The invention of electric arc welding revolutionized the steel construction industry, but also brought some problems when the welded region has inferior properties compared to the plate metal. A major cause of brittle failure was identi ed as hydrogen embrittlement of the weld zone, although a comprehensive understanding of this phenomenon is not, even now, available. Hydrogen in solution in the weld zone is found in arc welds, due to the intense conditions in the welding arc. There is invariably a sufficient source in the form of moisture and hydrocarbon residue to give a few parts-per-million (ppm) by mass of hydrogen in the weld pool, which is a sufficient concentration to bring the possibility of hydrogen cold cracking in the completed weld. Hydrogen is significantly mobile in steels at room temperature, which is certainly why a few ppm of hydrogen can concentrate on a microscopic scale and initiate cracks, but also means that on a macroscopic scale there is hydrogen dispersion, which can relieve the cracking risk or place hydrogen in hydrogen cracking susceptible regions. The understanding of solubility and mobility of hydrogen in steels of different compositions and microstructures is therefore paramount. The question investigated in this work is whether the characteristics of the weld hydrogen cracking tendency can be explained by the features of weld hydrogen transport, especially when steel selection is a variable. Plate steel ranging from a 0.22%C pearlitic steel to a 0.05%C thermo-mechanically controlled-rolled and accelerated-cooled (TMCR-AC) high strength low alloy (HSLA) steel with no pearlite, plus a 0.4%C non-plate steel, were included in the experimental program. Welds were made with rutile ux-cored-wire (R-FCW) at two hydrogen levels, together with rutile shielded-metal-arc (R-SMA) welds. In order to investigate the di usion rates, a novel experiment has been devised. The welded plate has been milled away at an angle from the underside of the weld to provide increasing distances between the fusion boundary and the plate under-surface. The formation of hydrogen bubbles in glycerol enabled the measurements of the time dependent diffusion distances. The results clearly show a square root time correlation, as expected from the Fickian mechanism and enabled the calculation of diffusion coefficients for different steels. A nearly four fold difference was found between the steels, with the fastest hydrogen movement in the TMCR-AC steel. To reveal the initial distribution of hydrogen some samples were frozen immediately after welding and machined under liquid nitrogen. This test ruled-out any signi cant hydrogen dispersion during the deposition of the weld and during the cooling down period. The experimental data were interpreted using a new numerical computer model, based on random jumps of hydrogen between equivalent lattice sites. It is shown that this numerical model gives identical results to the analytical Fickian approach, but has the advantage that it can be used for any boundary shape. When this model has been applied to the experimental data, some unexpected features have been found. The amount of hydrogen emerging at surfaces distant to the weld was higher than expected from a concentration-driven mechanism; suggesting that a di erent transport mechanism should be applied. The numerical model has also indicated a discontinuity in the hydrogen concentration at the fusion boundary. It is a consequence of the model that hydrogen solubilities and di usivities are inversely related properties of the metal; a feature which is supported by experimental evidence. The tendency of hydrogen cracking to appear in the weld metal rather than in the heat-a ected-zone (HAZ) can thus be explained by higher di usivity of hydrogen in the plate metal. It appears that there is a relationship between the diffusivity and the microstructure, particularly when the content and form of carbon is considered.

Abstract This thesis studies controlling factors that govern transverse hydrogen cracking in high-strength multipass weld metal (WM). The experiments were concerned with heavy-restraint Y- and U-Groove multipass cracking tests of shielded-metal arc (SMAW) and submerged-arc (SAW) weld metals. Results of tensile tests, hardness surveys, weld residual stress measurements and microstructural investigations are discussed. The analytical phase comprised numerical calculations for analysing the interactions between crack-controlling factors. The objectives were: (i) the assessment of WM hydrogen cracking risk by defining the Crack-No Crack boundary conditions in terms of 'safe line' description giving the desired lower-bound estimates, and (ii) to derive predictive equations capable of giving reliable estimates of the required preheat/interpass temperature T0/Ti for the avoidance of cracking. Hydrogen cracking occurred predominantly in high strength weld metals of Rp0.2 ≈ 580-900 MPa. At intermediate strengths of Rp0.2 ≈ 500-550 MPa, cracking took place in the cases where the holding time from welding to NDT inspection was prolonged to 7 days. Low strength WMs of Rp0.2 ≤ 480 MPa did not exhibit cracking under any conditions examined. Cracking occurrence was, above all, governed by WM tensile strength, weld diffusible hydrogen and weld residual stresses amounting to the yield strength. The appearance of cracking vanished when transferring from 40 to 6 mm thick welds. The implications of the holding time were more significant than anticipated previously. A period of 16 hrs in accordance with SFS-EN 1011 appeared much too short for thick multipass welds. Interpass time and heat input showed no measurable effect on cracking sensitivity, hence being of secondary importance. Equations were derived to assess the weld critical hydrogen content Hcr corresponding to the Crack-No Crack conditions as a function of either weld metal Pcm, yield strength Rp0.2 or weld metal maximum hardness HV5(max). For the calculation of safe T0/Ti estimates, a formula incorporating: (i) WM strength as a linear function of either weld carbon equivalent CET or weld HV5(max), (ii) weld build-up thickness aw in the form of tanh expression and (iii) weld diffusible hydrogen HD in terms of a combined [ln / power law] expression was found descriptive.

The repair welding of main steam pipelines, which involves the welding of new material onto service-exposed material, are investigated. This paper investigates the literature and experimental work surrounding this subject. The introduction provides a background to the applicable welding technology. In section two the heat-affected zone is discussed with emphasis on the residual stresses that develop in this zone. The mechanical properties of the heat-affected zone are also investigated. This includes the tensile, toughness and hardness properties as well as inspecting the relevant microstructures. The effect of post weld heat treatment on these properties is also investigated. Section three investigates the phenomenon of creep. Not only is this important due to the high temperatures at which these pipelines operate, but creep is also associated with some failures of these weld during post weld heat treatment. The creep properties of the heat-affected zone are investigated in detail with the use of weld simulation. Sections four and five detail reasons for weld failure after welding due to hydrogen and reheat cracking. Hydrogen cracking is investigated with the use of slow strain rate tensile tests during cathodical charging the specimen with hydrogen. The phenomenon of reheat cracking is investigated with the use of high temperature tensile tests as well as a novel approach in which the stress relief of a welded joint is simulated while measuring crack growth and stress relieved. / Dissertation (MEng)--University of Pretoria, 2014. / gm2014 / Materials Science and Metallurgical Engineering / Unrestricted

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