resistance to solidification cracking;
preventing the formation of a structure containing very hard brittle compounds;
preventing the formation of the sigma phase and, correspondingly, embrittlement in tempering;
the mechanical properties and fatigue strength satisfy the requirements of PNAE G-7-002–86.
Branch pipe welding without digging a paved road was enabled by the friction welding method.
The developed method does not require digging and repairing the paved road. Therefore, construction cost can be reduced, and traffic is not obstructed.
Tensile strength and elongation of branch pipe friction welded joint were 17.7–17.8 MPa and 520–525%, respectively. These were nearly equal to those of base polyethylene.
A large amount of mud was eliminated from the faying surface by friction. Therefore, the influence of mud on tensile strength and elongation was not recognized.
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During He‐TIG arc welding, the centre of the weldpool is characterised by a brighter, red‐tinged area being visible at its boundary. From the behaviour of the slag floating at this boundary, this area is named the inner zone, being the radial flow zone.
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The electromagnetic force distribution is determined on the assumption of a uniform current density on the anode area and on a cylindrical surface (thin plate) or hemispherical surface (thick plate) having the same radius R as the former. The absolute value of the electromagnetic force is maximum at the anode boundary and shows a singularity with an abrupt 90° direction change.
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On the assumption of Sozou and Pickering that the current flows from an infinitely remote location in another direction across the boundary path, the current is concentrated near the boundary, and the singularity becomes more pronounced. In the disc‐like and ring‐like inner zone, the current density appears to be uniform or concentrated.
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If the inner zone coincides with the anode area:
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At the centre, the centripetal electromagnetic force of the anode area is zero, and there is no obstruction of the surface flow.
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Nearer the boundary, the centripetal electromagnetic force intensifies, the radial flow is retarded, and the pressure is raised.
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The downward electromagnetic force concentrated at the anode boundary becomes the downward driving force.
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The liquid in the low‐temperature zone outside the boundary is adsorbed, with an inward directed flow being generated by desorption.
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Factors 1–4 jointly drive an internal‐radial and external‐axial type of double flow.
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The penetration change can be explained by the inner zone size.
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The paper also includes a discussion of the surface flow driving force and arc current concentration on the basis of the foregoing results.
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The impact properties of the high heat input weld metal produced under standard conditions (thickness of 40 mm) are generally such that ESW has a lower absorbed energy value (vE value) than SAW.
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The impact value of ESW high heat input weld metal is non‐uniform, and distinctive impact properties are found. That is to say, the vE value of the weld metal core is lower than that of the weld metal rim.
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ESW weld metal macrostructures have a non‐uniform morphology in both the core and rim. That is to say, a fine‐grained columnar zone is generated in the core and a coarse‐grained columnar zone in the rim.
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The results presented above in (1) and (2) suggest that the fine‐grained columnar zone in the ESW weld metal core has a low absorbed energy value (vE value) and that the coarse‐grained columnar zone in the rim conversely has a high one. This conflicts with what is conventionally stated about effects of grain size in otherwise identical microstructures, i.e. that the vE value decreases with an increasing grain size.
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The vE value of ESW weld metal tends to decrease in relation to the welding heat input Q. That is to say, it tends to have a low value at Q > 30.0 kJ/mm (up to around 80.0 kJ/mm).
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The analyses of the gas composition and five principal elements of the ESW weld metal at different Q values suggest that there is little change in relation to any heat input change. This suggests that the decrease in the vE value in relation to the welding heat input Q is not due to a change in the weld metal composition.
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The quality (joint interfacial strength and properties) of steel gas pressure welds is controlled by the area expansion factor of the gas pressure welds and gas pressure welding temperature. Within the range avoiding any melting of the surfaces being welded, the quality improves with increasing values of both factors.
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The quality of gas pressure welds can be evaluated by whether or not flat fracture occurs on the fracture surface under mechanical test and its proportion.
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The effects of the area expansion factor and gas pressure welding temperature on improvement of the quality of gas pressure welds vary depending on the carbon content of the base metal, their effects being more pronounced with a greater carbon content. This is due to reduction of the interfacial oxides by the carbon contained in the base metal.
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The electrode lives of HDG materials were changed at approximately 0.3–0.4 mass% Al content in Zn coating. Materials with low‐Al coating content showed over three times longer electrode lives than materials with high‐Al coating content.
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Although the thickness of Al oxide layers was in proportion to the Al content in Zn coatings, the obvious correlation between electrode life and thickness of Al oxide layers was not observed.
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In the case of low‐Al coating content, it was observed that Fe‐Zn alloy grew from the steel‐coating interface to the Zn coating. It was considered that, in the initial stage of welding, the content of Fe in Zn coating increased immediately.
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In the case of high‐Al coating content, Fe‐Al alloy was observed at coating‐steel interfaces instead of Fe‐Zn alloy. It was known that Fe‐Al alloy suppresses the Fe‐Zn alloying reaction. Zn coating was not alloyed with Fe in initial stages of welding.
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From these results, it was concluded that aluminium in coatings affected electrode life by changing the melting point of coating layers between the electrode and the steel. The melting point of low‐Al content coating layers rose because of the diffusion of Fe into the Zn coating. This phenomenon decreased electrode wear and electrode life was long. In contrast, the melting point of high‐Al content coating layers remained low. This phenomenon caused electrode alloying easily and also increased electrode wear. As a result, electrode life became shorter.
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Alloy 1 (Co-Cr-W-C) Alloy 4 (Fe-Cr-Mo-V)
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Alloy 2 (Ni-Cr-B-Si) Alloy 5 (Ni-Cr-C-Mo)
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Alloy 3 (Ni-Cr-C-W) Alloy 6 (Fe-Cr-Mo-Mn)
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The vE value of the weld metal core is lower than that of the weld metal rim. The absorbed energy transition temperature is also higher.
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The vE values of the weld metal core and rim altered with a change in the welding heat input Q (varied at six levels in the 10.1 ～ 126.7 kJ/mm range), generally decreasing with an increasing heat input. The vE value of the core is lower than that of the rim. In the weld metal produced at the maximum heat input (126.7 kJ/mm), however, the core and rim have much the same vE values.
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The vE value of the weld metal core shows little change with a change in the steel type (one type of TMCP and two types of SM490A), and remains at a low value. When there is a change in the surrounding gas (oxygen (O2), air, and argon gas), the vE value decreases in an Ar, air, O2 sequence.
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The weld metal core vE value is little dependent on any change in the flux type (three types of fused flux) and weight of flux used (0.5 or 0.7 N).
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The microstructural observations suggest that grain‐boundary ferrite (GBF) occurs at the coarse prior austenite grain boundaries of the rim, with fine acicular ferrite (AF) being mainly found trans‐granularly. In the core, grain‐boundary ferrite is formed at high density at the fine prior austenite grain boundaries, with massive polygonal ferrite (PF) being formed at high density transgranularly. There are thus distinct differences in the coarse ferrite morphology of the rim and core.
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The SEM observations of the fracture surfaces suggest that the core weld metal has a large grain size, a feature corresponding to its low vE value. The observations made at the fracture surface periphery indicate that numerous secondary cracks are initiated in the grain‐boundary ferrite of the core, suggesting that the ferrite has a low toughness.
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The results of the SEM simultaneous fractographic‐microstructural observations suggest that fracture selectively propagates along the grain‐boundary ferrite. This indicates that the low vE value of the core is due to the presence of high‐density grain‐boundary ferrite and massive transgranular polygonal ferrite.
Density, hardness and porosity data
Pore formation, evidence, film density values
Variation with processing parameters
Pore morphology, substrate effects
Depth of pore influencing surface layers
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The fatigue limit of smooth base metal which received strong peening cleaning at about 320 MPa was remarkably high in comparison with smooth base metal at about 245 MPa.
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The fatigue limit of a welded joint which received strong peening cleaning at about 300 MPa was remarkably high in comparison with a welded joint at about 170 MPa.
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The strong peening cleaning was highly efficient and the cleaning state was satisfactory.
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The cause of the remarkable rise of the fatigue limit (300 MPa) of the welded joint which received strong peening cleaning was because the fatigue limit (about 170 MPa, 57%) of the welded joint was improved (about 130 MPa, 43%) with peening cleaning. It was considered that improvement effects were: a rise (about 68 MPa, 23%) of hardness of the weld toe; relief (about 43 MPa, 14%) of stress concentration; increase (about 136 MPa, 45%) of compressive residual stress; and the decrease (about ? 96 MPa, ? 32%) by increase of surface roughness.
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The fatigue strength of the hot galvanized welded joint decreased remarkably. This was thought to be due to the decrease (about HV40) of hardness of the surface, the decrease (about 188 MPa) of the compressive residual stress and the influence of many factors which accompanied hot galvanizing.
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The results obtained during measurement of the stress relaxation behaviour in the rising‐temperature constant‐strain rate tests suggest that the bainite structures of both steels clearly show more stagnation or delay in their stress relaxation behaviour than the other HAZ structures at a PWHT temperature above 600 K. This implies that the matrix is resistant to softening. The non‐AcC type steel also exhibits more stagnation in the higher temperature range under the effect of alloy carbide precipitation at the grain boundaries than the AcC type steel.
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The results obtained in the rising‐temperature constant‐load tests run to determine the high‐temperature strength and inherent deformability of the HAZ structure suggest that the bainite structures of both steels tend to lose more ductility than the other HAZ structures, having a reduction of area of 35% at a fracture temperature of 850–900 K. The non‐AcC type steel also exhibits a greater loss of ductility in all HAZ structures than the AcC type steel.
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The results of the TEM observations made to determine the causes of this ductility loss suggest that a difference in the carbide precipitation behaviour near the grain boundaries in each HAZ structure in the PWHT process affects the plastic deformability of the matrix, and that the trend of reducing plasticity differs in each HAZ structure. These trends are more pronounced in the non‐AcC type steel containing alloying elements with a strong carbide‐producing tendency, such as e.g. Nb, Ti, etc.
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All HAZ structures of the AcC type steel show a trend of reducing ductility at a fracture temperature of 850–900 K. This feature is not found in conventional carbon steels with an identical composition and may well be due to the fact that this temperature range corresponds to the ductile‐brittle transition range. It is necessary to resort to a method of fabrication able to reduce the hardened structures as far as possible during welding, i.e. to ductility reducing counter‐measures in the PWHT process, such as e.g. welding heat input control, preheating, etc.
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To evaluate the ductility and brittleness of steels, it is important to gain a good understanding of their plastic deformability, and the paper proposes a method for evaluation of the ductility of the TMCP steels on the basis of the relationship between the amount of displacement produced in the rising‐temperature constant‐strain rate tests and the plastic deformability of each HAZ structure in the PWHT process as obtained in the rising‐temperature constant‐load tests. This method enables the risk of cracking and degree of embrittlement to be identified and proves effective in practical applications.