This report gives details of a new method to determine post heat (tempering) conditions for the spot welds in anti-corrosive high tensile steel sheets. Up to the present, selection of post heat conditions is normally based on the temperature measurement of the nugget. Since this can not be expected to give accurate indication of optimum conditions, the ductility of the weld must be checked by cruciform tension strength, micro-examination and hardness. In practical application a slight variation in welding and post heat conditions can have a singificant effect on weld ductility. Thus, as it is difficult to check the post heat effect by such a method the conditions selected are to be impracticable. The principle of the new method developed in this report is as follows : Electrical resistance of steels rises with temperature but the rate of its rise drops remarkably because of austenite transformation. When the weld periphery has been properly tempered, the center is heated above the level of austenitising temperature. Consequently, the resultant resistance in the weld during post heating increases first and then decreases slowly or steeply depending on post heat current level. Thus, by analysing the resistance characteristics, the optimum post heat conditions can be selected. Practically, since the resistance corresponds to voltage drops, post heat conditions to give the required heat treatment can be readily obtained from the records on an oscilloscope. Short post heat time and narrow allowable range of post heat time will be obtained for higher current while long post heat time and wide allowable range will be for small current. To check the validity of this method, fatigue strength test and hardness survey were carried out; the results are in fairly good agreement with post heat conditions selected. Plotted on the log-log scale, the post heat variable comes out in a straight line. This relation can be formulated by empirical equations. The determination of proper quench time was the most time-consuming operation but it can be made simply by analysing the voltage drop records. The rate of rise in electrical resistivity with temperature is measured and the average temperature in the weld during post heating is calculated.
It is generally known that the majority of fillet weld cracks are toe cracks, underbead cracks and so on. But we have found an unusual crack in the.fillet welds of 50 kg/mm2 grade high strength steels. The crack occurred at the heel-opposite side of the toe-of fillet weld. Therefore we refer to them by the name of "Heel-Crack" in this paper. We have carried out various experiments to make clear the characteristics of the crack and the causes of its formation, and established a practical preventive method for the crack based on these test results. The test results are summarized as follows: (1) Heel-cracks are located at the heel of fillt welds and they are hardly detectable by inspection methods such as visual observation and magnetic particle inspection. (2) It has been difficult to reproduce the heel-cracks by well-known weld cracking tests such as C.T.S. test and Reeve cracking test. We have developed a Non-restraint T type Cracking Test which can reproduce heel-cracks successfully. (3) Heel-crack is a cold crack which occurs at temperatures below 100°C. (4) Main factors which affect the heel-crack formation are the cooling process of the weld, rigidity and hydrogen in the weld. (5) Many 50 kg/mm2 grade high strength steels are susceptible to the formation of heel-cracks. But semikilled steels seems somewhat less susceptible to heel-cracks than killed steels. (6) Crack sensitivity is increased by increasing the carbon equivalent and/or the maximum hardness at underbead. (7) Heel-crack could be prevented by preheating or controlling the bead length. To prevent the heelcrack, higher preheating temperature and longer weld bead are desirable.
A new function mF0(δ) is introduced for the equation of the transient temperature rise, T=q/4πK·eX·mF0(δ) …(31) where mF0(δ), as well as m, δ are given in equ. (6), (32), (33). mF0(δ) has the following interesting characteristics mF0(δ)|m=1 = 1/2· mF0(δ)|m=∞ …(18) mF0(δ) + 1/mF0(δ) = ∞F0(δ) …(17) From equ. (17), the value of mF0(δ) for m>1 can be easily calculated from that of m<1. Fig. 3 shows the calculated results of mF0(δ) for range of m=0-1, in the form of ratio to ∞F0(δ), ∞F0(δ) being given in Table 1 for various values of δ. Equ. (18) suggests an interesting and useful understanding as to the transient temperature rise, namely when we neglect the radiation loss from the plate surface, the temperature of any point, whose distance from the arc point is equal to the distance of arc travel after arc starting, is just equal to the half of the quasi-stationary state value. See equ. (32). It is remarkable that the above mentioned relation expresed in equ. (18) holds true always independent of the thermal conduction constant of plate, arc traveling velocity and time elapsed after arc starting. The result holds under constant linear velocity with constant thermal input in thin plate.
The nucleation and growth processes of ferrite precipitation from austenite in low-carbon low-alloy high strength steels were investigated to make use of SH-CCT diagrams for welding. There were ferrite allotriomorphs and ferrite sawteeth which precipitated from austenite grainboundary, rodlike or needlelike ferrite which precipitated in austenite grain. Growth morphologies of these ferrites were independent of the strength level of testing steels and had the following growth processes : 1) In the case of slow cooling; grainboundary ferrite grew to massive ferrite, ferrite sawteeth and broad ferrite sideplate without carbide precipitation; intragranular ferrite grew to massive ferrite and rodlike ferrite without carbide precipitation; 2) in the case of rapid cooling, grainboundary ferrite grew to narrow ferrite sideplate, intragranular ferrite grew to needlelike ferrite without carbide precipitation. We defined the temperature region where ferrite grew to the configurations described above as the F transformation region in the SH-CCT diagrams for welding. As Zwischenstufengefuge (Zw) transformation proceeded continuously after ferrite transformation, it was difficult to accurately find the begining of transformation of Zw and the difference of Zw and ferrite by dilatation curve and microstructure. The starting line of Zw transformation following ferrite transformation, for these reasons, turns out as shown by broken line. The structual percentage of ferrite and Zw is exhibited as sum'of them under the Zw transformation area. Starting line of pearlite transformation does not correspond to the finishing line of ferrite transformation nor the starting line of Zw transformation corresponds to the finishing lines of ferrite and pearlite transformation.
Anew method for joining cast iron plates is developed utilizing a consumable electroslag welding. Four types of filler materials such as a mild steel wire, a silicon-containing wire, a cored wire containing white cast iron or gray cast iron powder and a meehanite cast iron rod, are used for comparison. By using a mild steel wire or a silicon-containing wire, it is possible to join cast iron, but a very hard zone due to formation of cementite is produced in the weld metal, and the strength of welded joint is unsatisfactory. On the other hand, the weld metal obtained by a cored wire or a cast iron rod process contains 3.05 to 3.38% carbon and 1.73 to 2.58% silicon, and is all satisfactory with respect to hardness, strength, colour tone and damped coefficient.