4.1 Mechanical Properties at Elevated Temperature
Austenitic stainless steels and nickel-based alloys are considered to have satisfactory ductility at elevated temperatures but with lower stress-carrying abilities. AISI 316L and AISI 310 showed a drop in ductility already at intermediate temperatures compared to RT. The increase in ductility at 973 K (700
\(^{\circ }\)C) for AISI 310 should be treated with care since it represents one specimen. Sanicro 25 and Alloy 800HT showed increasing ductility with increasing temperature until it dropped at 923 K and 873 K (650
\(^{\circ }\)C, 600
\(^{\circ }\)C), respectively, compared to RT. Alloy 617 showed increasing ductility with increasing temperature for all elevated temperatures compared to RT. The damage mechanisms found did not differ considerably between the alloys at elevated temperature, except for damage related to DRX in AISI 316L which is described later, and cannot explain the differences in ductility. All materials in the present study have similar crystallographic structure and are stable at elevated temperatures. Thus, the differences in ductility may depend on chemical composition. In terms of ductility, the alloys rank in order of increasing ductility from: AISI 316L/AISI 310
\(\rightarrow \) Alloy 800HT/Sanicro 25
\(\rightarrow \) Alloy 617. This correlates reasonably well with the alloys’ corresponding nickel content. Deformation twinning, which has been observed at elevated temperature for all alloys, except AISI 316L, may increase ductility by twinning-induced plasticity (TWIP).[
30] Low amounts of nickel and the absence of deformation twins partly explain the low ductility of AISI 316L at high temperature. However, AISI 310 showed deformation twinning but the ductility decreased as compared to AISI 316L. Hence, deformation twinning cannot alone explain the better ductility of Sanicro 25, Alloy 800HT and Alloy 617 at elevated temperature. The amount of nickel seems to be important for the high-temperature ductility. The marked drop in ductility for Sanicro 25 at 973 K (700
\(^{\circ }\)C), as seen in Figure
3(c), coincides with a sudden increase in the strain hardening rate, as seen in Figures
3(d) and
13(c). A possible explanation could be a rapid onset of nano-sized precipitate formation at this temperature that was found by Chai
et al.[
25] To be sure of the existence of nano-sized precipitate that forms rapidly at 973 K (700
\(^{\circ }\)C) in Sanicro 25, other microstructural analysis techniques are needed; a transmission electron microscopy (TEM) study is suggested for further investigation.
High-temperature strength may be achieved through solid solution hardening and by precipitation at high temperature. Given the chemical composition of the studied alloys, the following strengthening mechanisms may be expected: Precipitates, such as carbonitrides formed from elements such as carbon, nitrogen, niobium and titanium, increase the high-temperature strength. Elements such as molybdenum, tungsten, nitrogen and carbon are responsible for increasing the high-temperature strength by solid solution hardening. Chromium and nickel are known to increase the high-temperature strength,[
31] and the yield strength in Figure
3(a) and the tensile strength in (b) seem to, to some extent, follow that; when chromium and nickel contents are increased, the high-temperature strength also increases. The yield strengths of Sanicro 25 and Alloy 617 showed slight inclines, whereas AISI 310 showed a steeper incline with increasing temperature; this is attributed the higher amount of solid solution hardening element in Sanicro 25 and Alloy 617 compared to AISI 310. The reason for the higher yield strength with increasing temperature of AISI 310 compared to Alloy 800HT is believed to be the higher amount of chromium in AISI 310. Despite this, the tensile strengths with increasing temperature are rather the same, which is attributed to the higher amount of precipitation-hardening elements in Alloy 800HT compared to AISI 310. The microstructural analysis techniques used in this study could not reveal the assumed nano-scaled precipitates in Alloy 800HT; a TEM investigation is suggested for further investigation.
To further analyse the stress–strain curves in Figure
2 and the mechanical properties in Figure
3, the strain hardening rate as a function of true plastic strain is plotted in Figure
13. Figure
13(a) shows a comparison of all materials at 873 K (600
\(^{\circ }\)C) and Figures
13(b) through (d) gives examples of the strain hardening rates for different alloys. Figure
13(a) shows that AISI 316L and AISI 310 had high strain hardening rates initially, but they decreased during plastic deformation, and the other three alloys showed an increasing or nominally flat strain hardening rate during most of the plastic deformation. AISI 316L displayed a decreasing strain hardening rate with increasing true plastic strain at elevated temperatures, as seen in Figure
13(b). The strain hardening rate of AISI 316L was expected be low due to low amounts of solid solution-strenghtening and precipitation-hardening elements. Sanicro 25 first increased in strain hardening rate with increasing strain until a certain true plastic strain level was reached where it, from then on, decreased with further deformation, as shown in Figure
13(c). Alloy 617, as seen in Figure
13(d), showed an increase in strain hardening rate with increasing true plastic strain essentially all the way until fracture at elevated temperatures. The latter two alloys were expected to exhibit a high strain hardening rate due to strengthening mechanisms such as solid solution and precipitation stenghtening. In addition, DSA may increase the strain hardening rate in austenitic stainless steels, but since all alloys showed signs of DSA, even AISI 316L and AISI 310, it cannot be the single reason for the differences. Deformation twinning may also increase the strain hardening rate.[
30] However, AISI 310 showed no increase in strain hardening rate during plastic deformation even though it experienced deformation twinning.
In Figure
4, serrated yielding from some of the curves in Figure
2 are presented. Each curve was characterized with a PLC type and the serrated yielding was attributed to DSA. However, serrated yielding from deformation twinning cannot be precluded in the stress–strain curves, especially not at higher strain levels. Deformation twins were often found closer to the fracture surface and are related to higher stress levels which was experienced at high strain levels and closer to fracture. The differences of PLC types seems not to correlate to the differences of mechanical properties in Figure
3 or the active deformation mechanisms.
The investigated materials showed planar dislocation-driven deformation at elevated temperatures, as shown in Figures
12(a) and (b). Along the path indicated by arrow (B) in the EBSD map in Figure
14, a more continuous change in misorientation indicates dislocation activity, like planar slip and slip bands.
It is interesting that deformation twins were found at elevated temperatures. Deformation twinning usually does not occur in austenitic stainless steels at the tested temperatures.[
24,
29,
32] Deformation twins as shown in Figures
7 through
10, and the surrounding area, was further studied by EBSD; Figure
14 shows an EBSD map of a few twins in a specimen tested at 923 K (650
\({^\circ }\)C). The crystallographic orientation is displayed according to the colored stereographic triangle. The misorientation profiles in Figure
14 show variations in orientation along the directions marked by arrows (A) and (B). Arrow (A) shows an angular difference of 60 deg, indicating twin boundaries. To further analyse the deformation twining at elevated temperature in austenitic stainless steels and nickel-based alloys, a TEM investigation is suggested.
Deformation in austenitic stainless steel depends on the stacking fault energy (SFE). SFE is influenced by the alloying elements[
33] and temperature.[
34] Deformation twinning is active in the interval 18
\(\le \) SFE
\(\le \) 45 mJ m
−2.[
35] Under 18 mJ m
−2 , phase transformation from an austenitic phase to a martensitic phase is favored, and above 45 mJ m
−2 , the deformation process is controlled by dislocation glide.[
30] The low-alloyed steels can be assumed to have an SFE of 25± 5 mJ m
−2[
36] and the nickel-based alloys an SFE of 45± 5 mJ m
−2[
37] at RT. Since the SFE increases with increasing temperature,[
34] dislocation glide should be the favored deformation mechanism, at least in the nickel-based alloys. This was also supported by deformation mechanisms characterization using ECCI and EBSD. However, deformation twinning was active and often found closer to the fracture surface along with dislocation-driven deformation.
Deformation twinning is usually restricted in austenitic stainless steels at elevated temperature due to low stress levels.[
24,
29,
32,
38] Lin Peng
et al.[
39] reported that plastic deformation in austenitic stainless steels occurs by planar slip, followed by formation of stacking faults and eventually multi-directional slipping. Only at higher plastic strains does twinning become the dominant deformation mechanism as grains that favor slip become exhausted.[
39,
40] The deformation twins were found close to the fracture surface, which could mean that deformation twins were activated close to the fracture in the deformation process where the stresses were high. The high stress levels needed to activate deformation twinning at elevated temperatures and close to the fracture surface could also be supported by DSA, as proposed by Yapici
et al.[
32] However, AISI 316L did not show deformation twinning even though it experienced DSA at elevated temperatures; it, instead, showed DRX at elevated temperature and close to the fracture surface. The existence of deformation twins at elevated temperature and high SFE is probably due to preserved high-temperature strength from solid solution and precipitation strengthening supported by DSA. These strengthening mechanisms and the higher stresses close to the fracture in the deformation process, leads to stress levels high enough to activate deformation twinning, in all tested alloys except AISI 316L, which instead showed DRX.
At higher temperatures, when the recovery mechanisms are active, AISI 316L shows damage at grain boundaries, as can be seen in Figures
11(a) and (b). The dislocation density is increasing as the plastic deformation proceeds at the grain boundaries, as severe plastic deformation in Figures
11(a) and (b), suppressing further deformation. If the recovery occurs,
i.e. formation of subgrains and recrystallization at the grain boundaries, as shown in Figures
11(c) and (d), they will soften the grain boundary regions and dislocations can again multiply in these regions. If this process is repeated, it will initiate microcracks, as shown in Figures
11(a) and (b), at the grain boundaries which may lead to fracture and lower ductility.[
41]