Design and Characterization of a Heat-Resistant Ferritic Steel Strengthened by MX Precipitates
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The long-term performance of current heat-resistant ferritic steels used in steam generators is primarily limited by microstructure degradation as a result of coarsening or transformation of precipitates. To improve the microstructural stability, a heat-resistant ferritic steel was designed via computational thermodynamics to exclusively contain metal carbonitride (MX) precipitates as the primary means for strengthening at elevated temperatures. The volume fraction of precipitates is 0.35 vol pct, about one-seventh of P91. These MX precipitates are either V-rich or Nb-rich with plate-like or spheroidal morphology, respectively. The precipitate size remains almost constant at 973 K for up to 3000 hours aging. Electron diffraction analysis revealed a Baker-Nutting orientation relationship between the precipitate and the matrix. Consistent with the thermodynamics-based design, M23C6-, Laves-, and Z phase were not detected. The creep threshold stress, derived from high-temperature compressive creep tests, are evaluated to be 63 ± 1 and 43 ± 2 MPa at 923 K and 973 K, respectively, on par with or slightly better than P91. This study reveals that MX precipitates in ferritic steels coarsen slowly at temperatures up to 973 K and that a relatively small volume fraction of MX precipitates can provide effective long-term creep performance at elevated temperatures.
About 80 pct of the electricity generated worldwide today is produced by steam turbines. There has been a push in the past few decades to increase the operating temperature of steam generators to the ultra-supercritical (USC) regime, i.e., above 873 K, for improved thermal efficiency. More advanced USC power plants target temperatures at or above 923 K.[1, 2, 3, 4] This requires creep-resistant alloys that can operate at these temperatures for at least 100,000 hours. T/P91 and T/P92 steels are considered commercial standards for use in supercritical and USC steam generator applications. They rely on dispersion strengthening by MX precipitates, where M = (V,Nb,Ti) and X = (C,N), M23C6 precipitates, and Laves phase Fe2(W,Mo) and solid-solution strengthening by W or Mo.[5, 6, 7, 8, 9, 10, 11, 12] When exposed to these elevated temperatures for extended periods, there are concerns about microstructural degradation due to the metastability of MX, coarsening of M23C6 and the Laves phase, as well as Z phase formation (Cr(Nb,V)N).[13, 14, 15, 16, 17, 18, 19, 20, 21, 22, 23, 24] In particular, the Z phase has been held responsible for the reduction of creep strength and failure of these steels after extended operation, as its precipitation results in the dissolution of finely dispersed MX.[19,25] It has been suggested that reduction of Cr and N concentration can effectively restrict Z phase formation. Coarsening of M23C6 also causes decrease of creep strength due to the loss of effective grain boundary pinning. The coarsening resistance of MX vs M23C6 is illustrated by the work of Rojas et al. M23C6 precipitates grow from 78 to 103 nm after creep deformation at 923 K for about 8000 hours, while MX precipitates start at about 30 nm and remain almost constant after the same creep deformation.
Addition of alloying elements such as W, Co, and B has been used to improve the creep properties of ferritic steels.[27, 28, 29, 30, 31, 32, 33] For example, boron addition at certain alloy compositions can stabilize M23C6.[27, 28, 29] Low-carbon concentrations are known to suppress or eliminate M23C6 formation for improved microstructural stability.[34, 35, 36, 37, 38] For instance, Taneike and coworkers studied the effect of C concentration and found that lower C concentration gives rise to longer rupture time. Another approach is the use of coherent B2-NiAl precipitates for improved microstructure stability.[39,40] The microstructure in this class of steels is analogous to that of γ/γ′ Ni-based alloys. The hierarchical microstructure comprising of a network of B2-NiAl and L21-Ni2TiAl[41,42] has been proved to provide more effective load transfer with less significant diffusional flow along the matrix-precipitate interface compared with single-phase B2-NiAl structure, resulting in superior creep resistance. Threshold creep strength of these steels was shown to exceed 180 MPa at 973 K.
To improve the microstructural stability of heat-resistant ferritic steels, we explore the exclusive use of semi-coherent MX precipitates to provide the necessary creep strength at elevated temperatures, while eliminating the formation of M23C6, Laves phase, and Z phase. The intrinsic stability of MX precipitates and their possible pinning effect on grain and sub-grain boundaries could both contribute to the microstructural stability. The microstructure evolution and creep performance of the designed steel were studied. The steel was demonstrated to possess excellent microstructural stability and creep performance.
2 Alloy Design
Commercial thermodynamic modeling software Thermo-Calc was used with TCFE9 iron-alloy database to assist with the design of the alloy. We adopted the following design criteria: (1) existence of a fully austenitic phase field around 1273 K so as to obtain the tempered martensitic structure upon cooling, (2) thermodynamically stable precipitation of MX without M23C6, Laves phases, and Z phase between 923 K and 973 K, and (3) limiting the use of expensive elements such as Co. Apart from the usual C and Si additions, there are seven key elements included in this design: Cr, Mo, Nb, V, N, Mn, and Ni.
The general rule for precipitation control is to increase the concentration of alloying elements that enlarge the driving force for the formation of desirable precipitates, and reduce alloying elements that promote undesirable precipitates. Based on this principle, we limited Cr to 9 wt pct as a compromise between maintaining corrosion resistance and not promoting the formation of M23C6 and Z phase. Mo is generally added to provide solid-solution strengthening and to slow diffusion. In our case, we set Mo at a low level (0.13 wt pct) to minimize the driving force for the formation of M23C6 and Laves phase. C concentration was chosen to be 0.032 wt pct as a compromise between maximizing the volume fraction of MX and eliminating M23C6 in the temperature range of interest (between 923 K and 973 K). Nb concentration was set at 0.07 wt pct as in many creep-resistant ferritic steels.
Alloy Composition (Weight Percent)
3 Materials and Methods
A 100 g ingot was prepared by arc melting in an AM0.5 arc melter (Edmund Bühler GmbH, Germany). A prototype steel with composition of 0.032C-0.13Mo-0.19V-0.07Nb-1.21Mn-0.23Si (wt pct), designed by our group previously was used as the base material. Then, pure elements (except N, introduced by adding Si3N4) were added to achieve the target composition. The ingot was flipped over and re-melted six times during the process for homogeneity. The composition of the ingot measured by optical emission spectroscopy was listed in Table I. The ingot was then annealed at 1273 K for 1 h in air followed by air-cooling (~ 300 K/min from 1273 K to 673 K). Room temperature tensile test was conducted using a dog-bone specimen (aged at 973 K for 10 hours, with cross-section 3 × 1 mm and gauge length 13.7 mm) with a strain rate of 0.01/min. Appropriately sized samples were cut from the ingot, and polished for different investigations. Vickers hardness measurements at 973 K and rough vacuum (20 Pa) were performed on samples after they were aged at 973 K in air for 10 to 3000 hours. The microstructure was observed by optical and transmission electron microscopy (JEOL 2100F and JEOL ARM300F, JEOL Ltd., Japan). Polished metallographic samples were etched with Vilella’s reagent (picric acid 1 g, hydrochloric acid 5 ml, ethanol 100 ml). The sample for TEM characterization was mechanically polished to a thin foil about 50 µm thick. Disks of 3 mm diameter were punched off from the thin foil and double-jet thinned (TenuPol-5, Struers Inc.), in the electrolyte consisting of 10 pct perchloric acid and 90 pct ethanol at 248 K and a voltage of 30 V. Compressive creep tests were performed at 923 K and 973 K with stepwise load increase in the range of 50 to 100 MPa. Load increase was conducted only after reaching a constant creep rate. The total strain was < 10 pct for each test. All creep specimens, with dimensions of 5 mm diameter and 10 mm length, were heated to 973 K for 10 hours before creep testing.
4 Results and Discussion
4.1 Microstructure Characterization
4.2 Mechanical Properties
Models for threshold stress in creep have been discussed in terms of dislocation climb over precipitates, either in the form of local climb, in which the dislocation follows closely the matrix–precipitate interface, inducing a sharp bend in the dislocation line, or general climb, in which the climb portion of the dislocation extends smoothly from the matrix–precipitate interfacial region to the glide plane.[12,52, 53, 54, 55] The threshold stress normalized by the Orowan stress (σth/σOr) predicted by local climb is about 0.4-0.7, the correct order of magnitude for reported alloys,[9,40] as well as our designed steel. However, the dislocation configuration in local climb with the sharp bend is highly unstable and is thus unlikely to occur. On the other hand, the corresponding σth/σOr ratio for general climb was predicted to be about 0.03 to 0.06, not consistent with experimental findings.
Two refinements were proposed to address the shortcomings of the general climb model by incorporating attractive interaction between dislocations and incoherent precipitates[58,59] and elastic interaction between dislocations and coherent precipitates,[60,61] respectively. In the first case with incoherent precipitates, the attractive interaction is due to the reduction of dislocation line energy when the dislocation is attached to the matrix–precipitate interface, resulting in a detachment stress at the departure side of the precipitate. In the second case, with coherent precipitates, the elastic interaction comes from the stress field of the coherent precipitate induced by the modulus and lattice misfit between the precipitate and matrix. The stress field induced by the misfitting precipitate traps the dislocation at the departure side, resulting in a threshold stress. The incorporation of either the attractive or elastic interaction to general climb brings the normalized threshold stress (σth/σOr) to the correct order of magnitude, consistent with experimental results. This model also suggests the increase of lattice parameter mismatch of precipitates to be an effective strategy to improve creep resistance. The higher lattice mismatch of the MX (mainly VC) in this work might explain the slightly higher threshold stress than P91 where the MX is mainly VN. With our steel, this can be accomplished by increasing Nb concentration, although larger lattice mismatch of precipitates may result in loss of coherency, thereby accelerating precipitate coarsening.
The steel was designed to have MX precipitates as the primary means for strengthening; the volume fraction of these precipitates is only one-seventh of that in P91, a commercial heat-resistant steel used in power plants.
Two types of MX precipitates are predicted by thermodynamic modeling and observed experimentally: V-rich and Nb-rich precipitates, with plate-like or spheroidal morphology, respectively. The precipitate size is almost constant during isothermal aging at 973 K up to 3000 hours.
No M23C6, Laves phase, and Z phase were detected experimentally after aging at 973 K, consistent with the predictions from thermodynamic modeling.
The creep strain rate was found to be highly stress-sensitive. The creep threshold stresses were evaluated to be 63 ± 1 and 43 ± 2 MPa at 923 K and 973 K, respectively, comparable to or slightly better than P91.
These results demonstrate the excellent microstructural stability and creep resistance of our designed steel. Our approach, which centers on the formation of thermodynamically stable MX precipitates while eliminating M23C6-, Laves- and Z phase, has been proved to be effective to enhance the creep strength of the steel at elevated temperatures.
The authors would like to acknowledge the support from the US National Science Foundation (Grant no. CMMI-1462850) and the Chinese Scholarship Counsel (CSC). This work made use of the EPIC facility of Northwestern University’s NUANCE Center, which has received support from the Soft and Hybrid Nanotechnology Experimental (SHyNE) Resource (NSF ECCS-1542205), the MRSEC program (NSF DMR-1720139) at the Materials Research Center, the International Institute for Nanotechnology (IIN), the Keck Foundation, and the State of Illinois, through the IIN. This work used instrumentation at the Northwestern University Center for Atom-Probe Tomography (NUCAPT). NUCAPT received support from the NSF-MRI (DMR-0420532), ONR-DURIP (N00014-0400798, N00014-0610539, N00014-0910781, N00014-1712870) programs, the MRSEC program (NSF DMR-1720139) at the Materials Research Center, the SHyNE Resource (NSF ECCS-1542205), and the Initiative for Sustainability and Energy (ISEN) at Northwestern University. This work made use of the MatCI Facility which receives support from the MRSEC Program (NSF DMR- 1720139) of the Materials Research Center at Northwestern University. This work made use of the CLaMMP Facility at Northwestern University. We thank Prof. David Dunand of Northwestern University for providing access to the creep frame at his laboratory and for his critical reading of this manuscript and many helpful discussions. We would also like to thank Dr. Shrikant Bhat of ArcelorMittal for insightful comments on this manuscript.
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