summary
For alloys that undergo a solid phase transition when cooled, the lack of chemical segregation information at the time of solidification makes it challenging to eliminate hot cracks. One such material is C465, a high-strength martensitic aging steel that is prone to hot cracking. Here, we solve the problem of thermal cracking under the laser powder bed fusion process by adding titanium nitride (TiN) particles. During the solidification process, tin promotes grain refinement and reduces the formation of liquid films with low solidus temperatures. Under cooling conditions, the partial dissolution of tin lowers the starting temperature of martensite and produces more residual austenite. In the subsequent annealing, the dissolved titanium slows down the austenite recovery kinetics, while the dissolved nitrogen increases the yield strength. The material under tensile deformation follows a three-stage work hardening behavior, indicating that strain induces a martensitic transition. This work emphasizes that, in addition to the grain refinement effect of nucleating agents, the effects of partial dissolution during processing are critically examined when dealing with the problem of thermal cracking in alloys prone to phase transformation.
introduction
Conventional high-strength steels (HSSs), such as precipitation-hardening stainless steels 17–4ph (Fe-17Cr-4Ni-4Cu, wt. % and martensitic aging steel 18Ni-300 (Fe-18Ni-9Co-4Mo, wt. % has been widely researched and adopted by additive manufacturing (AM), but state-of-the-art HSS grades are not widely accepted. One of the main reasons for the difference in adoption rates is the presence of cracks during rapid solidification, as is the case with many other alloy systems. However, unlike most crack-sensitive materials such as aluminum and nickel alloys, steel typically undergoes multiple phase transitions during the manufacturing process, making the task of mitigating cracks in steel even more daunting. Custom 465 (C465, Fe-11Cr-11Ni-1.5Ti-1.0Mo, wt%) Since 2010, high-speed steels have attracted increasing attention from academia and industry due to their superior corrosion resistance compared to traditional high-strength stainless steels such as 18Ni-300 [4]. Since it can operate without the need for protective coatings, such as toxic cadmium and chromium, it provides a more sustainable environmental footprint. Among stainless steels, C465 is reported to have the highest strength. e. Approximately 50 % higher than 17–4 PH (1751 vs. MPa) with considerable corrosion resistance and toughness. If these characteristics are combined with the design freedom of AM, it is possible to envision several new applications and functions, such as shock-absorbing foams on high-performance vehicles, complex drill bits with internal cooling channels, and complex high-thrust ship propellers. The conventional manufacturing (CM) route for C465 consists of three processing steps after casting. (1) The alloy is first solution annealed at around 900 °C for 1 hour to obtain a fully austenitic centricity cubic (FCC) phase. (2) The material is then deeply quenched to 73 °C and held for 8 h to facilitate the martensite transition. (3) A final aging treatment is then undergo at a temperature between 480 and 650 °C to fine-tune the hexagonal close-packed (HCP) η-Ni3Ti to strengthen the precipitate and restore the properties of the austenite. In general, the volume fraction of recovered austenite increases with increasing aging temperature, and the maximum volume fraction of recovered austenite is about 14-19% under the peak condition that occurs at around 650°C. These recovered austenite grains preferentially nucleate at the pre-austenite or lath boundary of the parent martensite phase and typically have a lower dislocation density compared to their neighboring martensite. η-Ni3Ti precipitates nucleated mainly in martensite, following the orientation relationship of { 011 } α′/ { 0001 }η-Ni3Ti and 〈111α′/〈1120〉η-Ni3Ti. These precipitates have a rod-like morphology and range in length from about 5 to about 180 nanometers depending on the aging treatment employed. Reverting to changes in austenite content and precipitate morphology results in a variety of tensile, corrosion, and hydrogen resistance.
In this study, we attempted to fabricate C465 using laser powder bed fusion technology (LPBF). Our first results show that this alloy is highly susceptible to thermal cracking during LPBF. The presence of thermal cracks (also known as thermal tearing) in AM alloys is often attributed to a segmentation-induced liquid film at the end of solidification, where the solidus temperature is lower than that of its surrounding material. In previous AM projects, several thermal crack mitigation methods have been developed. Contis et al. prevented thermal cracking in nickel-base superalloys by reducing volumetric heat input, thereby reducing elemental partitioning and preventing the formation of a liquid film at low solidus temperatures. Sun et al. introduced secondary precipitates along grain boundaries during solidification to facilitate dendrite bridging and thus avoid thermal cracking. Sun et al. also eliminated hot cracking in nickel-based superalloys by controlling competitive solute partitioning with the help of thermodynamic calculations. Opprecht et al. addressed hot cracking in aluminum alloys through grain refinement.
Despite these early successes, as discussed below, the direct application of the hot crack mitigation methods described above is still problematic for steels that are prone to hot cracking. This is because modern high-strength stainless steels (mainly martensitic aging grades) undergo multiple phase transitions (unlike aluminum or nickel alloys that always have a face-centered cubic matrix), which can include primary and secondary phase formation during solidification, martensitic phase transitions during cooling, austenite recovery upon subsequent heating, and possible deformation-induced martensite phase transitions upon loading. Therefore, it is not straightforward to retrieve elemental assignment information experimentally immediately after curing. This, in turn, greatly hampers efforts to eliminate thermal cracks in the AM process through alloy design. Through the experimental work reported in this paper, we use C465 alloy as an example to highlight the complex nature and important aspects that need to be considered to solve hot cracks in high-strength martensitic aging steels. We will begin by covering the severity of thermal cracks in as-built materials and the limitations of several existing thermal crack elimination methods. A crack-free specimen was obtained after the introduction of tin particles into the precursor steel raw material powder. In addition, the effects of TiN on microstructure and tensile properties are studied and discussed in detail. We believe this work will not only help with the use of high-strength stainless martensitic aging steels for AM, but will also provide guidance for any alloy that undergoes a phase change in a production process involving rapid solidification.
Experimental setup
Commercial C465 alloy powders (size range: 15 to 45 microns) are obtained from Carpenter Additive. The chemical composition of the powder is listed in Table 1. A Trumpf TruPrint 1000 machine equipped with a laser with a wavelength of about 1070 nanometers (beam diameter of 30 microns) was used for sample production. Initial laser parameter screening was performed using a fixed laser power of 175 W, a powder bed thickness of 30 μm, and a hatch spacing of 100 μm. The laser scanning speed varies between 130 and 600 mm/s (8 uniform spacings), resulting in a volumetric energy density of 448 to 97 J/mm3. Cube specimens of 10 × 10 × 10 mm3 were fabricated to detect the microstructure and internal defects of the finished sample. Rectangular blocks with dimensions 10.6 × 26 × 10.6 mm3 are produced for subsequent tensile tests. Tin particles from Luoyang Tongrun Nanotechnology were then added to the C465 powder via a BMU-100–3 roller mill, which ran at 180 rpm for 12 hours. Scanning electron microscopy (SEM) imaging shows tin powder in the form of flakes with a size range of about 100 to 500 nm. All samples are produced using tool steel substrates.
The effects of austenite recovery and microstructure changes on mechanical properties were studied by heat treatment in QSXL-1616 chamber furnace. Heat the sample (at a rate of 10 °C/min) to the annealing temperature t a (range 360 to 630 °C), hold for 4 h, and then cool with the furnace. Microhardness was measured using an Innovatest Falcon 5000 machine at a load of 300 g and a hold time of 15 s. There are at least 12 indentations on each sample over the entire build height. Tensile specimens are processed into gauge lengths of 20 mm, gauge widths of 8 mm, and gauge thicknesses of 1 mmtensile tests performed on Instron Mechanical Tester 5569 machines with a nominal strain rate of 0.30 mm/min. The as-built sample was removed from the substrate using electrical discharge machining (EDM) and mounted in PloyFast carbon resin. The samples were then ground with silicon carbide paper with particle size numbers #220, #500, #1000, #2000, and #4000. This was followed by polishing with No. 3 and No. 1 micron diamond suspensions, followed by vibration polishing with colloidal silica suspensions prior to light microscopy (OM) and SEM observation. An additional electropolishing step (24 V and 0.7 A) was performed using Struers A2 electrolyte for electron backscatter diffraction (EBSD) studies. This procedure is performed to remove any possible grinding/polishing-induced martensite formed on the outer surface of the polished sample. By comparing the phase fraction between the EBSD results and the X-ray diffraction (XRD) data, it was determined that the electropolishing duration of 2 min was the optimal time for the process. EBSD data were obtained using a JEOL JSM-IT500HR microscope fitted with an Oxford symmetry detector. Overview scanning was performed with an accelerating voltage of 20 kV and a step size of 0.5 μm. For detailed scans with an area of less than 10 × 10 μm2, use a step size of 0.05 μm. All scans were scanned at 156 × 128 pixels at Ningbin size and acquired at 246.5 Hz. Oxford Aztec software is used for data collection. MTEX software was used for EBSD analysis and reconstitution of the original austenite grains. Backscattered electron (BSE) imaging and energy dispersive spectroscopy (EDS) mapping were also performed. XRD characterization was performed on a Bruker D8 discover diffractometer with Cu Kα radiation and a 2D vntec detector. The exposure time for each frame is set to 100 seconds. The phase fraction was estimated using the Rietveld fine method by TOPAS v5. Thin sections for scanning transmission electron microscopy (STEM) studies are prepared using a lifting technique using a focused ion beam (FIB, Zeiss Beam 540). Prior to the procedure, a 1 micron thick protective layer of platinum (Pt) was placed on the sample by ion beam-assisted chemical vapor deposition. Nanoscale microstructure characterization was performed at 200 kV on a FEI Titan 80–300 TEM. TEM EDS plotting was carried out using a Tecnai-20 operating under the same conditions. Thermodynamic calculations were performed using Thermo-Calc 2022b software and the TCFE12 database. The code written in-house is executed in console mode to calculate the partition coefficient for individual solute elements.
Results & Discussion
As mentioned earlier, all finished C465 specimens exhibited extensive cracking, despite the extensive laser processing parameters. Figure 1(a) shows a representative light micrograph of the as-built sample. The crack density found was 4.6 mm/mm2, which is comparable to that found in LPBF nickel-base superalloys [3]. As shown in Figures 1(a) and (b), most of these cracks are parallel to the direction in which the sample was built, and their lengths range from about 100 μm to about 1 mm. The surfaces of these cracks observed after tensile fracture are mostly smooth, and they mimic the shape of solidified dendrites, which confirms that these cracks are thermal cracks (also known as solidification cracks), Figure 1(c). The origin of this crack is usually due to the presence of a liquid film with low solidus temperature, which cracks under the combined influence of tensile residual stresses upon curing [2,3,12]. When viewed in BSE mode in Figure 1(d), there are bright contrast lines evenly distributed within the as-built material. These characteristics are most likely caused by significant changes in chemical composition resulting from the assignment or separation of elements. Nano-sized particles were also observed to be uniformly distributed across the entire cross-section.
By EBSD experiments, approximately 4 % residual austenite was detected near the 2 thermal cracks in the as-built C465 in Figure 2(a). In addition, slatted martensite is often observed in the inverse pole diagram (IPF) of the out-of-plane viewing direction in Figure 2(b). Various blocks and packages containing these slats are clearly observed. The iterative method proposed by Nyyss onen et al. was used to reconstruct the original austenite grains by MTEX software. In this method, the boundary orientation difference between the product martensitic variants is calculated sequentially. These individual orientation errors are then iteratively compared to the initially assumed orientation relationship (OR), such as Kurdjumov-Sachs (K-S) or Nishiyama-wasser Mann (N-W). The OR is constantly updated with the aim of obtaining the best match with the wrong orientation in the microstructure. There is always a deviation between the theoretical and experimental orientation error values, most likely due to local plastic deformation during the phase transition. Once the ideal OR has been determined, orientation analysis is performed by coloring all product martensitic variants based on their different phase transition variants, Figure 2(c). The orientation of the pole plots of all martensite variants within one parent austenite grain surrounded by a solid white line in Figure 2(c) with respect to the parent austenite is plotted in Figure 2(d). The colored data points correspond to the data captured experimentally, while the black hollow circles are theoretical calculations that follow K-S OR or . The high degree of agreement between the simulated and experimental data confirms that the current alloy follows the K-S OR during the martensitic transition. The reconstructed protoaustenite grain map is highlighted with a solid blue line within high-angle grain boundaries (HAGBs) above 15°, Figure 2(e). Thermal cracks were found to exist only along the original austenite grain boundaries, as indicated by the white arrows.
At higher magnifications, the residual austenite grains were found to be elongated and all appeared to have a common orientation of arrangement, as shown in Figure 3(a). Based on these characteristics, it is assumed that the bright contrast line observed under BSE conditions in Figure 1(d) is austenite grains. Most of the residual austenite grains are located on the boundary of the martensitic grains, see Figure 3(b). Some intracrystalline austenite grains circled by black dotted boxes were also observed. As expected, all residual austenite grains have the same orientation as the reconstituted original austenite, as shown in Figure 3(c).
Figure 4 identifies the dilution of iron and the enrichment of titanium in the vicinity of the thermal crack by means of an EDS profile. Chemical segregation of smaller precipitates in the range of hundreds of nanometers is not detected here, possibly due to insufficient spatial resolution of SEM-based characterization techniques. It is well known that thermal crack susceptibility is highly correlated with elemental assignment during solidification. Therefore, the Ti segregation observed here is inferred to be the main cause of crack occurrence.
A simple and effective way to reduce this harmful element distribution and prevent thermal cracking is to reduce the total heat input, which has been successfully demonstrated in nickel-based superalloys that are prone to thermal cracking fabricated using LPBF. However, this method does not work well for current alloys because cracks persist even when the laser power is reduced from 175 to 130 watts and the powder thickness is reduced from 30 microns to 10, while other manufacturing conditions remain constant, Figure 5(a). A total of 5 samples (130 W) with low laser power were fabricated with different scanning speeds, all of which had a high density of thermal cracking, suggesting that reducing heat input does not help to alleviate the cracking susceptibility of the current C465 alloy. Crack-free samples are introduced only at 1 wt. Figure 5(b) shows that 90% of the tin particles are added to the raw material by powder mixing prior to the manufacturing process. The amount % of 1 wt as the optimal amount of tin is inferred from the previous literature. 1.5 wt% additional experimental % of tin particles, the principle of selecting tin particles to eliminate cracks will be further elaborated in the discussion section.
All manufacturing alloys (with and without tin) are annealed over a wide range of temperatures. The samples were kept at a specific temperature for 4 hours, then cooled in a furnace and their microhardness values were measured, as shown in Figure 6. In as-built conditions, the material has similar microhardness of 302.6±7.7 HV (without tin) and 305.2±8.8 HV (with tin). For TA at 420°C, the heat-treated C465 alloy (without tin) achieves a peak hardness of 517±26.0 HV, and the high hardness remains relatively constant at TA up to 510°C. A further increase in TA led to a decrease in microhardness, which eventually became 337±12 after annealing at 630°C. 3 HV。 % of tin has a lower peak hardness of 470±21.1 HV, which is reached at a ta of 450 °C, which is slightly higher than the observed temperature of C465 without tin (420 °C). In addition, the hardness value of this material is more sensitive to temperature, and even the slightest change in TA can affect its hardness. The microhardness of the alloy continued to decrease until the lowest value (336.1±7.0 HV) was reached at TA = 630°C, when TA exceeded 450°C. Similarly, under over-annealing conditions, the microhardness of both alloys is similar.
The fractions of austenite and martensitic phases present in the microstructure of the C465 alloy without tin were investigated using the XRD technique in Figure 7(a) after annealing under different TAs. At a low t a of 360 °C, a sharp increase in the austenite volume fraction Vγ was observed, from about 3 % to about 17 % under as-built conditions. With the increase of ta, Vγ rises monotonically, reaching about 35% at a TA of 570 °C. After this point, the austenite recovery rate appears to increase again, reaching 55 % for Vγ at a TA of 630°C. Figure 7(b) shows representative XRD curves for the material under as-built conditions, peak microhardness conditions (420°C), and over-annealing conditions (630°C). The intensity of the austenite peak increases at the same time as the TA.
1 wt of phase fraction estimation of the material. Due to the overlap between the austenite (220) and tin (311) peaks, the peak around 74♀ is significantly broadened (not shown here). Therefore, we rely on EBSD to quantify the phase composition in tin-fortified samples. Figure 8 shows the IPF plot of C465 with and without tin. For each case, these diagrams show (1) reconstituted protoaustenite grains, (2) associated as-built microstructures, and microstructures after annealing at (3) 420°C and (4) 630°C, as well as associated phase diagrams. Table 2 gives statistics on phase fractions and average austenite grain sizes. Join 1 WT. % TiN, the reconstituted protoaustenite (about 40 μm diameter) has a smaller grain size compared to the pure C465 specimen (ø ̃ 75 μm), see Figures 8 (a1) and (b1). The morphology and size of the austenite phases in the two alloys are similar in the finished state, but the alloy containing tin particles has a higher residual austenite mass (9.8 %), compared to 4.1 % for finished pure C465, as shown in Table 2. In addition to affecting the original austenite grain size and residual austenite volume, the addition of tin appears to have slowed down the dynamics of austenite recovery in current steels. After annealing at 420°C and 630°C, the austenite content of pure C465 increased significantly to 20.2 % and 69.7 %, respectively. However, they only reach 17.2 % and 38.2 % in the Gatin group, although the austenite volume is slightly larger under as-built conditions. In addition, the average grain size (area) of pure C465 austenite was almost 5 times higher than that of the tin-added sample under over-annealing conditions (630°C) (89.6 μm2 vs. 19.0 μm2), indicating that tin severely limits the growth of revertant austenite grains.
In the construction state, the microstructure of C465 without the addition of tin consists of a high density of dislocations and sporadic distribution of particles, Fig. 9(a1). The average width of the martensitic slats is measured at about 500 nm. Near-spherical precipitates were found to form preferentially near the slats border, as indicated by the yellow arrows. The displacement properties of martensitic phase transitions require shear deformation and volume expansion. Due to the limitation of the surrounding austenite grains, the dislocation density in martensite is high. In addition, the repetitive thermomechanical cycling experienced by the deposited layer during LPBF also results in a high density of dislocations, Figure 9(a2). EDX point analysis of the near-spherical particles showed that they were predominantly titanium-based oxides with a composition of 50.1 titanium-35.3 iron-2.8 nickel-1.7 chromium (at. %), while the matrix remains 72.6 Fe-13Cr-12Ni-1.6 Titanium-0.8O(at. %), Figure 9 (a3). The presence of oxide particles in LPBF steels is not uncommon, as the outer surface of the raw metal powder usually contains a thin layer of residual oxide. After heat treatment at 630 °C, a uniformly distributed rod-like precipitate was detected within the martensite, Figure 9 (B1). The length of these precipitates is in the range of 20 to 100 nm, while their width varies between 5 and 15 nm, Figure 9 (B2). The selective electron diffraction pattern (SAEDP) obtained on a matrix with precipitate (Figure 9(B3)) shows a mixture of martensite, HCP η-Ni3Ti, and double diffraction spots, which matches well with the previous computer simulation diffraction pattern. The inferred orientation relationship between martensite and precipitate is (011)BCC//(0001)η, BCC//], η, which is very consistent with the reports of conventionally manufactured C465 after annealing.
Join 1 WT. % TiN, the constructed C465 also contains high-density dislocations, Figure 10(a). At the same time, nano-sized precipitates appear in the intra-grain and inter-grain regions. These precipitates appear to be remnants of those precursor tin particles, as they are much smaller in size (about 20 to 30 nanometers after fabrication and about 100 to 500 nanometers before melting). In addition, previous studies have shown that tin is unlikely to completely dissolve and reprecipitate near the melting temperature of the steel. After 4 h of heat treatment at 420°C, a high density of dislocations remained in the material, as shown in Figure 10(b). Figure 10(C) shows that by manually analyzing more than 20 individual precipitates, the average precipitate radius was found to be 22±5 nm (at as-built) and 31±7 nm (after heat treatment at 420°C). The area fraction of the precipitate decreased from 2.8±0.2% under as-built conditions to 1.7±0.1% after heat treatment, indicating that the particles in the matrix were partially dissolved during heat treatment. This hypothesis was confirmed by the TEM EDS results, as the N content in the matrix increased to approximately 0.83 wt. % after heat treatment and almost non-existent before annealing, see Table 3. During heat treatment, the size of the tin particles increases due to Ostwald curing, while the concentration of Ti and N in the precipitate increases due to the exclusion of other solutes into the matrix. It is likely that the dislocations observed in Figure 10(b) are dissolved at high densities however, direct experimental confirmation of this phenomenon is challenging and beyond the scope of this work. The EDS profile of the precipitate showed an enrichment of Ti and N, but the depletion of other elements confirmed that the particles were indeed tin.
Unsurprisingly, the tensile properties of LPBF C465 without tin were poor (yield strength (YS) of 149±20 MPa, ultimate tensile strength (UTS) of 234±23 MPa, and elongation at break values of 1.0±0.1% under as-built conditions, which remained unchanged after different heat treatments) due to the presence of a wide range of thermal deformation cracks, Fig. 11(a). Join 1 WT. % of tin, a significant improvement in tensile properties was noted, Figure 11(b). YS and UTS were 548±21 MPa and 952±11 MPa, respectively, with a total elongation of 11.8±0.5%. After annealing at 420 °C, the strength increased further (846±31 MPa for YS and 1207±35 MPa for UTS), while ductility decreased slightly (9.7±1.1%). A further increase in TA results in a decrease in strength and an increase in ductility at the same time, which can be attributed to the higher volume fraction of the reverted soft austenite. Now join 1 WT. Although the crack density is much lower than that of pure C465 (i.e., 0.1 mm/mm2 vs. 4.6 mm/mm2), % Tin vs. C465 does not completely eliminate crack formation during LPBF. In order to check that a greater tin content (1.5 wt%) will result in a completely crack-free sample, additional manufacturing was made. A high-density sample without thermal cracks was obtained, with a density value greater than 99.9 % (based on OM analysis) and the resulting tensile properties are shown in Figure 11(c). With 1.5 wt due to the complete elimination of thermal cracks. % of tin has better UTS and elongation values in its as-built condition (black curve), corresponding to 1 wt. Percentage of tin. At the same annealing temperature of 450 °C (1 wt. of the highest hardness value of the C465 alloy. % tin, blue curve), with 1.5 wt. % Tin has a slightly lower UTS (1111±8 MPa vs. 1186±37 MPa), but much higher ductility (20.6±2.1% vs. 12.0±0.8%). Compared to the peak aging C465 made from conventional process routes, the current 1.5 wt. After annealing at 450°C, the sample doped with % tin had lower strength (1100 MPa vs. MPa) but higher ductility (20% vs . 10%）。 It is assumed that the optimal TA conditions for the current alloy should be less than 450°C, and there is still room for improvement in its tensile strength value.
In the current work, TiN particles were added to the C465 HSS to address the problem of thermal cracking during AM production. While crack-free samples with excellent tensile properties can be obtained with this method, the addition of tin has several other effects. Here, we mainly aim to discuss three related questions, namely, (1) the mechanism of crack elimination due to the introduction of tin, (2) the effect of tin on the volume fraction of residual austenite and reverted austenite (after cooling and subsequent heating, respectively), and (3) the effect of microstructural changes on tensile behavior. The underlying mechanisms of thermal crack development are typically attributed to a combination of (1) low-melting liquid films at grain boundaries, initiated by solute partitioning during solidification, and (2) inherent tensile residual stresses within LPBF parts. A quick way to solve the problem of thermal cracking is to limit the extent of solute partitioning by reducing heat input (e.g., reducing laser power or increasing laser scanning speed). This method has been shown to be effective in several nickel-based superalloys, but it does not work well with current materials (Figure 5(a)). The difficulty in solving the hot crack problem, not only for this material, but also for other alloys that undergo a solid-state phase transition during cooling, is the lack of solidification state assignment information that can be retrieved experimentally. These data are critical for alloy design, both to remove hot crack promoters and to reduce their harmful effects by adjusting the content of other solute elements. It is recognized that certain computational methods (e.g., DICTRA) are useful in predicting this information, but they in turn require a highly accurate thermal history of the material, which is difficult to obtain and often yields lower prediction accuracy for low concentrations of elements. As a result, these simulation tools are largely used as qualitative guidance only. Figure 12(a) shows the solidification path of the current C465 alloy under Scharill-Gulliver conditions. This formula assumes that there is no diffusion in solids and that solutes in liquids redistribute infinitely and rapidly. Guided by the EBSD results, only austenite, ferrite, and liquid phases were enabled during the simulation (Figure 2). Most of the alloys (about 0.67 mole fractions) solidify as single-phase FCC, and the BCC phase is formed only near the end of solidification. The partition coefficient for a single element is calculated as k = XSXL, where XS and XL are the concentrations of that element in solids and liquids, respectively. Ti was found to have a very low partition coefficient in the current alloy, with values between 0.2 and 0.4 over the entire solidification range, Fig. 12(b). Therefore, it is believed that the strong distribution of Ti during LPBF is the root cause of the hot cracks in this alloy, which is consistent with our previous EDX plotting results, which showed Ti enrichment around the hot cracks, Figure 4. With such a high dispensing propensity, it is likely that simply reducing the heat input will not be sufficient to prevent the enrichment of residual liquids along the original austenite HAGBs. As a result, the solidus temperature of the residual liquid film decreases
This results in thermal cracking, as shown in Figure 5(a). In fact, after analyzing 13 possible alternative solutes in steels with a primary austenite solidification phase, titanium may be the most unfavorable candidate element considering minimizing thermal cracking, see Appendix 1. Another way to solve the problem of hot cracking is grain refinement, in which the composition of the original alloy remains unchanged. This method has proven successful in several LPBF aluminum alloys. The smaller grain size increases the density of HAGBs in the as-built sample, which is beneficial to reduce the solute enrichment of harmful elements and the residual stresses acting on a single GB, thereby improving the thermal cracking resistance of the material. Since the current alloy is mainly solidified into the BCC phase at the end of solidification (Fig. 12(a)), an effective ferrite grain refiner is required. The classical theory of grain refiner selection was originally based on the "lattice reorientation" approach, where the one-dimensional reorientation is defined as δ = δα0α0, which is the difference between the lattice constants of the low exponential planes of the crystal nucleus and the nucleating agent. Smaller values indicate a higher ability to refine grains. Bramfitt extended the method to two-dimensional space and ranked the effectiveness of several potential nucleating agents for δ-ferrite as ZrC, WC, ZrN, SiC, TiC, TiN. Subsequently, an edge-to-edge matching (E2EM) model was developed to illustrate a nucleating agent with a crystal structure that differs from that of the base alloy. Under this approach, the recommended nucleating agent for δ-ferrite is TiC"ce2o3"TiN"CeS"NbO, and NbO is the most effective candidate. Other possible nucleating agents, such as LaB6 and Ce2O2S, were also considered. Tin was chosen in this study mainly because of its high grain refinement capability and availability. Theoretically, however, the other compounds mentioned above should serve the same purpose of grain refinement and at least partially solve the problem of hot cracking in current alloys.
conclusion
In conclusion, this work highlights the potential factors that need to be considered when trying to solve the problem of thermal cracking in alloys prone to phase transformation during LPBF. The solid-state phase transition during cooling eliminates any pre-existing elemental assignment information, which is critical to solving the hot cracking problem. Based on the stainless steel martensitic aging steel C465 examined in this study, the main points are as follows.
(1) Although it helps to form the required η-Ni3Ti strengthening precipitate in steel, Ti has a strong distribution tendency during the solidification process, and the partition coefficient is close to 0.2. As a result, it splits along the HAGBs and forms a low-solidus liquid film, which initiates thermal cleavage.
Simple machining improvements, such as reducing heat input alone, are not enough to prevent it from splitting and eliminating thermal cracks. (2) The grain nucleating agent TiN can effectively reduce the grain size of the original austenite, thereby solving the problem of hot cracking of C465. However, unlike materials that do not undergo phase change during service, the grain refiner selection of phase change materials such as martensitic aging steel needs to consider several other factors besides grain refinement effectiveness.
(3) During the cooling process, the partial dissolution of Ti in the tin particles in the matrix reduces the martensitic starting temperature of the alloy, resulting in more residual austenite. In addition, this change in matrix composition reduces the equilibrium austenite content at high temperatures above 400°C. Coupled with the Zener pinning effect of nano-sized tin, the austenite recovery process is dynamically limited during heating.
(4) The introduction of tin also affects the tensile properties of the alloy. When annealing the finished sample with tin, the dissolved nitrogen produced by the partial dissolution of the tin particles helps to improve the yield strength. All tin-doped materials in the current study have undergone a typical three-stage work hardening behavior, indicating that strain-induced martensite phase transitions occur due to the low austenite stability. In order to ensure high elongation after uniform stretching, completely crack-free samples are required.
Thesis information
Laser powder bed fusion of crack-susceptible stainless maraging steel undergoing solid-state phase transformations
https://doi.org/10.1016/j.actamat.2023.119534
The copyright of this article belongs to the original author, only for communication and learning, and the final interpretation right belongs to this official account (laser manufacturing research).
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