The history of ferritic-martensitic (F-M) steels begins in 1912 with the first reported use of high chromium (Cr) steel being a 12wt% Cr and 2-5wt% Mo steel blade developed for use in steam turbines. The following year ‘stainless steel’ was developed when it was noticed that steels containing 13wt% Cr and 0.2wt% C did not rust, with the applications originally being as cutting edges due to the ability of the martensite to hold an edge. Although high Cr steels began being used in industry in the 1930’s, it was not until the jet age that a material with a high corrosion resistance, and a high operating temperature was required for use. This led to the development of high Cr steels with enhanced UTS, and creep rupture strength at temperatures of 550°C, and a desire to reduce both air pollution and operation costs of energy production in the 1990’s pushed the development of a steel with a creep rupture strength of 100 MPa at 105 hours while operating at 600° C [3].
2.1.3.1 Processing
Depending on the processing techniques, and alloying concentrations utilized, high Cr steels can be fully martensite, martensite and ferrite, martensite and austenite, or a combination of all three phases. Prior to quenching, a fully austenitic structure, or a combined austenite and δ-ferrite structure is engineered using a combination of alloy additions and heat treatments. In 12wt% Cr steels the alloying elements C, N, Ni, Mn, Cu, and Co promote the growth the austenite phase, while Cr, Mo, Nb, V, W, Si, Ti, and Al stabilize ferrite grains [3]. Chromium also promotes the formation of the ferrite phase, and if the concentration is too high then austenite will not precipitate out, and the
concentration of martensite will be limited [9].
Which elements are used depends on desired mechanical properties, and the type of operating environment the material will be exposed to [3]. A study by K. Hashimoto et al in 1983, found that for operation at high temperatures stabilizing elements should be added to replace carbon, and limit the carbide growth. As the carbides get larger, the precipitation hardening effect lowers as the precipitates coalesce into larger features, and strengthening elements are striped out of solid solution, removing some of the solid solution hardening. It was recommended that the concentration of carbon be less than 0.2wt% [10]. Abe, Araki, and Noda in 1991, and Shikakura et al in 1991 looked into the effect the W has on the final alloy, and found that it slows the recrystallization rate, and added to the long term thermal stability of the M23C6 precipitates that form during the tempering treatment [11], [12]. The mechanical effects are not always the most
an austenite forming element it doesn’t provide as much strengthening as carbon, but nickel is the preferred choice in nuclear applications as it is more difficult to activate [3].
As the combination of alloying elements is altered it is possible to predict the final phases that will be present based on nickel and chromium equivalents. Equation 2.8 and Equation 2.9 are valid for 12wt% Cr alloys, and are used along with a Schaeffler-
Schneider diagram, shown in Figure 2.11, to predict the microstructure that will result post processing. 𝐼𝑁𝑒𝑒𝑒𝑋𝑣𝑚𝑙𝑒𝑛𝑡(𝑤𝑤%) = (%𝐼𝑁) + (%𝐶𝐶) + 0.5(%𝐴𝑛) + 0.3(%𝐶𝐶) + 30(%𝐶) + 25 (%𝐼) Equation 2.8 [3] 𝐶𝑟𝑒𝑒𝑒𝑋𝑣𝑚𝑙𝑒𝑛𝑡(𝑤𝑤%) = (%𝐶𝑟) + 2(%𝐶𝑁) + 1.5(%𝐴𝐶) + 5(%𝑉) + 1.75 (%𝐼𝑏) + 0.75(%𝐶) + 1.5(%𝐴𝑁) + 5.5(%𝐴𝐴) + 1.2(%𝐴𝑎) + 1.2(%𝐻𝑓) + 1.0(% 𝐶𝐶) + 0.8(%𝑍𝑟) + 1.2(%𝐺𝐶) Equation 2.9 [3]
Relationships like these are verified against experimental results, and similar equations are developed for each type of material, with the relationships for a 9wt% Cr steel being found in High-Chromium Ferritic and Martensitic Steels for Nuclear Applications, by Klueh and Harries.
Although these are powerful tools, in actuality the percentage of austenite that transforms to martensite during cooling depends not only on the on the alloying elements, but also on the processing time and temperatures. Two temperatures of interest for high
chromium steels are the martensite start, MS, and martensite finish, MF, temperatures. Both of these temperatures are lowered by the addition of alloying components, and can prevent a full martensite conversion, leaving residual austenite in the matrix that
increases toughness while reducing hardness. The martensite start temperature can be estimated as: 𝐴𝑆(°𝐶) = 635−474[(%𝐶) + 0.86(%𝐼) −0.15(%𝐼𝑏+ %𝑍𝑟)) −0.066(%𝐴𝑎+ %𝐻𝑓)]−17(%𝐶𝑟) + 33(%𝐴𝑛) + 21(%𝐴𝐶) + 17(%𝐼𝑁) + 39(%𝑉) + 11(%𝐶) Equation 2.10 [3]
The rapid quenching process associated with F-M steels results in low carbon martensite laths whose hardness is a strong function of the carbon and nitrogen content. Prior to tempering, a double austenitizing treatment is applied to create uniformity between the prior austenite grains and the martensite structure, increasing the creep rupture strength.
To reduce brittleness, a tempering step is preformed below the α-iron to γ-iron transition temperature, which is altered by the alloying elements present, to prevent reaustenitization of the material. Table 2.1 shows the effect of different alloying
elements on the transition temperature, with the ferrite to austenite transition temperature ranging from 870° to 960°C for reduced activation steels. The evolution of the
microstructure throughout the tempering process is as follows:
• When T <350°C the M3C (Fe3C) precipitates form and grow using a branching dendritc structure to Widmanstätten ribbons. The precipitates
become chromium enriched, and M7C3 can form. The growth of these precipitates slows the rate of tempering.
• When T ≈ 450° - 500°C the Cr2CN (M2X) needles begin to form on dislocations in the martensitic laths, which slows the rate of tempering. • When T is between 500° and 550°C the precipitates grow in size, resulting
a drastic softening effect.
• When T > 550°C the M7C3 and M2X precipitates transform to M23C6 precipitates which are rich in chromium. This slows the rate of softening, and results in a reduction in dislocation density in the quenched
martensite.
• When T ≥ 650°C the M23C6 precipitates found at the martensite lath boundaries coarsen.
• When T ≥ 750°C, the grains within the martensitic laths become equal subgrains that may still have dislocations present. The M23C6 precipitates have removed the majority of carbon from solution as they continue to grow.
The rate of growth of the M3C, M2X, M7C3, and ultimately the M23C6 precipitates is heavily influenced by the alloying element concentration. Low concentrations of nitrogen, 0.02 to 0.03%, results in the preferential growth of Cr2N (M2X) over M7C3, which initially increases secondary hardening and overaged hardness by leaving higher concentrations of carbon in the matrix that drives the formation of additional precipitates, and raising the volume fraction of precipitates, which increases the overall hardening. Nickel accelerates precipitate growth, which lowers the tempering resistance [3], [13].
The final structure of both reduced activation and conventional high chromium martensitic steels contain martensitic laths with dislocations characterized with a Burgers vector of 1/2a0<111>, and M23C6 precipitates found both on prior austenite/ferrite grain boundaries, and within the martensite laths. The precipitates within the lath structure are finer than the precipitates found on the grain boundaries. M2X precipitates are also found within the martensite laths. This structure is shown in Figure 2.12 [3].
2.1.3.2 Effect of Chromium on Material Properties
The measured hardness, yield strength, and UTS in F-M steels trend with the chromium concentration within the alloy. As more chromium is added in solid solution, increased hardness is observed, with additional hardening observed if other solid solution strengthening alloys such as tungsten have been added [13]. Similar trends with the yield strength and UTS have been reported for F-M steels that have evolved to have higher creep-rupture strengths such as: T9, T91, HT9, HCM12, and HCM12A steels, which have chromium contents of 9wt%, 9wt%, 12wt%, 12wt%, and 12wt%, respectively. The mechanical properties are observed to reduce in magnitude as the testing temperature increases [13], [14]. The processing differences for the above mentioned steels are shown in Figure 2.13.