The testing of hot hardness is recently widely practised and drawing attention as a convenient way to find the mechanical properties of engineering materials, e.g. tensile strength, creep strength and creep rupture strength at an elevated temperature. There are two types of apparatuses for testing hot hardness that are now available in Japan to measure the Vickers hardness of materials, both metallic and non-metallic, and the measured values of the hot hardness are greatly variable depending on the measuring conditions and methods. In order to obtain the reliable values of hot hardness it is absolutely necessary, therefore, that the measuring conditions and method will be carefully studied. We have performed a number of experimental studies regarding the testing procedures and the conditions in connection therewith, and have been able, thanks to the cooperation of twelve research laboratories, to obtain a criterion for proper conditions and methods for the testing of hot hardness.
It is shown that the creep deformation resistance of a variety of metals is reduced by repeated reversed deformation at temperatures above 0.4Tm, where Tm is the absolute melting temperature of each metal. The reduction of creep deformation resistance due to stress reversals is most prominent at approximately 0.5Tm where the average creep rate may increase by as much as a factor of twenty due to 50-100 stress reversals. Metallographic studies show that the observed acceleration of creep at high temperatures is in part due to the enhancement of grain boundary sliding as a result of gradual grain boundary migration toward planes of maximum shear stress during reversed creep deformation.
It has been confirmed that the precipitated carbides in 18 Cr-8 Ni stainless steels play an important role in determining the creep rupture strength. The finer the dispersion of the carbides is, the higher rises the rupture strength. The small amounts of titanium and niobium, especially when both the elements coexist, have been found to be effective controlling elements to obtain a fine dispersion of the carbides. The addition of titanium and niobium of the order of 0.1% to 18 Cr-8 Ni stainless steel will considerably increase its creep rupture strength. To 18 Cr-10 Ni-2 Mo steel, also, small addition of titanium and niobium together with boron and nitrogen will have an effect of preventing abnormal decrease in creep rupture strength, and of giving the steel such high rupture strength as 18.5kg/mm2 at 700°C-1000hr.
The creep rupture strength of low alloy steels, 1/2Mo, 1Cr-1/2Mo, 21/4Cr-1Mo, 5Cr-1/2 Mo, 9Cr-1Mo was studied in relation to their microstructure. The results obtained are as follows; (1) The intersection of the stress-rupture time curves was observed in a steel with different heat treatment. This was observed also in 1Cr-1/2Mo and 21/4Cr-1Mo steels. (2) High Cr steels (5Cr-1/2Mo, 9Cr-1Mo) show weaker creep strength than low Cr steels (1/2Mo, LCr-1/2Mo, 21/4Cr-1Mo). (3) The Cr content in the residues does not increase by tempering. As for the amount of Mo in the residues, low Cr steels show remarkable increase in Mo content but high Cr steels show slight increase in Mo content. (4) Microscopical observation revealed that N. T. treated specimen contained finer precipitates in the ferrite than Ann. treated specimen. Tempered bainite which was observed in N. T. treated 21/4Cr-1Mo steel was apt to form coarse carbides. In fine the problems with the creep rupture strength of commercial low alloy steels will more clearly be boiled down to alternative between the following two questions: (1) whether or not the low alloy steels contain in their microstructure either bainite or martensite, or (2) whether or not the ferrite is adulterated with carbide particles.
The fatigue deformation of 18-8 austenitic steel and aluminum due to high temperature has been investigated mainly from the stand point of cross slip using the replica and transmission electron microscopy technique. The main results obtained are as follows. (1) Remarkable difference is found in the configuration of slip lines in both 18-8 austenitic steel and aluminum when fatigued either at room temperature or at elevated temperatures. In the former of low stacking fault energy, the morphology of surface slip lines formed at room temperature is fine and straight, while at elevated temperatures it is characterized by the thick and wavy appearance. In the latter of high stacking fault energy, on the other hand, the waviness increases with rise in temperature so strikingly that remarkable irregularity on specimen surface occurs. The morphology of deformation structure at high temperatures depends also on the stacking fault energy. The cell structure is more distinct at elevated temperatures than at room temperature. (2) The fatigue crack is initiated at grain boundary in fatigued 18-8 austenitic steel at 500°C, while at 300°C it is initiated at well-developed intrusion and at the region near the grain boundary, where the intensified slip lines including many cross slip lines intersect each other, and finely divided substructures are thus formed. (3) The well-developed substructures are formed near the crack in fatigued 18-8 austenitic steel at elevated temperatures. Therefore the fatigue crack might grow with the same mechanism as that in metals of high stacking fault energy at elevated temperatures. The clear cellular structures on the surface are observed at the crack tip in fatigued aluminum at high temperatures, and fatigue crack grows through the boundaries of these cellular structures of the surface. (4) Many pores were observed on the surface of the specimen of aluminum fatigued at elevated temperatures. It is considered that the pores are formed beneath the surface oxide film by the condensation of vacancies which were abundantly created during fatigue deformation at the elevated temperatures. (5) The characteristic feature in fatigue at elevated temperatures can be explained by considering cross slipping enhanced by high temperature.
In order to collect the results of micro-fractographic studies of austenitic stainless steel SUS 32 with fractured surfaces fatigued at high temperatures, and to find how far the number of cycles to fatigue failure is dependent on temperature and strain rate, the material was submitted to push-pull low cycle fatigue tests at high temperatures, and observation of the fractured surfaces was carried out by a scanning electron microscope. For fatigue test conditions the strain rates were set at 40, 4 and 0.4%/min, and the temperatures were set at room temperature, 450°, 600° and 700°C. The conclusions obtained are as follows; (1) Striations are observed in each fatigue condition. The striations observed clearly at high temperature are also corresponding to the respective strain cycles. (2) The striations are bent at twin boundaries. (3) Under the conditions of 600°C, 0.4%/min and large strain ranges, grain boundary fracture is observed clearly, which is of the same sort as that which is observed on the fracture surface at the creep rupture test made at 600°C. (4) Uuder the condition of 700°C, grain boundary fracture is also observed, but its surface is not always smooth. (5) Under the strain rate of 40%/min at 600° and 700°C, striations are observed more easily than grain boundary fracture.
In the present paper, the analytical and experimental studies on tensile creep of polycrystalline metal under hydrostatic pressure at elevated temperatures were presented, in a series of studies on the influence of hydrostatic stress on plasticity and fracture laws of the metals. From the tests performed of commercial pure aluminum at 200°C, the following conclusion has been derived. (1) The effect of hydrostatic pressure on minimum creep rate at elevated temperature is more intensive than that on tensile flow stress at the same temperature. This is considered to be mainly due to the pressure on the diffusion of vacancies in the creep process. (2) In order to predict the effect of hydrostatic stress on the creep rupture of the metal at elevated temperature from that on the steady state creep at the same temperature, it is required that the pressure effect on the structural change during the tertiary creep, such as inhibition of growth of void as the result of reduction in the diffusion of vacancies will be taken into consideration.