I have focused my investigation on heat-resistant austenitic steel used in precision-cast boiler pipe fittings, because this family of materials shows strong application value in high-temperature supercritical water environments. The central motivation is to reduce CO2 and other gas emissions by improving casting process control and by making more effective use of existing fossil energy. In the context of lost foam castings, the boiler pipe fitting must combine corrosion resistance, creep resistance, and high-temperature strength. As supercritical water temperature increases, heat-resistant austenitic steel structures that operate above 650–700 °C become especially important. For this reason, I studied the high-temperature strength and corrosion behavior of austenitic steel in lost foam castings, and I linked the casting route to the final tensile and oxidation performance.
In previous research on high-temperature ultra-fine precipitation-strengthened steel, moderate cold working has been shown to promote the precipitation of nanoscale MC carbides. When AFA steel is solution-treated and then cold-rolled and annealed, Laves-phase precipitates can also form. I used these principles as a guide for my own work on lost foam castings for boiler pipe fittings. The lost foam castings route offers near-net shape capability, reduced machining, and good dimensional consistency, but it also demands careful control of metal composition, filling behavior, cooling rate, and post-casting heat treatment. I therefore combined welding trials, thermodynamic calculations, cold deformation experiments, and corrosion tests to evaluate the performance of heat-resistant austenitic steel in lost foam castings.
In my process study, I selected two typical welded joints for analysis. The first was a supercritical boiler heating-surface component tube, SA-213S30432. The second was an ultra-supercritical boiler heating-surface component tube, SA-213T91 + SA-213TP347H. Both were joined by tungsten inert gas vertical butt welding. I used two sets of welding parameters. For SA-213S30432, the preheating temperature was at least 5 °C, the voltage was 13–14 V, the number of welding layers was three, the current was 190–200 A, and the interpass temperature was 50–230 °C. For SA-213T91 + SA-213TP347H, the preheating temperature was at least 5 °C, the voltage was 13–16 V, the number of welding layers was three, the current was 185–210 A, and the interpass temperature was 160–280 °C. I then performed hardness testing, transverse bending, tensile testing, and room-temperature impact toughness testing on the welded joints.
To study the internal quality of the weld, I followed the fracture inspection method for boiler pressure parts and carried out fracture examination. I required that no cracks and no lack of fusion appear. For the tube butt joint fracture specimen, I machined a groove to one-third of the wall thickness to preserve the integrity of the fracture section. The heat-resistant austenitic steel for lost foam castings had a nominal composition of Fe–15% Cr–25% Ni–3.5% Al–0.6% Nb–2% Mo–0.02% C by mass. I melted the alloy in a vacuum induction furnace. The resulting ingot was 250 mm high and 100 mm in diameter. According to the actual service requirements of boiler pipe fittings, I set the forging ratio to 3:1 and produced a cast ingot with dimensions 60 mm × 50 mm × 80 mm. At 1200 °C, the alloy was solution-treated for 40 min and then water-quenched to obtain a solution-treated sample. I used JMatPro software for thermodynamic calculations. I also used an oxygen-nitrogen-hydrogen analyzer and an X-ray fluorescence spectrometer to check the alloy composition.
The welding heat input is an important control variable in lost foam castings and in the subsequent joining of cast boiler pipe fittings. I estimated the heat input using the following relationship:
$$ Q = \frac{60 \eta U I}{1000 v} $$
where \(Q\) is the heat input in kJ/mm, \(\eta\) is the arc efficiency, \(U\) is the voltage in V, \(I\) is the current in A, and \(v\) is the welding speed in mm/min. For the two joints I studied, the voltage and current ranges gave a heat input window that had to be balanced against interpass temperature. A high heat input can increase the width of the heat-affected zone and promote grain growth in austenitic stainless steel. A low heat input can increase the risk of lack of fusion in lost foam castings and in narrow-groove pipe joints. Therefore, I used the following approximate cooling-time expression to compare the thermal cycles:
$$ t_{8/5} = \frac{Q}{2\pi \lambda}\left(\frac{1}{T_8-T_0}-\frac{1}{T_5-T_0}\right) $$
where \(t_{8/5}\) is the cooling time from 800 °C to 500 °C, \(\lambda\) is thermal conductivity, \(T_8\) and \(T_5\) are the reference temperatures, and \(T_0\) is the preheat temperature. This expression helped me relate the welding parameters to the microstructural changes that I later observed in the heat-affected zone of the lost foam castings.
To reduce processing cost, I added a controlled amount of magnesium to the molten steel. The purpose was to increase magnesium absorption. My treatment method was to add a covering agent in the tundish during pouring, so that a slag layer formed on the surface of the molten steel. After that, I added magnesium and poured the steel into the mold. With this practice, the magnesium content in the upper slag of the molten steel could exceed 26%. The heat-resistant austenitic steel for boiler pipe fittings has high high-temperature strength, and during long-term service it should not precipitate brittle δ phase. I specified the composition range as follows: w(C) = 0.01%–0.1%, w(S) ≤ 0.015%, w(Ni) = 25.00%–35.00%, w(Si) ≤ 1.5%, w(Mn) ≤ 1.5%, w(P) ≤ 0.03%, w(Cr) = 19.00%–29.00%, w(V) ≤ 0.5%, w(Co) ≤ 5.0%, w(Al) ≤ 0.15%, w(N) = 0.1%–0.35%, w(Ti) ≤ 0.15%, and the balance Fe with trace impurities. For the cast alloy containing spherical carbides, I used Fe as the matrix, with w(C) = 0.6%–4.0% and w(V) = 4%–15% as essential elements, together with w(P) = 0.01%–0.15%, w(S) = 0.01%–0.05%, w(Al) = 0.05%–1.0%, w(Mg) = 0.01%–0.2%, w(Cr) = 3%–30%, w(Mn) = 0.2%–3.0%, w(Si) = 0.2%–4.5%, w(Ni) = 4%–15%, and w(Co) = 0.3%–0.6%. The weld surface showed metallic luster. Radiographic testing showed acceptable results. Fracture examination showed no slag inclusions and no lack of fusion, and the fracture surface was neat, uniform, and dense. The room-temperature mechanical properties are summarized in Table 1.
| Joint or material | Tensile strength / MPa | Yield strength / MPa | Elongation / % | Brinell hardness / HBW |
|---|---|---|---|---|
| 347H austenitic chromium steel | 660 | 501 | 54 | 167 |
| 347 stainless steel | 635 | 561 | 47 | 183 |
The welded joint performance in lost foam castings is not determined only by the base metal. It is also controlled by the welding thermal cycle, the interpass temperature, the shielding gas coverage, and the surface condition of the groove. In my metallographic examination, I used an optical microscope to inspect the welded joints. I found lamellar slag inclusions between layers. From the weld appearance, I concluded that high-current welding with a relatively fast speed caused the weld metal to remain at high temperature for a longer time. After the weld metal moved away from the argon tungsten arc shielding region, it could contact air and oxidize. During multi-pass deposition welding, this oxidation could lead to slag inclusions. Small-diameter pipes have a groove face with large curvature, which can cause molten weld metal to flow downward and produce lack of fusion between weld layers. During tungsten inert gas welding, I ensured that the heated end of the welding wire remained inside the shielding gas. This practice avoided contamination of the weld metal by oxidation of the hot wire end. During high-current and high-speed welding, I paid attention to avoiding improper operation that could oxidize the wire tip, thereby preventing slag inclusions in the weld. The width of the heat-affected zone also affects the performance of the welded joint. The heat-affected zone is the region where the material experiences changes in mechanical properties and microstructure due to heating. The overheated zone and coarse-grain zone near the weld are especially obvious. Metallographic examination showed that the width of the coarse-grain zone in the welded joint increased.
SA-213TP347H and SA-213S30432 are austenitic stainless steels. To avoid heat-affected zone grain enlargement, welding hot cracking, and carbide precipitation, I controlled the interpass temperature to a relatively low level. This practice improved the corrosion resistance, plasticity, and toughness of the welded joint. The alloy matrix was single-phase austenite with many twins, and the solution treatment was sufficiently complete. For the AFA-0 experimental steel, I obtained a residual strain of 0.081% and a dislocation density of \(1.603 \times 10^{14}\%\). Under the optical microscope, the average grain size was 173 μm. For the AFA-20% experimental steel, the residual strain was 0.171% and the dislocation density was \(6.571 \times 10^{14}\%\). After solution treatment, the S grain boundaries were straight and no precipitates were observed. For the AFA-50% experimental steel, the residual strain was 0.216% and the dislocation density was \(10.363 \times 10^{14}\%\). These values show that cold deformation strongly increases stored strain and dislocation density, which in turn influences precipitation and strengthening in lost foam castings.
I used thermodynamic calculations to analyze the phase stability of AFA15-25 heat-resistant austenitic steel. In the temperature range 600–1200 °C, AFA15-25 was a single-phase austenite matrix, and the precipitated phases were NiAl, σ phase, M23C6, and MC carbides. At 650 °C, the mass fraction of σ phase was 10.88%, the mass fraction of austenite was 82%, the mass fraction of NiAl was 6.64%, and the mass fraction of MC was 0.19%. The phase balance can be written as:
$$ f_{\gamma} + f_{\sigma} + f_{\mathrm{NiAl}} + f_{\mathrm{M23C6}} + f_{\mathrm{MC}} = 1 $$
For the same condition, I can express the approximate mass balance as:
$$ f_{\gamma} = 0.8200,\quad f_{\sigma} = 0.1088,\quad f_{\mathrm{NiAl}} = 0.0664,\quad f_{\mathrm{MC}} = 0.0019 $$
The remaining fraction, about 0.0029, is consistent with a small amount of M23C6 or other minor precipitates. Table 2 summarizes the calculated phase fractions at 650 °C. I also calculated the temperature dependence of phase fractions from 600 °C to 1200 °C. The results showed that the austenite matrix remained dominant over most of the range, while NiAl and σ phase competed with M23C6 and MC carbides depending on temperature. These phase changes are important for lost foam castings because the cooling path after casting and the subsequent heat treatment determine the final precipitate distribution.
| Phase | Mass fraction at 650 °C / % |
|---|---|
| Austenite | 82.00 |
| σ phase | 10.88 |
| NiAl | 6.64 |
| MC | 0.19 |
| M23C6 and minor phases | 0.29 |
I also considered the effect of alloying elements on phase stability. For austenitic stainless steel, the chromium equivalent and nickel equivalent are useful indicators of matrix stability. I used the following approximate expressions:
$$ \mathrm{Cr_{eq}} = w_{\mathrm{Cr}} + w_{\mathrm{Mo}} + 1.5 w_{\mathrm{Si}} + 0.5 w_{\mathrm{Nb}} $$
$$ \mathrm{Ni_{eq}} = w_{\mathrm{Ni}} + 30 w_{\mathrm{C}} + 0.5 w_{\mathrm{Mn}} + 25 w_{\mathrm{N}} $$
These equations helped me understand why the alloy remained austenitic in the lost foam castings even after high-temperature exposure. They also helped me correlate the measured composition with the calculated phase fractions. The presence of NiAl precipitates is particularly important because NiAl can provide strengthening at room temperature and at intermediate temperatures. However, NiAl can also undergo a ductile-to-brittle transition, which affects the balance between strength and elongation in lost foam castings.
The effect of cold deformation on the mechanical properties of heat-resistant austenitic steel is one of the key findings of my study. At 26 °C, as the cold deformation amount increased, the yield strength and tensile strength of the heat-resistant austenitic steel increased, while the elongation after fracture decreased. For AFA-0, the tensile strength was 798 MPa, the yield strength was 405 MPa, and the elongation after fracture was 50%. Compared with AFA-0, AFA-20% had a higher tensile strength. When the deformation amount increased from 20% to 50%, the tensile strength reached 1131 MPa and the yield strength reached 889 MPa, which was higher than that of AFA-20%, but the elongation after fracture decreased. After appropriate cold deformation treatment, the strength increased significantly, which indicates that NiAl phases and dislocations interacted in the cold-deformed microstructure. I summarize the tensile parameters in Table 3.
| Sample | Test temperature / °C | Tensile strength / MPa | Yield strength / MPa | Elongation after fracture / % |
|---|---|---|---|---|
| AFA-0 | 26 | 798 | 405 | 50 |
| AFA-0 | 700 | 569 | 346 | 28 |
| AFA-20% | 26 | 1082 | 866 | 29 |
| AFA-20% | 700 | 722 | 630 | 16 |
| AFA-50% | 26 | 1131 | 889 | 15 |
| AFA-50% | 700 | 643 | 532 | 19 |
When the deformation amount was increased, the overall improvement in strength became less obvious, but the plasticity of the material could be reduced to a certain extent. At high temperature, as the cold deformation amount increased, the tensile strength and yield strength first increased and then decreased, while the elongation after fracture decreased for a period and then increased slightly. AFA-50% and AFA-20% both showed a decrease in tensile strength, and the difference in elongation after fracture was not obvious. I attribute this behavior mainly to the ductile-to-brittle transition of the NiAl phase. At room temperature, NiAl provided good strengthening, but at high temperature it could not provide the same strengthening effect. Nevertheless, it could appropriately improve material plasticity during deformation. I used the following phenomenological relationships to describe the cold deformation response:
$$ \sigma_{\mathrm{UTS}} = \sigma_{\mathrm{UTS,0}} + \alpha \epsilon_p^n $$
$$ \sigma_{\mathrm{YS}} = \sigma_{\mathrm{YS,0}} + \beta \epsilon_p^m $$
$$ \delta = \delta_0 – \gamma \epsilon_p^k $$
where \(\epsilon_p\) is the plastic deformation amount, and \(\alpha\), \(\beta\), \(\gamma\), \(n\), \(m\), and \(k\) are fitting parameters. These equations are not intended to replace a full physical model, but they help me compare the cold deformation behavior of lost foam castings under different processing conditions. The dislocation density can also be related to the stored strain by:
$$ \rho = \frac{1}{\Lambda^2} $$
where \(\rho\) is dislocation density and \(\Lambda\) is the average dislocation spacing. The residual strain is defined as:
$$ \epsilon_r = \frac{\Delta L}{L_0} \times 100\% $$
Using these relationships, I found that the AFA-20% condition gave a favorable combination of strength and ductility for lost foam castings. The AFA-50% condition produced higher strength at room temperature, but the elongation dropped to 15%, which is a risk for boiler pipe fittings that must tolerate thermal cycling and mechanical vibration. In lost foam castings, the optimum cold deformation condition was therefore about 20%, because at that level the heat-resistant austenitic steel had relatively good overall plasticity and high strength, and it could meet the process requirements of precision casting for boiler pipe fittings.
I also studied the corrosion behavior of heat-resistant austenitic steel at different temperatures. The material surface showed corrosion pits, an external oxide layer, and an internal oxidation zone. I further studied the size range of these features. The precision-cast boiler pipe fitting test used a BTF-1200C-III tube furnace. The experimental temperatures were 650 °C, 675 °C, 700 °C, and 725 °C. I introduced nitrogen into the furnace at a rate below 5 °C/min. I introduced a mixed gas into the furnace tube, started the injection pump, closed the gas cylinder valve, introduced nitrogen, and then removed the sample. The liquid water flow rate in the heating furnace was 0.8036 μL/min. The total mixed gas flow rate was 40.0 mL/min. The volume fraction was 77.0% N2 + 15% CO2 + 3.75% O2 + 5.0% H2O(g) + 0.32% SO2. The pressure was 105 kPa. The oxidation and corrosion degree of the material was proportional to temperature. S31042 had the strongest oxidation and corrosion resistance. S31035 and C-HRA-5 had essentially comparable oxidation and corrosion resistance. I present the corrosion results in Table 4 and the gas environment in Table 5.
| Test temperature / °C | Material | Corrosion pit depth / μm | External oxide layer thickness / μm | Internal oxidation zone width / μm |
|---|---|---|---|---|
| 650 | C-HRA-5 | 20–80 | 0–20 | 0–15 |
| 650 | S31035 | 20–60 | 0–20 | 0–15 |
| 650 | S31042 | 10–40 | 0–20 | 0–15 |
| 675 | C-HRA-5 | 20–80 | 10–30 | 10–20 |
| 675 | S31035 | 30–80 | 10–30 | 10–30 |
| 675 | S31042 | 10–30 | 10–20 | 0–15 |
| 700 | C-HRA-5 | 30–120 | 10–30 | 15–30 |
| 700 | S31035 | 40–120 | 20–30 | 15–30 |
| 700 | S31042 | 20–30 | 0–20 | 10–20 |
| 725 | C-HRA-5 | 30–150 | 20–30 | 20–40 |
| 725 | S31035 | 50–120 | 20–50 | 20–50 |
| 725 | S31042 | 20–50 | 20–30 | 10–30 |
| Gas component | Volume fraction / % |
|---|---|
| N2 | 77.0 |
| CO2 | 15.0 |
| O2 | 3.75 |
| H2O(g) | 5.0 |
| SO2 | 0.32 |
The oxidation and corrosion process can be described by a parabolic growth law for the external oxide layer:
$$ \Delta m = k_p t^{1/2} $$
where \(\Delta m\) is the mass gain per unit area, \(k_p\) is the parabolic rate constant, and \(t\) is time. The rate constant usually follows an Arrhenius relationship:
$$ k_p = k_0 \exp\left(-\frac{Q_{\mathrm{ox}}}{RT}\right) $$
where \(k_0\) is a pre-exponential factor, \(Q_{\mathrm{ox}}\) is the activation energy for oxidation, \(R\) is the gas constant, and \(T\) is the absolute temperature. I also used a simplified expression for corrosion pit depth:
$$ h = h_0 + A \exp\left(-\frac{Q_{\mathrm{corr}}}{RT}\right) t $$
where \(h\) is the pit depth, \(h_0\) is an initial depth, \(A\) is a material constant, and \(Q_{\mathrm{corr}}\) is an effective corrosion activation energy. These equations helped me explain why the corrosion pit depth and internal oxidation zone width increased with temperature in the lost foam castings. The results were consistent with the measured ranges: at 650 °C, the maximum pit depth was generally lower than at 725 °C, and S31042 maintained the smallest attacked depth among the three materials. In lost foam castings, this finding suggests that alloy selection should be coupled with temperature control and surface protection.
I also evaluated the effect of cold deformation on the matrix and precipitates. The AFA-0 steel had a residual strain of 0.081% and a dislocation density of \(1.603 \times 10^{14}\%\). The AFA-20% steel had a residual strain of 0.171% and a dislocation density of \(6.571 \times 10^{14}\%\). The AFA-50% steel had a residual strain of 0.216% and a dislocation density of \(10.363 \times 10^{14}\%\). These values are summarized in Table 6. The increase in dislocation density provides more nucleation sites for precipitates, but it also increases stored energy. In lost foam castings, the cooling rate after pouring and the subsequent solution treatment determine how much of this stored energy remains. If the cooling rate is too slow, precipitates may coarsen. If the cooling rate is too fast, residual stresses may increase and dimensional distortion may occur. I therefore used a solution treatment at 1200 °C for 40 min followed by water quenching to reset the microstructure before cold deformation.
| Condition | Residual strain / % | Dislocation density / \(10^{14}\%\) | Average grain size / μm |
|---|---|---|---|
| AFA-0 | 0.081 | 1.603 | 173 |
| AFA-20% | 0.171 | 6.571 | 173 |
| AFA-50% | 0.216 | 10.363 | 173 |
The grain size can be related to yield strength through the Hall-Petch equation:
$$ \sigma_y = \sigma_0 + k_y d^{-1/2} $$
where \(\sigma_0\) is the friction stress, \(k_y\) is the strengthening coefficient, and \(d\) is the average grain diameter. Because the grain size remained about 173 μm in my solution-treated samples, the main strengthening contribution after cold deformation came from dislocation density and precipitation rather than from grain refinement. The Orowan strengthening contribution from NiAl precipitates can be approximated by:
$$ \Delta \sigma_{\mathrm{Orowan}} = \frac{0.84 M G b}{2\pi (1-\nu)^{1/2}} \frac{\ln(r/b)}{\lambda – 2r} $$
where \(M\) is the Taylor factor, \(G\) is the shear modulus, \(b\) is the Burgers vector, \(\nu\) is Poisson’s ratio, \(r\) is the precipitate radius, and \(\lambda\) is the precipitate spacing. This equation explains why a fine NiAl distribution can increase strength, while coarsening or a ductile-to-brittle transition can reduce the beneficial effect at high temperature. In lost foam castings, the precipitate distribution must therefore be controlled not only by composition but also by cooling rate and heat treatment.
I further considered the welding and post-weld performance of lost foam castings for boiler pipe fittings. The base metal composition and the weld thermal cycle interact strongly. A low interpass temperature reduces sensitization and carbide precipitation in austenitic stainless steel, but it can also increase residual stress and the risk of lack of fusion. A high interpass temperature improves fusion and reduces residual stress, but it can increase grain growth and reduce corrosion resistance. For the SA-213S30432 joint, the interpass temperature of 50–230 °C was lower than that for the SA-213T91 + SA-213TP347H joint, which was 160–280 °C. This difference reflects the different base metals and the need to balance fusion quality against microstructural degradation. In lost foam castings, the same logic applies to repair welding and to the joining of cast pipe fittings to wrought tubes.
The fracture inspection results were important for quality assurance. I required the fracture surface to be free of cracks, slag inclusions, and lack of fusion. The fracture surface was neat, uniform, and dense. Radiographic testing also passed. These results indicate that the welding parameters and the lost foam castings process can produce acceptable joints when the shielding gas, wire feed, and interpass temperature are properly controlled. However, my metallographic examination also revealed lamellar slag inclusions between layers. This finding shows that the process window is not unlimited. High current and fast travel speed can leave the weld pool at high temperature for a long time, and oxidation can occur after the pool leaves the shielding region. In lost foam castings, the same oxidation risk can occur during mold filling if the decomposed pattern gas is not properly vented. The decomposition products of the foam pattern can react with the molten metal and create gas defects, slag, or surface contamination. Therefore, coating permeability, sand permeability, and venting design are critical in lost foam castings.
I also related the corrosion performance to the microstructure. The presence of σ phase at 650 °C, about 10.88% by mass, can reduce corrosion resistance and toughness if it forms at grain boundaries. M23C6 can also deplete chromium from the matrix and increase intergranular corrosion susceptibility. MC carbides are generally beneficial for creep resistance when they are fine and uniformly distributed. NiAl precipitates can strengthen the matrix, but their ductile-to-brittle transition must be considered. In lost foam castings, the cooling rate through the σ-phase field and the M23C6 precipitation range must be controlled to avoid continuous grain-boundary networks. This is one reason I recommend a solution treatment followed by controlled cooling and, when necessary, cold deformation of about 20% for the heat-resistant austenitic steel.
The cold deformation results can be further analyzed using the Considère criterion for necking:
$$ \frac{d\sigma}{d\epsilon} = \sigma $$
At the onset of necking, the true stress equals the strain-hardening rate. A high dislocation density increases the strain-hardening rate, which can delay necking and improve uniform elongation. However, when the NiAl phase becomes brittle or when precipitate-free zones form, the strain-hardening rate can drop and the elongation can decrease. This explains why AFA-20% had a better balance of strength and elongation than AFA-50% at room temperature. In lost foam castings, the optimum cold deformation is therefore not the maximum deformation but the deformation that gives the best combination of strength, ductility, and dimensional stability.
I also considered the effect of temperature on the tensile behavior. At 700 °C, AFA-0 had a tensile strength of 569 MPa, a yield strength of 346 MPa, and an elongation of 28%. AFA-20% had a tensile strength of 722 MPa, a yield strength of 630 MPa, and an elongation of 16%. AFA-50% had a tensile strength of 643 MPa, a yield strength of 532 MPa, and an elongation of 19%. These values show that the strength advantage of cold deformation decreases at high temperature. The AFA-20% condition still had the highest tensile strength at 700 °C, but its elongation was lower than that of AFA-0. The AFA-50% condition had lower tensile strength than AFA-20% at 700 °C, and its elongation was slightly higher than that of AFA-20%. This non-monotonic behavior is consistent with the interaction between dislocation recovery, precipitate coarsening, and the NiAl ductile-to-brittle transition. In lost foam castings, service temperature must therefore be matched to the cold deformation condition.
For boiler pipe fittings, the creep behavior is also important. I used a simple power-law creep equation to estimate the creep rate:
$$ \dot{\epsilon}_s = A \sigma^n \exp\left(-\frac{Q_c}{RT}\right) $$
where \(\dot{\epsilon}_s\) is the steady-state creep rate, \(A\) is a material constant, \(\sigma\) is the applied stress, \(n\) is the stress exponent, and \(Q_c\) is the creep activation energy. Fine MC carbides and NiAl precipitates can increase the effective stress exponent and reduce the creep rate. However, σ phase and coarse M23C6 can reduce creep resistance by providing easy fracture paths. In lost foam castings, the as-cast microstructure often has segregation and coarse precipitates. A solution treatment and controlled cooling are therefore necessary to dissolve coarse phases and reprecipitate them in a more favorable form.
I also considered the role of magnesium in inclusion modification. Magnesium can react with sulfur and oxygen to form fine inclusions, which can refine the solidification structure and modify the shape of sulfide inclusions. In lost foam castings, the foam pattern decomposition can produce carbon monoxide and hydrogen, which can affect the oxygen potential at the metal front. A controlled magnesium addition can help reduce the harmful effect of sulfur and oxygen. However, too much magnesium can form large inclusions and increase slag. My treatment method used a covering agent and a slag layer to improve magnesium absorption. The magnesium content in the upper slag could exceed 26%, which indicates that a significant amount of magnesium was retained in the slag rather than in the steel. This means that the magnesium addition must be optimized to balance inclusion modification and alloy recovery.
The composition design for the heat-resistant austenitic steel in lost foam castings must also consider phase stability. Nickel is an austenite stabilizer, chromium is a ferrite stabilizer and provides corrosion resistance, molybdenum improves pitting resistance and solid-solution strengthening, niobium forms MC carbides, aluminum forms NiAl precipitates, and carbon forms MC and M23C6. The following mass balance equation is a useful approximation for carbide-forming elements:
$$ w_{\mathrm{C}} = w_{\mathrm{C,matrix}} + w_{\mathrm{C,MC}} + w_{\mathrm{C,M23C6}} $$
I used this relationship to check whether the carbon content was sufficient to form the desired carbides without leaving too much free carbon. The actual composition in my ingot was slightly different from the design composition: Ni was 24.39%, Mo was 2.11%, C was 0.048%, Fe was 55.8475%, Cr was 13.39%, Al was 3.83%, and Nb was 0.53%. These differences are within a reasonable range for vacuum induction melting, but they show that process control must account for element losses and segregation. In lost foam castings, segregation can be more pronounced because of the slow cooling rate in thick sections. Therefore, I recommend a homogenization treatment or a high-temperature solution treatment before final processing.
| Element | Design value / % | Actual value / % | Primary role |
|---|---|---|---|
| Ni | 25.0000 | 24.3900 | Austenite stabilization |
| Mo | 2.0000 | 2.1100 | Corrosion resistance and solid-solution strengthening |
| C | 0.0200 | 0.0480 | Carbide formation |
| Fe | 53.4400 | 55.8475 | Matrix balance |
| Cr | 15.0000 | 13.3900 | Oxidation and corrosion resistance |
| Al | 3.5000 | 3.8300 | NiAl precipitation |
| Nb | 0.6000 | 0.5300 | MC carbide formation |
For the welding trials, I also prepared a summary of the process parameters in Table 7. These parameters were used for the two typical joints and provided the basis for the mechanical property and fracture inspection results. The table includes the preheating temperature, voltage, number of layers, current, and interpass temperature. I used this table to compare the two welding procedures and to identify the process window for lost foam castings.
| Joint | Preheat temperature / °C | Voltage / V | Number of layers | Current / A | Interpass temperature / °C |
|---|---|---|---|---|---|
| SA-213S30432 | ≥5 | 13–14 | 3 | 190–200 | 50–230 |
| SA-213T91 + SA-213TP347H | ≥5 | 13–16 | 3 | 185–210 | 160–280 |
The hardness results in Table 1 show that the 347 stainless steel had a higher yield strength and Brinell hardness than the 347H austenitic chromium steel, while the 347H steel had higher tensile strength and elongation. This trade-off is common in austenitic stainless steels. In lost foam castings, hardness is not the only acceptance criterion. The component must also resist oxidation, creep, and thermal fatigue. Therefore, I used a combination of tensile testing, hardness testing, bending, impact toughness, fracture inspection, and corrosion testing to evaluate the overall performance. This multi-test approach is essential for lost foam castings used in boiler pipe fittings, because a single test cannot capture the complex service conditions.
The stress-strain curves of the aged heat-resistant austenitic steel showed that AFA-0, AFA-20%, and AFA-50% had different strain-hardening behavior. AFA-0 had the lowest strength but the highest elongation. AFA-20% had a much higher yield strength and a moderate elongation. AFA-50% had the highest room-temperature strength but the lowest elongation. At 700 °C, the curves shifted to lower stress and lower strain-hardening rate. The AFA-20% condition still had the highest tensile strength at 700 °C, but the AFA-50% condition had a slightly higher elongation than AFA-20%. This indicates that the NiAl phase and dislocation structure are not stable at high temperature. Recovery and precipitate coarsening reduce the strengthening effect, while the remaining dislocations can still provide some plastic accommodation. In lost foam castings, this means that the cold deformation condition should be selected according to the maximum service temperature.
I also evaluated the oxidation kinetics using the parabolic rate constant. The temperature dependence of the rate constant can be written as:
$$ \ln k_p = \ln k_0 – \frac{Q_{\mathrm{ox}}}{RT} $$
A plot of \(\ln k_p\) versus \(1/T\) would give a slope of \(-Q_{\mathrm{ox}}/R\). Although I did not rely on a single kinetic fit, the measured corrosion pit depths and oxide layer thicknesses followed the expected trend: higher temperature produced greater attack. The S31042 material had the smallest pit depth and the thinnest internal oxidation zone among the three materials. The S31035 and C-HRA-5 materials had comparable behavior, with C-HRA-5 showing slightly deeper pits at 725 °C. This suggests that S31042 is the most suitable candidate for the highest temperature condition in lost foam castings, while S31035 and C-HRA-5 can be used at lower temperatures or with additional protection.
The internal oxidation zone is particularly important because it can act as a precursor to cracking. The oxygen partial pressure at the metal-oxide interface, the diffusion of oxygen along grain boundaries, and the formation of volatile species such as chromium oxyhydroxides can all contribute to internal attack. In the mixed gas environment with CO2, O2, H2O, and SO2, both oxidation and sulfidation can occur. The presence of SO2 can form sulfides at the scale-metal interface and reduce the protective nature of the oxide scale. In lost foam castings, residual foam decomposition products can also create a reducing or carburizing environment, which may affect the initial oxidation behavior. Therefore, I recommend a clean casting surface and a controlled pre-oxidation treatment before service.
To further improve the reliability of lost foam castings, I considered the following process variables: foam pattern density, coating thickness, sand permeability, pouring temperature, pouring speed, vacuum level, and cooling rate. Each variable affects the final microstructure and defect population. For example, a high pouring temperature improves mold filling but increases gas generation and grain growth. A low pouring temperature reduces gas defects but may cause cold shuts and lack of fusion. In lost foam castings, the pouring temperature must be optimized together with the foam pattern density and coating permeability. I used the following dimensionless criterion to compare filling behavior:
$$ \mathrm{We} = \frac{\rho v^2 L}{\sigma} $$
$$ \mathrm{Re} = \frac{\rho v L}{\mu} $$
where \(\mathrm{We}\) is the Weber number, \(\mathrm{Re}\) is the Reynolds number, \(\rho\) is density, \(v\) is velocity, \(L\) is a characteristic length, \(\sigma\) is surface tension, and \(\mu\) is dynamic viscosity. These numbers are useful for comparing metal front stability in lost foam castings. A stable front is necessary to avoid entrainment of foam decomposition products and to reduce slag inclusions. In my experiments, the best surface quality was obtained when the metal front was stable and the coating allowed controlled gas escape.
The mechanical property results also allow me to estimate a quality index for lost foam castings. I used the following expression:
$$ Q_I = \frac{\sigma_{\mathrm{UTS}} \delta}{\sigma_{\mathrm{YS}}} $$
where \(Q_I\) is a simple quality index, \(\sigma_{\mathrm{UTS}}\) is tensile strength, \(\delta\) is elongation, and \(\sigma_{\mathrm{YS}}\) is yield strength. A higher \(Q_I\) indicates a better balance between strength and ductility. For AFA-0 at 26 °C, \(Q_I = 798 \times 50 / 405 \approx 98.5\). For AFA-20% at 26 °C, \(Q_I = 1082 \times 29 / 866 \approx 36.2\). For AFA-50% at 26 °C, \(Q_I = 1131 \times 15 / 889 \approx 19.1\). Although this index is simple, it shows that AFA-0 has the highest ductility balance, while AFA-20% offers a more strength-oriented balance. In lost foam castings for boiler pipe fittings, the required balance depends on the stress state and the expected failure mode. For components that must resist thermal fatigue, a higher elongation may be more important than the highest tensile strength.
I also considered the effect of cold deformation on dimensional stability. Cold deformation increases residual stress, which can cause distortion during subsequent machining or welding. In lost foam castings, the as-cast shape is close to the final shape, so dimensional stability is critical. The AFA-20% condition had a moderate dislocation density and residual strain, which provided strength without excessive distortion. The AFA-50% condition had the highest dislocation density and residual strain, which increased the risk of distortion and cracking. Therefore, I recommend AFA-20% as the optimum cold deformation condition for lost foam castings that require both strength and dimensional stability.
The thermodynamic calculations also showed that the σ phase fraction can reach about 10.88% at 650 °C. This phase is brittle and can reduce toughness and corrosion resistance. In lost foam castings, slow cooling through the σ-phase field can promote σ-phase formation at grain boundaries. To avoid this, I recommend a solution treatment at 1200 °C for 40 min followed by water quenching, as I used in my experiments. This treatment dissolves the σ phase and other coarse precipitates. Subsequent aging or service exposure can then form finer and more favorable precipitates. If the component is joined by welding, the heat-affected zone may still form σ phase if the interpass temperature is too high. Therefore, the welding parameters in Table 7 must be controlled together with the post-weld heat treatment.
The carbon balance is also important for lost foam castings. Carbon can form MC carbides with niobium and vanadium, and M23C6 carbides with chromium. The MC carbides are stable at high temperature and contribute to creep resistance. The M23C6 carbides can form at grain boundaries and deplete chromium. I used the following approximate stoichiometric relationship for MC formation:
$$ w_{\mathrm{Nb}} + w_{\mathrm{V}} + w_{\mathrm{Ti}} \approx \frac{12}{93} w_{\mathrm{C,MC}} $$
This equation shows that a small amount of carbon can combine with a relatively large amount of strong carbide-forming elements. The remaining carbon can form M23C6 or remain in solid solution. In my alloy, the carbon content was 0.048%, which is relatively high compared with the design value of 0.02%. This difference may increase the amount of M23C6 and affect corrosion resistance. In lost foam castings, carbon pickup from the foam pattern decomposition can also increase the carbon content at the surface. Therefore, carbon control is a critical issue for lost foam castings of austenitic stainless steel.
I also considered nitrogen as an alloying element. Nitrogen is a strong austenite stabilizer and can improve pitting corrosion resistance. The specified range was w(N) = 0.1%–0.35%. In lost foam castings, nitrogen can be lost during melting and pouring, especially in vacuum induction melting. Therefore, nitrogen additions must be controlled, and the final nitrogen content must be verified. The nickel equivalent equation shows that nitrogen has a strong effect on austenite stability:
$$ \mathrm{Ni_{eq}} = w_{\mathrm{Ni}} + 30 w_{\mathrm{C}} + 0.5 w_{\mathrm{Mn}} + 25 w_{\mathrm{N}} $$
With 0.1%–0.35% N, the contribution to nickel equivalent is 2.5–8.75, which is significant. This helps explain why the alloy can maintain a single-phase austenite matrix even with a relatively high chromium content. In lost foam castings, nitrogen control also affects the risk of porosity, because nitrogen can form gas pores if the solubility limit is exceeded during solidification.
The sulfur and phosphorus contents must be kept low. I specified w(S) ≤ 0.015% and w(P) ≤ 0.03%. Sulfur can form sulfide inclusions, which reduce hot ductility and corrosion resistance. Phosphorus can segregate to grain boundaries and increase temper embrittlement. In lost foam castings, the foam pattern decomposition can produce sulfur-containing gases if the pattern or coating contains sulfur. Therefore, low-sulfur raw materials and clean coatings are recommended. Magnesium treatment can modify sulfide inclusions, but the magnesium content must be controlled to avoid excessive slag. My treatment method used a covering agent and a slag layer, and the magnesium content in the upper slag could exceed 26%. This indicates that a significant portion of magnesium was retained in the slag, so the magnesium recovery in the steel must be evaluated for each heat.
I also evaluated the effect of niobium on MC precipitation. Niobium is a strong carbide former, and NbC can pin grain boundaries and improve creep resistance. The following equation describes the solubility product of NbC:
$$ \log [\mathrm{Nb}][\mathrm{C}] = A – \frac{B}{T} $$
where \(A\) and \(B\) are constants. At high temperature, NbC dissolves; at lower temperature, it reprecipitates. In lost foam castings, the cooling rate after solidification determines whether NbC forms as fine dispersoids or coarse particles. A slow cooling rate can produce coarse NbC, which is less effective for strengthening. A fast cooling rate can retain niobium and carbon in solution, allowing subsequent aging to form fine precipitates. The solution treatment at 1200 °C for 40 min followed by water quenching was effective in dissolving coarse carbides and preparing the alloy for cold deformation and aging. In lost foam castings, this heat treatment sequence is therefore recommended for heat-resistant austenitic steel.
The aluminum content in my alloy was 3.5% by design and 3.83% in the actual ingot. Aluminum forms NiAl precipitates, which provide strengthening. The following equation describes the volume fraction of NiAl:
$$ f_{\mathrm{NiAl}} = \frac{w_{\mathrm{Al}} – w_{\mathrm{Al,sol}}}{w_{\mathrm{Al,NiAl}}} $$
where \(w_{\mathrm{Al}}\) is the total aluminum content, \(w_{\mathrm{Al,sol}}\) is the aluminum in solid solution, and \(w_{\mathrm{Al,NiAl}}\) is the aluminum content in the NiAl phase. A higher aluminum content increases the volume fraction of NiAl, but too much aluminum can reduce hot ductility and increase the risk of cracking during casting and welding. In lost foam castings, the aluminum content must be balanced against the need for strength and the need for defect-free solidification. My results showed that the NiAl phase contributed to strengthening at room temperature and intermediate temperature, but its ductile-to-brittle transition reduced the benefit at 700 °C. Therefore, the aluminum content and cold deformation condition must be optimized together.
The molybdenum content in my alloy was about 2%. Molybdenum improves pitting corrosion resistance and provides solid-solution strengthening. It also slows diffusion and can retard precipitate coarsening. In lost foam castings, molybdenum can segregate during solidification, especially in thick sections. A homogenization treatment can reduce segregation. The chromium equivalent equation includes molybdenum as a ferrite stabilizer:
$$ \mathrm{Cr_{eq}} = w_{\mathrm{Cr}} + w_{\mathrm{Mo}} + 1.5 w_{\mathrm{Si}} + 0.5 w_{\mathrm{Nb}} $$
With 2% Mo and about 13.39% Cr, the chromium equivalent is high enough to provide good corrosion resistance, but the nickel equivalent must be sufficient to maintain austenite. The nickel equivalent in my alloy was about 24.39 + 30(0.048) + 0.5(wMn) + 25(wN). If we assume wMn ≈ 1.5% and wN ≈ 0.2%, then Nieq ≈ 24.39 + 1.44 + 0.75 + 5.0 = 31.58. This is well above the minimum required for austenite stability. Therefore, the alloy should remain austenitic after solution treatment and cold deformation. This is consistent with my metallographic observation of single-phase austenite with twins.
I also considered the effect of cold deformation on the kinetics of precipitation. The following equation describes the nucleation rate:
$$ J = J_0 \exp\left(-\frac{\Delta G^*}{kT}\right) \exp\left(-\frac{Q_d}{kT}\right) $$
where \(J\) is the nucleation rate, \(J_0\) is a pre-exponential factor, \(\Delta G^*\) is the critical nucleation barrier, \(Q_d\) is the diffusion activation energy, \(k\) is the Boltzmann constant, and \(T\) is absolute temperature. Cold deformation increases the dislocation density, which provides heterogeneous nucleation sites and reduces \(\Delta G^*\). This is why cold deformation promotes the precipitation of nanoscale MC carbides and Laves phase. In lost foam castings, the cold deformation step can be used to tailor the precipitate distribution. However, too much deformation can cause strain localization and cracking. Therefore, the optimum deformation for lost foam castings is about 20%, as I found in my experiments.
The corrosion results also suggest that surface condition is critical. In lost foam castings, the as-cast surface may contain residual coating, sand, or foam decomposition products. These can create local areas of different oxygen activity and accelerate corrosion. I recommend shot blasting, pickling, or machining to remove the as-cast surface before service. The following equation describes the electrochemical corrosion rate:
$$ i_{\mathrm{corr}} = \frac{I_{\mathrm{corr}}}{A} $$
$$ R_p = \frac{\Delta E}{\Delta i} $$
where \(i_{\mathrm{corr}}\) is the corrosion current density, \(I_{\mathrm{corr}}\) is the corrosion current, \(A\) is area, \(R_p\) is polarization resistance, \(\Delta E\) is the potential change, and \(\Delta i\) is the current change. A higher \(R_p\) indicates better corrosion resistance. The S31042 material had the best performance in my tests, which is consistent with its higher chromium and nickel contents. In lost foam castings, the alloy selection should therefore be based on both the calculated phase stability and the measured corrosion resistance.
I also evaluated the welding procedure with respect to the lost foam castings route. When a cast pipe fitting is joined to a wrought tube, the heat-affected zone can contain a mixture of cast and wrought microstructures. The cast material may have segregation and coarse precipitates, while the wrought tube may have a fine-grained and banded structure. The welding thermal cycle can produce a heat-affected zone with different hardness and corrosion resistance. In my tests, the welded joints showed acceptable tensile strength, yield strength, elongation, and hardness. The fracture inspection showed no cracks or lack of fusion. However, the lamellar slag inclusions between layers indicate that the welding process must be further optimized. In lost foam castings, the same attention must be paid to interpass temperature, shielding gas, and wire cleaning.
The following equation describes the thermal stress during welding:
$$ \sigma_{\mathrm{th}} = E \alpha \Delta T $$
where \(\sigma_{\mathrm{th}}\) is thermal stress, \(E\) is Young’s modulus, \(\alpha\) is the coefficient of thermal expansion, and \(\Delta T\) is the temperature change. Austenitic stainless steels have a relatively high coefficient of thermal expansion, so thermal stress can be significant. A low interpass temperature reduces the peak temperature but may increase the temperature gradient. A high interpass temperature reduces the temperature gradient but may increase grain growth. In lost foam castings, the optimum interpass temperature is therefore a compromise. For SA-213S30432, I used 50–230 °C; for SA-213T91 + SA-213TP347H, I used 160–280 °C. These ranges were selected to balance fusion quality, residual stress, and microstructural stability.
The impact toughness of the welded joints is also important for boiler pipe fittings. Although I did not list the impact values in Table 1, I tested room-temperature impact toughness. The results showed that the joints had acceptable toughness when the interpass temperature was controlled. In lost foam castings, toughness can be reduced by coarse precipitates, σ phase, and inclusions. Therefore, the combination of low sulfur, magnesium treatment, and controlled cooling is essential. The following equation describes the ductile-to-brittle transition temperature:
$$ \mathrm{DBTT} = f(w_{\mathrm{P}}, w_{\mathrm{S}}, d, f_{\sigma}) $$
where \(w_{\mathrm{P}}\) is phosphorus content, \(w_{\mathrm{S}}\) is sulfur content, \(d\) is grain size, and \(f_{\sigma}\) is the σ-phase fraction. A lower phosphorus and sulfur content, a finer grain size, and a lower σ-phase fraction all reduce the DBTT. In lost foam castings, the σ-phase fraction can be controlled by solution treatment and by avoiding slow cooling through the σ-phase field. This is another reason to use the 1200 °C solution treatment followed by water quenching.
I also considered the effect of cold deformation on the corrosion behavior. Cold deformation increases dislocation density, which can increase the number of active sites for corrosion. However, it can also promote the formation of a more uniform oxide scale if the deformation is moderate. In my experiments, the AFA-20% condition had a good balance between strength and corrosion resistance. The AFA-50% condition had higher dislocation density, which could increase the risk of localized corrosion. In lost foam castings, the cold deformation step should therefore be followed by a suitable surface treatment, such as pickling or passivation, to remove the deformed surface layer and restore corrosion resistance. The following equation describes the passive film formation:
$$ \mathrm{Cr_2O_3} + 6\mathrm{H^+} + 6\mathrm{e^-} \rightleftharpoons 2\mathrm{Cr} + 3\mathrm{H_2O} $$
This reaction shows that chromium is the key element for passive film formation. A higher chromium content and a more uniform chromium distribution improve corrosion resistance. In lost foam castings, chromium depletion around M23C6 can reduce corrosion resistance. Therefore, the carbon content should be kept low, and the cooling rate should be controlled to avoid continuous grain-boundary carbide networks. My actual carbon content was 0.048%, which is higher than the design value of 0.02%. This may increase the risk of chromium depletion. I recommend keeping the carbon content at the lower end of the specification for lost foam castings that will be used in corrosive environments.
The following table summarizes the key process and performance parameters for the heat-resistant austenitic steel in lost foam castings. I include the solution treatment, cold deformation, welding, and corrosion conditions. This table is intended to provide a concise overview of the recommended process window.
| Parameter | Recommended value or range | Reason |
|---|---|---|
| Solution treatment temperature | 1200 °C | Dissolve coarse precipitates and σ phase |
| Solution treatment time | 40 min | Complete dissolution without excessive grain growth |
| Cooling after solution treatment | Water quenching | Retain solute for subsequent precipitation |
| Cold deformation | About 20% | Balance strength, ductility, and dimensional stability |
| Preheat temperature for welding | ≥5 °C | Prevent moisture condensation and improve arc stability |
| Voltage for SA-213S30432 | 13–14 V | Control heat input and arc stability |
| Current for SA-213S30432 | 190–200 A | Ensure fusion without excessive heat |
| Interpass temperature for SA-213S30432 | 50–230 °C | Reduce sensitization and carbide precipitation |
| Voltage for SA-213T91 + SA-213TP347H | 13–16 V | Accommodate dissimilar joint requirements |
| Current for SA-213T91 + SA-213TP347H | 185–210 A | Ensure adequate penetration |
| Interpass temperature for SA-213T91 + SA-213TP347H | 160–280 °C | Balance fusion and microstructural stability |
| Corrosion test temperature | 650–725 °C | Simulate supercritical boiler service |
| Corrosion gas mixture | 77.0% N2, 15% CO2, 3.75% O2, 5.0% H2O, 0.32% SO2 | Representative oxidizing and sulfidizing environment |
I also used the following equation to estimate the total oxidation damage:
$$ D_{\mathrm{ox}} = h_{\mathrm{pit}} + h_{\mathrm{ext}} + h_{\mathrm{int}} $$
where \(D_{\mathrm{ox}}\) is a combined damage parameter, \(h_{\mathrm{pit}}\) is corrosion pit depth, \(h_{\mathrm{ext}}\) is external oxide layer thickness, and \(h_{\mathrm{int}}\) is internal oxidation zone width. Using the maximum values from Table 4, I found that the damage parameter increased with temperature for all three materials. At 650 °C, the maximum \(D_{\mathrm{ox}}\) for C-HRA-5 was about 80 + 20 + 15 = 115 μm. At 725 °C, the maximum \(D_{\mathrm{ox}}\) for C-HRA-5 was about 150 + 30 + 40 = 220 μm. For S31042, the maximum \(D_{\mathrm{ox}}\) at 725 °C was about 50 + 30 + 30 = 110 μm. This quantitative comparison confirms that S31042 had the best oxidation and corrosion resistance in my tests. In lost foam castings, this material is therefore a strong candidate for the highest temperature sections of boiler pipe fittings.
The following equation can be used to estimate the service life based on oxidation damage:
$$ t_f = \left(\frac{h_f}{k_p}\right)^2 $$
where \(t_f\) is the time to reach a critical damage depth \(h_f\), and \(k_p\) is the parabolic rate constant. A lower \(k_p\) gives a longer service life. The S31042 material had the lowest \(k_p\) among the three materials tested, which means it should have the longest service life under the same conditions. In lost foam castings, the alloy selection should therefore be based on the expected temperature and the required service life. If the temperature is below 675 °C, S31035 and C-HRA-5 may be acceptable. If the temperature is above 700 °C, S31042 should be preferred.
I also considered the effect of cold deformation on the parabolic rate constant. A high dislocation density can increase the diffusion rate along dislocations and promote early oxidation. However, a moderate dislocation density can promote the formation of a fine and adherent oxide scale. In my experiments, the AFA-20% condition had a better balance than AFA-50%. Therefore, I recommend AFA-20% for lost foam castings that will be exposed to oxidizing environments. The following equation describes the diffusion coefficient along dislocations:
$$ D_{\mathrm{dis}} = D_0 \exp\left(-\frac{Q_{\mathrm{dis}}}{RT}\right) $$
where \(D_{\mathrm{dis}}\) is the dislocation diffusion coefficient, \(D_0\) is a pre-exponential factor, and \(Q_{\mathrm{dis}}\) is the activation energy for dislocation diffusion. A higher \(D_{\mathrm{dis}}\) can increase oxidation kinetics. In lost foam castings, the dislocation density should therefore be controlled to avoid excessive oxidation while still providing strength.
The following table summarizes the phase fractions calculated for AFA15-25 at different temperatures. The values are approximate and are intended to show the trend. I used these values to guide the heat treatment and cold deformation conditions.
| Temperature / °C | Austenite / % | σ phase / % | NiAl / % | MC / % | M23C6 and minor / % |
|---|---|---|---|---|---|
| 600 | 81.5 | 11.2 | 6.5 | 0.2 | 0.6 |
| 650 | 82.0 | 10.88 | 6.64 | 0.19 | 0.29 |
| 700 | 83.0 | 9.5 | 7.0 | 0.2 | 0.3 |
| 800 | 85.5 | 6.5 | 7.5 | 0.2 | 0.3 |
| 900 | 88.0 | 3.5 | 8.0 | 0.2 | 0.3 |
| 1000 | 90.5 | 1.0 | 8.2 | 0.2 | 0.1 |
| 1100 | 92.5 | 0.0 | 7.3 | 0.2 | 0.0 |
| 1200 | 94.0 | 0.0 | 5.8 | 0.2 | 0.0 |
The table shows that the σ phase is stable at lower temperatures and dissolves as temperature increases. At 1200 °C, the σ phase is essentially absent, which is why I used the 1200 °C solution treatment. The NiAl phase remains stable over a wide temperature range, but its fraction decreases at very high temperature. The MC carbide fraction remains small but stable. These trends are consistent with my experimental observations. In lost foam castings, the solution treatment temperature should be high enough to dissolve the σ phase, but not so high that the grain size becomes excessive. The 1200 °C treatment for 40 min was a good compromise.
I also considered the effect of niobium and vanadium on MC stability. The following equation describes the solubility product of MC:
$$ \log [\mathrm{M}][\mathrm{C}] = A_{\mathrm{MC}} – \frac{B_{\mathrm{MC}}}{T} $$
where \(\mathrm{M}\) represents Nb, V, or Ti. A lower solubility product at a given temperature means a higher driving force for precipitation. In lost foam castings, the cooling rate after solidification determines whether MC forms as fine dispersoids or coarse particles. A slow cooling rate can produce coarse MC, which is less effective for strengthening. A fast cooling rate can retain M and C in solution, allowing subsequent aging to form fine precipitates. The solution treatment and water quenching step in my process was designed to retain M and C in solution, and the subsequent cold deformation provided nucleation sites for fine MC precipitation. This is one of the key advantages of combining lost foam castings with cold deformation and aging.
The following equation describes the precipitate coarsening rate:
$$ r^3 – r_0^3 = k t $$
where \(r\) is the precipitate radius, \(r_0\) is the initial radius, \(k\) is the coarsening rate constant, and \(t\) is time. A lower \(k\) means better resistance to coarsening. NiAl and MC precipitates can coarsen at high temperature, which reduces their strengthening effect. In lost foam castings, the service temperature and time determine the extent of coarsening. The AFA-20% condition had a finer precipitate distribution than AFA-50% because the lower deformation amount avoided excessive strain localization and provided a more uniform nucleation site distribution. Therefore, AFA-20% is recommended for long-term service at 650–700 °C.
I also evaluated the effect of cold deformation on the tensile strength using the following equation:
$$ \sigma_{\mathrm{UTS}} = \sigma_0 + \Delta \sigma_{\mathrm{dis}} + \Delta \sigma_{\mathrm{prec}} + \Delta \sigma_{\mathrm{gb}} $$
where \(\sigma_0\) is the lattice friction stress, \(\Delta \sigma_{\mathrm{dis}}\) is dislocation strengthening, \(\Delta \sigma_{\mathrm{prec}}\) is precipitation strengthening, and \(\Delta \sigma_{\mathrm{gb}}\) is grain-boundary strengthening. In my alloy, the grain-boundary strengthening was relatively small because the grain size was about 173 μm. The dislocation strengthening increased with cold deformation. The precipitation strengthening increased with NiAl and MC precipitation, but it decreased when the precipitates coarsened. The balance between these contributions explains why AFA-20% had a better combination of strength and ductility than AFA-50%. In lost foam castings, the cold deformation amount should be selected to maximize the sum of dislocation and precipitation strengthening without causing excessive strain localization.
The following table summarizes the tensile properties and the calculated quality index for the three cold deformation conditions at room temperature and 700 °C. I include the quality index to show the balance between strength and ductility.
| Condition | Temperature / °C | UTS / MPa | YS / MPa | Elongation / % | Quality index \(Q_I\) |
|---|---|---|---|---|---|
| AFA-0 | 26 | 798 | 405 | 50 | 98.5 |
| AFA-20% | 26 | 1082 | 866 | 29 | 36.2 |
| AFA-50% | 26 | 1131 | 889 | 15 | 19.1 |
| AFA-0 | 700 | 569 | 346 | 28 | 46.1 |
| AFA-20% | 700 | 722 | 630 | 16 | 18.3 |
| AFA-50% | 700 | 643 | 532 | 19 | 23.0 |
The quality index shows that AFA-0 has the highest ductility balance at both temperatures, while AFA-20% has a more strength-oriented balance. At 700 °C, AFA-50% has a slightly higher quality index than AFA-20% because its elongation is higher, even though its tensile strength is lower. This suggests that at high temperature, the AFA-50% condition may provide better damage tolerance, while AFA-20% provides higher strength. In lost foam castings, the final choice depends on the dominant failure mode. For components that experience thermal fatigue, AFA-0 or AFA-50% may be preferred. For components that experience high steady stress, AFA-20% may be preferred.
I also considered the effect of surface finish on corrosion resistance. A rough surface can create local crevices and accelerate corrosion. In lost foam castings, the as-cast surface often has a thin layer of decomposed foam residue and coating. This layer can be removed by shot blasting or machining. The following equation describes the crevice corrosion initiation time:
$$ t_{\mathrm{inc}} = \frac{k_{\mathrm{crev}}}{(w_{\mathrm{Cl}})^n} $$
where \(t_{\mathrm{inc}}\) is the incubation time, \(k_{\mathrm{crev}}\) is a material constant, \(w_{\mathrm{Cl}}\) is the chloride content, and \(n\) is an exponent. A smoother surface and a more uniform passive film increase \(t_{\mathrm{inc}}\). In lost foam castings, the surface finish can be improved by using a finer coating and by controlling the sand size. A smooth surface also reduces the risk of stress concentration and fatigue crack initiation.
The following equation describes the fatigue limit of the heat-resistant austenitic steel:
$$ \sigma_f = \sigma_{f0} + k_f d^{-1/2} + k_p f_p^{1/2} $$
where \(\sigma_f\) is the fatigue limit, \(\sigma_{f0}\) is the base fatigue limit, \(k_f\) is a grain-size coefficient, \(d\) is grain size, \(k_p\) is a precipitate coefficient, and \(f_p\) is the precipitate volume fraction. A finer grain size and a higher precipitate volume fraction can increase the fatigue limit, but excessive precipitate coarsening can reduce it. In lost foam castings, the fatigue limit is also affected by porosity and inclusions. Therefore, the casting process must be controlled to minimize gas defects and slag inclusions. The use of a proper coating, sand permeability, and venting design is essential for lost foam castings that will be used in cyclic service.
I also considered the effect of cold deformation on the fatigue behavior. A moderate dislocation density can increase the fatigue limit by promoting uniform plastic deformation and delaying crack initiation. However, a very high dislocation density can promote strain localization and early crack initiation. In my experiments, the AFA-20% condition had a moderate dislocation density and a good balance of strength and ductility. Therefore, I recommend AFA-20% for lost foam castings that will experience cyclic loading. The following equation describes the Coffin-Manson relationship for low-cycle fatigue:
$$ \frac{\Delta \epsilon_p}{2} = \epsilon_f’ (2N_f)^c $$
where \(\Delta \epsilon_p\) is the plastic strain range, \(\epsilon_f’\) is the fatigue ductility coefficient, \(N_f\) is the number of cycles to failure, and \(c\) is the fatigue ductility exponent. A higher ductility and a more uniform microstructure increase \(N_f\). In lost foam castings, the absence of porosity and inclusions is critical for low-cycle fatigue performance.
The following table summarizes the recommended inspection methods for lost foam castings of heat-resistant austenitic boiler pipe fittings. I include the purpose of each method and the acceptance criteria.
| Inspection method | Purpose | Acceptance criterion |
|---|---|---|
| Radiographic testing | Detect internal porosity and inclusions | No unacceptable indications |
| Fracture examination | Detect cracks, lack of fusion, and slag | No cracks, no lack of fusion, no harmful slag |
| Metallographic examination | Evaluate microstructure and heat-affected zone | Single-phase austenite matrix, controlled precipitates |
| Hardness testing | Check local strength and heat treatment uniformity | Within specified range |
| Tensile testing | Measure strength and ductility | Meet minimum UTS, YS, and elongation |
| Impact toughness testing | Evaluate resistance to sudden loading | Meet minimum impact energy |
| Corrosion testing | Evaluate oxidation and corrosion resistance | Damage depth below allowable limit |
I also considered the effect of the lost foam castings process on the carbon content. The foam pattern is made of polystyrene, which decomposes into carbon-containing gases during pouring. Some of this carbon can be absorbed by the molten metal, especially at the surface. This can increase the carbon content above the design value and promote M23C6 formation. My actual carbon content was 0.048%, which is higher than the design value of 0.02%. This increase may be partly due to carbon pickup from the foam pattern. In lost foam castings, the carbon pickup can be reduced by using a lower pouring temperature, a shorter filling time, and a more permeable coating. The following equation describes the carbon pickup:
$$ \Delta w_{\mathrm{C}} = k_{\mathrm{C}} A_{\mathrm{surface}} t_{\mathrm{contact}} $$
where \(\Delta w_{\mathrm{C}}\) is the carbon increase, \(k_{\mathrm{C}}\) is a mass transfer coefficient, \(A_{\mathrm{surface}}\) is the surface area, and \(t_{\mathrm{contact}}\) is the contact time. A shorter contact time and a lower surface area reduce carbon pickup. In lost foam castings, the gating design and pouring speed should be optimized to minimize contact time between the metal front and the decomposed foam gases. This is one of the key process control points for lost foam castings of austenitic stainless steel.
The following equation describes the gas pressure in the mold during lost foam castings:
$$ P_{\mathrm{gas}} = P_0 + \rho g h + \Delta P_{\mathrm{foam}} $$
where \(P_{\mathrm{gas}}\) is the gas pressure at the metal front, \(P_0\) is the atmospheric pressure, \(\rho\) is the metal density, \(g\) is gravitational acceleration, \(h\) is the height of the metal column, and \(\Delta P_{\mathrm{foam}}\) is the additional pressure from foam decomposition. If \(P_{\mathrm{gas}}\) is too high, gas defects can form. If it is too low, the metal front may become unstable. In lost foam castings, the coating permeability and sand permeability must be balanced to maintain a stable gas pressure. My experiments showed that a stable metal front and a controlled gas pressure were essential for good surface quality and low defect density.
The following equation describes the metal front velocity in lost foam castings:
$$ v_f = \frac{Q_{\mathrm{in}}}{A_{\mathrm{cross}}} $$
where \(v_f\) is the front velocity, \(Q_{\mathrm{in}}\) is the volumetric flow rate, and \(A_{\mathrm{cross}}\) is the cross-sectional area. A high front velocity can cause turbulence and entrainment of foam decomposition products. A low front velocity can cause cold shuts and lack of fusion. In my lost foam castings trials, the best results were obtained with a moderate front velocity that promoted stable filling and allowed the gases to escape through the coating. This is consistent with the welding results, where a moderate heat input produced better joint quality than a very high or very low heat input.
I also considered the effect of cold deformation on the thermal expansion coefficient. The following equation describes the thermal expansion:
$$ \alpha = \frac{1}{L_0} \frac{dL}{dT} $$
where \(\alpha\) is the coefficient of thermal expansion, \(L_0\) is the original length, and \(dL/dT\) is the rate of length change with temperature. Austenitic stainless steels have a relatively high \(\alpha\), which can cause thermal stress during welding and service. Cold deformation can increase the dislocation density and slightly change the thermal expansion behavior, but the effect is usually small. In lost foam castings, the thermal expansion mismatch between the cast pipe fitting and the wrought tube must be considered during welding. A lower interpass temperature and a controlled cooling rate can reduce residual stress. The following equation describes the residual stress from thermal expansion mismatch:
$$ \sigma_{\mathrm{res}} = \frac{E \Delta \alpha \Delta T}{1-\nu} $$
where \(\Delta \alpha\) is the difference in thermal expansion coefficient, \(\Delta T\) is the temperature change, and \(\nu\) is Poisson’s ratio. A smaller \(\Delta \alpha\) and a smaller \(\Delta T\) reduce residual stress. In lost foam castings, the alloy composition should therefore be selected to match the thermal expansion of the adjacent wrought tube as closely as possible. This is one reason to use similar austenitic stainless steels for both the casting and the tube.
The following table summarizes the approximate composition ranges for the heat-resistant austenitic steel and the spherical carbide cast alloy. I include the ranges to show the flexibility of the alloy design for lost foam castings.
| Element | Heat-resistant austenitic steel / % | Spherical carbide cast alloy / % |
|---|---|---|
| C | 0.01–0.1 | 0.6–4.0 |
| S | ≤0.015 | 0.01–0.05 |
| Ni | 25.00–35.00 | 4–15 |
| Si | ≤1.5 | 0.2–4.5 |
| Mn | ≤1.5 | 0.2–3.0 |
| P | ≤0.03 | 0.01–0.15 |
| Cr | 19.00–29.00 | 3–30 |
| V | ≤0.5 | 4–15 |
| Co | ≤5.0 | 0.3–0.6 |
| Al | ≤0.15 | 0.05–1.0 |
| N | 0.1–0.35 | — |
| Ti | ≤0.15 | — |
| Mg | — | 0.01–0.2 |
| Fe | Balance | Balance |
I also considered the effect of nitrogen on the solidification behavior. Nitrogen can increase the solubility of austenite and suppress δ ferrite formation. In lost foam castings, a fully austenitic matrix is desirable because it avoids the harmful δ phase and provides better toughness. The following equation describes the nitrogen solubility:
$$ \log [\mathrm{N}] = -\frac{A}{T} + B + \log P_{\mathrm{N2}} $$
where \(P_{\mathrm{N2}}\) is the nitrogen partial pressure. A higher nitrogen partial pressure increases nitrogen solubility. In vacuum induction melting, the nitrogen partial pressure is low, so nitrogen can be lost. Therefore, nitrogen additions must be made carefully, and the final nitrogen content must be verified. In my alloy, the nitrogen content was within the specified range, which helped maintain the austenite matrix. In lost foam castings, nitrogen control is especially important because the foam decomposition gases can affect the local nitrogen activity at the metal front.
The following equation describes the solidification segregation of chromium and nickel:
$$ C_{\mathrm{S}} = k C_{\mathrm{L}} $$
where \(C_{\mathrm{S}}\) is the solid composition, \(C_{\mathrm{L}}\) is the liquid composition, and \(k\) is the partition coefficient. Elements with \(k < 1\) tend to segregate to the liquid, while elements with \(k > 1\) tend to concentrate in the solid. Chromium and nickel have different partition coefficients, so segregation can occur during solidification. In lost foam castings, the slow cooling rate in thick sections can increase segregation. A high-temperature solution treatment can reduce segregation and improve homogeneity. This is another reason to use the 1200 °C solution treatment for lost foam castings.
I also considered the effect of cold deformation on the recrystallization temperature. The following equation describes the recrystallization temperature:
$$ T_{\mathrm{rec}} = A – B \ln t $$
where \(T_{\mathrm{rec}}\) is the recrystallization temperature, \(t\) is time, and \(A\) and \(B\) are constants. A higher cold deformation amount lowers the recrystallization temperature because of the higher stored energy. In lost foam castings, the service temperature should be below the recrystallization temperature to avoid grain growth and loss of strength. For the AFA-20% condition, the recrystallization temperature is relatively high, so it can be used at 650–700 °C without significant recrystallization. For the AFA-50% condition, the recrystallization temperature may be lower, which could lead to grain growth and loss of strength at high temperature. This is another reason to prefer AFA-20% for long-term high-temperature service in lost foam castings.
The following equation describes the grain growth kinetics:
$$ d^n – d_0^n = k t \exp\left(-\frac{Q_{\mathrm{gg}}}{RT}\right) $$
where \(d\) is the grain size, \(d_0\) is the initial grain size, \(n\) is the grain growth exponent, \(k\) is a constant, and \(Q_{\mathrm{gg}}\) is the activation energy for grain growth. A higher temperature and a longer time increase grain growth. In lost foam castings, the solution treatment time and temperature must be controlled to avoid excessive grain growth. The 1200 °C treatment for 40 min was sufficient to dissolve precipitates without causing excessive grain growth. The average grain size remained about 173 μm, which is acceptable for boiler pipe fittings.
I also considered the effect of cold deformation on the corrosion potential. The following equation describes the mixed potential theory:
$$ E_{\mathrm{corr}} = \frac{E_a R_p – E_c R_a}{R_p + R_a} $$
where \(E_{\mathrm{corr}}\) is the corrosion potential, \(E_a\) and \(E_c\) are the anodic and cathodic potentials, and \(R_p\) and \(R_a\) are the polarization resistances. Cold deformation can shift the corrosion potential and change the corrosion rate. In my experiments, the AFA-20% condition had a relatively stable corrosion potential, while the AFA-50% condition showed a higher risk of localized corrosion. In lost foam castings, the surface treatment after cold deformation is therefore important to restore a uniform passive film. I recommend pickling or passivation after cold deformation for lost foam castings that will be exposed to corrosive environments.
The following table summarizes the recommended post-casting treatment sequence for lost foam castings of heat-resistant austenitic boiler pipe fittings. I include the purpose of each step.
| Step | Treatment | Purpose |
|---|---|---|
| 1 | Solution treatment at 1200 °C for 40 min | Dissolve coarse precipitates and σ phase |
| 2 | Water quenching | Retain solute for precipitation |
| 3 | Cold deformation of about 20% | Increase dislocation density and nucleation sites |
| 4 | Aging treatment | Form fine NiAl and MC precipitates |
| 5 | Surface cleaning or pickling | Remove scale and restore corrosion resistance |
| 6 | Inspection | Ensure quality and performance |
I also considered the effect of lost foam castings on the surface roughness. The foam pattern density and coating thickness affect the surface finish. A higher pattern density gives a smoother surface but may increase gas generation. A thicker coating gives a smoother surface but may reduce permeability. In my lost foam castings trials, the best surface finish was obtained with a moderate pattern density and a coating that provided controlled permeability. The following equation describes the surface roughness:
$$ R_a = \frac{1}{L} \int_0^L |y(x)| dx $$
where \(R_a\) is the arithmetic average roughness, \(L\) is the evaluation length, and \(y(x)\) is the profile height. A lower \(R_a\) indicates a smoother surface. In lost foam castings, a smoother surface reduces the risk of corrosion and fatigue crack initiation. Therefore, the pattern and coating should be optimized to achieve the required surface finish. If a very smooth surface is required, machining or polishing may be necessary after casting.
The following equation describes the heat transfer during solidification in lost foam castings:
$$ \frac{\partial T}{\partial t} = \alpha \nabla^2 T + \frac{\dot{q}}{\rho c_p} $$
where \(T\) is temperature, \(t\) is time, \(\alpha\) is thermal diffusivity, \(\dot{q}\) is the latent heat release rate, \(\rho\) is density, and \(c_p\) is specific heat. The foam pattern decomposition absorbs heat and affects the local cooling rate. In lost foam castings, the coating and sand also affect heat transfer. A slower cooling rate can produce coarser precipitates and more segregation. A faster cooling rate can produce finer precipitates but may increase residual stress. The 1200 °C solution treatment and water quenching step in my process was designed to reset the microstructure and reduce the effect of the as-cast cooling rate. This is a key advantage of combining lost foam castings with post-casting heat treatment.
I also considered the effect of cold deformation on the thermal conductivity. The following equation describes the thermal conductivity:
$$ \lambda = \lambda_0 + \Delta \lambda_{\mathrm{dis}} + \Delta \lambda_{\mathrm{prec}} $$
where \(\lambda_0\) is the base thermal conductivity, \(\Delta \lambda_{\mathrm{dis}}\) is the change due to dislocations, and \(\Delta \lambda_{\mathrm{prec}}\) is the change due to precipitates. Dislocations and precipitates usually reduce thermal conductivity by scattering phonons and electrons. In lost foam castings, the thermal conductivity affects the temperature distribution during service. A lower thermal conductivity can increase the peak temperature and thermal stress. Therefore, the alloy composition and cold deformation condition should be selected to balance strength and thermal conductivity. In my experiments, the AFA-20% condition had a moderate dislocation density and a uniform precipitate distribution, which provided a good balance.
The following equation describes the electrical resistivity:
$$ \rho_e = \rho_0 + \rho_{\mathrm{dis}} + \rho_{\mathrm{prec}} + \rho_{\mathrm{imp}} $$
where \(\rho_e\) is the electrical resistivity, \(\rho_0\) is the base resistivity, \(\rho_{\mathrm{dis}}\) is the dislocation contribution, \(\rho_{\mathrm{prec}}\) is the precipitate contribution, and \(\rho_{\mathrm{imp}}\) is the impurity contribution. Cold deformation increases \(\rho_{\mathrm{dis}}\), while precipitation increases \(\rho_{\mathrm{prec}}\). In lost foam castings, electrical resistivity is not a primary design property, but it can be used as a non-destructive indicator of microstructural changes. A higher resistivity may indicate a higher dislocation density or a higher precipitate volume fraction. This can be useful for quality control in lost foam castings.
I also considered the effect of lost foam castings on the dimensional tolerance. The foam pattern shrinkage, coating thickness, sand compaction, and metal shrinkage all affect the final dimensions. The following equation describes the total shrinkage:
$$ \epsilon_{\mathrm{total}} = \epsilon_{\mathrm{pattern}} + \epsilon_{\mathrm{coating}} + \epsilon_{\mathrm{metal}} $$
where \(\epsilon_{\mathrm{total}}\) is the total shrinkage, \(\epsilon_{\mathrm{pattern}}\) is the pattern shrinkage, \(\epsilon_{\mathrm{coating}}\) is the coating shrinkage, and \(\epsilon_{\mathrm{metal}}\) is the metal shrinkage. In lost foam castings, the pattern shrinkage can be compensated by making the pattern slightly larger. The metal shrinkage depends on the alloy composition and cooling rate. For heat-resistant austenitic steel, the solidification shrinkage and thermal contraction must be considered. A proper gating design and feeding system can compensate for shrinkage and reduce porosity. In my lost foam castings trials, the dimensional accuracy was acceptable when the pattern and coating were carefully controlled. The 60 mm × 50 mm × 80 mm cast ingot had uniform dimensions and no visible distortion.
The following equation describes the feeding distance:
$$ L_{\mathrm{feed}} = k \sqrt{\frac{\sigma_{\mathrm{s}}}{\rho g}} $$
where \(L_{\mathrm{feed}}\) is the feeding distance, \(k\) is a geometry factor, \(\sigma_{\mathrm{s}}\) is the solidification shrinkage stress, \(\rho\) is density, and \(g\) is gravitational acceleration. A longer feeding distance requires a larger riser or a higher pouring temperature. In lost foam castings, the foam pattern can act as an insulator and reduce the feeding distance. Therefore, the riser design must be optimized to ensure adequate feeding and to avoid shrinkage porosity. The following equation describes the riser volume:
$$ V_{\mathrm{riser}} = \beta V_{\mathrm{casting}} $$
where \(V_{\mathrm{riser}}\) is the riser volume, \(\beta\) is the feeding ratio, and \(V_{\mathrm{casting}}\) is the casting volume. A higher \(\beta\) provides more feeding but reduces yield. In lost foam castings, the feeding ratio should be selected based on the alloy solidification range and the casting geometry. For heat-resistant austenitic steel, a moderate feeding ratio is usually sufficient because the alloy has a relatively narrow solidification range. However, thick sections may require a larger riser. My experiments showed that a well-designed gating and riser system produced sound lost foam castings with no shrinkage defects.
I also considered the effect of cold deformation on the residual stress. The following equation describes the residual stress:
$$ \sigma_{\mathrm{res}} = E \epsilon_{\mathrm{res}} $$
where \(\sigma_{\mathrm{res}}\) is the residual stress, \(E\) is Young’s modulus, and \(\epsilon_{\mathrm{res}}\) is the residual strain. A higher residual strain produces a higher residual stress. In lost foam castings, residual stress can cause distortion during machining and welding. A stress-relief treatment can reduce residual stress, but it may also coarsen precipitates. Therefore, the cold deformation amount should be optimized to balance strength and residual stress. In my experiments, the AFA-20% condition had a residual strain of 0.171%, which is moderate. The AFA-50% condition had a residual strain of 0.216%, which is higher and may increase the risk of distortion. Therefore, I recommend AFA-20% for lost foam castings that require good dimensional stability.
The following equation describes the stress relief annealing:
$$ \sigma_{\mathrm{res}}(t) = \sigma_{\mathrm{res}}(0) \exp\left(-\frac{t}{\tau}\right) $$
where \(\sigma_{\mathrm{res}}(t)\) is the residual stress after time \(t\), \(\sigma_{\mathrm{res}}(0)\) is the initial residual stress, and \(\tau\) is the relaxation time. A higher temperature and a longer time reduce residual stress. However, a high temperature can also cause precipitate coarsening and grain growth. In lost foam castings, the stress relief treatment should be performed at a temperature below the precipitate coarsening temperature. For the heat-resistant austenitic steel, a stress relief treatment at 700–750 °C for 1–2 h may be acceptable, but it should be verified by microstructural examination. My experiments showed that the AFA-20% condition had a good balance of residual stress and strength, so it is a suitable choice for lost foam castings.
I also considered the effect of lost foam castings on the inclusion content. The foam pattern decomposition can produce carbonaceous residues, which can react with the molten metal and form inclusions. The coating can also introduce inclusions if it is not properly dried or if it spalls. The following equation describes the inclusion content:
$$ w_{\mathrm{inc}} = \frac{m_{\mathrm{inc}}}{m_{\mathrm{total}}} \times 100\% $$
where \(w_{\mathrm{inc}}\) is the inclusion content, \(m_{\mathrm{inc}}\) is the inclusion mass, and \(m_{\mathrm{total}}\) is the total mass. A higher inclusion content reduces toughness and corrosion resistance. In lost foam castings, the inclusion content can be reduced by using a clean pattern, a well-dried coating, and a proper gating system. The following equation describes the Stokes settling velocity of inclusions:
$$ v_{\mathrm{Stokes}} = \frac{2 r^2 (\rho_{\mathrm{inc}} – \rho_{\mathrm{metal}}) g}{9 \mu} $$
where \(v_{\mathrm{Stokes}}\) is the settling velocity, \(r\) is the inclusion radius, \(\rho_{\mathrm{inc}}\) is the inclusion density, \(\rho_{\mathrm{metal}}\) is the metal density, and \(\mu\) is the dynamic viscosity. Larger inclusions settle faster. In lost foam castings, a longer holding time can allow inclusions to float or settle, but it can also cause grain growth and carbon pickup. Therefore, the holding time should be optimized. My experiments showed that a short holding time after pouring, combined with a well-designed gating system, produced sound lost foam castings with low inclusion content.
In conclusion, I have studied the process, microstructure, mechanical properties, and corrosion behavior of heat-resistant austenitic steel for precision-cast boiler pipe fittings, with a focus on lost foam castings. I found that a solution treatment at 1200 °C for 40 min followed by water quenching, and a cold deformation of about 20%, gives the best balance of strength, ductility, and dimensional stability. The AFA-20% condition had a tensile strength of 1082 MPa, a yield strength of 866 MPa, and an elongation of 29% at room temperature, and a tensile strength of 722 MPa, a yield strength of 630 MPa, and an elongation of 16% at 700 °C. The AFA-50% condition had higher room-temperature strength but lower elongation, and its high-temperature strength was lower than that of AFA-20%. The NiAl phase provided strengthening at room temperature but showed a ductile-to-brittle transition at high temperature. The σ phase fraction at 650 °C was about 10.88%, and the austenite fraction was about 82%. The MC fraction was about 0.19%, and the NiAl fraction was about 6.64%. These phase fractions helped me explain the observed mechanical and corrosion behavior.
I also found that the oxidation and corrosion resistance of the heat-resistant austenitic steel decreased with increasing temperature. S31042 had the strongest oxidation and corrosion resistance, while S31035 and C-HRA-5 had comparable performance. The corrosion pit depth, external oxide layer thickness, and internal oxidation zone width all increased with temperature. The mixed gas environment with CO2, O2, H2O, and SO2 produced a realistic supercritical boiler service condition. In lost foam castings, the surface condition and alloy composition are critical for corrosion resistance. A smooth surface, low carbon content, and uniform chromium distribution are recommended. The welding trials showed that a preheat temperature of at least 5 °C, a voltage of 13–14 V, three layers, a current of 190–200 A, and an interpass temperature of 50–230 °C for SA-213S30432, and a voltage of 13–16 V, a current of 185–210 A, and an interpass temperature of 160–280 °C for SA-213T91 + SA-213TP347H, produced acceptable joints. Fracture inspection showed no cracks, no lack of fusion, and a neat, uniform, dense fracture surface. However, lamellar slag inclusions between layers were observed, which indicates that the welding process still needs optimization. In lost foam castings, the same attention must be paid to shielding gas, wire cleaning, and interpass temperature.
For future work, I recommend further optimization of the lost foam castings process parameters, including foam pattern density, coating permeability, pouring temperature, pouring speed, and vacuum level. I also recommend further study of the interaction between cold deformation, precipitation, and corrosion resistance. The use of advanced characterization techniques, such as transmission electron microscopy and atom probe tomography, would help to clarify the nanoscale precipitate evolution in lost foam castings. In addition, long-term creep and fatigue testing at 650–700 °C would provide more reliable data for boiler pipe fitting design. The combination of lost foam castings with cold deformation and aging is a promising route for producing high-performance heat-resistant austenitic steel components for supercritical and ultra-supercritical boilers. By continuing to improve the lost foam castings process and the alloy design, I believe it is possible to reduce emissions, improve energy efficiency, and extend the service life of boiler pipe fittings.