Carbon Solid Solution Effect on Microstructures of Laser Powder Bed Fusion Prepared Ti Alloys

Abstract

In this study, titanium (Ti) alloy with carbon (C) solid solution was fabricated by laser powder bed fusion (LPBF) from Ti-TiC mixture powder to investigate the effect of C solid solution on microstructure and tensile properties of LPBF Ti alloy. XRD and TEM analysis showed that part of TiC particles was decomposed from the surface during melting process of LPBF and C was solid soluted in α-Ti for Ti-(0.09∼0.4 wt%) C. The solid solution of carbon changed the microstructure of the LPBF Ti alloy to fine acicular microstructure due to the martensitic phase transformation. This was also observed in Ti-0.4 wt% C, where solid solution of carbon was not confirmed, suggesting that in Ti-0.4 wt% C, C was once solid solution and precipitated as TiC after phase transformation. LPBF Ti-C alloys showed good strength-ductility, UTS and elongation at break for Ti-0.2 wt% C were 746 MPa and 26.3%, respectively.

 

This Paper was Originally Published in Japanese in J. Jpn. Soc. Powder Powder Metallurgy 71 (2024) 517–523, https://doi.org/10.2497/jjspm.23-00079. The citation in reference 22 is corrected.

1. Introduction

Titanium (Ti) is characterized by high specific strength, high corrosion resistance due to its dense oxide film, and high biocompatibility due to osseointegration (bonding between Ti and bone) [1, 2]. The use of these materials is progressing in a wide range of industrial fields, including medical devices [3–6] and transportation equipment [7], where there have been problems with the application of light metals such as aluminum and magnesium, and highly corrosion-resistant metals such as stainless steel. In such cases, Ti alloys are often used instead of pure Ti, owing to their mechanical properties. Among the current general-purpose Ti alloys, α + β type alloys represented by Ti-6Al-4V composition (ASTM B348 Gr5, Ti 64 alloy) are the mainstream, but the addition of rare metals such as V, Cr, and Mo is essential for stabilizing the β phase [8]. The production of these elements is limited to a few countries [9], and depending on the type of alloying elements and their addition, there are geopolitical risks such as an increase in the price of materials and supply instability.

In contrast, oxygen (O), nitrogen (N), and carbon (C) are inexpensive and ubiquitous elements that are known as interstitial element of Ti to elongate the α-Ti crystal lattice and significantly increase the strength of Ti alloys [10]. However, in conventional Ti castings, the tensile ductility is considered to decrease significantly with the O content above 0.6 wt% [10–12], and the O content in Ti is strictly limited to a maximum of 0.40 wt% in CP-Ti (ASTM Gr. 4) and 0.20 wt% in a general-purpose Ti-6Al-4V alloy (ASTM Gr. 5). In contrast, it has been reported that Ti materials with high concentrations of O (0.4 to 1.0 wt%), N, and C solid solution prepared by powder metallurgy exhibit excellent tensile strength while maintaining high ductility [13, 14]. Ti-O/N alloys fabricated by laser powder bed fusion (LPBF), which have attracted much attention in recent years, have been reported to exhibit excellent strength-ductility balance [15, 16]. On the other hand, there have been several reports on LPBF Ti-C alloys from the viewpoint of using TiC particles as a dispersion strengthening phase, but there are no reports on microstructural analysis or evaluation of mechanical properties from the viewpoint of C solid solution [17–21].

In this study, TiC powder was selected as the C source, and the decomposition of TiC and the solid solution of C atoms in the Ti crystal were investigated through local heating melting and ultra-rapid cooling solidification phenomena using the LPBF method. The effects of solute C atoms on the microstructure and tensile properties of LPBF Ti-C alloys were analyzed in detail.

2. Experimental Procedure

Pure Ti powder (purity 99.7%, D50 = 26.2 µm, TILOP-45, Osaka Titanium Technologies Ltd.) and TiC powder (purity 99.9%, D50 = 3.4 µm, TII02PB, Kojundo chemical laboratory Co., Ltd.) were used as starting material. These powders were mixed to Ti-x wt% TiC (x = 0, 0.25, 0.5, 1.0, 2.0) using a rocking mill (RM-05, Seiwa Giken, Ltd.). The C contents of the LPBF Ti-C alloys prepared from these mixed powders were 0.02, 0.09, 0.16, 0.25, and 0.41 wt%, respectively, and as shown in the Ti-C equilibrium phase diagram in Fig. 1, the C content is within the solid solution limit for β-Ti up to 0.15 wt%. These powder mixtures were fabricated using an LPBF system (TruPrint1000, Trumpf Corporation) with a power output of 160 W, a hatch width of 0.11 mm, a powder layer thickness of 0.02 mm, and a scanning speed of 535 mm s-1. A chessboard pattern consisting of two types of squares with orthogonal laser scanning directions (3.96 mm per side, 2.73 mm and 3.22 mm shifts in the X-axis and Y-axis directions, respectively, for each layer) was used as the laser scanning pattern. Ar gas was used as the shielding gas to maintain the O concentration in the chamber at less than 100 ppm during fabrication. Subsequently, the specimens were vacuum heat-treated at 773 K for 3.6 ks to remove the residual stress.

Fig. 1

Binary phase diagram of Ti-C system.

The LPBF Ti-C alloys were analyzed by X-ray diffraction (XRD, XRD-6100, Shimadzu Corporation), scanning electron microscopy (SEM, JSM-6500F, JEOL Ltd.), electron beam backscatter diffraction (EBSD, Digi Viewer IV detector, TSL Solutions Inc. TSL Solutions, Inc.), and TEM-EDS (TEM, JEOL Ltd., JEM-2100F). For the evaluation of tensile properties, plate-like tensile specimens (parallel section width, 2 mm; thickness, 1 mm; length, 10 mm) were machined from the LPBF sample, and tensile tests were conducted using an autograph (AG-X 50 kN, Shimadzu Corporation) at a strain rate of 5.0 × 10⁻⁴ s⁻¹.

3. Results and Discussions

First, the O, N, and C contents of the LPBF Ti-C alloys and the powder mixture used for their fabrication were analyzed, and the results are shown in Fig. 2. The C content of the LPBF Ti-C alloy increased with TiC addition, confirming a proportional relationship with a coefficient of determination of 0.99. This result indicates that the TiC powder was not separated from the powder mixture by the rising airflow at the heat input and shielding gas during the LPBF process and acted as a source of C. The C content of the LPBF alloy could be controlled with high accuracy using the amount of TiC added. O and N also showed an increasing trend with the amount of TiC added. This is mainly due to the fact that the TiC powder contains a large amount of impurities (O: 0.66 wt% and N: 0.58 wt%), while the O and N contents in the pure Ti powder are 0.12 wt% and 0.001 wt%, respectively. In addition, during the LPBF process, the O₂ and N₂ in the atmosphere are absorbed [22], which is also considered to be the cause of the increase in O and N in the LPBF material.

Fig. 2

(a) Dependence of C, O, N contents of LPBF Ti-C alloy on TiC addition and (b) relationship between C content of Ti-TiC mixture powder and LPBF Ti-C alloy.

Next, to confirm the reaction between the Ti matrix and TiC particles during the LPBF process, XRD analysis was performed on four samples: each raw powder (pure Ti and TiC powders), a Ti-2.0 wt% TiC mixed powder (C content 0.41 wt%), and an LPBF Ti-0.41 wt% C alloy fabricated from the mixed powder. As shown in Fig. 3, the intensity of the TiC diffraction peak at approximately 41.7° detected in the mixed powder was significantly reduced in the LPBF Ti-0.41 wt% C alloy, suggesting that the TiC powder was decomposed during the LPBF process.

Fig. 3

XRD profiles of Ti powder, TiC powder, Ti-2 wt% TiC (0.41 wt% C) mixture powder and LPBF Ti-0.41 wt% C alloy. (online color)

Therefore, XRD analysis of LPBF Ti-C alloys was performed to investigate the phase composition in LPBF Ti-C alloys and to quantitatively analyze the solid solution state of C atoms from the changes in the lattice constants of them. The XRD profiles of the LPBF Ti-C alloys are shown in Fig. 4. The α-Ti peak was detected in all samples, and the TiC diffraction peak was observed only in the Ti-0.41 wt% C sample. Therefore, it was found that the LPBF Ti-C alloy is composed of α-Ti matrix and TiC. In general, it is known that lattice elongation in the c-axis direction occurs when C atoms and other intrusive solid solution atoms are in solid solution in α-Ti crystals [23]. We investigated the change in diffraction peaks corresponding to the α-Ti {0002} and {10$\bar{1}$0} in LPBF Ti-C alloys, and found that the peak at 38.4° corresponding to the α-Ti {0002} peak shifted to a lower angle with increasing TiC addition up to 1 wt% (Ti-0.25 wt% C), but the peak at 35.1° corresponding to the {10$\bar{1}$0} peak around 35.1° remained unchanged. These results suggest that the lattice of α-Ti crystals elongated only in the c-axis direction due to the solid solution of C accompanying the decomposition of TiC powder in the range of TiC addition up to 1.0 wt% (Ti-0.25 wt% C).

Fig. 4

XRD profiles of LPBF Ti-x wt% C alloy (x = (a) 0.02, (b) 0.09, (c) 0.16, (d) 0.25 and (e) 0.41). (online color)

Based on the above results, we investigated the relationship between the lattice constant of the α-Ti crystals and the C content. Figure 5 shows the relationship between the lattice constant of α-Ti and the C content in LPBF Ti-C alloys. The a-axis length was almost constant (2.954–2.956 Å), whereas the c-axis length decreased at Ti-0.41 wt% C. The increase rate of the c-axis lattice constant at Ti-0.02∼0.25 wt% C was calculated to be 0.0044 Å/at%, which was lower than the increase rate (0.0145 Å/at%) reported in a previous study [24] for C solid solution Ti sintered extruded material. The solid solution of C was maintained in a supersaturated solid solution state in the case of Ti-0.02∼0.25 wt% C, while in the case of Ti-0.41 wt% C, the solid solution C was precipitated as TiC and the supersaturated solid solution state could not be maintained and the amount of C solid solution decreased.

Fig. 5

Relationship between lattice constant of α-Ti and C content.

Therefore, we investigated the distribution of the TiC partcles in the LPBF Ti-0.09∼0.41 wt% C alloy through SEM observation. In the XRD analysis results described above, no TiC diffraction peaks were detected in the samples other than the Ti-0.41 wt% C alloy, but the SEM images shown in Fig. 6 confirm the presence of TiC particles in all samples. The number of dispersed TiC particles increased with an increase in the C content. To clarify the formation mechanism of these TiC particles, the structures of the raw TiC particles and the dispersed TiC particles were analyzed by TEM. The results are shown in Fig. 7. As shown in Fig. 7(a), the crystallographic orientation of the raw TiC particles differed from place to place, indicating that the TiC particles used in this study had a polycrystalline structure with multiple crystal grains within a single particle. It has been reported that reprecipitated TiC particles due to the decreased solid solution limit during the cooling process have a single crystal structure [25]. The crystal structure of the dispersed TiC particles observed in Ti-0.41 wt% C shows that the left part of the particle shows ⟨33$\bar{2}$⟩ while the right part shows ⟨1$\bar{1}\bar{1}$⟩, indicating that there are two or more orientations within the particle. These results indicate that the TiC phase in LPBF Ti-C alloys, similar to the raw TiC particles, has a polycrystalline structure with multiple orientations within a single particle. Therefore, the results suggest that the TiC particles remained in the LPBF Ti-C alloy is origin to added TiC particles. However, the TiC particles in the LPBF Ti-C alloy is rounded, whereas the raw TiC particles have an angular shape owing to the milling process. This suggests that the TiC particles was partially, although not completely, decomposed during the melting process by the LPBF method. In addition to the above results, the XRD analysis confirmed the change in the lattice parameter accompanying the solid solution of C. Although some of the solid solution C could precipitate as TiC, the intended decomposition of TiC particles and the accompanying solid solution of C in the LPBF process were achieved.

Fig. 6

SEM images of LPBF Ti-x wt% C alloy (x = (a) 0.09, (b) 0.16, (c) 0.25 and (d) 0.41). Black arrow shows TiC particles.

Fig. 7

(-1) TEM image and (-2, 3) SAD patterns of (a) TiC raw particle and (b) TiC particle dispersed in LPBF Ti-C alloy.

An inverse pole figure (IPF) maps of the LPBF Ti-C alloy is shown in Fig. 8. The Ti-0 wt% C (pure Ti material) consists of coarse columnar grains, whereas Ti-0.09∼0.41 wt% C consists of fine acicular grains, whose grain size decreases with increasing C content. Such changes have also been reported for LPBF Ti alloys with solid solutions of O and N [15, 16]. Figure 9 shows the relationship between the average grain diameter and average width of acicular grains and the C content in the LPBF Ti-C alloys. The grains of Ti-0.25 wt% C became finer, and the average grain size was approximately 2.1 µm. On the other hand, Ti-0.41 wt% C exhibited a microstructure composed of fine acicular grains and coarse grains of 20 to 40 µm, with a slightly larger average grain size of 3.5 µm. The mode of grain size was 2.1 µm for Ti-0.25 wt% C and 1.9 µm for Ti-0.41 wt% C. Therefore, there was no significant difference in the size of acicular grains in LPBF Ti-C alloys.

Fig. 8

IPF maps of LPBF Ti-x wt% C alloys (x = (a) 0.02, (b) 0.09, (c) 0.16, (d) 0.25, and (e) 0.41). IPF maps are colored along building direction. (online color)

Fig. 9

Relationship between grain size and C content.

Figure 10 shows the IPF of the crystal orientation of the LPBF Ti-C alloys in the building direction. Ti-0.02 wt% C has a strong ⟨0001⟩ orientation, while Ti-0.09∼0.41 wt% C has strong ⟨2$\bar{1}\bar{1}$0⟩ and ⟨10$\bar{1}$2⟩ orientations. According to Burgers’ orientation relation [26], during phase transformation, β-Ti transforms to α-Ti to satisfy (0001)α-Ti//(101)β-Ti and ⟨11$\bar{2}$0⟩α-Ti//⟨1$\bar{1}\bar{1}$⟩β-Ti; thus, ⟨2$\bar{1}\bar{1}$0⟩α-Ti and ⟨10$\bar{1}$2⟩α-Ti correspond to ⟨001⟩β-Ti. The orientation of β-Ti was evaluated by reconstructing prior β-Ti grain from the crystal orientation of α-Ti, and the results are shown in Fig. 11. It has been reported that LPBF β-Ti alloys form ⟨001⟩ and ⟨110⟩ orientations depending on the laser scan strategy, and the ⟨001⟩ orientation is characteristic of β-Ti alloys fabricated with a laser scan strategy of the X-Y scan [27]. The laser scan strategy employed in this study is called a chessboard pattern, in which X and Y scans with orthogonal laser scan directions are periodically repeated. Therefore, it can be concluded that β-Ti with a strong orientation of ⟨001⟩β-Ti in the building direction was formed with the chessboard pattern owing to the same microstructure formation as in the X-Y scan, and that the phase transformation according to Burgers’ orientation relationship resulted in a strong orientation of ⟨2$\bar{1}\bar{1}$0⟩ and ⟨10$\bar{1}$2⟩. Such microstructural changes are observed even in Ti-0.09 wt% C, where the dispersion of TiC phase is limited, suggesting that they are not due to TiC dispersion. As mentioned above, it has been reported that fine acicular grain structures are formed in LPBF Ti-O/N alloys because of the solid solution of trace amounts of O and N [15, 16]. Therefore, it can be concluded that the similar microstructural changes observed in this LPBF Ti-C alloy are also due to solid solution elements and that fine acicular grains are formed by a solid solution of C atoms generated by the decomposition of TiC particles. The formation of fine acicular grains was also observed in Ti-0.41 wt% C, where no elongation of the crystal lattice in the c-axis direction was observed. This is considered to be due to the fact that the C component was supersaturated after the acicular microstructure formation (phase transformation) and the TiC particles were precipitated in α-phase due to decreased C solubility in α-Ti.

Fig. 10

IPF along building direction of Ti-x wt% C alloys (x = (a) 0.02, (b) 0.09, (c) 0.16, (d) 0.25, and (e) 0.41). (online color)

Fig. 11

IPF maps of reconstructed β phase of LPBF Ti-x wt% C alloys (x = (a) 0.09, (b) 0.16, (c) 0.25 and (d) 0.41). IPF maps are colored along building direction. (online color)

Finally, tensile tests were conducted at room temperature to investigate the effect of varying the C content on the tensile properties of the LPBF Ti-C alloy. The obtained stress-strain curves and tensile properties are shown in Fig. 12. A proof stress (YS) of 328 MPa and ultimate tensile stress (UTS) of 392 MPa were obtained for Ti-0.02 wt% C. The tensile strength increased with increasing C content in Ti-0.25 wt% C, 0.2% YS and UTS were 661 MPa and 746 MPa, respectively, and in Ti-0.41 wt% C, where the maximum C content was added, 0.2% YS was 647 MPa and UTS was 770 MPa, respectively. These higher strengths can be attributed to grain refinement, solid-solution strengthening of C, and dispersion strengthening of TiC. The elongation at break ranged from 24.5 to 28.4% for Ti-0∼0.25 wt% C. High elongation at break was maintained even as the C content increased. The Ti-0.41 wt% C showed a decrease in ductility, but exhibited a sufficient elongation at break of 16.3%. Figure 13 shows the SEM observation results of the fracture surfaces of Ti-0.25 and 0.41 wt% C. Ti-0.25 wt% C exhibited a ductile fracture surface with equiaxed dimples with several micrometers on the entire surface. On the other hand, in Ti-0.4 wt% C, equiaxed dimples similar to those in Ti-0.25 wt% C and quasicleavage fracture surfaces indicating brittle fracture were observed. These quasicleavage fracture are 30–40 µm in size, which is equivalent to the coarse grains observed in the Ti-0.41 wt% C. It has also been reported that in LPBF Ti-O/N alloys, no reduction in ductility associated with solid solution strengthening was observed until the UTS exceeded 1000 MPa [15, 16]. On the other hand, because the UTS of Ti-0.41 wt% C is 770 MPa, it can be concluded that the reduction in ductility associated with the solid solution strengthening of C is limited. Therefore, the quasicleavage surfaces observed in Ti-0.41 wt% C are due to the coarse grains, and the poor plastic deformation ability of these coarse grains have reduced the ductility of Ti-0.41 wt% C.

Fig. 12

(a) Stress-strain curves and (b) tensile properties of LPBF Ti-C alloys. (online color)

Fig. 13

Fractography of (a) LPBF Ti-0.25 wt% C and (b) LPBF Ti-0.41 wt% C alloys.

4. Conclusion

In this study, we attempted to produce a Ti alloy with C solid solution by applying the LPBF method to Ti-TiC mixed powder. The decomposition of TiC during the LPBF process was comfirmed, and the effects of the C solid solution on the microstructure and mechanical properties were investigated. The following conclusions were obtained:

1. (1)    Microstructural analysis using XRD and TEM revealed that although the TiC powder was decomposed in the LPBF process, a part of the TiC powder remained without decomposed, and that the C generated by the decomposition of TiC was in solid solution, and the crystal lattice expanded in the c-axis direction up to 0.25 wt% C addition. However, no crystal lattice expansion was observed for Ti-0.41 wt% C, and the expansion rate per C content was about 1/3 of that reported in previous studies. Therefore, although TiC decomposes during the melting process in LPBF, some of them have re-precipitated as TiC.
2. (2)    The LPBF Ti-C alloy exhibits a fine acicular microstructure resulting from martensitic transformation, and the grain becomes finer with the amount of TiC added, with an average grain size of 2.1 µm for Ti-0.25 wt% C. This change in microstructure can be attributed to solid solution of C. Fine acicular microstructures were also observed in Ti-0.41 wt% C, where no lattice expansion from Ti-0.25 wt% C was observed, suggesting that solid solution C reprecipitated as TiC after the phase transformation.
3. (3)    The LPBF Ti-C alloy exhibited excellent strength and ductility properties, with a UTS of 392 MPa and elongation at break of 25.8% for Ti-0 wt% C, compared with UTS 746 MPa and elongation at break of 26.3% for Ti-0.25 wt% C and UTS 770 MPa and elongation at break of 16.3% for Ti-0.41 wt% C. The reduction in ductility was limited even when large amounts of TiC were added, which may be due to the smaller solid-solution strengthening capacity of C compared to O and N.

Acknowledgments

The authors acknowledge funding from OU Master Plan Implementation Project of Osaka University and the International Joint Research Promotion Program of Osaka University.

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• This Paper was Originally Published in Japanese in J. Jpn. Soc. Powder Powder Metallurgy 71 (2024) 517–523, https://doi.org/10.2497/jjspm.23-00079. The citation in reference 22 is corrected.