best of the authors' knowledge, there have been no comprehensive studies of the laser melting of TiC/H13 nanocomposites. The development of MMNCs by AM technology is still in its infancy; currently, no company or academic institution is able to process complex near-net shaped MMNC's by AM techniques to meet the requirements of user industries [26]. The physical and mechanical properties that can be obtained with TiC/H13 nanocomposites via SLM, discussed above, will make them very attractive candidate materials for future tooling applications. Therefore, in the light of the above-mentioned advantages, this work focuses on the processing of TiC/H13 nanocomposites by the SLM technique. Comprehensive microstructural observations were performed through scanning electron microscopy (SEM), electron backscatter diffraction (EBSD), and transmission electron microscopy (TEM), while the mechanical properties were assessed by hardness, friction and wear testing. 2. Experimental procedures 2.1. Powder preparation H13 steel powder was used as the metal matrix in this study. The H13 powder had a spherical shape and particle size distribution: d10 = 23.84 μm, d50 = 45.29 μm, d90 = 142 μm (Fig. 1A). The reinforcement powder was TiC nanopowder (99 +% purity) with a nearspherical shape and a mean particle size of 50 nm (Fig. 1B). The nanocomposite powders were prepared by a mechanical alloying method using high-energy ball milling (Pulverisette 4 vario-planetary ball mill) of H13 with 15% vol of TiC powder added. The ball-to-powder weight ratio was kept at 5:1, with a main disc rotation speed of
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200 rpm for milling times of 2, 4, 6, and 8 h. The balls and the inner surface of the mill were made of a stainless steel, and the milling process was carried out under a protective atmosphere of Ar. In order to avoid excessive temperature rise within the grinding bowl, there was one 15-min pause in the middle of each 2-h cycle. Clusters of nanoparticles resulting from insufficient mixing or electrostatic phenomena degrade the final microstructure and properties; therefore, it is expected that ball milling should enhance the distribution of the added nanoparticles, resulting in lower microporosity and fewer cracks. 2.2. SLM process The SLM parameters were fixed as follows: laser power, 100 W; spot size, 0.18 mm; scan speed, 250 mm/s; and hatch spacing, 0.12 mm. The powder layer thickness was only 0.05 mm, since higher layer thicknesses result in insufficient melting of the particles and lead to poor bonding between layers. The atmosphere inside the chamber was high purity argon to avoid any oxidation. Afterwards, nanocomposite powder that had been ball milled for 8 h was deposited on the substrate by a layering mechanism. The laser beam scanned the powder bed surface to form a layered profile according to the CAD data for the part, using a “cross-hatching” scanning method. The process was repeated in a layer-by-layer manner until multi-layer cylindrical parts with dimensions of 8 mm × 6 mm were constructed. 2.3. Microstructural observation Phase identification of the ball-milled powders and SLM-processed specimens was performed by X-ray diffraction (XRD) using a PANalytical X'Pert PRO X-ray powder diffractometer with Cu Kα radiation at 45 kV and 40 mA, using a continuous scan mode at 5°/min. Transverse and longitudinal sections of the fabricated samples were removed with a water-cooled cutoff wheel and then mounted on bakelite. The surfaces to be viewed were ground with progressively finer grits of silicon carbide paper and then polished with 1 and 0.3 μ alumina suspensions. Finally, the samples were etched with a 3% Nital solution for 50 s. Microstructures were observed using the Nova 230 SEM. The microstructure was further characterized by EBSD to study the crystal orientations, texture, and grain boundary misorientations for the nanocomposite sample from both the top and side views. A beam-scanning step-size of 40 nm was used and the EBSD data were collected using an EDAX system. The EBSD scans were performed at low magnification to improve the statistical reliability of the data. A clean-up procedure was applied to the EBSD images to remove rogue points. A Philips 420 TEM was also used at 120 kV to examine the internal microstructure of the nanocomposite samples that had been prepared using a focused-ion beam.
Fig. 1. Microstructures of the starting material. (A) H13 powder (the inset in the upper-right corner shows equiaxed grains of a single particle); (B) TiC powder.
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2.4. Hardness and wear behavior The Vickers hardness was measured on the polished sections of SLM parts using a microhardness tester (Leco, LM800AT) at a load of 200 g and an indentation time of 10 s for a minimum of 15 indentations. Also, nanoindentation tests were performed at room temperature, using an MST nanoindenter to measure the nanohardness and elastic modulus. The wear behavior of the fabricated samples was estimated by dry rotative wear tests in a T50 ball-on-disk tribometer at room temperature following the ASTM G99 standard. The surfaces of the specimens were ground and polished to a uniform standard surface finish prior to wear tests. The counter material was a 52100 bearing-steel ball with a diameter of 3 mm. The friction unit was rotated at a speed of 840 rpm for 20 min with a testing load of 3 N and a rotation radius of 2 mm. The coefficient of friction (COF) of the specimens was recorded during the wear tests. The volume of materials lost (V) was measured with the ST400 white light profilometer, which has a vertical resolution and accuracy of 8 and 80 nm respectively, and a lateral resolution of 2 μm. The wear rate (w) was calculated by: w ¼ FVL (mm3/Nm), where F is the contact load, and L is the total sliding distance. The morphologies of the worn surfaces were observed by SEM.
3. Results and discussion 3.1. Characterization of MA processed powders Fig. 2. XRD spectra of the powders. (a) Pure H13 before milling; (b), (c), (d), and (e) the mixture of H13 and 15% TiC after milling for 2, 4, 6, and 8 h, respectively.
3.1.1. Phase evolution Fig. 2 illustrates the XRD spectra of the H13 powder by itself, and also after mixing with 15% TiC for four different milling times. Before and
Fig. 3. SEM images showing characteristics of the milled 15% TiC/H13 nanocomposite powders after various milling times: (A) 2 h, (B) 4 h, (C) 6 h, (D) 8 h. Higher-magnification insets are included in the upper-left corners of (C) and (D) to show additional detail.
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Fig. 4. Size distribution of milled TiC/H13 nanocomposite powders after different milling times: (A) 2 h, (B) 4 h, (C) 6 h, (D) 8 h.
after milling, the strong diffraction peaks for α-Fe corresponding to H13 steel were clearly detected. After 2–8 h of milling, the TiC peaks were observed. Broadening of the peaks increased gradually with increased milling times because of the impact forces imposed on the powder particles by the grinding medium, which led to severe plastic deformation of the particles. The peak broadening indicates the formation of considerably smaller crystallites in the milled powders (Scherrer's equation). In addition to showing peak broadening, the α-Fe peaks generally shifted to lower angles with increased milling times, suggesting the generation of tensile stresses.
3.1.2. Microstructural development Fig. 3 shows the influence of the applied milling time on the morphology evolution of the TiC/H13 milled powders. Fig. 4 depicts the particle size distribution of the milled powders as a function of the milling time, revealing that the microstructural characteristics of the milled powders (e.g., particle shape, particle size and its distribution, and reinforcement dispersion state) were significantly influenced by the milling times. At the early stages of milling, i.e., the first two to four hours, the particles almost retained their spherical shape. As the milling time increased, part of the powder volume fractured intensively to form fine particles, resulting in a broad range of particle sizes. This is because a portion of the milled powders becomes work-hardened and has a limited ability to accept further plastic deformation, resulting in
Fig. 5. XRD spectra of SLM-processed materials. (a) Pure H13; and (b) 15% vol TiC/H13 nanocomposite.
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the absence of strong agglomerating forces [27]. The H13 particles tend to undergo plastic deformation, while the brittle TiC particles tend not to deform plastically. Overall, because of the severe plastic deformation experienced by powder particles during milling, particles tend to become cold-welded to each other, resulting in an increase in the average particle size. A comparative study of Fig. 3A and C reveals that when the milling time was increased to 6 h, a change occurred from a loose particle surface to a flattened and irregular morphology because of the competition between fracturing and cold-welding processes. Traces of fine TiC reinforcements could be easily seen on the rough surface of H13 particles, suggesting enhancement of the atomic diffusion effect and coldwelding mechanism. However, the high magnification image (corner of Fig. 3D) suggests that a longer milling time of 8 h might be optimal to ensure uniform dispersion of the TiC powder around the matrix. When the ball milling was prolonged to 8 h, a dynamic balance was eventually achieved between cold-welding, which tends to increase the particle size, and fracturing, which tends to decrease the mean
composite particle size. It was noted that the average particle size tends to increase with increasing milling time, but we predict that milling for more than 8 h may lead to further fracturing of the particles and a sharp decrease in the particle size [28]. Therefore, the above ball-milling times were carefully limited to 8 h or less in order to maintain a reasonable approximation of the original spherical structure of the H13 powder and thus obtain high flowability of the mixed powder during the SLM processing. 3.2. Characterization of SLM processed parts 3.2.1. Phases Fig. 5 presents the XRD patterns of SLM-processed pure H13 and of TiC/H13 nanocomposites. As expected, the fast cooling leads to a strong and complete martensite corresponding to an α-Fe phase. Similar observations were reported in some previous work [20], but other work [29] reported formation of retained austenite in addition to the martensitic phase. The difference could be due to the processing
Fig. 6. SEM images showing the characteristic morphologies of etched SLM-processed samples. (A) The top view of a pure H13 sample showing three distinct regions, denoted by “A,” “B,” and “C.” (B) A high-magnification top view of region “A”, revealing homogenous and fine equiaxed grains due to the rapid solidification. (C) A side view of pure H13 revealing elongated columnar grains toward the building direction (BD). (D) TiC/H13 nanocomposite illustrating the limited densification level. (E) TiC/H13 nanocomposite viewed from the top. (F) TiC/H13 nanocomposite viewed from the side.
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parameters—mainly the scanning speed that controls the solidification rate. Also, no peaks involving carbides were detected; this could be attributed to the relatively small amount of carbon (e.g., compared to M2 steel) [30]. After the addition of 15% TiC, diffraction peaks appear for the TiC-containing phase. This confirms the formation of TiCreinforced H13 based nanocomposites after SLM. The low intensity of the TiC peaks can be associated with the small initial size of the nanocrystallites and their further refinement after milling, according to Scherrer's formula [31]. 3.2.2. Microstructural observation by SEM The SEM observation of the pure H13 sample from the top reveals three distinct regions suggesting that the final microstructure was a thermally modified solidified structure (Fig. 6A). A continuous grain boundary with fine carbide networks is shown in region A, while region B is composed of cellular dendrite grain structures that grow along the direction of higher temperature gradient. Region C is similar to region
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B, but with a finer structure. The fine secondary dendrite cell spacing is indicative of high strength and hardness [32]. The difference in dendrite growth between regions “B” and “C” is due to the cyclic heat dissipation. The secondary-cell spacing of the dendrite depends on the laser scanning speed, preheating temperature, and the location of the region with respect to the scanning tracks. The cooling rates of the melting pool in the SLM process are generally dependent on the thermal gradients experienced in the build direction and the laser scan direction [33, 34]. The high solidification rate limits the complete growth of the dendrite structure and the formation of lath martensite [35]. The chemical concentration or temperature gradients in the molten pool may generate a surface tension gradient and resultant Marangoni convection, making the solidification a non-steady-state process. These result in a variety of crystal orientations with localized regularity. Fig. 6B is a high magnification image of region “A” in Fig. 6A, showing uniform and fine equiaxed grains with an average size of ~1.5 μm. It is expected that the temperature of the melting pool is high enough to
Fig. 7. (A,B) Unique grain color map of the 15% TiC/H13 nanocomposite; (C,D) grain boundary misorientation maps superimposed on image quality maps of the H13 matrix; (E,F) Kernel average misorientation (KAM) maps. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
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dissolve chromium carbides phases (M7C3/M23C6), which are known to be easily soluble in tool steels at temperatures of ~ 900 °C [36]. In the side view (Fig. 6C), the additive manufacturing nature of SLM can be seen through the elongated columnar grains that grow along the building direction (BD). This epitaxial growth is based on re-melted grains on the previously solidified material, which is influenced by the cooling rate, thermal gradient, and growth velocity. With the addition of TiC particles, significant differences occur in the wettability and flowability of the liquid present in the melting pool, leading to some porosity and cracks (Fig. 6D). During SLM, the TiC/ H13 nanocomposite powder is melted line-by-line via the laser beam, forming a continuous and mobile molten pool. The amount of the liquid formed in the molten pool influences thermo-capillary characteristics such as the viscosity, wettability, and liquid-solid rheological properties. The presence of TiC nanoparticles in the H13 matrix liquid tends to increase the viscosity of the melt significantly, limiting the flow and decreasing the overall rheological performance of the composite melt. It has also been reported that the relatively high C content of the H13 steel (compared to 316 L stainless steel, for example) leads to severe balling of the molten pool [37], which interferes with the processing. The high melting viscosity and limited wetting characteristics caused by insufficient laser energy are the key factors in producing the balling effect and interlayer pores, thereby decreasing the densification level of SLM-processed TiC/H13 nanocomposite parts. The average relative
density for the nanocomposite is ~91.1% of bulk in comparison to that of the pure H13 with ~95.46% of bulk. Further optimization of the SLM processing parameters is necessary in order to reduce the porosity level as was shown by the study of Read et al. [38]. Fig. 6E and f are high-magnification images of the TiC/H13 nanocomposite, viewed from the top and side, respectively. A series of TiC nanoparticles formed uniformly in the H13 matrix. Their homogenous and interconnected distribution must be attributed to the strong convection caused by the rapid heating and solidification [39]. Interestingly, the side view (Fig. 6F) did not reveal a clear layered microstructure but showed a microstructure that was similar to the top view: heterogeneous and continuous grain boundaries rich with TiC reinforcement. Yuan et al. [40] found out that increasing the energy density up to a certain limit enhances the distribution of the TiC reinforcements within the matrix. 3.2.3. Microstructural observation by EBSD Fig. 7A and B show the unique grain color maps, obtained by EBSD, from the top and side views, respectively, of the TiC/H13 nanocomposite. The top view shows a structure of fine equiaxed grains oriented parallel to the scanning direction (SD), with an average size ~ 1 μm as a result of the rapid solidification that was discussed earlier. The microstructure when viewed from the side, shows epitaxial columnar grains oriented along the SLM building direction. These grains formed during the solidification of the previous layer elongated toward the building
Fig. 8. (A,B) Crystal orientation maps of the 15% TiC/H13 nanocomposite; (C,D) Inverse pole figures for the α-Fe phase (top view) indicating random texture; (E,F) Pole figures for the α-Fe phase (top view).
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Fig. 9. TEM images for the 15% TiC/H13 nanocomposite at two different regions revealing (A) spherical TiC precipitates at the interface; and (B) an increased dislocation density in the H13 steel matrix. The single arrow in (A) indicates a grain boundary; the triple arrows indicate TiC precipitates near a grain boundary. The arrows in (B) indicate the TiC particle size and shape; (C) selected area diffraction pattern (SADP) showing the existence of TiC and BCC-Fe constitution phases.
direction of heat conduction through the laser scan with the same orientation. Rotation of the scanning direction by 90° usually breaks up the defined epitaxial columnar structure [33]. It was also noticed that the presence of unmelted particles, porosity, or TiC nanoparticles results in a breakup of the epitaxial growth.
The distribution of grain boundary misorientation angles are presented in Fig. 7C and D. Clearly, the grains from the top and side views were separated by a majority of high-angle grain boundaries distribution ([N 15°] [blue color]). This could be due to the high scanning speed implemented in this study, which leads to elongation of the
Table 1 Vickers and nanohardness of SLM-processed samples.
Vicker's and nanohardness (GPa) Elastic modulus (GPa) Anisotropy (%)
Pure H13 (top view)
Pure H13 (side view)
Nanocomposite (top view)
Nanocomposite (side view)
748.04 ± 27.61 (8.36 ± 0.17) 230.71 ± 8.41 ~8.35
810.51 ± 40.11 (8.56 ± 0.11) 236 ± 5.5
811.23 ± 44.18 (9.41 ± 0.30) 255.83 ± 8.14 ~5.82
858.45 ± 36.56 (9.71 ± 0.24) 259 ± 6.47
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crystallographic orientation may be deviated because of repeated solidification and annealing in the narrow-scanned builds.
Table 2 Coefficient of friction (COF) and wear rates of SLM-processed samples.
Average COF Max COF Wear rate (mm3/Nm)
Pure H13
Nanocomposite
0.550 0.989 4.906 × 10−6
0.526 0.822 3.623 × 10−6
molten pool [41]. The small zones containing low angle boundaries (2–5°) that are visible at the borders of the tracks are probably not grain boundaries, but rather sub-grain boundaries formed by dislocation rearrangement. The Kernel average misorientation (KAM), shown in Fig. 7E and F, calculates the average misorientation between a data point and all of its neighbors. No significant differences were noted between the side view and the top view. The higher KAM portion in the grains (i.e., green color) indicates a higher degree of misorientation, which implies high stored energy in the grain. Generally, KAM is high in deformed grains because of higher dislocation density. The crystal orientation maps shown in Fig. 8A and B, demonstrate clear random grain orientations. This was expected because the “cross-hatching” scanning strategy (i.e., the scanning direction is altered by 90° between consecutive layers) results in a reduction of texture along the build and scanning directions, which minimizes the anisotropic properties of such structures. This agrees with the previous observation reported by Boegelein and co-workers [42]. However, looking closely at the inverse pole figures, shown in Fig. 8C and D, a preferential orientation of grain growth in the 〈100〉 direction is observed, which is favored for cubic crystals [43]. When comparing the side view of the SLM build with the top view from the pole figure maps, shown in Fig. 8E and F, one can see some differences in the preferential alignment. This was expected because the same area was melted several times as a result of the narrow scan spacing that was used. Therefore, the
3.2.4. Microstructural observation by TEM Fig. 9 shows the bright-field TEM images of the SLM-processed nanocomposite with selected area diffraction pattern (SADP). Fig. 9A presents the H13 matrix grain boundary with a “metal-ceramic” interface in which the H13 matrix has wet the TiC particles, and small TiC spherical precipitates have formed in the side of the interface (indicated by the triple arrows). Bolton et al. [44] reported minor diffusion of iron from the steel matrix into the TiC particles (limited solubility), which may have caused the formation of Fe-rich precipitates at the interface. It is generally believed that a very thin precipitate-rich interface layer (b200 nm) can be beneficial for the strength of the interface. The reinforcements were not just located at the grain boundaries, which hinders the grain growth at elevated temperature, but also within the grain (indicated by the single arrow). More importantly, TiC nanoparticles with a particle size ~ 50 nm were uniformly dispersed in the grain boundaries of the H13 matrix retaining its size and initial spherical shape (Fig. 9B). The high solidification rate ~106–108 K/s [45], inhibits the grain growth of the TiC nanoparticles because of insufficient time for grain coarsening, hence remaining as the favorable nanostructured TiC reinforcement after SLM. The sample exhibited a high level of dislocations (Fig. 9B), as a result of the fast solidification rate and also as a result of the differences in coefficient of thermal expansion between the matrix and the TiC reinforcement. The TiC nanoparticles act as obstacles to dislocation glide and grain growth. The SADP taken at the grain boundaries of the nanocomposite sample revealed the existence of TiC and BCC-Fe constitution phases (Fig. 9C). No evidence of any amorphous regions was found anywhere within the sample. To confirm that, we tilted the regions that we
Fig. 10. SEM images showing morphologies of the worn surfaces of the nanocomposites. (A,C) Pure H13; (BD) 15% TiC/H13.
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believed might be amorphous, but the tilting of the regions showed a contrast change, suggesting full crystallinity.
3.3. Hardness and wear behavior The effects of TiC reinforcement on the micro- and nano-hardness of SLM-processed H13 steel and nanocomposite are shown in Table 1. The SLM-processed samples demonstrate superior hardness compared to conventionally processed H13 (e.g., casting), with a typical hardness of ~750 HV0.2. This high initial hardness is attributed to the significant grain refinement effect resulting from laser rapid solidification and the presence of a completely martensitic phase that favors high hardness. With the addition of 15% vol TiC, the hardness and elastic modulus were considerably improved by ~ 15% (rule of mixtures), because of the homogenous incorporation of nanoscale TiC reinforcing particles throughout the matrix. The SLM-processed parts are relatively isotropic, as would be expected from the EBSD results, indicating that the orientations of the grains were random and not associated with any preferred directions. Similar results were obtained by Walker et al. [46] who showed that a “cross-hatching” scanning strategy minimizes the anisotropic properties.
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The variation of friction coefficients and wear rates of the pure H13 and TiC/H13 are given in Table 2. Compared to the unreinforced H13 steel, the coefficient of friction of TiC/H13 nanocomposite decreased, and the resultant wear rate was lowered by ~ 26%. Fig. 10 shows the morphology of the worn surfaces of the pure H13 and the nanocomposites. The material loss is believed to result from the removal of material chips from the matrix phase because of micro-cutting by the abrasive particles [7]. The presence of irregularly shaped fragments reveals that local deformation and delamination of the worn surface occurred. With the addition of TiC nanoparticles, the mechanism of material removal during sliding was changed from partial abrasion or a localized tribolayer to a further strain-hardened tribolayer (i.e., shallower grooves and less abrasive fragments). Clearly, the strain-induced tribolayer portion was increased after the incorporation of TiC nanoparticles (Fig. 10C, D). Such a transition is expected to reduce the wear volume after the wear tests, although there was still some minor microplowing. No particle cracking or separation of particles from the matrix can be observed. The improvement in wear performance can thus be attributed to the uniformly dispersed TiC nanoparticle reinforcement and low interfacial stress with the matrix, leading to less breakage of the nanocomposites during sliding. Moreover, the improvement in
Fig. 11. SEM micrographs of the worn surface of the nanocomposites, and the corresponding elemental maps. (a) Pure H13; (b) 15% TiC/H13.
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hardness (Table 1) helped to increase the area of the strain-hardened tribolayer, which prevented further wearing. Ghosh et al. [47] have shown that when the size of reinforcement decreases the wear resistance of the composite improves, with no further enhancement in wear resistance after 20 vol%. Elemental maps of the worn surfaces of the pure steel and the nanocomposites are shown in Fig. 11a and b, respectively. The elemental mapping of the pure steel shows that the worn surface was rich with iron, carbon, and chromium. For the worn surface of the nanocomposite, the map indicates the presence of more iron, but less carbon, suggesting that a small amount of TiC was lost. The TiC particles were present on the worn surface; thus, they protected the matrix during the wear test. The only damage on the surface seems to be some pits from spalled M7C3 carbides and the corresponding mini-grooves produced by them. No presence of oxides after sliding was seen, suggesting that there was no significant temperature rise at the sample surface. Some fine reinforcing particles were observed at the edges of the grooves on the worn surface. 4. Conclusion 1. The mechanical intermixing caused by the high-energy ball milling produces H13 steel reinforced with 15% vol TiC nanocomposite powder. With increased milling time, the powder, which initially had a spherical shape, underwent successive changes in its morphology and particle size distribution. This microstructural evolution is determined by the competition between cold-welding and fracturing of the particles. 2. The microstructure of the SLM-processed nanocomposite viewed from the top reveals equiaxed fine grains as a result of rapid solidification, while the side view clearly shows elongated columnar grains caused by the additive nature of the SLM process. 3. Overall, it was shown that the “cross-hatching” scanning strategy (i.e., the scanning direction is altered 90° between consecutive layers) results in a reduction of texture along the build and scanning directions, which consequently minimizes the anisotropic properties of such structures. 4. In comparison to the unreinforced H13 steel, the average hardness and elastic modulus of the nanocomposite part increased, resulting in enhanced wear resistance. 5. The benefits of AM can be explained by the changes in the microstructure and mechanical properties when compared to cast materials. With a better understanding of these differences, the technology can be further developed and AM methods will be more commonly used. Acknowledgment One of the authors, Bandar AlMangour, gratefully appreciates the financial support from the Saudi Arabia Basic Industries Corporation (SABIC). References [1] J.F. Shackelford, W. Alexander, CRC Materials Science and Engineering Handbook, CRC Press, 2015. [2] G. Krauss, Steels: Processing, Structure, and Performance, ASM International, 2015. [3] S. Hashmi, Comprehensive Materials Processing, Elsevier, 2014. [4] I. Ibrahim, F. Mohamed, E. Lavernia, Particulate reinforced metal matrix composites—a review, J. Mater. Sci. 26 (1991) 1137–1156. [5] S.C. Tjong, Novel nanoparticle-reinforced metal matrix composites with enhanced mechanical properties, Adv. Eng. Mater. 9 (2007) 639–652. [6] E. Pagounis, V. Lindroos, M. Talvitie, Influence of reinforcement volume fraction and size on the microstructure and abrasion wear resistance of hot isostatic pressed white iron matrix composites, Metall. Mater. Trans. A 27 (1996) 4171–4181. [7] F. Akhtar, Microstructure evolution and wear properties of in situ synthesized TiB2 and TiC reinforced steel matrix composites, J. Alloys Compd. 459 (2008) 491–497. [8] W. Jiang, P. Molian, Nanocrystalline TiC powder alloying and glazing of H13 steel using a CO2 laser for improved life of die-casting dies, Surf. Coat. Technol. 135 (2001) 139–149. [9] J. Campbell, Complete Casting Handbook: Metal Casting Processes, Metallurgy, Techniques and Design, Butterworth-Heinemann, 2015.
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