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The Design of Submerged Arc Weld Deposits for High-strength Steels
1.8 The Design of Submerged Arc Weld Deposits for High-strength Steels L-E. Svensson* and H. K. D. H. Bhadeshia** *ESAB AB, Gothenburg, Sweden "Cambridge University, UK
ABSTRACT The physical metallurgy of high-strength steel welds is studied both theoretically and experimentally with a view to establishing a basis for the further development and optimisation of arc welding consumables. The primary microstructure of such welds is unique, consisting of a mixture of acicular ferrite, bainite and low-carbon martensite. Since this is very resistant to tempering, and since the alloys necessarily have a very low Ae3 temperature, the effect of reheating is found to be minimal. Indeed, the microstructure of multirun welds is found to be fairly homogeneous and not very different from the primary microstructure. These and other factors combine to give very good toughness levels in spite of the high yield strength of the weld deposits. KEYWORDS High-strength steel welds; acicular ferrite; low-carbon martensite; microstructure; microstructure modelling; physical metallurgy of welds, ; bainite; thermodynamics. INTRODUCTION The level of alloying additions in most weld deposits is usually kept low in order to prevent cold cracking and other defects associated with the formation of brittle phases within the weld itself. The hardenability of the weld has to be low enough to avoid the transformation of residual austenite to relatively high carbon, untempered martensite (α'), as the weld cools towards ambient temperature. The problem can be made worse by the partitioning of carbon into the austenite as a consequence of other transformations which precede the martensite transformation; the carbon-enriched austenite then transforms into even harder martensite. At the same time, the austenite hardenability has to be sufficiently high to promote the formation of desirable phases such as acicular ferrite, at the expense of weaker phases such as allotriomorphic ferrite and Widmanstàtten ferrite. In a multirun weld deposit, the alloy composition may also have a strong effect on the extent of microstructural change in the material already deposited, due to the additional heat input as further layers are deposited. For these reasons, the yield strength (σ γ ) of conventional welds has in general been kept rather low (~ 350-550 MPa) with occasional excursions to higher strength experimental welds but usually at the expense of toughness. The use of low-carbon high-strength steels based on a microstructure of bainite or a mixed microstructure of bainite and martensite has recently assumed prominence in the fabrication of pipelines (Nakasngi, Matsuda and Tamehiro, 1983) and of ships and submarines (Deb, Challenger and Themen,
71
1987). To meet simultaneously the requirements of strength and toughness, a class of rather heavily alloyed, complex steel welding consumables have been developed, a typical chemical composition being Fe-2.2Ni-1.3Mn-0.5Mo-0.5Cr-0.1C wt.%. In spite of the high alloy content, these consumables give tough weld deposits which are extremely strong (σ ν « 730 MPa) and yet are resistant to hydrogen cracking and even are insensitive to certain specification requirement tests involving explosive deformation. The physical metallurgy of these alloy welds is not well understood either theoretically or experimentally. For example, metallography using light microscopy has proven impossible to interpret in terms of the numerous microstructure classification schemes that exist in the welding industry (Deb, Challenger and Themen, 1987). Even experiments using transmission electron microscopy have failed to characterise the microstructure with the sort of clarity needed for systematic experimentation or calculation. The purpose of this work was to attempt an understanding of the physical metallurgy of the highstrength steel welding alloys in order to provide a basis for systematic research and development towards the optimisation of such welds. The work also includes an experimental assessment of the microstructure in order to clarify the nature of the transformation products present in the various regions of multirun welds. EXPERIMENTAL The weld was fabricated using a submerged arc technique, as an X joint using 60 mm thick plate with 36 runs on each side. The welding current was 625 Amps at 30V, the weld being deposited in the flat position at a speed of « 40m/h (0.0111 m/s) with the interpass temperature (TQ) being kept within the range 150-200DC. The welding consumables used have the proprietary ESAB designations OK 10.62 (flux) and OK Autrod 13.43 (4 mm diameter wire); the base plate and deposit compositions are given in Table 1. TABLE 1 Chemical Compositions (wt.%) Weld Plate
C 0.071 0.110
Si 0.26 0.27
Weld Plate
Al 0.024 0.076
Cu 0.13 0.17
Mn 1.23 0.86 Sn 0.011 0.007
Ni 2.20 1.09 Pb 0.003 0.003
Cr 0.55 0.51
Mo 0.44 0.40 As 0.010 0.005
Sb 0.001 0.001
P 0.012 0.015 B 0.0004 0.0013
S 0.007 0.005 N 0.0083 0.0041
Nb 0.005 0.017
Ti 0.004 0.004
O 0.0371 0.0010
Tensile tests were carried out on standard cylindrical specimens of 12.8mm diameter with a gauge length of 65mm. The specimens were degassed at 100°C for 48h prior to testing. Charpy tests were carried out on standard 10mm square V-notch specimens at -18 and -51 °C. Transmission electron microscopy samples were prepared from 3mm diameter discs machined from known positions in the weld; by defining the extraction positions using light microscopy, it was possible to obtain discs containing just the primary microstructure or those just containing the reheated microstructure. The discs were mechanically ground down to a thickness of 0.08mm on 1200 grit SiC paper before electropolishing. The specimens were electropolished using a Tenupol jet polisher at a voltage of 55V in an electrolyte of 5% perchloric acid, 25% glycerol in ethanol, at ambient temperature. It was found that by reducing the polishing voltage (« 50V), the acicular ferrite could be attacked preferentially with respect to the martensite matrix, so that it could be distinguished clearly from the matrix phase. RESULTS AND DISCUSSION Mechanical Properties The results from mechanical tests are given in Table 2. The strain hardening coefficient 'n' and strengthening coefficient 'Κ' are crude estimates: if the [true stress o]/[true strain ε] curve is of the form σ = Κε11, and if o u is the tensile strength, then it can be shown that (o/a u ) = η" η ε η . Hence, values of
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K and n can be estimated using the proof stress data (σοο02, σ001 measured at strains of 0.002 and 0.01 respectively), together with the measured tensile strength, using the criterion that a plot of (o/cu) versus (η"ηεη) should have a slope of unity. The results give an idea of the good toughness even though the yield strength is so high. The Vickers hardness of the top bead, which contains the microstructure unaffected by other beads, was found to be 295 whereas the mean hardness of the entire weld was found to be 281, indicating that the weld is extremely homogeneous in strength in spite of its multirun character. As will be discussed later, this is because of the high alloy content and temper resistance of its microstructure. TABLE 2 Mechanical Property Data
οοι' ΜΡ 3 Gu, MPa Reduction of Area, % Elongation, % K, MPa n
Test 1 761 786 854 62 21 980 0.044
Test 2 765 779 826 62 22 930 0.034
Charpy tests Energy absorbed, J
-51°C 75, 69, 71
-18°C 82, 79, 82
Tensile Properties
^
Brief Review of Transformation Products It is necessary to consider briefly the nature of the transformation products in steel welds, in order to set the scene for the discussion of theoretical and experimental results. Allotriomorphic ferrite (a) is the first phase to form during the cooling of austenite in low-alloy steels. It nucleates at the austenite grain boundaries and grows by a diffusional transformation mechanism which involves at the very least, reconstructive diffusion, which is the diffusion necessary to achieve the lattice change with a minimum of strain. The growth of a is anisotropie, the rate being highest along the γ boundaries which soon become decorated with layers of a. Pearlite or degenerate pearlite is the other product which grows from austenite by diffusional transformation. Detailed discussions of transformation mechanisms can be found in some recent reviews (Bhadeshia, 1985; Grong and Matlock, 1986; Abson and Pargeter, 1986; Bhadeshia, 1987). As the transformation temperature decreases, diffusion becomes sluggish and displacive transformations are kinetically (though not thermodynamically) favoured; these transformations do not involve any reconstructive diffusion so that substitutional or iron atoms do not diffuse during reaction. Consequently, the transformations are accompanied by macroscopic displacements reflecting the coordinated movement of atoms during the lattice change, the displacements having the characteristics of invariant-plane strains. Their morphology is thus dominated by the need to minimise strain energy. Plates of Widmanstätten ferrite (ο^) grow at relatively low undercoolings by a displacive paraequilibrium mechanism which involves the simultaneous growth of pairs of mutually accommodating plates of ferrite, so that the strain energy of transformation is reduced (Bhadeshia, 1981). Widmanstätten ferrite and bainite ( α ^ both nucleate by the same mechanism and are both represented, at least initially, by the same ' C curve on the TTT diagram. Bainite, which grows in the form of sheaves of small platelets, can only form at a higher undercooling because it deviates further from equilibrium, forming by diffusionless transformation from austenite (the carbon redistributing into γ subsequent to transformation), the strain energy term also being higher since the plates do not form in an accommodating manner (Bhadeshia, 1987). Lenticular plates of acicular ferrite (a a ) form by the same mechanism as bainite, but have a different morphology because nucleation occurs intragranularly at inclusions (rather than at γ grain boundaries) and the development of sheaves is stifled by hard impingement between plates nucleated at neighbouring sites
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(Yang and Bhadeshia, 1987; Strangwood and Bhadeshia, 1987). Theory Transformations from austenite can therefore be categorised into diffusional and displacive reactions, each category beingrepresentedby a 'C curve on the time-temperature-transformation diagram for the steel. The continuous lines on Fig. la show the time-temperature-transformation (TTT) diagram calculated (Bhadeshia, 1982) assuming that the weld is chemically homogeneous. We have noted earlier that the TTT diagram should consist of just two 'C curvesrepresentingthe initiation of diffusional and of displacive transformation. However, for the present alloy, a complication arises for the displacive 'C curve, since nucleation becomes impossible at an intermediate temperature, so that the curve splits into two halves, the higher temperature one for Widmanstätten ferrite and the other one for bainite and acicular ferrite. This is illustrated in Fig. lb where the F n curve represents the free energy needed to obtain a detectable nucleation rate of o^, a b or a a for any steel. The AFty-r^+a} curve is however specific to the present alloy, and represents the free energy available at any temperature; when AF{Y-^Y1+a}
Fig. 1. (a) Calculated TTT diagrams; (b) free energy curves. Weld deposits are not in fact chemically homogeneous since nonequilibrium solidification leads to coring. Regions lean in austenite stabilising alloying elements will have a lower hardenability and will influence the transformation behaviour of the entire weld. The approximate procedure for taking this into account has been discussed elsewhere (Gretoft, Bhadeshia and Svensson, 1986), but the TTT curve for the solute-depleted regions of the weld (calculated composition Fe-0.33C-0.37Si-0.93Mn-0.95Ni-0.13Mo0.49Cr as compared with Fe-0.33C-0.51Si-l.25Mn-2.09Ni-0.26Mo-0.59Cr at.% for the homogeneous alloy) is given in Fig. la as dashed lines. Because of the higher free energy change accompanying the γ-»α transformation in the solute-depleted regions, the two parts of the displacive *C* curve unite, although because of their different growth mechanism and stored energy, the bainite and acicular ferrite transformations can still only occur below the indicated temperature of 586°C. TTT curves can be used to estimate reactions occurring during continuous cooling, using the Scheil
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method, in which the cooling curve is divided into a series of isothermal steps. If vx is the incubation time as given by the TTT curve for the ith step, then transformation is considered initiated when S(tj/Tj)=l, where tj is the time for the ith isothermal heat treatment. Values of IXtj/Xj) obtained by calculating the cooling conditions of the weld (for the welding parameters presented earlier) and using the diffusional ' C curve for the solute depleted regions were always found to be less than unity, so that the the microstructure was predicted to consist of just acicular ferrite, bainite and low-carbon martensite. Widmanstätten ferrite was also ruled out because Z(tj/ij) for the displacive curve had a value of just 0.14 by the time the bainite-start temperature for the solute-depleted regions was reached. In fact, Ι^/τ^ became unity at = 480°C, only about 10°C above the martensite-start temperature of these regions, so that the volume fraction of acicular ferrite and bainite should be rather low, most of the microstructure being low-carbon martensite of a carbon concentration not very different from the starting level. Estimation of the volume fractions awaits the development of theory for predicting the nucleation rate of acicular ferrite. Theoretical Assessment of Alloy and Welding Effects A model for the prediction of microstructure in low-alloy steel weld deposits (Bhadeshia, Svensson and Gretoft, 1985, 1987) was used to assess the effect of alloy chemistry, welding current and speed on the microstructure of the primary regions of the weld. Calculations were done for 98 different cases; the results are presented in Tables 3 (changes in composition with respect to weld 1 have been emphasised using bold typeface) and Table 4. The latter includes only the cases where phases other than acicular ferrite and martensite were predicted, although the calculations were repeated for all the weld compositions listed in Table 3. The experimental weld is numbered 14. TABLE 3 The Influence of Composition Variations
c
Si
Mn
Ni
Mo
Cr
V
V
0.070 0.040 0.100 0.070 0.070 0.070 0.070 0.070 0.070 0.070 0.070 0.070 0.070 0.071
0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.26
1.25 1.25 1.25 1.75 0.75 1.25 1.25 1.25 1.25 1.25 1.25 1.25 1.25 1.23
2.20 2.20 2.20 2.20 2.20 3.00 1.40 2.20 2.20 2.20 2.20 2.20 2.20 2.20
0.45 0.45 0.45 0.45 0.45 0.45 0.45 0.65 0.25 0.45 0.45 0.45 0.45 0.44
0.55 0.55 0.55 0.55 0.55 0.55 0.55 0.55 0.55 0.80 0.30 0.55 0.55 0.55
0.0 0.0 0.0 0.0 0.0 0.0 0.0 0.0 0.0 0.0 0.0 0.1 0.2 0.0
0.00 0.00 0.00 0.00 0.17 0.00 0.13 0.00 0.00 0.00 0.00 0.00 0.00 0.00
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a
v
w
0.00 0.00 0.00 0.00 0.06 0.00 0.08 0.00 0.00 0.00 0.00 0.00 0.00 0.00
1 2 3 4 5 6 7 8 9 10 11 12 13 14
TABLE 4 The Influence of Other Variables Comment
v
Welding speed reduced (0.0083 m/s) Welding speed reduced (0.0083 m/s) Welding speed reduced (0.0083 m/s) Welding speed reduced (0.0083 m/s) Welding speed increased (0.0139 m/s) Interpass temperature reduced (150°C) Interpass temperature (250°C) Interpass temperature (250°C) Welding current reduced (525 A) Welding current increased (725 A) Welding current increased (725 A)
0.20 0.20 0.13 0.13 0.14 0.15 0.19 0.18 0.15 0.19 0.17
a
V
w
0.06 0.07 0.04 0.04 0.06 0.06 0.06 0.07 0.06 0.06 0.07
No. 5A (low Mn) 7A (low Ni) 9A (low Mo) 11A (lowCr) 5B (low Mn) 5C (low Mn) 5D (low Mn) 7B (low Ni) 5E (low Mn) 5F (low Mn) 7C (low Ni)
The calculations show clearly the role of Mn and to a lesser extent Ni, in preventing the formation of detrimental phases such as allotriomorphic ferrite and Widmanstàtten ferrite. A reduction in welding speed or an increase in the interpass temperature or an increase in welding current all cause a decrease in weld cooling rate and consequently can in some alloys enhance the formation of a and o^. It is evident the experimental weld (no. 14) is very tolerant to quite large variations in welding conditions in the sense that the primary microstructure remains a mixture of acicular ferrite and martensite for all the conditions used. There are of course other welds which behave similarly (nos. 1-4, 8, 10, 12, 13) and these compositions need to be investigated experimentally, since the mechanical properties could then be varied without changing the mixed acicular ferrite and martensite microstructure. There is as yet no model allowing the calculation of the strength of multirun submerged arc welds, but we use here a model by Svensson and co workers (1988), which is based on multirun manual metal arc welds, in order to obtain some crude estimates of yield strength as a function of composition. Note that the model takes account of solid solution strengthening and microstructural strengthening due to individual phases, both for the primary and reheated regions of the weld, the volume fraction of these regions depending on the Ae3' temperature (i.e., a point on the paraequilibrium γ/γ+α phase boundary for the alloy concerned). The results are presented in Table 5 (V p is the calculated volume fraction of the primary region) and with the earlier data of Table 3,4, can be used in order to design weld deposits of the desired strength and microstructure. Note that the calculated yield strength of weld 14 is about 40MPa lower than the measured strength. TABLE 5 Estimates of Mechanical Properties No. 1 2 3 4 5 6 7 8 9 10 11 12 13 14
Ae3\ °C 699 699 691 672 723 667 725 699 695 691 699 694 693 700
σ , MPa 728 729 739 783 642 795 644 732 730 740 727 735 737 723
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V
P
0.72 0.73 0.76 0.83 0.64 0.84 0.63 0.73 0.74 0.76 0.73 0.74 0.75 0.72
Transmission Electron Microscopy Examination of weld 14 revealed a microstructure consistent with the calculations and mechanical property data presented above. Its primary regions were found to consist of an estimated 40% of intragranularly nucleated, lenticular acicular ferrite plates, a small amount of γ grain boundary nucleated bainite, the remainder being low-carbon dislocated martensite, with thin films of retained austenite between the martensite plates (Fig. 2). The reheated regions of the weld were not in fact very different, since most of them had been fully reaustenitised due to the low Ae3' temperature of the alloy; due to space restrictions, the detailed microstructural study will be presented elsewhere. Both the primary and reaustenitised microstructure seem resistant to tempering, probably because of the very low-carbon content of the alloy. All this explains the relatively small observed variation in hardness; such a mechanically homogeneous microstructure is believed to be beneficial to toughness (Tweed and Knott, 1987). Since most of the microstructure is martensite, and since acicular ferrite is also a strong phase (Sugden and Bhadeshia, 1988; Svensson and coworkers, 1988), the strain hardening coefficient is expected to be low. The data of Table 2 are about the same as reported values of strain hardening coefficients for lowcarbon martensite in steels (Davies, 1978).
Fig. 2.
The primary microstructure; (a) a a in a martensite matrix; (b) α^ growing from a γ grain boundary.
CONCLUSIONS A primary weld microstructure consisting mainly of low-carbon martensite as the major phase, and acicular ferrite has a good combination of strength and toughness. However, to ensure high-strength in multirun weld deposits, it is also necessary to alloy in a manner which stabilises the austenite (reduces the Ae3 temperature), so that any additional heating due to the deposition of subsequent runs causes much of the reheated region to be reaustenitised. The hardenability of the alloy has to be high enough to permit these reaustenitised regions to transform to just acicular ferrite, low-carbon bainite and low-carbon martensite during cooling to ambient temperature. The tempering resistance of the alloy also has to be high so that those reheated regions which are not substantially reaustenitised do not lose their strength. These conditions all ensure that the multirun weld, is mechanically homogeneous with little variation in strength, an important requirement for the achievement of optimum toughness.
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ACKNOWLEDGMENTS The authors are grateful to ESAB AB (Sweden) for the provision of laboratory facilities and financial support for this research and to Professor D. Hull for the provision of laboratory facilities at the University of Cambridge.
REFERENCES Abson, D. J. (1986). Review on microstructure and toughness of arc welds. International Metals Rev., 3Λ, 141-194. Bhadeshia, H. K. D. H. (1981). Rationalisation of shear transformations in steels. Acta Metall., 29, 1117-1130. Bhadeshia, H. K. D. H. (1982). Thermodynamic analysis of TTT curves. Metal Science, 16, 159-165. Bhadeshia, H. K. D. H. (1985). Diffusional formation of ferrite in iron and its alloys. Prog, in Mat. Sci., 29, 321-386. Bhadeshia, H. K. D. H. (1987). Bainite in steels. Phase Transformations 87, Institute of Metals, London. Bhadeshia, H. K. D. H. and Edmonds, D. V. (1980). The mechanism of bainite formation in steels. Acta Metall., 28, 1265-1273. Bhadeshia, H. K. D. H., Svensson, L.-E. and Gretoft, B. (1985). Model for the development of microstructure in steel welds. Acta Metall., 33, 1271-1283. Bhadeshia, H. K. D. H., Svensson, L.-E. and Gretoft, B. (1987). Prediction of the microstructure of the fusion zone of multicomponent steel welds. Advances in Welding Science and Technology, ASM, Ohio, pp. 225-229. Davies, R. G. (1978). Mechanical properties of zero-carbon ferrite and martensite structures. Metall. Trans. A, 9 451-455. Deb, P., Challenger, K. D. and Themen, A. E. (1987). Structure-property correlation of SA and GMA weldments in HY100 steel. Metall. Trans. A, 18A. 987-999. Gretoft, B., Bhadeshia, H. K. D. H. and Svensson, L.-E. (1986). Development of microstructure in the fusion zone of steel weld deposits. Acta Stereologica, 5, 365-371. Grong, O. and Matlock, D. K. (1986). Microstructural development in mild and low-alloy steel welds. International Metals Rev., 31., 27-48. Nakasugi, N., Matsuda, H. and Tamehiro, H. (1983). Ultra-low carbon bainitic steel for line pipes. Steels for line pipe and pipeline fittings, Metals Society, London, pp. 90-95. Strangwood, M. and Bhadeshia, H. K. D. H. (1987). Mechanism of the acicular ferrite transformation. Advances in Welding Science and Technology, ASM, Ohio, pp. 209-213. Sugden, A. A. B. and Bhadeshia, H. K. D. H. (1988). Strength of steel welds, submitted to Metall. Trans. A. Svensson, L.-E., Gretoft, B., Sugden, A. A. B. and Bhadeshia, H. K. D. H. (1988). Computer-aided design of electrodes for arc welding, II. Computer Technology in Welding, Welding Institute, U.K. Tweed, J. and Knott, J. (1987). Micromechanisms of failure in C-Mn weld metals. Acta Metall., 35. 1401-1414. Yang, J. R. and Bhadeshia, H. K. D. H. (1987). Thermodynamics of the acicular ferrite transformation. Advances in Welding Science and Technology, ASM, Ohio, pp. 187-191.
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ABSTRACT The physical metallurgy of high-strength steel welds is studied both theoretically and experimentally with a view to establishing a basis for the further development and optimisation of arc welding consumables. The primary microstructure of such welds is unique, consisting of a mixture of acicular ferrite, bainite and low-carbon martensite. Since this is very resistant to tempering, and since the alloys necessarily have a very low Ae3 temperature, the effect of reheating is found to be minimal. Indeed, the microstructure of multirun welds is found to be fairly homogeneous and not very different from the primary microstructure. These and other factors combine to give very good toughness levels in spite of the high yield strength of the weld deposits. KEYWORDS High-strength steel welds; acicular ferrite; low-carbon martensite; microstructure; microstructure modelling; physical metallurgy of welds, ; bainite; thermodynamics. INTRODUCTION The level of alloying additions in most weld deposits is usually kept low in order to prevent cold cracking and other defects associated with the formation of brittle phases within the weld itself. The hardenability of the weld has to be low enough to avoid the transformation of residual austenite to relatively high carbon, untempered martensite (α'), as the weld cools towards ambient temperature. The problem can be made worse by the partitioning of carbon into the austenite as a consequence of other transformations which precede the martensite transformation; the carbon-enriched austenite then transforms into even harder martensite. At the same time, the austenite hardenability has to be sufficiently high to promote the formation of desirable phases such as acicular ferrite, at the expense of weaker phases such as allotriomorphic ferrite and Widmanstàtten ferrite. In a multirun weld deposit, the alloy composition may also have a strong effect on the extent of microstructural change in the material already deposited, due to the additional heat input as further layers are deposited. For these reasons, the yield strength (σ γ ) of conventional welds has in general been kept rather low (~ 350-550 MPa) with occasional excursions to higher strength experimental welds but usually at the expense of toughness. The use of low-carbon high-strength steels based on a microstructure of bainite or a mixed microstructure of bainite and martensite has recently assumed prominence in the fabrication of pipelines (Nakasngi, Matsuda and Tamehiro, 1983) and of ships and submarines (Deb, Challenger and Themen,
71
1987). To meet simultaneously the requirements of strength and toughness, a class of rather heavily alloyed, complex steel welding consumables have been developed, a typical chemical composition being Fe-2.2Ni-1.3Mn-0.5Mo-0.5Cr-0.1C wt.%. In spite of the high alloy content, these consumables give tough weld deposits which are extremely strong (σ ν « 730 MPa) and yet are resistant to hydrogen cracking and even are insensitive to certain specification requirement tests involving explosive deformation. The physical metallurgy of these alloy welds is not well understood either theoretically or experimentally. For example, metallography using light microscopy has proven impossible to interpret in terms of the numerous microstructure classification schemes that exist in the welding industry (Deb, Challenger and Themen, 1987). Even experiments using transmission electron microscopy have failed to characterise the microstructure with the sort of clarity needed for systematic experimentation or calculation. The purpose of this work was to attempt an understanding of the physical metallurgy of the highstrength steel welding alloys in order to provide a basis for systematic research and development towards the optimisation of such welds. The work also includes an experimental assessment of the microstructure in order to clarify the nature of the transformation products present in the various regions of multirun welds. EXPERIMENTAL The weld was fabricated using a submerged arc technique, as an X joint using 60 mm thick plate with 36 runs on each side. The welding current was 625 Amps at 30V, the weld being deposited in the flat position at a speed of « 40m/h (0.0111 m/s) with the interpass temperature (TQ) being kept within the range 150-200DC. The welding consumables used have the proprietary ESAB designations OK 10.62 (flux) and OK Autrod 13.43 (4 mm diameter wire); the base plate and deposit compositions are given in Table 1. TABLE 1 Chemical Compositions (wt.%) Weld Plate
C 0.071 0.110
Si 0.26 0.27
Weld Plate
Al 0.024 0.076
Cu 0.13 0.17
Mn 1.23 0.86 Sn 0.011 0.007
Ni 2.20 1.09 Pb 0.003 0.003
Cr 0.55 0.51
Mo 0.44 0.40 As 0.010 0.005
Sb 0.001 0.001
P 0.012 0.015 B 0.0004 0.0013
S 0.007 0.005 N 0.0083 0.0041
Nb 0.005 0.017
Ti 0.004 0.004
O 0.0371 0.0010
Tensile tests were carried out on standard cylindrical specimens of 12.8mm diameter with a gauge length of 65mm. The specimens were degassed at 100°C for 48h prior to testing. Charpy tests were carried out on standard 10mm square V-notch specimens at -18 and -51 °C. Transmission electron microscopy samples were prepared from 3mm diameter discs machined from known positions in the weld; by defining the extraction positions using light microscopy, it was possible to obtain discs containing just the primary microstructure or those just containing the reheated microstructure. The discs were mechanically ground down to a thickness of 0.08mm on 1200 grit SiC paper before electropolishing. The specimens were electropolished using a Tenupol jet polisher at a voltage of 55V in an electrolyte of 5% perchloric acid, 25% glycerol in ethanol, at ambient temperature. It was found that by reducing the polishing voltage (« 50V), the acicular ferrite could be attacked preferentially with respect to the martensite matrix, so that it could be distinguished clearly from the matrix phase. RESULTS AND DISCUSSION Mechanical Properties The results from mechanical tests are given in Table 2. The strain hardening coefficient 'n' and strengthening coefficient 'Κ' are crude estimates: if the [true stress o]/[true strain ε] curve is of the form σ = Κε11, and if o u is the tensile strength, then it can be shown that (o/a u ) = η" η ε η . Hence, values of
72
K and n can be estimated using the proof stress data (σοο02, σ001 measured at strains of 0.002 and 0.01 respectively), together with the measured tensile strength, using the criterion that a plot of (o/cu) versus (η"ηεη) should have a slope of unity. The results give an idea of the good toughness even though the yield strength is so high. The Vickers hardness of the top bead, which contains the microstructure unaffected by other beads, was found to be 295 whereas the mean hardness of the entire weld was found to be 281, indicating that the weld is extremely homogeneous in strength in spite of its multirun character. As will be discussed later, this is because of the high alloy content and temper resistance of its microstructure. TABLE 2 Mechanical Property Data
οοι' ΜΡ 3 Gu, MPa Reduction of Area, % Elongation, % K, MPa n
Test 1 761 786 854 62 21 980 0.044
Test 2 765 779 826 62 22 930 0.034
Charpy tests Energy absorbed, J
-51°C 75, 69, 71
-18°C 82, 79, 82
Tensile Properties
^
Brief Review of Transformation Products It is necessary to consider briefly the nature of the transformation products in steel welds, in order to set the scene for the discussion of theoretical and experimental results. Allotriomorphic ferrite (a) is the first phase to form during the cooling of austenite in low-alloy steels. It nucleates at the austenite grain boundaries and grows by a diffusional transformation mechanism which involves at the very least, reconstructive diffusion, which is the diffusion necessary to achieve the lattice change with a minimum of strain. The growth of a is anisotropie, the rate being highest along the γ boundaries which soon become decorated with layers of a. Pearlite or degenerate pearlite is the other product which grows from austenite by diffusional transformation. Detailed discussions of transformation mechanisms can be found in some recent reviews (Bhadeshia, 1985; Grong and Matlock, 1986; Abson and Pargeter, 1986; Bhadeshia, 1987). As the transformation temperature decreases, diffusion becomes sluggish and displacive transformations are kinetically (though not thermodynamically) favoured; these transformations do not involve any reconstructive diffusion so that substitutional or iron atoms do not diffuse during reaction. Consequently, the transformations are accompanied by macroscopic displacements reflecting the coordinated movement of atoms during the lattice change, the displacements having the characteristics of invariant-plane strains. Their morphology is thus dominated by the need to minimise strain energy. Plates of Widmanstätten ferrite (ο^) grow at relatively low undercoolings by a displacive paraequilibrium mechanism which involves the simultaneous growth of pairs of mutually accommodating plates of ferrite, so that the strain energy of transformation is reduced (Bhadeshia, 1981). Widmanstätten ferrite and bainite ( α ^ both nucleate by the same mechanism and are both represented, at least initially, by the same ' C curve on the TTT diagram. Bainite, which grows in the form of sheaves of small platelets, can only form at a higher undercooling because it deviates further from equilibrium, forming by diffusionless transformation from austenite (the carbon redistributing into γ subsequent to transformation), the strain energy term also being higher since the plates do not form in an accommodating manner (Bhadeshia, 1987). Lenticular plates of acicular ferrite (a a ) form by the same mechanism as bainite, but have a different morphology because nucleation occurs intragranularly at inclusions (rather than at γ grain boundaries) and the development of sheaves is stifled by hard impingement between plates nucleated at neighbouring sites
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(Yang and Bhadeshia, 1987; Strangwood and Bhadeshia, 1987). Theory Transformations from austenite can therefore be categorised into diffusional and displacive reactions, each category beingrepresentedby a 'C curve on the time-temperature-transformation diagram for the steel. The continuous lines on Fig. la show the time-temperature-transformation (TTT) diagram calculated (Bhadeshia, 1982) assuming that the weld is chemically homogeneous. We have noted earlier that the TTT diagram should consist of just two 'C curvesrepresentingthe initiation of diffusional and of displacive transformation. However, for the present alloy, a complication arises for the displacive 'C curve, since nucleation becomes impossible at an intermediate temperature, so that the curve splits into two halves, the higher temperature one for Widmanstätten ferrite and the other one for bainite and acicular ferrite. This is illustrated in Fig. lb where the F n curve represents the free energy needed to obtain a detectable nucleation rate of o^, a b or a a for any steel. The AFty-r^+a} curve is however specific to the present alloy, and represents the free energy available at any temperature; when AF{Y-^Y1+a}
Fig. 1. (a) Calculated TTT diagrams; (b) free energy curves. Weld deposits are not in fact chemically homogeneous since nonequilibrium solidification leads to coring. Regions lean in austenite stabilising alloying elements will have a lower hardenability and will influence the transformation behaviour of the entire weld. The approximate procedure for taking this into account has been discussed elsewhere (Gretoft, Bhadeshia and Svensson, 1986), but the TTT curve for the solute-depleted regions of the weld (calculated composition Fe-0.33C-0.37Si-0.93Mn-0.95Ni-0.13Mo0.49Cr as compared with Fe-0.33C-0.51Si-l.25Mn-2.09Ni-0.26Mo-0.59Cr at.% for the homogeneous alloy) is given in Fig. la as dashed lines. Because of the higher free energy change accompanying the γ-»α transformation in the solute-depleted regions, the two parts of the displacive *C* curve unite, although because of their different growth mechanism and stored energy, the bainite and acicular ferrite transformations can still only occur below the indicated temperature of 586°C. TTT curves can be used to estimate reactions occurring during continuous cooling, using the Scheil
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method, in which the cooling curve is divided into a series of isothermal steps. If vx is the incubation time as given by the TTT curve for the ith step, then transformation is considered initiated when S(tj/Tj)=l, where tj is the time for the ith isothermal heat treatment. Values of IXtj/Xj) obtained by calculating the cooling conditions of the weld (for the welding parameters presented earlier) and using the diffusional ' C curve for the solute depleted regions were always found to be less than unity, so that the the microstructure was predicted to consist of just acicular ferrite, bainite and low-carbon martensite. Widmanstätten ferrite was also ruled out because Z(tj/ij) for the displacive curve had a value of just 0.14 by the time the bainite-start temperature for the solute-depleted regions was reached. In fact, Ι^/τ^ became unity at = 480°C, only about 10°C above the martensite-start temperature of these regions, so that the volume fraction of acicular ferrite and bainite should be rather low, most of the microstructure being low-carbon martensite of a carbon concentration not very different from the starting level. Estimation of the volume fractions awaits the development of theory for predicting the nucleation rate of acicular ferrite. Theoretical Assessment of Alloy and Welding Effects A model for the prediction of microstructure in low-alloy steel weld deposits (Bhadeshia, Svensson and Gretoft, 1985, 1987) was used to assess the effect of alloy chemistry, welding current and speed on the microstructure of the primary regions of the weld. Calculations were done for 98 different cases; the results are presented in Tables 3 (changes in composition with respect to weld 1 have been emphasised using bold typeface) and Table 4. The latter includes only the cases where phases other than acicular ferrite and martensite were predicted, although the calculations were repeated for all the weld compositions listed in Table 3. The experimental weld is numbered 14. TABLE 3 The Influence of Composition Variations
c
Si
Mn
Ni
Mo
Cr
V
V
0.070 0.040 0.100 0.070 0.070 0.070 0.070 0.070 0.070 0.070 0.070 0.070 0.070 0.071
0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.30 0.26
1.25 1.25 1.25 1.75 0.75 1.25 1.25 1.25 1.25 1.25 1.25 1.25 1.25 1.23
2.20 2.20 2.20 2.20 2.20 3.00 1.40 2.20 2.20 2.20 2.20 2.20 2.20 2.20
0.45 0.45 0.45 0.45 0.45 0.45 0.45 0.65 0.25 0.45 0.45 0.45 0.45 0.44
0.55 0.55 0.55 0.55 0.55 0.55 0.55 0.55 0.55 0.80 0.30 0.55 0.55 0.55
0.0 0.0 0.0 0.0 0.0 0.0 0.0 0.0 0.0 0.0 0.0 0.1 0.2 0.0
0.00 0.00 0.00 0.00 0.17 0.00 0.13 0.00 0.00 0.00 0.00 0.00 0.00 0.00
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a
v
w
0.00 0.00 0.00 0.00 0.06 0.00 0.08 0.00 0.00 0.00 0.00 0.00 0.00 0.00
1 2 3 4 5 6 7 8 9 10 11 12 13 14
TABLE 4 The Influence of Other Variables Comment
v
Welding speed reduced (0.0083 m/s) Welding speed reduced (0.0083 m/s) Welding speed reduced (0.0083 m/s) Welding speed reduced (0.0083 m/s) Welding speed increased (0.0139 m/s) Interpass temperature reduced (150°C) Interpass temperature (250°C) Interpass temperature (250°C) Welding current reduced (525 A) Welding current increased (725 A) Welding current increased (725 A)
0.20 0.20 0.13 0.13 0.14 0.15 0.19 0.18 0.15 0.19 0.17
a
V
w
0.06 0.07 0.04 0.04 0.06 0.06 0.06 0.07 0.06 0.06 0.07
No. 5A (low Mn) 7A (low Ni) 9A (low Mo) 11A (lowCr) 5B (low Mn) 5C (low Mn) 5D (low Mn) 7B (low Ni) 5E (low Mn) 5F (low Mn) 7C (low Ni)
The calculations show clearly the role of Mn and to a lesser extent Ni, in preventing the formation of detrimental phases such as allotriomorphic ferrite and Widmanstàtten ferrite. A reduction in welding speed or an increase in the interpass temperature or an increase in welding current all cause a decrease in weld cooling rate and consequently can in some alloys enhance the formation of a and o^. It is evident the experimental weld (no. 14) is very tolerant to quite large variations in welding conditions in the sense that the primary microstructure remains a mixture of acicular ferrite and martensite for all the conditions used. There are of course other welds which behave similarly (nos. 1-4, 8, 10, 12, 13) and these compositions need to be investigated experimentally, since the mechanical properties could then be varied without changing the mixed acicular ferrite and martensite microstructure. There is as yet no model allowing the calculation of the strength of multirun submerged arc welds, but we use here a model by Svensson and co workers (1988), which is based on multirun manual metal arc welds, in order to obtain some crude estimates of yield strength as a function of composition. Note that the model takes account of solid solution strengthening and microstructural strengthening due to individual phases, both for the primary and reheated regions of the weld, the volume fraction of these regions depending on the Ae3' temperature (i.e., a point on the paraequilibrium γ/γ+α phase boundary for the alloy concerned). The results are presented in Table 5 (V p is the calculated volume fraction of the primary region) and with the earlier data of Table 3,4, can be used in order to design weld deposits of the desired strength and microstructure. Note that the calculated yield strength of weld 14 is about 40MPa lower than the measured strength. TABLE 5 Estimates of Mechanical Properties No. 1 2 3 4 5 6 7 8 9 10 11 12 13 14
Ae3\ °C 699 699 691 672 723 667 725 699 695 691 699 694 693 700
σ , MPa 728 729 739 783 642 795 644 732 730 740 727 735 737 723
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V
P
0.72 0.73 0.76 0.83 0.64 0.84 0.63 0.73 0.74 0.76 0.73 0.74 0.75 0.72
Transmission Electron Microscopy Examination of weld 14 revealed a microstructure consistent with the calculations and mechanical property data presented above. Its primary regions were found to consist of an estimated 40% of intragranularly nucleated, lenticular acicular ferrite plates, a small amount of γ grain boundary nucleated bainite, the remainder being low-carbon dislocated martensite, with thin films of retained austenite between the martensite plates (Fig. 2). The reheated regions of the weld were not in fact very different, since most of them had been fully reaustenitised due to the low Ae3' temperature of the alloy; due to space restrictions, the detailed microstructural study will be presented elsewhere. Both the primary and reaustenitised microstructure seem resistant to tempering, probably because of the very low-carbon content of the alloy. All this explains the relatively small observed variation in hardness; such a mechanically homogeneous microstructure is believed to be beneficial to toughness (Tweed and Knott, 1987). Since most of the microstructure is martensite, and since acicular ferrite is also a strong phase (Sugden and Bhadeshia, 1988; Svensson and coworkers, 1988), the strain hardening coefficient is expected to be low. The data of Table 2 are about the same as reported values of strain hardening coefficients for lowcarbon martensite in steels (Davies, 1978).
Fig. 2.
The primary microstructure; (a) a a in a martensite matrix; (b) α^ growing from a γ grain boundary.
CONCLUSIONS A primary weld microstructure consisting mainly of low-carbon martensite as the major phase, and acicular ferrite has a good combination of strength and toughness. However, to ensure high-strength in multirun weld deposits, it is also necessary to alloy in a manner which stabilises the austenite (reduces the Ae3 temperature), so that any additional heating due to the deposition of subsequent runs causes much of the reheated region to be reaustenitised. The hardenability of the alloy has to be high enough to permit these reaustenitised regions to transform to just acicular ferrite, low-carbon bainite and low-carbon martensite during cooling to ambient temperature. The tempering resistance of the alloy also has to be high so that those reheated regions which are not substantially reaustenitised do not lose their strength. These conditions all ensure that the multirun weld, is mechanically homogeneous with little variation in strength, an important requirement for the achievement of optimum toughness.
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ACKNOWLEDGMENTS The authors are grateful to ESAB AB (Sweden) for the provision of laboratory facilities and financial support for this research and to Professor D. Hull for the provision of laboratory facilities at the University of Cambridge.
REFERENCES Abson, D. J. (1986). Review on microstructure and toughness of arc welds. International Metals Rev., 3Λ, 141-194. Bhadeshia, H. K. D. H. (1981). Rationalisation of shear transformations in steels. Acta Metall., 29, 1117-1130. Bhadeshia, H. K. D. H. (1982). Thermodynamic analysis of TTT curves. Metal Science, 16, 159-165. Bhadeshia, H. K. D. H. (1985). Diffusional formation of ferrite in iron and its alloys. Prog, in Mat. Sci., 29, 321-386. Bhadeshia, H. K. D. H. (1987). Bainite in steels. Phase Transformations 87, Institute of Metals, London. Bhadeshia, H. K. D. H. and Edmonds, D. V. (1980). The mechanism of bainite formation in steels. Acta Metall., 28, 1265-1273. Bhadeshia, H. K. D. H., Svensson, L.-E. and Gretoft, B. (1985). Model for the development of microstructure in steel welds. Acta Metall., 33, 1271-1283. Bhadeshia, H. K. D. H., Svensson, L.-E. and Gretoft, B. (1987). Prediction of the microstructure of the fusion zone of multicomponent steel welds. Advances in Welding Science and Technology, ASM, Ohio, pp. 225-229. Davies, R. G. (1978). Mechanical properties of zero-carbon ferrite and martensite structures. Metall. Trans. A, 9 451-455. Deb, P., Challenger, K. D. and Themen, A. E. (1987). Structure-property correlation of SA and GMA weldments in HY100 steel. Metall. Trans. A, 18A. 987-999. Gretoft, B., Bhadeshia, H. K. D. H. and Svensson, L.-E. (1986). Development of microstructure in the fusion zone of steel weld deposits. Acta Stereologica, 5, 365-371. Grong, O. and Matlock, D. K. (1986). Microstructural development in mild and low-alloy steel welds. International Metals Rev., 31., 27-48. Nakasugi, N., Matsuda, H. and Tamehiro, H. (1983). Ultra-low carbon bainitic steel for line pipes. Steels for line pipe and pipeline fittings, Metals Society, London, pp. 90-95. Strangwood, M. and Bhadeshia, H. K. D. H. (1987). Mechanism of the acicular ferrite transformation. Advances in Welding Science and Technology, ASM, Ohio, pp. 209-213. Sugden, A. A. B. and Bhadeshia, H. K. D. H. (1988). Strength of steel welds, submitted to Metall. Trans. A. Svensson, L.-E., Gretoft, B., Sugden, A. A. B. and Bhadeshia, H. K. D. H. (1988). Computer-aided design of electrodes for arc welding, II. Computer Technology in Welding, Welding Institute, U.K. Tweed, J. and Knott, J. (1987). Micromechanisms of failure in C-Mn weld metals. Acta Metall., 35. 1401-1414. Yang, J. R. and Bhadeshia, H. K. D. H. (1987). Thermodynamics of the acicular ferrite transformation. Advances in Welding Science and Technology, ASM, Ohio, pp. 187-191.
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