Effect of heat-treatment temperature on microstructures and mechanical properties of Co–Cr–Mo alloys fabricated by selective laser melting

Author’s Accepted Manuscript Effect of heat-treatment temperature on microstructures and mechanical properties of Co– Cr–Mo alloys fabricated by selective laser melting Yuka Kajima, Atsushi Takaichi, Nuttaphon Kittikundecha, Takayuki Nakamoto, Takahiro Kimura, Naoyuki Nomura, Akira Kawasaki, Takao Hanawa, Hidekazu Takahashi, Noriyuki Wakabayashi

PII: DOI: Reference:

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S0921-5093(18)30545-8 https://doi.org/10.1016/j.msea.2018.04.048 MSA36368

To appear in: Materials Science & Engineering A Received date: 18 January 2018 Revised date: 10 April 2018 Accepted date: 11 April 2018 Cite this article as: Yuka Kajima, Atsushi Takaichi, Nuttaphon Kittikundecha, Takayuki Nakamoto, Takahiro Kimura, Naoyuki Nomura, Akira Kawasaki, Takao Hanawa, Hidekazu Takahashi and Noriyuki Wakabayashi, Effect of heattreatment temperature on microstructures and mechanical properties of Co–Cr– Mo alloys fabricated by selective laser melting, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2018.04.048 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Effect of heat-treatment temperature on microstructures and mechanical properties of Co–Cr–Mo alloys fabricated by selective laser melting

Yuka Kajimaa,b, Atsushi Takaichib*, Nuttaphon Kittikundechab, Takayuki Nakamotoc, Takahiro Kimurac, Naoyuki Nomuraa, Akira Kawasakia, Takao Hanawad, Hidekazu Takahashie, Noriyuki Wakabayashib

a

Department of Materials Processing, Graduate School of Engineering, Tohoku University, 6-6-02 Aramaki-aza Aoba, Aoba-ku, Sendai, Miyagi 980-8579, Japan

b

Removable Partial Prosthodontics, Oral Health Sciences, Graduate School of Medical and Dental

Sciences, Tokyo Medical and Dental University, 1-5-45 Yushima, Bunkyo-ku, Tokyo 113-8549, Japan c

Research Division of Machining and Molding, Osaka Research Institute of Industrial Science and Technology Izumi Center, 2-7-1 Ayumino, Izumi, Osaka 594-1157, Japan

d

Metallic Biomaterials, Biomedical Materials, Institute of Biomaterials and Bioengineering, Tokyo Medical and Dental University, 2-3-10 Kanda-surugadai, Chiyoda-ku, Tokyo 101-0062, Japan

e

Department of Oral Biomaterials Engineering, Graduate School of Medical and Dental Sciences, Tokyo

Medical and Dental University, 1-5-45 Yushima, Bunkyo-ku, Tokyo 113-8549, Japan

*

Corresponding author: Removable Partial Prosthodontics, Oral Health Sciences, Graduate School of

Medical and Dental Sciences, Tokyo Medical and Dental University, 1-5-45, Yushima, Bunkyo-ku, Tokyo 113-8549, Japan, Tel.:/fax: +81-3-5803-5842; [email protected]

Abstract Selective laser melting (SLM) has attracted considerable attention as an advanced method for the fabrication of biomedical devices. However, SLM-manufactured parts easily accumulate large amounts of residual stress due to rapid heating and cooling, which negatively affects their mechanical properties. In this study, Co–Cr–Mo alloy specimens were fabricated by SLM and then heat-treated at various temperatures (750, 900, 1050, or 1150 °C) to relieve the residual stress and improve their mechanical properties. The alloy microstructure was analyzed via confocal laser scanning microscopy, scanning electron microscopy combined with energy dispersive X-ray spectroscopy, electron backscattered diffraction, and X-ray diffraction techniques, whereas the mechanical properties of the produced specimens were evaluated by tensile and Vickers hardness tests. The results showed that increasing the heat-treatment temperature from 750 °C to 1150 °C increased the ductility of the alloy and decreased its 0.2% offset yield strength and Vickers hardness. Both γ and ε phases formed in all heat-treated specimens, and the volume fraction of the ε phase decreased with increasing heat-treatment temperature. After the specimens were heated to 750–1050 °C, a recovery process was initiated, which proceeded as the temperature increased; however, the residual stress in the studied specimens was not sufficiently relieved. In contrast, after heating to 1150 °C, the formation of equiaxed grains and the drastic relief of the residual stress were observed simultaneously, accompanied by an increase in the elongation of the specimen and a decrease in its strength (as compared to those of the other heat-treated specimens), indicating that the specimen completely recrystallized and that the residual stress was the driving force of this recrystallization. Thus, heat-treating at 1150 °C for 6 h is an effective method for eliminating the residual stress, leading to a homogenized microstructure and satisfactory ductility. Keywords: Selective laser melting; Heat treatment; Mechanical property; Microstructure; Co–Cr–Mo alloy; Recrystallization

1. Introduction Co–Cr–Mo alloys are widely used to manufacture orthopedic implants such as artificial hips, knee joints, and dental prostheses owing to their excellent biocompatibility and good mechanical properties as well as their high wear and corrosion resistances [1-4]. In recent years, selective laser melting (SLM), an additive manufacturing process, has attracted significant attention as an advanced method for fabricating biomedical devices, and it has begun replacing conventional manufacturing methods [5-7]. SLM can create three-dimensional metal parts via the layer-by-layer melting of metal powders with a laser beam according to a computer-aided design [7], and it enables the rapid, low-cost, and precise fabrication of custom-made medical devices [8]. In particular, SLM-processed Co–Cr alloys have been used to manufacture dental implants, crowns, and bridges [8]. However, a well-known drawback of SLM is the large amount of residual stress that accumulates owing to the rapid heating and cooling of the manufactured parts, which is detrimental to the resulting mechanical properties [9-12]. For example, Edwards et al. [9] reported that SLM specimens exhibited high values of the residual tensile stress on their surfaces, which led to poor fatigue performance [9]. Therefore, SLM-fabricated metal parts usually require a post-fabrication heat treatment to relieve the residual stress and further improve their microstructures and mechanical properties [10-13]. Conventionally cast Co–Cr–Mo alloys are composed of two phases: the γ phase (face-centered cubic (FCC) structure) and the ε phase (hexagonal close-packed (HCP) structure), which are stable at high and low temperatures, respectively [14]. To improve the mechanical properties of these alloys, heat treatments have been successfully applied to vary the volume fraction of these microstructurally distinct phases, thus producing a more homogeneous structure with enhanced mechanical properties [15-18]. A similar approach can potentially be applied to SLM-processed Co–Cr–Mo alloys. Lu et al. [13] investigated the microstructures, mechanical properties, and metal release behaviors of SLM-processed Co–Cr–W alloys under different solution heat-treatment conditions and concluded that an appropriate heat-treatment procedure could significantly increase the strength and tensile ductility of SLM-fabricated

parts [13]. However, the as-built Co–Cr–W alloys show duplex phases of γ and ε phases, which is different from the as-built Co–Cr–Mo alloys (the as-built Co–Cr–Mo alloys were exclusively composed of the γ (FCC) cobalt phase in the previous study). In addition, the micrograph of as-built Co–Cr–W alloys exhibits a distinctive diamond-like network pattern consisting of a border structure and square-like structure, which were not observed in as-built Co–Cr–Mo alloys in the previous study. Mengucci et al. [11] investigated the effects of thermal treatments on the microstructure and mechanical properties of a Co–Cr–Mo–W alloy produced by laser sintering. However, their heat-treatment conditions were especially intended for porcelain fused to metal crowns, i.e., for dental devices; thus, the specimens in that study were subjected to a complex firing procedure specifically for veneering with a dental ceramic material [11]. In addition, these studies did not refer to the effects of the heat treatment on the resulting residual stress relief. Therefore, the effects of the heat treatment on the residual stress relief and the changes in microstructures and mechanical properties should be clarified in as-SLM Co– Cr–Mo alloys. In this work, the microstructures and mechanical properties of Co–Cr–Mo alloy specimens fabricated by SLM before and after heat treatments for 6 h at 750, 900, 1050, and 1150 °C were examined, and an optimal stress-relieving procedure was determined. The microstructural change was discussed along with previous studies.

2. Materials and methods 2.1. Specimen preparation Commercially available Co–Cr–Mo alloy powders (MP1, EOS, Krailling, Germany) were used in this study, and their chemical compositions are listed in Table 1. Dumbbell-shaped specimens with diameters of 3 mm and gage lengths of 18 mm were prepared using an SLM machine equipped with a fiber laser (EOSINT M280, EOS, Krailling, Germany). The SLM instrument was operated using the standard deposition parameters for MP1 in a nitrogen atmosphere. The longitudinal direction of the specimens

was oriented parallel to the build direction (BD). To examine the microstructures of the prepared specimens, some of them were cut perpendicularly to the longitudinal direction to produce disc specimens with approximate diameters of 3 mm and heights of 1 mm using a slow-speed diamond cutter. For the heat treatment, specimens were heated at a ramp rate of 60 °C/min to 750, 900, 1050, or 1150 °C and held at this temperature for 6 h in a furnace in an Ar atmosphere. The heat-treated specimens were slowly cooled to room temperature (approximately 20 °C) inside the same furnace. When the temperature reached approximately 300 °C, the furnace door was opened.

Table 1 Chemical composition of the Co–Cr–Mo alloy powder (MP1) in mass%.

MP1

Co 60–65

Cr 26–30

Mo 5–7

Si <1.0

Mn <1.0

Fe <0.75

C <0.16

Ni <0.1

2.2. Microstructural observations and analysis After the heat treatment, the disc specimens were polished with waterproof emery paper (up to 1000 grit), 9-µm diamond, and a 0.04-µm colloidal silica suspension followed by electropolishing in a H2SO4/CH3OH (5:95) solution at a voltage of 16–20 V and a temperature of 268–273 K. Their microstructures were investigated via confocal laser scanning microscopy (CLSM; OLS4000, OLYMPUS, Tokyo, Japan) and scanning electron microscopy combined with energy dispersive X-ray spectroscopy (SEM-EDS; S-3400NX, Hitachi, Tokyo, Japan) at an accelerating voltage of 15 kV and an emission current of 68 mA. Spot analyses were conducted to examine the compositional segregation in the specimens by SEM-EDS.

In addition, the elemental distributions of the specimens were determined

using field-emission electron probe microanalysis (FE-EPMA; JXA-8530F, JEOL Ltd., Tokyo, Japan) at an acceleration voltage of 10 kV. Their crystallographic orientations were determined via electron backscattered diffraction (EBSD; TexSEM Laboratories Inc., USA) using a field-emission scanning electron microscope (XL-30FEG, Philips, The Netherlands). Phases were identified by X-ray diffraction (XRD; D8 Advance, Bruker-AXS, Germany) using Cu Kα radiation at a voltage of 40 kV and a current

of 40 mA in the 2θ scanning range between 40 and 55.

2.3. Mechanical properties The Vickers hardness values of the cylindrical specimens were measured with a load of 100 g for 30 s (n = 3 for each condition). At least five independent measurements were collected for each specimen. Uniaxial tensile tests were performed on the dumbbell specimens (n = 4 for each condition) at an initial strain rate of 1.1×10−3 s−1 using an Instron universal testing machine (AG-2000B, Shimadzu, Kyoto, Japan). Strain values were measured using a noncontact optical strain gauge. The 0.2% offset yield strengths (0.2% YS), ultimate tensile strengths (UTS), and elongations of the tested specimens were obtained from the resulting stress–strain curves. After the tensile tests, the fracture surfaces were observed by SEM to determine the fracture mode.

2.4. Statistical analysis The data obtained from the Vickers hardness and tensile stress measurements were statistically analyzed by performing one-way analysis of variance (ANOVA) and Tukey’s multiple comparison tests. The statistical significance was set to p = 0.01.

3. Results 3.1. Microstructures Figs. 1(a) and (b) show the CLSM and SEM images of the as-built specimens, respectively. On the macroscale, molten pool boundaries were clearly aligned, and the distance was 70 m, which corresponds to the hatching distance. On the other hand, numerous fine cellular dendrites with diameters of approximately 0.5–1 µm were detected on the microscale.

(a)

50 µm

BD

Fig. 1. Microstructures of the horizontal cross section to the BD of the SLM-processed alloy specimens: (a) CLSM and (b) SEM images.

The CLSM images of the heat-treated specimens are shown in Fig. 2. The specimens heated to 750, 1050, and 1150 °C exhibit well-defined grain boundaries (GBs), which were not as distinct in the specimens treated at 900 °C. According to Fig. 2(a), striations were observed inside the grains of the alloy specimens heated to 750 °C, whereas the grains of the specimens heated to 1150 °C were larger than the grains of the specimens heated to different temperatures.

(a)

(b)

(c)

(d)

BD Fig. 2. CLSM images of the horizontal cross section to the BD of the Co–Cr–Mo alloy heated to (a)

50 µm

750 °C, (b) 900 °C, (c) 1050 °C, and (d) 1150 °C.

Fig. 3 shows SEM images of the specimens heat-treated at various temperatures. Numerous fine cellular dendrites, which were observed in the as-built specimens, were also detected in the specimens heated to 750 °C; however, their boundaries were not very distinct. Multiple precipitates formed inside the grains of the specimens heated to 750, 900, and 1050 °C, but not in the specimens heated to 1150 °C. In addition, the precipitates grew along the GBs under all conditions.

Fig. 3. SEM images of the horizontal cross section to the BD of the alloy specimens heat-treated at (a) 750 °C, (b) 900 °C, (c) 1050 °C, and (d) 1150 °C.

Fig. 4 shows the backscattered electron (BSE) images of the various heat-treated specimens and their Co, Cr, and Mo contents, as determined by EDS. Point analyses were performed for both the matrix and precipitates of all specimens (represented by the dark and bright regions, respectively). The matrix was Co-rich after treating under all conditions. The specimens heated to 750, 900, and 1050 °C contained very fine precipitates with diameters ranging from several tens of nanometers to a few hundred nanometers both inside the grains and at the GBs. In addition, larger particles with diameters of approximately 1–1.5 μm were observed at the GBs of the specimens heated to 1050 °C (indicated by the white arrows in Fig. 4c). In contrast, after heating to 1150 °C, the long and narrow precipitates grew along the GBs. Alloy elements tended to segregate similarly in the fine precipitates (formed at 750, 900, and 1050 °C) and the long and narrow precipitates (formed at 1150 °C). Their Cr and Mo contents were higher than those of the alloy matrix. In contrast, the larger precipitates (formed at 1050 °C) contained

higher concentrations of Mo and lower concentrations of Co and Cr.

*2 *1

*5

*8 *7

*6

*3 *4

*9

Fig. 4. BSE images of the horizontal cross section to the BD of the alloy specimens heated to (a) 750 °C, (b) 900 °C, (c) 1050 °C, and (d) 1150 °C, and their Co, Cr, and Mo contents as determined by EDS. The asterisks denote the analyzed locations, whereas the white arrows indicate large precipitates.

Figs. 5 and 6 show SEM-BSE images and the corresponding EPMA maps of the specimens treated at

750, 900, and 1150 °C (Fig. 5) and at 1050 °C (Fig. 6). The results revealed that C was also segregated in the precipitates under all conditions (Figs. 5a, b, and c and Fig. 6). Moreover, two types of precipitates were identified in the specimens heated to 1050 °C (Fig. 6). One type was the bright precipitates in the SEM-BSE image (indicated by the yellow arrows), which were enriched with Mo and C; they were depleted of Cr and Co, as with the compositional segregation of larger precipitates (formed at 1050 °C) determined by EDS (Fig. 4c). The other was the precipitates with high concentrations of Cr, Mo, and C and a low concentration of Co (indicated by the red arrow).

BSE

C

BSE

C

BSE

C

BD

Low

High

Fig. 5. SEM-BSE images and EPMA elemental maps of the horizontal cross section to the BD of the

Co–Cr–Mo alloy heated to (a) 750 °C, (b) 900 °C, and (c) 1150 °C.

Mo

Cr

Co

C

BSE

5 µm BD

Low

High

Fig. 6. SEM-BSE images and EPMA elemental maps of the horizontal cross section to the BD of the Co–Cr–Mo alloy heated to 1050 °C.

The XRD patterns (Fig. 7) of the alloy specimens treated at various temperatures show that the as-built specimen is exclusively composed of the γ (FCC) cobalt phase, whereas the heat-treated specimens contain both γ (FCC) and ε (HCP) phases. Fig. 8 displays the volume fractions of the ε phase of the alloy specimens treated under various conditions, which were calculated from the intensities of their XRD peaks using the expression developed by Sage and Gillaud [19]:

The obtained results indicate that the volume fraction of the ε phase decreases with the increasing heat-treatment temperature, approaching zero at 1150 °C.

Fig. 7. XRD spectra of the specimens heat-treated at (a) 1150 °C, (b) 1050 °C, (c) 900 °C, and (d) 750 °C and (e) the as-built specimen.

Fig. 8. Volume fractions of the ε phase in the alloy specimens heated to different temperatures.

Fig. 9 shows the EBSD data for the alloy specimens heated to different temperatures, which were used to construct the corresponding GB, inverse pole figure (IPF), kernel average misorientation (KAM), and misorientation angle (MA) distribution maps. In the obtained GB maps, the high-angle boundaries (HABs) with misorientation angles greater than 15°, low-angle boundaries (LABs) with misorientation angles in the range of 2–15°, and Σ3 boundaries are represented by the black, red, and blue lines, respectively. Both the GB and MA distribution maps exhibited similar trends under all conditions, except for the specimens heated to 1150 °C, which produced distinct peaks corresponding to LABs with a numerical fraction of approximately 40%. In addition, multiple subboundaries formed inside the grains of this specimen, whereas the fraction of HABs was very low. On the other hand, the specimens heated to 1150 °C exhibited equiaxed grains, and approximately 60% of the GBs have a clear misorientation angle of approximately 60°. Thus, it was concluded that these GBs are Σ the fraction of

LABs in these specimens was close to zero, and their KAM values were much lower than the magnitudes obtained for the specimens treated at different temperatures.

Fig. 9. Results of EBSD studies: GB (a, e, i, m, q), IPF (b, f, j, n, r), KAM (c, g, k, o, s), and MA (d, h, l, p, t) distribution maps of the horizontal cross section to the BD of the alloy specimens heated for 6 h at different temperatures: (a, b, c, d) as-built; (e, f, g, h) 750 °C; (i, j, k, l) 900 °C; (m, n, o, p) 1050 °C; and (q, r, s, t) 1150 °C. The areas surrounded by thick black lines indicate the ε phases (HCP) in the IPF maps.

3.2. Mechanical properties Table 2 lists the Vickers hardness (HV) values of the alloy specimens heated to different temperatures, indicating that the HV magnitude generally decreases with increasing temperature. In particular, the HV values of the specimens heated to 1050 and 1150 °C were significantly lower than those of the as-built specimens as well as those of the alloys treated at 750 and 900 °C. However, the HV values of the untreated alloys and those treated at 750 and 900 °C were not significantly different (p > 0.01).

Table 2. Vickers hardness values of the alloy specimens heated to different temperatures. Standard deviations are listed in parentheses. The values marked with identical letters exhibited no significant differences (p > 0.01). Specimen As-built heated to 750 °C heated to 900 °C heated to 1050 °C heated to 1150 °C

Vickers hardness (HV) 477 (9)a 498 (2)a 495 (10)a 428 (9)b 365 (11)c

Fig. 10 shows the typical stress–strain curves obtained for the specimens heat-treated at various temperatures, and their mechanical properties are summarized in Table 3. Relative to those of the as-built specimens, the 0.2% YS values increased after heating to 750 °C but decreased with higher heat-treatment temperatures. The UTS of the as-built specimens (1173 MPa) was the highest among all

specimens, and this value decreased after all heat treatments. The heat-treated specimens all exhibited similar UTS magnitudes; however, a significant difference was observed between the UTSs of specimens heated to 750, 900, and 1050 °C and those of the specimens heat-treated at 1150 °C. With the increasing heat-treatment temperature, the sample elongation increased from 3.1% to 16.3%, and the highest elongation was observed for the specimens heated to 1150 °C.

Fig. 10. Typical stress–strain curves obtained for the alloy specimens heat-treated at (a) 750 °C, (b) 900 °C, (c) 1050 °C, and (d) 1150 °C.

Table 3. Mechanical properties of the alloy specimens heated to different temperatures. Standard deviations are listed in parentheses. The values marked with identical letters exhibited no significant differences (p > 0.01). Specimen

UTS (MPa)

0.2% YS (MPa)

Elongation

As-built heated to 750 °C heated to 900 °C heated to 1050 °C heated to 1150 °C

1173 (25)a 1097 (11)b 1071 (14)b 1075 (11)b 1007 (32)c

839 (32)a 953 (16)b 793 (25)a,c 738 (4.0)c 614 (11)d

12.3 (2.8)a,c 3 (0.5)b 5.8 (0.9)b 11.4 (0.7)a 16.3 (1.4)c

Fig. 11 shows the fracture surfaces of different heat-treated specimens after the tensile test. Stepped cleavage fractures were observed with wedge-type cracks for the as-built specimens and specimens heat-treated at 750, 900, and 1050 °C. On the other hand, for the specimens heat-treated at 1150 °C, it is seen that cracks propagated along the GBs, and distinct facets were not observed at low magnification. Therefore, the fractures of the samples appeared to be intergranular. At high magnification, dimple-type fracture surfaces were clearly observed in the specimens heated to 1150 °C.

Fig. 11. Fracture surface morphologies of specimens heated to different temperatures after tensile testing: (a) as-built, (b) 750 °C, (c) 900 °C, (d) 1050 °C, (e) 1150 °C, and (f) 1150 °C at higher magnification.

4. Discussion The results of the tensile and Vickers hardness tests conducted after all heat treatments revealed that as the heat-treatment temperature increased, the ductility of the alloy specimens increased, whereas their 0.2% YS and Vickers hardness decreased. The observed variations in the mechanical behavior of the alloy are strongly related to the microstructural changes induced by the heat treatment.

4.1. Effects of the heat-treatment temperature on the microstructure Co–Cr–Mo alloys are well-known to contain two main phases, HCP (ε) and FCC (γ) [20], and their ductility increases with the increasing volume fraction of the γ phase because this phase contains a larger number of independent slip systems [21]. In this study, the volume fraction of the ε phase formed by the heat treatment decreased with increasing heat-treatment temperature (Fig. 8), which could contribute to the observed increase in the sample elongation. According to the phase diagram, the HCP phase is more stable at room temperature [21]. Nevertheless, the FCC-to-HCP transformation rarely occurs under normal cooling conditions; as a result, the FCC structure is usually retained at room temperature [20]. This transformation can be induced athermally (by quenching from temperatures corresponding to the FCC region of stability) [21-23], isothermally (by aging at temperatures between 650 °C and 950 °C) [16,17,21,24,25], or through plastic strain [26,27], which is known as the ε

heated to 750 °C in this study was the presence of intragranular striations, presumably corresponding to the ε martensite phase [15,21,28]. They resulted from the extremely low stacking fault energy (SFE) of Co–Cr–Mo alloys, which assume negative values at temperatures below 850 °C [29] and favor the formation of ε martensite [30]. Some researchers reported that the ε martensite phase formed within the original γ grains through the accumulation of stacking faults, and the grains contained many lattice defects such as dislocations [21,30,31]. In general,

ε

phase exhibits a relatively high strength; however,

this phase negatively affects the ductility of Co–Cr-based alloys [28,32-34]. Therefore, the specimens heated to 750 °C were characterized by the highest Vickers hardness and 0.2% YS values but the lowest elongations among all heat-treated specimens. On the other hand, after heating to 900 °C or 1050 °C, intragranular striations were not apparent in the corresponding CLSM images, although the ε ε martensitic transformation occurred in the lower temperature region (approximately 700 °C), while a massive transformation was predominant at higher temperatures (approximately 900 °C) [21]. These two-phase transformations are considered to compete with each other, depending on the temperature. In the time–temperature–transformation diagram, the nose temperature of the start of the massive transformation process is at approximately 900 °C, and at temperatures below 800 °C, the isothermal martensitic transformation replaces the massive transformation with decreasing temperature [21]. Therefore, in this study, the γ ε

900 °C and 1050 °C.

Furthermore, we identified a variety of precipitates. Precipitates were observed inside the grains and at the GBs after the heat treatments at 750, 900, and 1050 °C. In contrast, after heating to 1150 °C, precipitates were not observed inside the grains, indicating that any formed precipitates dissolved into the γ matrix, thus increasing the elongation of the specimens and decreasing their hardness [13]. The XRD patterns obtained for all of the studied alloys do not exhibit any clear peaks indicating precipitates of the σphase, Laves phase, carbides, or other components. The absence of these peaks suggests that the volume fraction of these precipitates was not sufficiently high to distinguish them from the background. The bright, fine (at 750, 900, and 1050 °C) and long, narrow (at 1150 °C) precipitates contained in the SEM-BSE images were rich in Cr and Mo and poor in Co atoms from the EDS results. This observation was in good agreement with the results reported by Takaichi et al. on the chemical compositions of the precipitates corresponding to the σ σ

γ

γ

he precipitates are composed of large quantities of C, Cr, and Mo, and 23C6

elemental

carbide identified by a previous report [37].

On the basis of the above results, the precipitates are considered to be M(Cr, Mo)23C6 . Ramírez et al. investigated the precipitation mechanism in the Co–25.5Cr–5.5Mo–0.26C biomedical alloy and suggested that M23C6 carbide precipitates formed from the σ → M23C6

σ σ

23C6

carbide (M= Mo, Cr) °C can be assumed to belong to the Laves phase because their Cr concentration was lower than that of the surrounding matrix [30,39]. Precipitates in Co–Cr alloys are well-known to affect the mechanical properties [30,39], wear [40,41] and corrosion resistance [42], and grain refinement characteristics [43] of the alloys. The presence of carbide precipitates at both the GBs and inside the grains can promote the alloy strengthening process [38,40]. The presence of precipitates at the GBs prevents the GBs from sliding and migrating, and a skeletal network of precipitates can resist the imposed stress to some extent [38,40]. At the same time, precipitation inside the grains strengthens the alloy matrix by providing obstacles to the movement of dislocations, thus inhibiting crystallographic slip. However, the formation of coarse M23C6 carbide particles at higher carbon concentrations was found to be detrimental to the ductility of the alloy [44]. On the other hand, Laves phase precipitates are plastically deformable at elevated temperatures and generally act as obstacles to grain growth [39]. However, the relationship between the alloy characteristics and those of the formed precipitates would need to be investigated in more detail in future studies.

4.2. Influence of the heat-treatment temperature on the recovery and recrystallization behavior The KAM maps constructed for the specimens heated to 750, 900, and 1050 °C (which show a difference in the crystal orientations of the adjacent measurement points and are typically used to characterize the residual plastic strain [45,46]) indicate that the magnitudes of the plastic strain inside the grains were relatively high. In addition, the studied specimens contained a large number of LABs inside the grains (subgrains), which are often observed in the deformed microstructure before recrystallization [47]. Some of these regions with high KAMs overlapped with LABs, suggesting that the recovery process occurred after heating to 750, 900, and 1050 °C. The recovery process is defined as the stage during which the deformed grains can reduce their stored energy through the removal or rearrangement of defects in their crystal structure [48]. As recovery proceeds, various structural changes occur. Specifically, dislocations are annihilated and/or rearranged, and subgrains surrounded by LABs subsequently form [49]. In the present study, the recovery process was likely initiated by some defects (primarily dislocations) that were introduced during SLM. The excess dislocations that rearranged during recovery were subsequently transformed into LABs, leading to the formation of subgrains. Since the recovery process reduces the dislocation density, it is normally accompanied by a reduction in the material strength and a simultaneous increase in the ductility. In this work, the 0.2% YS and Vickers hardness values of the alloy specimens were found to decrease, whereas the elongation increased when the heat-treatment temperature increased from 750 to 1050 °C, indicating that the recovery process could proceed at elevated temperatures. In the specimens heated to 1150 °C, equiaxed grains containing a large number of Σ3 annealing twin boundaries were detected, and the residual plastic strain inside the grains was drastically relieved. The strain in a metal structure can be significantly reduced through the stress relief process, which represents the release of internal energy related to the presence of defects in the microstructure (including intergranular boundaries and dislocations) [10]. The results obtained in this study suggest that the alloy homogenization and recrystallization processes were accompanied by the relief of the residual stress

primarily introduced by the formation of dense dislocations during SLM processing. It was shown previously that alloy recrystallization tends to convert LABs (which are mostly composed of dislocations) into HABs [47,48,50], and this trend was also observed in the present study. In general, a large number of Σ3 annealing twin boundaries form in Co–Cr alloys during recrystallization owing to the extremely low SFE of these alloys [51]. It was previously reported that higher temperatures (corresponding to a higher thermal driving force) promoted the formation of Σ3 annealing twin boundaries from the stacking fault bands [51]; therefore, in the present study, temperatures below 1050 °C may not have been sufficient to enhance the recrystallization process. In the MA distribution, the small peak centered at approximately 40° corresponds to Σ Σ

Σ

Σ

Σ

In general, the initiation of the recrystallization process requires the driving force of the stored dislocation energy, which is supplied by the prior plastic deformation process [48]. However, in this work, the SLM-fabricated components were not deformed during SLM. Some researchers reported the accumulation of thermal residual stress due to the repeated rapid heating and cooling of the molten pool during the SLM process, which thus could be considered the driving force of the recrystallization of the SLM-fabricated components during annealing [12,53,54]. Therefore, in the present study, the extremely rapid solidification of the alloy during SLM could also serve as the driving force of the recrystallization process. Recrystallization is usually accompanied by reductions in the strength and hardness of a material and a simultaneous increase in its ductility, which explains why the 0.2% YS drastically decreased and why the alloy elongation increased after the heat treatment at 1150 °C in this study.

4.3. Effects of the heat-treatment temperature on the mechanical properties and fracture mechanisms The SEM images of the fracture surface after the tensile test (Fig. 11) clearly demonstrated that the fracture mechanism was similar among the as-built specimens and specimens heated to 750, 900, and 1050 °C. Takaichi et al., reported that the strain-induced martensite transformation (SIMT) from the γ

phase to the ε phase occurred in the as-SLM Co–Cr–Mo alloy after a tensile test and worked as a dominant deformation mechanism [35]. In this study, the same phenomenon was observed for the specimens heated to 750, 900, or 1050 °C. Cracks propagated along the interface between the γ-FCC matrix and the SIMT on {1 1 1} planes, resulting in the occurrence of a quasicleavage fracture. Kubota et al., reported that this mechanism contributes to work hardening because of the formation of a sessile dislocation by the partial dislocation reaction on {1 1 1} [55]. Accordingly, the specimens heated to 750, 900, or 1050 °C showed limited elongation. On the other hand, for the specimens heated to 1150 °C, the fracture surfaces were covered by dimples, indicative of highly ductile fracture. These observations are in good agreement with the results that the specimens heated to 1150 °C increased their ductility compared to the other heat-treated conditions. The specimens heat-treated at 1150 °C exhibited a lower UTS and 0.2% YS than the as-built specimens. In addition, although the elongation of the specimens heated to 1150 °C (16.3%) was higher than that of the as-built specimens (12.3%), no significant differences between them were noted (p = 0.055). Therefore, in this study, the specimens heat-treated at 1150 °C for 6 h could not sufficiently improve the mechanical properties compared to the as-built specimen. The as-built specimen is exclusively composed of the γ phase, and the specimen heated to 1150 °C consists almost entirely of the γ phase; thus, the reasons for these results are attributed to the differences in the grain size and the precipitate distribution of each specimens. In the as-built specimens, numerous fine cellular dendrites were observed, and small precipitates were widely distributed at the cell boundaries in the grains. These microstructures reduced the dislocation mobility, which contributed to the high UTS and 0.2% YS as well as the hardness values. On the other hand, for the specimens after the heat treatment at 1150 °C, recrystallization occurred, coarse grains without cellular substructures were formed, and precipitates were not observed inside the grains, which significantly reduced the UTS and 0.2% YS. The cracks in these specimens seemed to initiate at or near the GBs and propagate along the GBs decorated with carbides. It was reported that carbides decorating the GBs increased the incompatibility of plastic

deformation across the GBs [44]; thus, intergranular fractures were observed in the specimens heated to 1150 °C. As shown in the CLSM image in Fig. 2d, heating to 1150 °C for 6 h led to grain coarsening. Normally, longer heat-treatment times promote grain growth [10], which in turn decreases the material strength and ductility; thus, the heat-treatment time in the present study could be longer than the optimum time. Hence, additional studies must be conducted to elucidate the relationship between the heat-treatment time and the progression of the recrystallization and grain growth processes (which can potentially affect the mechanical properties of Co–Cr–Mo alloys) to optimize the parameters of the heat-treatment procedure and improve the mechanical properties.

5. Conclusions The microstructures and mechanical properties of Co–Cr–Mo alloy specimens fabricated by SLM were studied before and after heat treatments for 6 h at 750, 900, 1050, or 1150 °C with a primary focus on the stress relief process. We found that increasing the heat-treatment temperature from 750 to 1150 °C enhanced the ductility of the alloy specimens and decreased their 0.2% YS and Vickers hardness values. The as-built specimen was exclusively composed of the γ (FCC) cobalt phase, whereas the heat-treated specimens contained both the γ (FCC) and ε (HCP) phases. The volume fraction of the ε phase decreased as the heat-treatment temperature increased from 750 to 1150 °C and reached almost zero after heating to 1150 °C (i.e., complete transformation), which could contribute to the observed increase in the sample elongation. After heat treatment at 750, 900, or 1050 °C, a large number of LABs and relatively high KAM values were observed, indicating that the recovery process occurred without significant stress relief. On the other hand, the specimens heated to 1150 °C exhibited equiaxed grains containing a large number of Σ3 annealing twin boundaries and drastic relief of the residual stress, suggesting that recrystallization occurred, which was responsible for the observed drastic increase in the alloy elongation and the decrease in strength (as compared to the magnitudes obtained at other temperatures). However, no significant differences between the specimens heated to 1150 °C and the as-built specimens

were noted (p = 0.055), although the elongation of the specimens heated to 1150 °C (16.3%) was higher than that of the as-built specimens (12.3%). We believe that long heat-treatment times promote grain growth excessively, which in turn decreases the material strength and ductility; thus, further investigations are needed to optimize the heat-treatment conditions for improving the mechanical properties.

Acknowledgments This work was partially supported by a Grant-in-Aid for Fundamental Scientific Research (Nos. 17K17152 and 17K17157) from the Ministry of Education, Culture, Sports, Science, and Technology of Japan and a Research Promotion Grant (No. AF-2016231) from the Amada Foundation.

Declarations of interest: none

Fig. 1. Microstructures of the horizontal cross section to the BD of the SLM-processed alloy specimens: (a) CLSM and (b) SEM images. Fig. 2. CLSM images of the horizontal cross section to the BD of the CoCrMo alloy heated to (a) 750 °C, (b) 900 °C, (c) 1050 °C, and (d) 1150 °C. Fig. 3. SEM images of the horizontal cross section to the BD of the alloy specimens heat-treated at (a) 750 °C, (b) 900 °C, (c) 1050 °C, and (d) 1150 °C. Fig. 4. BSE images of the horizontal cross section to the BD of the alloy specimens heated to (a) 750 °C, (b) 900 °C, (c) 1050 °C, and (d) 1150 °C, and their Co, Cr, and Mo contents as determined by EDS. The asterisks denote the analyzed locations, whereas the white arrows indicate large precipitates. Fig. 5. SEM-BSE images and EPMA elemental maps of the horizontal cross section to the BD of the Co–Cr–Mo alloy heated to (a) 750 °C, (b) 900 °C, and (c) 1150 °C. Fig. 6. SEM-BSE images and EPMA elemental maps of the horizontal cross section to the BD of the Co–Cr–Mo alloy heated to 1050 °C. Fig. 7. XRD spectra of the specimens heat-treated at (a) 1150 °C, (b) 1050 °C, (c) 900 °C, and (d) 750 °C and (e) the as-built specimen. Fig. 8. Volume fractions of the ε phase in the alloy specimens heated to different temperatures. Fig. 9. Results of EBSD studies: GB (a, e, i, m, q), IPF (b, f, j, n, r), KAM (c, g, k, o, s), and MA (d, h, l, p, t) distribution maps of the horizontal cross section to the BD of the alloy specimens heated for 6 h at different temperatures: (a, b, c, d) as-built; (e, f, g, h) 750 °C; (i, j, k, l) 900 °C; (m, n, o, p) 1050 °C; and (q, r, s, t) 1150 °C. Fig. 10. Typical stress–strain curves obtained for the alloy specimens heat-treated at (a) 750 °C, (b) 900 °C, (c) 1050 °C, and (d) 1150 °C. Fig. 11. Fracture surface morphologies of specimens heated to different temperatures after tensile testing: (a) as-built, (b) 750 °C, (c) 900 °C, (d) 1050 °C, (e) 1150 °C, and (f) 1150 °C at higher magnification.

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