Methods which measure hardness by indentation are widely used in engineering. The size of indent imposed in a material varies with the load used and the shape of the tool tip used to make the indent. Standard hardness measurements may introduce indents some hundreds of μm in diameter or diamond diagonal length. Such test methods return results which measure the average hardness of fine grained polycrystalline metals. Mixed phase materials have been shown to display hardness relative to the properties of each discreet phase. Where micro-indents smaller than the materials grain structure are made it can be possible to separate the hardness contribution of each phase. However, the effects of the surrounding material on the measured results from isolated grains can lead to inaccurate assessment of the materials properties.
Sawada et al. [239] performed nano-indentation tests on ASME T91 alloy which had been subjected to a variety of thermal treatments and creep tests. The aim of their work was to quantify the contribution of microstructural features to the hardness exhibited by the alloy.
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Quantification of dislocation densities, precipitation populations, lath, block and packet size and inter-grain angles (grain boundary miss-orientations) were made. This data was correlated to hardness measurements recorded by both micro-hardness and nano-indentation results. Their findings showed that, in agreement with results in this work (fig. 7.8), the overall hardness of the material reduced as a function of time at temperature. They attributed the loss of strength to an increase in inter-particle spacing, a decreased dislocation density and the increased distance between high angled boundaries. They conclude that prior austenite grain boundaries do not contribute to strengthening and the effect of lath boundaries is only minor. They quote high angle boundaries as contributing significantly to strength as does dislocation and precipitation hardening. They were unable to separate contributions of dislocation and high angle boundaries but the decreased dislocation densities and coarsening of subgrains and precipitation which was encountered during creep testing resulted in a significant loss of hardness.
The hardness of recrystallised P22 weld alloy was measured between 120 and 150HV. Once recrystallised, no further degradation of hardness was apparent, even in crept specimens at equivalent times at 730°C approaching 25 hours. In contrast, continued exposure of bainitic and martensitic microstructures, to high temperatures, resulted in an unremitting loss of hardness. It was concluded that recrystallisation resulted in the instantaneous removal of strengthening mechanisms; whereas, continued tempering of bainitic and martensitic material resulted in the progressive removal of strengthening mechanisms. Results show the hardness of P23, P24 and P91 alloys tending toward the inherent strength of the alloy. Recrystallised 8 hour PWHT P22, P23 and P24 material returned averaged hardness values of 144, 168 and 162HV respectively. The values recorded for recrystallised P23 and P24 material was based on only 5 readings some of which were made within 50μm of sample
edges. The greater hardness of MX forming alloys could be the result of mechanical constriction provided by surrounding bainitic material or could be a result of the retention of MX precipitation. Tsuchiyama et al. [187] report an inherent hardness of 120HV for recrystallised low carbon 9Cr-1Mo alloy.
The high initial hardness of P91 HAZ material which ranged approximately between 400 and 450HV corresponds to that expected for similar martensitic material [246] and was comparable to results reported in literature [234].
After 2 hours PWHT at 730°C, P23 alloy displayed an average hardness of approximately 225HV. The hardness of material adjacent to the fusion line, in a region which was most probably depleted of carbon, showed some softening but hardness was not significantly lower than that of the bulk weld material.
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The response of welds made with P24 alloy to post weld heat treatment was quite different to that of P23. Although the application of a 2 hour PWHT at 730°C softened the weld alloy by approximately 136HV, the average hardness remained comparatively high at 266HV (std. dev. 10.5HV). Multiple hardness tests were carried out on P24/P91 welds in 2 hour PWHT condition, all of which gave similar results. Whether the difference between P23 and P24 in their response to tempering was a result of differing rates of softening or a secondary hardening effect on P24 was not clear. Tempering of carbide forming steels often involve complex responses to exposure to high temperatures. The results presented in this work correspond to 2 and 8 hour PWHT and a number of creep tested specimens. The hardness of materials corresponding to the intervals between 2 and 8 hours are not known and may have involved secondary hardening effects. Studies carried out by Mohyla et al. [122] on P23 and P24 welds revealed multiple secondary hardening peaks in the case of P24 alloy, whereas, P23 displayed only one. This indicates a difference in precipitation response to PWHT but other microstructural strengthening contributions must also be considered. However, this does not explain the anomalous behavior observed of 2 hour PWHT P91 HAZ in joints made with P24. Again the results were consistent for the multiple tests performed but similar behavior was not seen in joints made with P22 and P23 weld alloy in 2 hour PWHT condition. Tests on the same welds were also carried out at Doosan Babcock of Renfrew; their results were in good agreement with those presented here.
Accuracy and scatter of hardness measurements made using the Matsuzawa Seiki micro- indentation hardness tester were calculated using alloy calibration blocks of known hardness. A mean hardness of 212.5HV with a standard deviation of 8HV was recorded using 50gr loads on a 220HV calibration block; the graphs have not been corrected accordingly. Assuming the inaccuracies in indentation measurement remained constant; the relative error was expected to increase when measuring harder regions of material due to the smaller indentations. The opposite would be true for softer material.
Further work is required to substantiate the mechanisms responsible for the response of P24 and P91 HAZ material that is welded to P24 alloy to post weld heat treatment. There is also scope for further work regarding the separation and quantification of the effects of individual strengthening mechanisms which contribute to the mechanical properties of ferritic, martensitic and bainitic materials.
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CHAPTER EIGHT
Precipitation in 2.25 and 9 % Cr Creep
Resistant Steels
8.1
Introduction
Recent decades have seen the development of ferritic steels in which the control of carbide, nitride and carbonitride precipitation has resulted in enhancement of their creep performance. Models which predict the thermodynamics and kinetics of precipitation, precipitation sequences and coarsening rates can be used to predict stable phases in these alloys [111, 247-250]. By careful selection of alloying additions and the application of stringent thermo-mechanical treatments it has become possible, in many ways, to control the species, size and distribution of particles and tailor alloys accordingly [56]. Exposure to high temperatures and stresses, during service, results in the aging of alloys, which manifests as precipitate phase changes and coarsening. Precipitation of carbides of increasing thermodynamic stability frequently occurs at the expense of metastable phases. These metastable phases precipitate, not because they represent the lowest free energy state but because they nucleate easily. Changes in matrix composition which occur as a result of preferential growth of specific carbides can also affect the stability of other pre-existing precipitates.
In the case of dissimilar metal welds made between 9% and 2.25%Cr creep resistant steels an added complication in the control of precipitation comes in the form of compositional changes at the weld interface.
This section of experimental work was carried out in order to determine the types and distributions of precipitation present in different areas of dissimilar metal welds. Examination of welds, in as received and post weld heat treated condition, was completed. Phase identification and characterisation was performed by a combination of SEM, TEM, EDS and XRD. Efforts were made to substantiate the formation of, and retention of, MX type precipitation in P23 and P24 welds. In particular, confirmation of their presence in weld metal, which had become depleted of carbon during PWHT, was sought.
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