In the tube design of boilers and other high temperature and high pressure installations, most of the tubes under internal pressure are designed on the basis of creep or creep rupture strength of simple tension bar specimens. In practice, these are used under multiaxial stresses. It is required therefore that the most representative formula will be determined experimentally which is to be applied to the designing of tubes subject to internal pressure at high temperature.
In the present study, the creep tests on tubular specimens of 18-8 Mo steel (thin-walled tube), 21/4 Cr-1 Mo steel (thick-walled tube) and a low carbon steel (electric resistance welded tube) under internal pressure were carried out. The data of the stress rupture tests on tubular specimens were compared with those of simple tension bar specimens cut from the same tubes. The results are summarized as follows:
(1) A reasonable estimation of creep rupture life of the tube subject to internal pressure may be made from the uniaxial creep rupture data by using the Mean diameter formula. This agreement is independent of the steels, the size of the tubes, and of the temperature and time. The reason for this agreement is regared as the redistribution and the levelling of circumferential stress that occurred in the creep stage over the tube wall.
(2) The creep rupture test on the welded tube of a low carbon steel under internal pressure was carried out. The tube was the stretch reduced at 800°C after welding. In this test, the welded tube showed almost equivalent creep rupture strength as compared with the seamless tube.
(3) From this result, the welded tube can be used in high temperature service by giving proper treatment to eliminate the weld structure.

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It is not possible yet at present to relate completely the mechanical properties of materials with the knowledges of physical process. As the deformation process of creep begins to be explained by the dislocation theory and the electronmicroscopic observation, it is expected that the estimation of structural changes in creep and the extrapolation of long-time rupture strength will possibly be made by clarifying the relations between the mechanical factors, such as stress, rupture time, elongation in creep rupture test, and physical factors.
In this paper some information obtained regarding structural changes with regard to creep rupture strength are described, applying Siegfried's method to creep rupture data, which modified the method of O.D. Sherby and J.E. Dorn by introducing temperature dependence for the activation energy. That is, from the curves of stress/log Z(=H/2.3RT-log t_b) for ferritic steels containing 0.5% Mo, 1% Cr 0.5% Mo, 1% Cr 1% Mo 0.25% V, 1.25% Cr 0.5% Mo 0.25% V, 1.25% Cr 0.5% Mo 0.75% Si, 2.25% Cr 1% Mo, 5% Cr 0.5% Mo and 9% Cr, 1% Mo it was shown that a tempering which tends to spheroidize carbides was undesirable to the creep rupture strength, and that the addition of silicon will have an injurious effect on the creep rupture strength while vanadium is available at higher temperature. And 10⁵hr rupture strengths for the above-mentioned steels were estimated from the rupture data within 10⁴hr through the application of this method to the extrapolation of long-time rupture strength. The obtained values agreed approximately with the extrapolated values on ASTM report.
On applying this method to these ferritic steels, the values ΔH which are determined by W. Siegfried were used. In an attempt to evaluate the creep rupture data more specifically, the relationships between a value of ΔH and metallurgical factors-chemical composition, structure and grain size-for various materials must be explained precisely. For ferritic steels, it is considered that a mass effect will practically become an important problem in future.
In the case of applying this method to 18-8 Ti steel and 18-8 Cb steel, it was shown that the break points on the curve of log Z/stress corresponded to the rupture with certain changes of creep deformation process which is caused by the precipitation of carbides and σ phase. And there were some cases that a simple curve of log Z/stress was not fit to be drawn. It was difficult, therefore, to extrapolate the long-time rupture strength. The creep deformation process can be explained as a complex process which consists of various phenomena, i.e. dislocation movement, diffusion of vacancy, viscous flow at grain boundary and grain boundary migration. It is necessary to make some corrections in the values of ΔS and ΔH in the thermodynamical function based on the phenomenon which is found as most prominent. On the extrapolation of long-time rupture strength, it is necessary that further investigations must be carried out to promote the accuracy of the correction by the observation of structural and physical changes in the materials.

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Recent developments in high temperature technology require advanced knowledge of high temperature mechanical properties of graphite. Quite a lot of information has recently become available concerning the high temperature behavior of the mechanical properties of certain types of commercial graphites. However, there has been but little information available on the dependence of the elastic modulus on temperature for carbon materials made from various types of carbon and heat treated at various temperatures. In this paper, the following studies on static elastic modulus of carbons in bending at high temperatures have been carried out as follows; (1) dependence of Young's modulus on temperature for four basic types of carbons heat treated at 3000°C (soft filler-soft binder, soft filler-hard binder, hard filler-soft binder); (2) a complete system of temperature dependence for these four types of carbons and for glassy carbon and low density carbon has been obtained as these carbons progressively heat treated to higher temperature and then cooled back to room temperature; (3) the dependence of Young's modulus on temperature to 2500°C was measured for pyrolytic and ZTA graphite. It was found that for soft filler materials all the curves of Young's modulus vs. temperature shows a maximum below the heat-treatment temperature. Hard filler bodies show a continuous decrease in modulus with increase of temperature. The ZTA graphite, the low density carbon and the pyrolytic graphite behavior are respectively similar to that of soft filler bodies, whereas the glassy carbon behavior similar to that of hard filler materials.

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Stainless steel-castings Type 316 and Type 316L were subjected to high-temperature tensile test, creep-rupture test up to 10000hr at 650°C, and aging treatment under various levels of stress at 650°C and stress-free aging treatment up to 6400hr at 650°C, and the structural changes during these test were observed in order to study the relations between the high-temperature properties and the structural changes of both alloys.
Carbide and σ-phase were indentified by the following means.
(1) The techniques of electrolytic etching by using concentrated strong Hydrooxide solutions (10N, KOH, to color the σ-phase, concentrated NH₄OH, to color the carbide)
(2) The X-ray diffraction analysis of the residues obtained by the electrolytic extraction with 10% HCl alcohol solution and HOC₆H₂ (NO₂)₃ 5% HCl alcohol solution.
(3) Distribution of Fe, Ni and Cr in carbide and σ-phase decomposed from δ-ferrite, examined by X-ray microanalyzer.
The short and long time high-temperature strength of Type 316 is slightly superior to that of Type 316L. The superiority of Type 316 is attributed to the solid solution hardening effect of carbon and the more uniform distributions of fine carbide and σ-phase in the austenite matrix than in the case of Type 316L. But the high-temperature long time load-carrying ability of both alloys seems to be fairly good.
There has been a little difference observed between Type 316 and 316L in their structural changes during the tests, but they take place generally in the following sequences;
(1) The δ-ferrite decomposes into γ-phase, carbide and σ-phase.
(2) The carbide and σ-phase precipitate at austenite grain boundary.
(3) Carbide and σ-phase precipitate in austenite matrix.
(1) Decomposition of δ-ferrite and (2) precipitation at austenite grain boundary are accomplished within 1000hr but (3) the precipitation in austenite matrix gradually proceeds for a long time.
Therefore it is probable that the precipitation of fine carbide and σ-phase in austenite matrix influences mainly the long time high-temperature strength of these alloys, and that Type 316 is a little superior in the high temperature strength to Type 316L, since the precipitates of fine carbide and σ-phase are dispersed uniformly in the austenite matrix of Type 316, while the precipitation in the austenite matrix of Type 316L is limited only round the regions which have been formerly occupied by δ-ferrite.

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It is a well known fact that austenitic stainless steels such as 18 Cr-8 Ni type or high Ni base alloys Inconel (70 Ni-15 Cr alloy). They are quite useful materials as the chemical plant fabricated materials used in a high temperature or corrosion atmosphere service.
Incoloy (32 Ni-15 Cr-45 Fe) is also a useful material as intermediate between the two above mentioned materials. But its high temperature properties, especially creep rupture data are not accurately defined as yet.
We have made high temperature creep rupture tests on domestic Incoloy mother plates and weld joints, and attempted to clarify their high temperature characteristics which has so far been left ambiguous.
The results can be summarized as follows.
1) The high temperature creep rupture strength on the mother plates increases as the grain size grows to the grain number 3-4 by the solution heat treatment.
2) The creep rupture strength on the weld joints {mother plate: coarse-grained material, welding rod: Inco weld-A rod (70 Ni-15 Cr)} is more or less superior to that of the coarse grained mother plate.
3) We dwelt upon the relation which exists among the optional temperature, stress and creep rate. The results are formal that a fine-grained material behavior is superior to a coarse-grained one below 720°C and vice versa above 720°C, and generally speaking any fine-grained material tends to decrease its strength in comparison to any coarse-grained materials as the temperature increases.
4) The high temperature creep rupture strength level of Incoloy is almost identical with those of AISI 304, 310, 330 type or Inconel, which are generally considered to be heat resistant alloys.

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In the present experiment the effect of active nitrogen contents on creep rupture strength of low carbon steels was investigated. The materials were smelted in an atomosphere-controlled induction furnace of 50kg capacity, and the nitrogen contents from 0.004% to 0.03% were covered. Various tests were performed, tensile, bending, charpy impact, recrystallization and creep rupture, The results are as follows:
1. Tensile strength increased with increased active nitrogen content, while elongation and reduction of the area decreased. However, even if as much as 0.03% nitrogen (uncombined) was contained, the elongation and reduction of the area at room temperature remained quite adequate. namely, 30% and 60% respectively.
2. Hot tensile strength in the range of blue shortness temperatures was found to be remarkably higher as compared with commercial steel. The strength increased with increased nitrogen content, while the elongation remained almost unaffected.
3. The transition temperature in charpy impact test became higher with increased nitrogen content as generally recognized.
4. The recrystallizetion temperature shifted slightly to the higher temperature side, as the nitrogen content increased.
5. The short-time rupture strength was found to be definitely improved effectively by nitrogen, thougth the strengthening effect of nitrogen appeared to decline at temperature above 500°C.
Longer-time rupture tests are in progress, so that more comprehensive effects of nitrogen will be obtained.

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To establish the influence of heat treatments and notch sharpness of specimens upon the notchedbar creep rupture strength of a Cr-Mo-V steel for steam turbine rotor forgings, creep rupture tests were conducted with variously heat treated and notched bar specimens. Three ways of austenitization were employed, i.e., heating at 950°C for 20hr (Sample A), heating at 1000°C for 20hr (Sample B), and heating at 1000°C for 20hr, followed by cooling to 950°C and holding at the temperature for 1hr (Sample C). All the specimens were tempered at 675°C for 24hr. Two kinds of notch sharpness were used, one had a notch with elastic stress concentration factor k_t of 6.5 and another of 2.5. The creep rupture tests were conducted with these specimens for about 4000hr at 550°C.
The results obtained were as follows.
(1) The sample B showed the highest plain bar creep rupture strength of theses three samples, the sample C was second to the sample B, and the sample A was the poorest.
(2) Both the sample A and the sample C were notch strengthened within the range of this investigation. However the sample B austenitized at 1000°C was notch weakened at about 1300 hr for kt of 2.5 and about 500hr for k_t of 6.5.
(3) The plain bar creep rupture properties of the sample C were similar to those obtained with the sample B, while for the notched bar specimen the sample C showed a similar behavior to the sample A. From these results it was found that austenitizing at 1000°C was effective for the plain bar creep rupture strength, and cooling from 950°C was effective for the notched bar creep rupture strength.
(4) The rupture time ratios of notched bar specimens for k_t of 2.5 were higher than those for k_t of 6.5 in all the three samples.
(5) From the results of observation of cracks at the notch root, it was found that in specimens for k_t of 6.5 cracks were formed in comparatively earlier stage of test, i.e., within 300 of rupture life, while in the specimens for k_t of 2.5 no crack could be observed within 900 of rupture life.
Consequently, it was supposed that in the specimens for k_t of 6.5 almost all of rupture lives were consumed by the propagation of cracks, while rupture lives of the specimens for k_t of 2.5 were decided by initiation of cracks.

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The cantilever-bending creep tests of 21/4% Cr-1% Mo steel were carried out under steady or interrupted loads at 500°C to 580°C. The tests have revealed the following facts: In the transient creep stage the log-log plots of stress versus creep rate at the outermost fibre of the specimen lie on a straight line or two at any duration of time, and their slopes decrease with the time, indicating the time-dependency of the non-linear stress-distribution over the cross-section of the specimen. The creep rate ε during the transient stage is related to the time t according to ε=at^m or ε=at^m+k, where a, m and k are constants sensitive to temperature and load, similar to those in the conventional tensile creep test. The creep strain is partly recovered during the subsequent period of stress removal, and the strain recovery is characterized by the straight-line relationship between log (rate of strain recovery) and log (time), and by the recovery percentage decreasing with the increasing creep strain.

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The present investigation has been carried out on a water-quenched 18-8 stainless steel in order to find the correlations of torsion or bending creep strength to tensile creep strength. And the analytical consideration due to von Mises criterion was added on the experimental results obtained. Tension creep tests were conducted on 8mmφ cylindrical specimen, torsion creep test on 8mmφ hollow and solid cylindrical specimens, and pure bending creep test on 6mm^t×10mm plate specimen. All the test of the creep were made for 400 hours at 600° and 650°C. The machines employed in the torsion and bending creep test were made in our laboratory. The results obtained can be summarized as follows:
1) The ratio of the torsion creep strength of hollow cylindrical specimen to tension creep strength for a mean creep rate of 2×10^-4∼5×10^-3%/hr during 300 to 400 hours were 45∼46% for 600°C, and 38∼42% for 650°C. On the other hand, the ratio derieved from the calculation based on von Mises criterion was 52% for 600°C, and 53% for 650°C.
2) The ratio of the torsion creep strength of solid cylindrical specimen to the tension creep strength was 54∼56% for 600°C. This magnitude of ratio was smaller than the ratio of 64% obtained by calculation.
3) The ratio of the bending creep strength to the tension creep strength was 141∼158% for 600°C, and 145∼160% for 650°C. On the other hand, the ratio obtained by calculation was 134% for 600°C, and 139% for 650°C.

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A number of experimental works have been carried out on the multiaxial creep problem, but most of them are concerned with the multiaxial creep under steady stress conditions. It is difficult to find creep data for non-steady multiaxial stress conditions, which will provide the basis for the analysis of multiaxial creep problem under the general stress state including the non-steady one. In view of such a situation, the authors have conducted a study of dynamic creep under combined static tension and alternating torsion, and have shown that a strain-hardening type stress-strain rate equation provides a good basis for the analysis of this kind of multiaxial dynamic creep, in the same way as it does for the case of uniaxial dynamic or non-steady stress creep.
This report deals with further extension of the previous work to the more general case of multiaxial dynamic creep, dynamic creep under combined static tension and repeated torsion. In this case, there appear biaxial creep strains, tensile and torsional, while the creep strain was uniaxial in the previous study, torsional creep being zero due to the alternating torsional stress.
The tests were conducted with a low carbon steel at the temperature of 450°C, with a fixed ratio 0.75 of alternating torsional stress to mean torsional stress and with various ratios of mean torsional stress to static tensile stress in the range from zero to infinity. Static creep tests under combined tension and torsion were also made with various ratios of torsional stress to tensile stress from zero to infinity for the purpose of comparison.
The results showed that the tensile creep was greatly accelerated when a repeated torsional stress was superimposed on a static tensile stress. Alternatively, torsional creep was considerably increased by superposing a static tensile stress on a repeated torsional stress. The figure of creep curves was not altered appreciably by combining a tensile stress with a repeated torsional stress as compared with the creep curve for simple tension, except that the period of transient creep was somewhat longer.
These results were discussed from the standpoint of the multiaxial creep theory which was employed successfully in the previous study. In doing this, an effective stress was derived from the information of multiaxial static creep tests under combined tension and torsion, in the same way as was done in the previous study. This effective stress, together with the strain hardening type stress-strain rate equation, was applied successfully also to the dynamic creep of this time, whereas the Mises or Tresca effective stress failed to give the results that agree with the experimental data. Slight discrepancy was observed between the theory and the experiments mainly as the result of the difference in transient period between the static and the dynamic case. However, when a prediction of creep under multiaxial dynamic stress is to be made from the information of static creep data, the results predicted is on the safe side for design, making this method of prediction applicable to the practical design purposes.

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Studies were made of thermal fatigue strength of metals with reference to the relation between the plastic strain range ε_p and a number of cycles to fracture _N. The relation was first presented by S.S. Manson in the from ε_pNα=C, and the values of the constants α and C were given by L.F. Coffin as α=1/2 and C=1/2 ln{1/(1-φ)}, where φ was the reduction of area. In the first part of this paper, these values of the constants are examined experimentally by means of several results of thermal fatigue tests on various metals presented by N. Kato, R.W. Swindeman & D.A. Douglas and others. It is found that in most cases α=0.5∼0.6 and the value of C is smaller than the value given by Coffin's equation and is much affected by thermal cycle conditions, such as upper temperature and mean temperature of the cycle. As the upper temperature T_max and mean temperature T_m of the thermal cycle become high, the value of C decreases rapidly, and the life up to fracture is much reduced, so that Coffin's equation might give a thermal fatigue life on unsafe side.
In order to examine the effect of thermal cycle conditions on the value of C, the authors carried out a series of thermal fatigue tests on 12-Cr steel in the condition that during each series of test the upper and lower temperatures were kept constant and a variable mechanical strain was combined to a constant thermal strain, so that a different resultant strain could be imposed on each specimen under the same thermal cycle condition. The result was that the value of α was not much varied from the mean value α=0.55 by variation of thermal cycle condition, but the value of C decreased as T_max and T_m became high, and was far lower than Coffin's value. From the test results the authors reduced an experimental equation, which expressed the relation between plastic strain range ε_p and the life N, encluding the effect of thermal cycle condition, i.e.
ε_pN^α=K·exp(Q₁/T_m+Q₂/T_max) (1)
In the second part of the paper, the relation is reexamined from another standpoint by the aid of dislocation theory. The proposed model for the mechanism of thermal fatigue is that during the thermal cycling the density of supersaturated vacancies, which are produced by repeated plastic strain, increases gradually and reaches the critical value, then a number of voids are created suddenly by segregation of vacancies, and the density of supersaturated vacancies decreases to a stable state. Thenceforth, the surfaces of the created voids act as absorption saurces of vacancies, that is, the most part of the supersaturated vacancies produced by the subsequent repeated plastic strain is absorbed into the voids. Thus with repetition of cycles, the voids grow and reach a dimention that the adjacent voids touch together, then the voids expand suddenly and microcracks are created and extended to macro-cracks.
By means of this model, some theoretical calculations were carried out, and the relation of the form of eq. (1) was again obtained. The value of the constant α was estimated to be 0.5∼0.6 by the theory, so that it agreed with the above experimental values quite well.

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It is well known that the tool of extruding and the die of casting, of which the materials are Cr-Mo-V steel, are frequently broken by fire crack. It is believed that this is the result of thermal cycling during the hot working. In order to obtain basic information for a reasonable design and working of such material, a series of thermal fatigue tests have been made under large amplitude of temperature. In the present study, particular attention has been directed to the effect that the rapid change of deformability of the present material owing to the temperature level will have on the thermal fatigue characteristics. Attention has been paid also to microscopic observations of nucleation and propagation of thermal crack.
From the present experiments the following conclusions are marked:
(1) Thermal crack nucleates at the crevice of the scale on the surface of the specimen, in case of rich ductility and of fine structure (grain) of the material. The crack at early period grows by repetition of oxidation and at the latter period it propagates under thermal stress cycling. In case of coarse structure (grain) and of poor ductility, the crack propagates rapidly along the austenite grain boundary.
(2) Thermal fatigue life of the material, of which ductility is much dependent on the temperature level, is approximately estimated from the information of static tensile tests of the same material at various temperature levels. This is expressed as
[ΣNfN=1(Δεp)2]1/2=1/√2εf*=1/√2∫T2T1εf(T)dT/T2-T1,
or
Nf1/2(Δεp)=1/√2εf*=1/√2∫T2T1εf(T)dT/T2-T1,
where Δεp is plastic strain range, Nf the number of cycles to fracture, T2 and T1 the upper and lower temperature levels, εf(T) tensile elongation at temperature level T. The latter equation holds in case of constant range of plastic strain during the thermal fatigue test.

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