F transformation region in SH-CCT diagram for welding of high strength steels was investigated to make use of the dilation curves and A→F transformation curves. The starting line of F transformation as measured by the dilatation and the thermal analysis method has departed slightly from the 1% precipitation lines of ferrite. The starting line of F transformation has departed maximum ±10°C from the 1% precipitation line of ferrite. In the case of slow cooling, F transformation curve was lower than the 1% precipitation line of ferrite. On the other hand, F transfer ma tion curve, in the case of rapid cooling, was higher than the 1% precipitation line of ferrite.
When the dilatation curve possesed a two stage expansion process, the growth of ferrite did not cease in the temperature region between the start of secondary contruction and the start of secondary expansion of the dilatation curve during the cooling process, but continued even in this temperature region. For this reason, we considered that no nontransformation region existed between the F and Zw transformation regions in SH-CCT diagram for welding. It was considered that the two stage expansion process of dilatation curve above mentioned was the effects of Cr, Mo and the other elements which were a strong carbide former. Moreover, the F transformation region of SH-CCT diagram for welding was defined.

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溶着金属が150℃に冷却した時点での拡散性水素量を測定することを提案した.しかして予熱,直後熱など溶接条件の変化によるこれらの拡散性水素量を測定し,定ひずみ割れ試験における溶接の遅れ割れとの関係を調査した.さらにこれらの拡散性水素量を用い低温で発生する溶接割れ防止におよぼす予熱の効果について考察を加えた.また405ステンレス鋼の溶接割れの潜伏期間が炭素鋼や高張力鋼に比べ長期化する理由について,拡散性水素の拡散係教から考察した.それらの結果をまとみると次のごとくである.
(1)新しく提案した測定法による拡散性性水素量は予熱によって減少する.したがって予熱が溶接割れ防止に有効なのは溶接部の軟化による伸び,じん性の改善よりもむしろ拡散性水素量の減少にあることが明らかになった.
(2)405ステンレス鋼と炭素鋼の拡散性水素の常温における拡散教を計算して405ステンレス鋼は炭素鋼のほぼ1/10～1/25になることを示した.したがって405ステンレス鋼の溶接における割れの潜伏期間が炭素鋼のそれに比べてきわめて長期化する大きな理由は拡散性水素の拡散係教が非常に異なるためであることを示した。
終りに本研究の実施に当り,適切な御教示を賜わった渡辺正紀先生,井川博先生に厚くお礼申し上げます.

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In our previous paper, the nature of the precipitate in 17-4 PH stainless steel was described. In this paper, we report on the weld heat-affected zone of 17-4 PH stainless steel.
The main conclusions obtained in the study are as follows;
(1) The region within the HAZ somewhat was softened by weld thermal cycles and this region was not hardened by the aging treatment at 450°C for I hr after welding.
(2) When the heating rate was increased to that in welding thermal cycles, A_s (martensite to austenite reversion temperature) increased from 600 to 750°C.
(3) In austenite matrix, precipitates didn't nucleate but grew into rods. Resolution of the precipitates was not observed at relatively low austenitizing temperatures. With higher temperature (900°C) and prolonged holding time, precipitates became spheroidal.
(4) Specimens which had been rapid- heated above A_s point and air- cooled contained 10-20 vol% retained austenite. This retained austenite was one of the causes of softening in HAZ.

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In order to study the effect of air on the fatigue strength of steel, mild steel plate specimens with and without round or vee shape notch were tested by the laboratory--made fatigue testing machine in the atmosphere controlled tank.
The results obtained are as follows.
1) The fatigue strength of mild steel is decreased by oxygen or water vapour in atmosphere and if both exist in atmosphere, the influence becomes more remarkable. Nitrogen has little effect on it.
2) Fatigue crack initiation and endurance limit of mild steel are little affected by oxygen and water vapour in atmosphere.
3) The effect of atmosphere on the fatigue strength is not affected by the profile of specimen.
4) Oxygen or water vapour accelerates the fatigue crack propagation.

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It is previously known that the original weld interface of carbon steel pressure weld made in air or vacuum is apt to consist with the austenite grain boundaries at 1200°C and that this phenomenon is probably resulted from existence of micro-cavities on the original weld interface. It is the purpose of this study to clear the effect of microscopic causes in the weld interface, which constitute of the ununiform structure in pressure weld, on the mechanical properties of the pressure welded joints in carbon steel. The following results were obtained.
1. The microscopic causes which impede the migration of austenite grain boundaries could not to be removal from the pressure weld by annealing. The cyclic-heat treatment between 1200°C and 500°C is a successful method for elimination of the ununiform micro-structure in pressure weld.
2. It was found that the mechanical properties of the annealed joints of S1OC and S45C carbon steel pressure welds made at 1200°C for 5 min were just the same to those of parent metals, respectively. Therefore the microscopic causes on the weld interface, which retard the migration of grain boundary, does not always deteriorate on the mechanical properties of carbon steel pressure welded joints at room temperature.

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The influence of shield gas compositions in argon-nitrogen, argon-oxygen, argon-carbon dioxide and argon-hydrogen systems on the arc temperature was studied by means of the spectral line intensity method.
The findings in this work are as follows:
1) The arc temperature increases slightly with an increase of welding current.
2) The arc temperature does not undergo a noticeable change with arc voltage.
3) The arc temperature increases with an increase in nitrogen, oxygen, carbon dioxide and hydrogen additions.
The arc temperature in argon-hydrogen gas mixtures is highest, and the one in argon-carbon dioxide gas mixtures is lowest.
4) The arc temperature seems to be influenced by electrode materials.

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The cause of generation of cracking in the bead end part in one side high speed automatic arc welding of heavy plate is studied. It is well known that the base metal tends to bend as shown in Fig. 3-b due to the heat delivered to the plate. The actual rotational deformation might be restricted by some restraining devices. When these devices are constituted only by tack weld beads along the weld line, cracking may occur because the welding arc will melt off these tack weld metals one by one as the arc proceeds, and at the terminal end when the last tack weld is melted off, the constraining force also vanishes and the rotational deformation takes place abruptly in large amount which results in a cracking of the bead at the part whose temperature is about 1200-1300°C. See Fig. 8.
Fig. 8-10 and Fig. 15 show the deflection of point N on the base metal shown in Fig. 6 under the welding condition shown in Table 1. The abrupt increase of deflection observable in Fig. 8 causes cracking shown in Fig. 2 and Fig. 13, and it is noticeable that in case of Fig. 10 no cracking is observed.
The point R in these figures corresponds to the time when the arc proceeds into the tab plate passing over the border line of the base metal. The increase in displacement of the measuring point N in RS part in Figs. 9, 10 is caused by the rotational moment of the tab plate side as illustrated in Fig. 12-b.
In Fig. 15, the tab is separated into two parts in order to eliminate the above mentioned tab actions.
Fig. 21 shows the deflection of points N₁-N₅ along the bead, Z being the tack weld. Figs. 21, 23 are the results of explanational experiment to illustrate the action of the tack weld at root.

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