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Weld HAZ embrittlement of Nb containing C–Mn steels
International Journal of Pressure Vessels and Piping 75 (1998) 661–677
Weld HAZ embrittlement of Nb containing C–Mn steels Roelof J. Hattingh a,*, Gert Pienaar b b
a Itec R&D, Iscor Ltd, Roger Dyason Road, Pretoria, South Africa Materials Science and Metallurgical Engineering Department, University of Pretoria, Pretoria, South Africa
Received 12 June 1998; accepted 22 June 1998
Abstract The adverse influence of Nb on weld HAZ properties is still an active issue of discussion between construction companies and steel manufacturers. Some controversy exists in the literature concerning the influence of Nb on HAZ properties under certain conditions, and this investigation was subsequently performed over a range of C and Nb compositions, typical of the steels concerned, and for three different single cycle heat inputs ranging from 1.5 to 6 kJ/mm. Simulated thermal cycles were employed, using a Gleeble 1500 thermomechanical simulator, followed by CVN testing. It is shown that Nb additions can have a detrimental or beneficial effect at low heat inputs, depending on the C level, but a severe detrimental effect of Nb on HAZ toughness is observed at high heat inputs and C levels. q 1998 Elsevier Science Ltd. All rights reserved Keywords: CGHAZ embrittlement; Nb; C–Mn steels; Simulated HAZ; CVN toughness; FATT 50; High heat input welding
1. Introduction Normalised carbon–manganese steels are still being used for the manufacturing of large constructions such as containment tanks for liquids and gases. These tanks can be several storeys high, with very large diameters. They are manufactured by welding the steel plates together in a staggered brick-like pattern along the circumference of the tank. Plates thicker than 50 mm are used for larger vessels, depending on the specific size and application. The welding of these plates constitutes a large fraction of the total cost of manufacturing due to the thickness of the weld, and the total length of the welds. Extremely high heat input welding processes are used for this reason, as a higher deposition rate can be obtained this way. Processes such as electro-slag and electro-gas welding are used with heat inputs as high as 35 kJ/mm. Due to the safety hazard of such massive storage tanks which can hold millions of liters of liquid or gas, the construction companies must ensure that material with sufficient strength and toughness is used, and that the construction as a whole has sufficient mechanical properties to prevent rupture. This implies that the weld metal and base metal heat affected zones of the welds must have sufficient strength and toughness. * Corresponding author. Tel.: +27-12-307-3335; Fax: +27-12-325-6769
The steel manufacturer, however, is mainly concerned with the mechanical properties of the steel, and would like to consistently obtain the minimum specified properties. Obtaining the required through-thickness mechanical properties for thick plate, normalised carbon–manganese steels can become borderline. The steel manufacturer would like to add micro-alloying elements, such as Nb, to consistently obtain the minimum specified properties to ensure a high yield. Addition of Nb up to a certain level, depending on the application, is allowed by certain construction companies, but the steel manufacturer would like to exceed these specified limits to obtain the specified strength and toughness properties more easily. The typical steels concerned are BS 1501 grades 223 and 225, BS 4360 grade 50D, and ASTM A516 grades 60, 65 and 70. This investigation was initiated as a result of the above mentioned conflicting interests between construction companies and steelmakers. The construction companies indicated that the limit on Nb additions is due to references in the literature concerning the adverse influence of Nb on weld metal and HAZ toughness. There are, however, conflicting reports in the literature on the influence of Nb, but in consideration of the catastrophic consequences associated with the failure of large storage tanks, construction companies prefer a more conservative approach. It is the objective of this investigation to fundamentally quantify the
0308-0161/98/$ - see front matter q 1998 Elsevier Science Ltd. All rights reserved PII: S0 30 8 -0 1 61 ( 98 ) 00 0 66 - 0
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Fig. 1. Schematic diagram of the experimental procedure.
influence of Nb on HAZ properties during high heat input welding, in order to assist resolution of the issue.
2. Experimental procedure The experimental procedure is schematically represented in Fig. 1. Each of the procedure steps is subsequently discussed in detail, including motivation for its inclusion in the procedure.
2.1. Laboratory casts To isolate the influence of Nb, and to fundamentally study its influence, as many of the variables as possible must be excluded or kept constant. The C content has a major influence on the behaviour of Nb and must also be accommodated. All other alloying elements are kept constant between tight tolerances. The C contents are typical of the relevant steels, while the Nb contents have been chosen inside as well as outside acceptable limits dictated by construction
Table 1 Chemical compositions of laboratory casts with the aim values in brackets [wt%] Cast no.
C
Nb
Mn (1.2)
Si (0.3)
P (0.005 max)
S (0.005 max)
Al (0.03)
N (0.007)
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.071 (0.06) 0.057 (0.06) 0.062 (0.06) 0.124 (0.12) 0.133 (0.12) 0.120 (0.12) 0.20 (0.19) 0.187 (0.19) 0.192 (0.19)
0.004 (0) 0.033 (0.03) 0.055 (0.06) 0.003 (0) 0.032 (0.03) 0.059 (0.06) 0.003 (0) 0.032 (0.03) 0.065 (0.06)
1.27 1.14 1.24 1.18 1.19 1.28 1.19 1.18 1.16
0.26 0.33 0.28 0.32 0.34 0.28 0.34 0.33 0.33
0.017 0.021 0.017 0.020 0.018 0.008 0.020 0.020 0.021
0.009 0.010 0.010 0.010 0.010 0.003 0.013 0.011 0.010
0.033 0.029 0.033 0.032 0.035 0.021 0.035 0.037 0.019
0.0069 0.0062 0.0079 0.0053 0.0052 0.0066 0.0050 0.0046 0.0065
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companies and material specifications. The aim and obtained chemical compositions are given in Table 1. 2.2. Reheating Initially it was thought to use the same reheating temperature as used in practice for Nb microalloyed steels, namely 12508C. However, to prevent excessive austenite grain growth in the non-Nb containing steels, it was decided to use a slightly lower reheat temperature of 12008C as most of the Nb carbides in the steels investigated should be dissolved at this temperature [1]. The soaking time at the reheating temperature is determined by the Nb(CN) precipitate size in the billets. In the case of laboratory casts these precipitates will be relatively small due to the fast cooling rate. The precipitate size is not expected to be larger than obtained in practice during continuous casting. It has been shown during a previous investigation [2] that a soaking time of 1 h is sufficient to dissolve all the Nb(CN) precipitates in continuously cast BS 4360 grade 50B containing 0.028%Nb and 0.16%C. 2.3. Rolling Rolling instead of forging down to the final thickness of 11 mm is preferred in order to obtain similar directional properties as in practice and to obtain controlled deformation, resulting in strain-induced Nb(CN) precipitation, resembling practical conditions. A total strain of 1.5 is obtained by rolling 50 mm billets, forged from 20 kg laboratory casts, down to 11 mm, thereby simulating practical rolling of a plate with a final thickness of 50 mm. The reduction in temperature per pass allows only four passes to finish at 9008C, which is the recommended finishing temperature for these steels [2]. 2.4. Normalising The material was normalised after rolling to equalise properties and because 50 mm plate material is usually normalised in practice. Thick plate HSLA material is normalised to obtain the specified through-thickness properties. 2.5. Reference specimens Nine CVN specimens and two tensile specimens were removed from each cast to serve as reference specimens of the mechanical properties of the base material. Hardness measurements were carried out on these specimens.
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influence of Nb on the toughness properties obtained in specific microstructures generated by a weld thermal cycle. As many variables as possible had to be excluded, of which the welding process has the most. The purpose therefore, was not the exact replication of practical welding processes or practical welding conditions. (2) Weld thermal simulation is an inexpensive, simple and rapid test method to investigate fundamental phenomena in a large number of specimens. (3) CVN impact testing is a valid indication of impact toughness in simulated specimens due to the much larger homogenous HAZ microstructure obtained. (4) A Gleeble simulator was available. The simulation route was chosen to obtain the most adverse HAZ toughness properties. This is obtained by a single thermal cycle with a peak temperature of 13508C, representing the coarse-grained HAZ (CGHAZ) at a position of 0.4 mm from the fusion line. The thermal cycles were varied to represent welding of a 50 mm thick plate at a low, medium and high heat input of 1.5, 3 and 6 kJ/mm. Analytical Rosenthal equations were used with refinements as suggested by Ion et al. [3] to predict thermal cycles as a function of plate thickness, preheat and heat input. 2.7. Mechanical testing Impact transition curves were constructed from CVN data and hardness measurements were made. 2.8. TEM analysis and optical metallography The microstructures were analyzed by optical metallography, and TEM analyses were performed in order to investigate precipitation behaviour and microstructures where they could not easily be characterised by optical methods. 2.9. Mechanisms The TEM, optical metallography and mechanical testing results were used to identify embrittling mechanisms for each of the chemical compositions and thermal routes. 2.10. Operating window From the identified mechanisms and data generated, an operating window was constructed. This window presents data on the HAZ toughness behaviour of Nb-containing steels in terms of the Nb and C content, heat input and position in the HAZ (0.4 mm from the fusion line). 3. Results and discussion
2.6. Thermal simulation
3.1. Base material properties
The thermal simulation route, using 11 mm 3 11 mm CVN blank specimens, was chosen for the following reasons. (1) This is a fundamental investigation into the
3.1.1. Chemical composition The obtained chemical compositions are listed in Table 1. LC signifies low C, MC medium C and HC high C.
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3.1.2. Microstructural analysis A definite grain refinement effect is noticed with an increase in Nb level. The microstructures consist of ferrite and pearlite, with an increase in pearlite fraction with increased C content, as expected. A TEM analysis was performed on material from selected compositions for two reasons, namely to explain some unexpected mechanical property behaviour, and to determine the size and distribution of Nb(CN) precipitates to confirm proper dissolution during reheating before rolling. For a 0.19%C–0.06%Nb steel, practically all the Nb carbides and nitrides will be in solution at 12008C [1]. It was on this basis that the reheat temperature before rolling was chosen as 12008C, at which temperature excessive grain growth in the lower alloyed steels is also not expected to occur. The combined effect of C and nitrogen was not taken into account at that stage, and it is expected that Nb(CN) precipitates will be more stable at a particular temperature than either NbC or NbN [1]. Some undissolved precipitates are subsequently expected in the higher alloyed steels, and the TEM analysis was performed to confirm this. TEM micrographs of the 0.12%C–0.06%Nb and 0.19%C–0.03%Nb rolled and normalised base material are shown in Fig. 2. Some relatively large undissolved Nb rich precipitates are indeed observed, which could have an effect on lowering the strength as observed in Table 3 (MC3). A network of extremely fine Nb-rich precipitates is also observed in the matrix, which is an indication that enough Nb and C were taken into solution. The properties obtained conform to the specified properties for these steels. For instance, the specified strength and toughness values for a normalised 50 mm BS4360 grade 50D steel plate are 490– 620 MPa tensile strength, 340 MPa minimum yield strength, and a CVN impact toughness of 27 J at ¹208C. The composition of this steel compares well with that of the 0.19%C–0.03%Nb steel, which delivered similar or better properties in the as-rolled and normalised condition, namely a tensile strength of 546 MPa, a yield strength of 367 MPa, and CVN impact toughness in excess of 60 J at ¹208C. The precipitation of aluminium nitride also plays a significant role in the behaviour of Nb micro-alloyed steels. It has been shown that the ductile-to-brittle transition temperatures of Nb-containing steels are further reduced with the addition of aluminium. An optimum aluminium addition, which determines the Al/N ratio for a particular N content in the steel, is required for maximum advantage [4]. The effect of aluminium on ductile-to-brittle transformation temperature is due to grain refinement following precipitation, and the removal of free nitrogen from solution, which is the most severe interstitial solid solution hardening element. Numerous AlN precipitates were observed during the TEM analysis of the base material. The ferrite grain size and pearlite fraction of the base materials were measured by image analysis, and are shown in Table 2. A general grain refining effect is obtained
Fig. 2. TEM micrographs showing large Nb(CN) precipitates, but also a fine distribution of Nb(CN) precipitates.
with an increase in Nb and C levels, and an increase in pearlite fraction is obtained with an increase in C level, as expected. 3.2. Mechanical properties The base material mechanical properties are listed in Table 3.
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R.J. Hattingh, G. Pienaar / International Journal of Pressure Vessels and Piping 75 (1998) 661–677 Table 2 Ferrite grain sizes and pearlite fractions of the as-rolled and normalised base material Cast no.
Cast composition
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.06C–0Nb 0.06C–0.03Nb 0.06C–0.06Nb 0.12C–0Nb 0.12C–0.03Nb 0.12C–0.06Nb 0.19C–0Nb 0.19C–0.03Nb 0.19C–0.06Nb
Ferrite grain size [mm] 13.1 10.5 9.4 11.5 9.1 5.2 8.5 7.8 6.7
d ¹1/2 [mm ¹1/2]
ASTM no. 10 10 11 10 11 12 11 11 12
8.7 9.7 10.3 9.3 10.5 13.9 10.9 11.4 12.3
Pearlite fraction [%] 16 7.5 10.6 19 28 18 23 24 36
exception of the 0.12%C casts, where the increase in Nb content had a significant increase in hardness. The FATT 50 was plotted as a function of hardness with the purpose of correcting the toughness for hardness, but a poor relationship was obtained. The trend does indicate a general increase in FATT 50 with increased hardness. However, this is not necessarily always the case, as lower FATT 50 values are obtained with increased Nb content, resulting in increased hardness, for some of the lower C steels. A decrease in grain size with increased Nb content is obtained due to more efficient pinning of grain boundaries by Nb(CN). More Nb at a particular C content, will result in more stable precipitates and more efficient pinning. The exceptional and poor toughness behaviour of the 0.06%C–0.03%Nb material cannot be explained at this stage. The cast was repeated, followed by the exact reheating, rolling, cooling and normalising route that were used as previously. Very much the same mechanical properties were obtained as with the original material. Further work is required in this regard and the matter will not be pursued further here.
3.2.1. CVN toughness properties The influence of different strengthening mechanisms on toughness and yield strength is thoroughly discussed by Gladman [1] and Pickering [5], of which a graphical representation is shown in Fig. 3. It can be seen that increased pearlite fraction has no influence on yield strength, but has an adverse influence on toughness. Refined grain size and the addition of manganese and aluminium are the only options which contribute positively to strength, while having a positive influence on toughness as well. The positive influence of aluminium on toughness is due to the removal of nitrogen from solution, while both elements have a grain refining effect. The observed properties will be analysed with reference to this model. The rolled and normalised base material CVN toughness properties, as a function of C and Nb content, are shown in Figs 4 and 5. FATT 50 and CVN upper shelf energy behaviour are shown. Three-dimensional plots show the overall trends in properties as a function of all the parameters very clearly. Additional points were added inbetween the existing data points, and were obtained from two-dimensional second-order regression curves. It follows that the influence of C on toughness properties is dominant over that of Nb, with one exception at 0.06%C– 0.03%Nb. As shown in Table 3, a significant increase in hardness is obtained with the increase in C content, but only a slight increase is obtained with the increase in Nb content, with the
3.3. Simulated CGHAZ properties 3.3.1. Metallography The optical micrographs are attached in Appendix A. The following aspects must be kept in mind during interpretation of these microstructures, and in the discussion of the
Table 3 Mechanical properties of the rolled and normalised base material Tensile properties
Toughness properties
Cast no.
Cast composition
Hardness [HV 30]
Rm [MPa]
Re [MPa]
El [%]
RA [%]
FATT 50 [8C]
US [J] a
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.06C–0Nb 0.06C–0.03Nb 0.06C–0.06Nb 0.12C–0Nb 0.12C–0.03Nb 0.12C–0.06Nb 0.19C–0Nb 0.19C–0.03Nb 0.19C–0.06Nb
124 129 (133) 132 134 145 151 153 153 160
438 444 (447) 457 475 510 497 543 546 556
305 318 (325) 330 313 367 347 350 367 380
40 37 (35) 38 37 35 36 34 34 32
60 53 (55) 57 47 48 52 44 45 43
¹87 ¹40 (¹53) ¹100 ¹28 ¹40 ¹56 ¹24 ¹34 ¹25
266 121 (132) 283 109 97 128 91 90 63
Note: Values in bracketts are those of the repeated LC2 cast. a US–CVN upper shelf energy.
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Fig. 3. The influence of strengthening mechanisms on yield strength and impact transition temperature [1,5].
mechanical properties of the simulated CGHAZ at the different weld heat inputs. There are primarily two aspects that determine the microstructure of a steel, namely the characteristics of the thermal profile during reheating and cooling, and the particular chemical composition. The parameters during the thermal cycle, such as the heating rate, maximum temperature, time at maximum temperature and cooling rate, determines
Fig. 4. Base material FATT 50 as a function of C and Nb content.
the behaviour of precipitates, phase transformations and grain growth or refinement at a particular chemical composition. A faster heating rate will increase transformation temperatures and precipitate dissolution temperatures with a subsequent influence on the grain growth behaviour. Time at some maximum temperature will determine the precipitate dissolution behaviour and once again the subsequent
Fig. 5. Base material CVN upper shelf energy as a function of C and Nb content.
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grain growth obtained. The cooling rate will determine the precipitation behaviour and transformation behaviour, as observed on the CCT diagram for a particular chemical composition and grain size, cooled from a given reheat temperature. Faster cooling rates will depress transformation temperatures, resulting in lower transformation temperature products being formed. Precipitation reactions will also be suppressed. Materials of different chemical composition will also deliver different microstructural compositions when subjected to the same thermal cycle. Several mechanisms are responsible for this behaviour, namely the depression of transformation temperatures by virtually all alloying elements in solid solution, differences in precipitation behaviour, the influence of alloying elements in solid solution on transformation and other kinetics, resulting in changes in hardenability and diffusion-dependent phenomena. Some alloying elements, such as C and N, have more significant influences on the transformation behaviour and the resultant microstructure than others such as Nb. The relative effect and quantity of a particular alloying element must therefore be taken into account. Less precipitates, which could pin grain boundaries and prevent austenite grain growth, are expected at low C levels. Furthermore, plain cementite precipitates are less stable at high temperatures than Nb(CN) precipitates, and an increase in the Nb level at a particular C level is expected to prevent austenite grain growth more efficiently. A microstructural refinement is therefore expected with Nb additions. However, during fast cooling rates, insufficient time may be available for complete precipitation to occur, and some alloying and interstitial elements may remain in solid solution. Slower cooling rates obtained during high heat input welding are expected to result in more complete
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precipitation. The relative influences of precipitation strengthening, elements in solid solution, the specific microstructure and the resultant grain size creates a complex situation which is rather difficult to quantify. At low alloying element levels, martensite and bainite start temperatures can be relatively high, and a larger fraction of these phases can form than at higher alloying element contents. Due to the higher transformation temperature, autotempering can take place, which will deliver structures with superior mechanical properties. Large austenite grain sizes will increase the hardenability of the material, resulting in a larger fraction of lower temperature transformation products. An increased final grain size and higher bainite fraction will have an adverse effect on toughness. A significant increase in alloying elements, and C in particular, will depress transformation temperatures and increase hardenability. This will result in low temperature transformation products such as bainite and martensite being formed at low temperatures, without autotempering being possible, during fast cooling rates. These microstructures will have high strength, but poor toughness properties. A decreased cooling rate and longer times at peak temperature, as obtained during higher heat input welding, will result in increased austenite grain size. The transformation temperatures are then depressed less and more complete precipitation is obtained during cooling. Relatively coarse, higher transformation temperature products are expected at higher heat inputs, with accompanying poor strength and toughness. A qualitative representation of the typical microstructures formed as a function of transformation temperature and C content, as well as the relative toughness properties of these microstructures, are shown in Fig. 6. Lower bainite is often generally classified as having better toughness properties than upper bainite. This,
Fig. 6. The toughness properties of microstructures formed at varying C levels and transformation temperatures [6].
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however, is only true for lower C steels. In the vicinity of 0.2%C, lower bainite exhibits poorer toughness than upper bainite. The other important observation is the formation of martensite at higher C levels which occurs at lower transformation temperatures without significant autotempering. A relatively high FATT 50 is subsequently obtained, as compared to autotempered martensite forming at higher transformation temperatures in low C steels. 3.3.2. Mechanical properties The strengthening mechanism obtained in pure iron is described by the Hall–Petch equation as follows [1,5]: jy ¼ ji þ ky d ¹ 1=2
(1)
where: j y ¼ lower yield stress (MPa); j i ¼ Peirls friction stress (MPa); k y ¼ constant; and d ¼ grain diameter (mm). The Hall–Petch equation has been extended for steels to include terms for dislocation density strengthening, solid solution strengthening and precipitation strengthening, but the exact relationship of these to each other (additivity) is still under discussion [1]. A similar equation exists for ductile-to-brittle transition temperature of pure iron (DBTT) [5] DBTT (8C) ¼ ln B ¹ ln C ¹ ky d ¹ 1=2
(2)
where: B, C and k y are all constants and d ¼ grain diameter (mm). This equation has also been extended for steels to include the effects of substitutional and interstitial solid solution, dislocation strengthening and precipitation strengthening [1]. This is schematically represented in Fig. 3. It must, however, be kept in mind that more subtle influences could be responsible for changes in the strength-totoughness ratio. Precipitation, for instance, removes alloying elements out of solid solution, and the effect of precipitation on strength and toughness must then be weighed up against that of reducing elements in solid solution. The influence of the particular microstructure is described in terms of grain size, precipitation and dislocation density. Lower transformation temperature products such as martensite have significantly higher dislocation densities than that of undeformed polygonal ferrite. Other deviations from the relationships above are concerned with the definition of grain size in non-polygonal microstructures such as acicular ferrite, bainite and martensite. It has been shown by Naylor [7] that the lath diameter and lath packet sizes become the important parameters to describe mechanical property behaviour, in particular toughness, and that prior austenite grain size is not the determining factor in lath type microstructures as was always accepted. The morphology and size of secondary phases such as carbides and other precipitates has also been shown to influence fracture behaviour. The FATT 50 was plotted as a function of HV 30 hardness in order to use the relationship to correct the toughness for
Fig. 7. FATT 50 as a function of C and Nb content for the 1.5 kJ/mm weld HAZ.
hardness (strength). Relatively poor correlations exist for the 1.5 and 3 kJ/mm heat inputs, but it is acceptable for the 6 kJ/mm heat input. The gradients obtained differ widely between the different heat inputs, where it is 0.19 for the 1.5 kJ/mm, 0.064 for the 3 kJ/mm, and 0.65 for the 6 kJ/mm simulated weld heat affected zones. This is understandable in the light of the widely varying microstructures obtained from one situation to another. It was with these reasons in mind that the decision was made not to correct the FATT 50 values for hardness. The individual hardness values will be kept in mind during the interpretation of the results. Graphs of FATT 50 and DFATT 50 as a function of C and Nb content for the various simulated weld heat inputs are
Fig. 8. DFATT 50 as a function of C and Nb content for the 1.5 kJ/mm weld HAZ.
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Fig. 9. FATT 50 as a function of C and Nb content for the 3 kJ/mm weld HAZ.
Fig. 11. FATT 50 as a function of C and Nb content for the 6 kJ/mm weld HAZ.
shown in Figs 7–12. During the interpretation of the change in property graphs, the shape of the base material curve (Fig. 4) must be kept in mind. The FATT 50 behaviour of the 0.06%C–0.03%Nb discussed earlier is referred to in particular. The high FATT 50 of this material in the rolled and normalised condition is clearly reflected in the DFATT 50 graphs of the weld HAZ material, which might create the impression that this material has superior resistance to toughness deterioration during welding. In fact, it is only a reflection of poor base material toughness properties. The above mentioned aspects have been considered during the interpretation of the results obtained for each condition.
3.4. Mechanical properties of the simulated 1.5 kJ/mm weld HAZ material
Fig. 10. DFATT 50 as a function of C and Nb content for the 3 kJ/mm weld HAZ.
Fig. 12. DFATT 50 as a function of C and Nb content for the 6 kJ/mm weld HAZ.
The mechanical properties of the simulated weld HAZ at a heat input of 1.5 kJ/mm are listed in Table 4. Taking into account the various factors which determine the specific microstructure, namely composition, thermal cycle and austenite grain size, the following attempt is made at describing the formation of these microstructures, and to explain their influence on the obtained mechanical properties. The factors described in the Hall–Petch equations for strengthening and DBTT will be taken into account, namely precipitation, solid solution, dislocation
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Table 4 Mechanical properties of the 1.5 kJ/mm simulated weld HAZ material Toughness properties Cast no.
Cast composition
Hardness [HV 30]
DHardness [HV 30]
FATT 50 [8C]
DFATT 50 [8C]
US [J]
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.06C–0Nb 0.06C–0.03Nb 0.06C–0.06Nb 0.12C–0Nb 0.12C–0.03Nb 0.12C–0.06Nb 0.19C–0Nb 0.19C–0.03Nb 0.19C–0.06Nb
233 199 215 225 310 246 433 410 445
109 70 83 91 165 95 280 257 285
8 3 13 20 ¹19 ¹26 62 49 30
95 43 113 48 21 30 86 83 55
120 107 115 37 33 68 29 29 40
DUS [J] ¹146 ¹14 ¹168 ¹72 ¹64 ¹60 ¹62 ¹61 ¹23
density (specific microstructure) and grain size. The toughness properties are shown in Figs 7 and 8.
3.5. Mechanical properties of the simulated 3 kJ/mm weld HAZ material
3.4.1. Interpretation of properties obtained during the 1.5kJ/mm weld haz simulation The increase in C from 0.06% to 0.19% dominates the behaviour of the material. Insufficient pinning of grain boundaries at low C and Nb levels are obtained. Higher bainite–ferrite type microstructures are obtained at low C content with effective grain sizes smaller than that of the base material, but with a higher dislocation density and more alloying and interstitial elements in solid solution due to the relatively fast cooling rate. The best strength and toughness properties are obtained at the intermediate C level with a Nb addition of 0.03% due to the formation of autotempered martensite. Grain refinement due to increased precipitate stability and an associated decrease in elements in solid solution with a Nb addition of 0.06% delivers a softer material with the lowest FATT 50. Untempered bainite and martensite are obtained at the high C level, associated with a high hardness of 410HV 30 to 445HV 30, and poor toughness properties. A gradual and consistent increase in hardness is observed with the increase in C content, while inconsistent hardness behaviour is obtained with the increase in Nb at the various C levels.
The mechanical properties of the simulated weld HAZ at a heat input of 3 kJ/mm are listed in Table 5. Figs 9 and 10 show the toughness properties. 3.5.1. Interpretation of properties obtained during the 3 kJ/ mm weld HAZ simulation The increase in C from 0.06% to 0.19% once again dominates the behaviour of the material. The overall behaviour of the material at a heat input of 3 kJ/mm is very much the same as that obtained for the 1.5 kJ/mm heat input, except at the high C range where autotempered martensite and bainite are obtained with better toughness properties and a lower hardness. Ferrite–bainite type microstructures are obtained at lower C levels. The microstructures are coarser than that obtained at 1.5 kJ/ mm due to more time at high temperatures, which means that more complete dissolution of precipitates will occur, leading to poorer grain boundary pinning. However, more complete precipitation will be obtained during cooling due to the slower cooling rate. The best strength and toughness properties are obtained at the intermediate C level with a Nb addition of 0.06% due to the formation of autotempered martensite and the refinement of the
Table 5 Mechanical properties of the 3 kJ/mm simulated weld HAZ material Toughness properties Cast no.
Cast composition
Hardness [HV 30]
DHardness [HV 30]
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.06C–0Nb 0.06C–0.03Nb 0.06C–0.06Nb 0.12C–0Nb 0.12C–0.03Nb 0.12C–0.06Nb 0.19C–0Nb 0.19C–0.03Nb 0.19C–0.06Nb
175 191 231 273 341 364 366 365 372
51 62 99 139 196 213 213 212 212
FATT 50 [8C] 9 17 6 19 8 ¹3 35 30 44
DFATT 50 [8C]
US [J]
96 57 106 47 48 53 59 64 69
110 74 116 66 43 81 41 38 33
DUS [J] ¹156 ¹47 ¹167 ¹43 ¹54 ¹47 ¹50 ¹52 ¹30
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Cast composition
Hardness [HV 30]
DHardness [HV 30]
FATT 50 [8C]
DFATT 50 [8C]
US [J]
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.06C–0Nb 0.06C–0.03Nb 0.06C–0.06Nb 0.12C–0Nb 0.12C–0.03Nb 0.12C–0.06Nb 0.19C–0Nb 0.19C–0.03Nb 0.19C–0.06Nb
195 195 206 223 232 302 308 285 326
71 66 74 89 87 151 155 132 166
¹14 1 1 46 42 68 53 83 83
73 41 101 74 82 124 77 117 108
278 128 178 54 88 98 60 32 43
structure due to the grain boundary pinning effect of more stable precipitates. 3.6. Mechanical properties of the simulated 6 kJ/mm weld HAZ material The mechanical properties of the simulated weld HAZ at a heat input of 6 kJ/mm are listed in Table 6. Figs 11 and 12 show the toughness properties. 3.6.1. Interpretation of properties obtained during the 6 kJ/ mm weld HAZ simulation The toughness behaviour over the whole C range differs significantly from that obtained for heat inputs of 1.5 and 3 kJ/mm. Contrary to expectation, the lowest FATT 50 values for all of the conditions investigated are obtained at the low C range, with a rather sharp increase to high FATT 50 values at intermediate and high C levels. The highest overall FATT 50 values in the vicinity of 70–808C are obtained at intermediate-C and high-Nb levels, and at high C and intermediate and high Nb levels. This is a clear indication that Nb has an adverse influence on toughness properties at high heat inputs, typically above 3 kJ/mm. It is further shown that C has a very significant influence on the general toughness behaviour of steels, and that an increase in C, associated with Nb additions has a significantly adverse effect on HAZ toughness of high heat input welds. The long times at high temperatures and the slow cooling rates obtained at a simulated weld heat input of 6 kJ/mm result in excessive grain growth, leading to the coarse microstructures observed. The slow cooling rate results in more complete precipitation, particularly at higher C and Nb levels. Coarse precipitates have a significant adverse effect on toughness. Larger prior austenite grain sizes mean higher hardenability, and at high C and Nb levels, the hardenability is high enough to deliver a mainly lower bainitic microstructure with poor toughness properties as shown in Fig. 6. Due to the lath-like microstructures obtained in the simulated heat affected zones, the findings of Naylor [7] are more applicable in this case, where a parameter describing lath packet geometry (size) is used to describe structure
DUS [J] 12 7 ¹105 ¹55 ¹9 ¹30 ¹31 ¹58 ¹20
property relationships. The lath packet size and geometry are, however, functions of prior austenite grain size, meaning that the same tendency of prior austenite grain size on toughness will also hold true for the lath packets. The TEM analysis showed that solute enriched areas transform to brittle untempered martensite at high C and Nb levels, which has a further adverse effect on toughness. The extremely poor toughness properties at high C and Nb levels are due to a combination of these microstructural effects, namely a coarse prior austenite grain size, a combination of high-C lower bainite and solute enriched untempered martensite areas with poor toughness properties, and coarse precipitates.
4. General discussion With reference to the three-dimensional graphs, it is clear that the toughness behaviour of the 1.5 and 3 kJ/mm simulated weld HAZ material performed in much the same way, except at high C levels where the low heat input simulation resulted in the formation of brittle untempered martensite. The best toughness properties are in both cases obtained at an intermediate C level of 0.12%C and at intermediate to high Nb levels. The toughness behaviour of the 6 kJ/mm weld HAZ simulation is clearly different to that of the lower heat input simulations. The best toughness properties are obtained at low C levels with a significant decrease in toughness at intermediate and high C levels. Nb additions at the higher C levels have a further significant adverse effect on toughness at a heat input of 6 kJ/mm, to deliver the poorest toughness for all the conditions investigated. These findings are summarised in Fig. 13 for each of the simulated welding conditions. The optimum operating windows for each condition is evident from these contour plots. From Fig. 13 it can be seen that the best toughness properties in the simulated HAZ at a heat input of 1.5 kJ/mm are obtained in the Nb region of 0.03–0.06%, and at a C level of 0.12%. In this case, minimum Nb and C levels are required
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Fig. 13. FATT 50 contour plots (8C) as a function of C and Nb content for weld HAZ simulation at 1.5, 3 and 6 kJ/mm.
for optimum toughness. It can further be seen that the tolerance for C increases with increased Nb additions above a C level of 0.12%. The positive gradients of the contours mean that the same FATT 50 can be obtained at higher C levels, with the addition of Nb. Almost the same behaviour is observed for the simulated HAZ at a heat input of 3 kJ/mm, with the best toughness properties at a Nb level of 0.06% and a C level of 0.12%. The lowest FATT 50 obtained in this area is, however, higher than that obtained at a heat input of 1.5 kJ/mm, and the tolerance for C decreases instead of increases at Nb levels above 0.03%. A completely different picture emerges at a simulated heat input of 6 kJ/mm. The best toughness properties are obtained at low C and Nb levels. The negative gradients
of the contour lines further indicate less tolerance towards C with increased levels of Nb. This intolerance is more severe at higher C levels, and extremely high FATT 50 values of the order of 808C are obtained at high C and Nb levels. It is shown that the behaviour of C–Mn steels with Nb additions can vary quite significantly depending on the particular chemical composition and weld thermal cycle. The adverse influence of Nb at high heat inputs has been confirmed, which supports the conservative attitudes of the construction companies. However, it is also shown that Nb additions can have a beneficial effect on HAZ properties at lower heat inputs over plain C–Mn steels. An operating window in terms of composition and heat input is established. It must, however, be verified by actual welding
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tests, as weld thermal simulations do not necessarily deliver the same results as obtained during actual welds.
5. Conclusions The adverse influence of Nb on HAZ toughness at high heat inputs is confirmed. The highest FATT 50 values of 838C for all the investigated conditions are obtained at a simulated weld HAZ heat input of 6 kJ/mm at intermediate to high C levels (0.12%C up to 0.19%C), and intermediate to high Nb levels (0.03%Nb up to 0.06%Nb). This confirms the fears of construction companies concerning the adverse influence of Nb on high heat input weld HAZ toughness properties, as reported in the literature. A combination of a few microstructurally observed phenomena is given as explanation for the observed poor toughness behaviour in this area, namely: (1) an overall coarser prior austenite grain size at a simulated heat input of 6 kJ/mm as compared to the other simulated heat inputs. The extended Hall–Petch [1,5] relationship for DBTT shows a negative effect of grain size on toughness (DBTT). Due to the lath-like microstructures obtained in the simulated heat affected zones, the findings of Naylor [7] are more applicable in this case, where a parameter describing lath packet geometry (size) is used to describe structure property relationships. The lath packet size and geometry is, however, a function of prior austenite grain size, meaning that the same tendency of prior austenite grain size on toughness will also hold true for the lath packet geometry, and (2) the larger prior austenite grain size increases hardenability, and although the cooling rates at a simulated weld thermal cycle of 6 kJ/mm are lower than at the other heat inputs, hardenability is sufficient to cause transformation to coarse upper and lower bainite and pockets of martensite at the 0.19%C level with additions of Nb. The slower cooling rate at a simulated heat input of 6 kJ/mm also causes more complete precipitation and coarsening of
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carbides. These coarse carbides, combined with brittle upper bainite and untempered martensite, are considered to be the main contributors towards the poor toughness properties. The C content is shown to determine the toughness properties in general and tends to be particularly detrimental to HAZ toughness at higher C levels (0.19%C) in combination with Nb. Nb does not have a significant influence on HAZ toughness at low C levels (0.06%C) during high heat input thermal cycling. Good toughness properties can be obtained at intermediate C levels of 0.12%C with intermediate to high Nb additions at lower heat inputs of 1.5 and 3 kJ/mm. High C levels (0.19%C) combined with a low heat input result in the formation of untempered brittle martensite and lower bainite with poor toughness properties. An operating window is established in terms of chemical composition and heat input, which shows the influence of these variables on HAZ toughness and strength properties. The strengthening and embrittling mechanisms has been identified to better understand the behaviour of Nb in C– Mn steels during weld thermal cycling. The operating window needs to be verified by actual weld runs.
Acknowledgements This work was financially sponsored by the Product Technology Department, Iscor Vanderbijlpark Works. Their contribution is gratefully acknowledged.
Appendix A Simulated weld HAZ microstructures at different heat inputs Figs A1–A3
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Fig. A1. 1.5kJ/mm Weld simulation (micrographs 5003 magnification)
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Fig. A2. 3kJ/mm Weld simulation (micrographs 5003 magnification)
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Fig. A3. 6kJ/mm Weld simulation (micrographs 5003 magnification)
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References [1] Gladman T. The physical metallurgy of microalloyed steels. London: The Institute of Materials, 1997. [2] Buirski J, Verdoorn FA, Whitfield JJ. Influence of plate rolling parameters and normalising on the microstructure and mechanical properties of BS 4360 Gr 50B. Iscor Research and Development Report, January 1994. [3] Ion JC, Easterling KE, Ashby MF. A second report on diagrams of microstructure and hardness for heat-affected zones in welds. Acta Metallurgica, 1984;32(11):1949–1962.
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[4] Kirkwood PR. Welding niobium containing microalloyed steels. In: Proceedings International Symposium: Niobium, San-Francisco, 8– 11 November 1981:761–802. [5] Pickering FB. Physical metallurgy and the design of steels. London: Applied Science Publishers, 1978. [6] Batte AD, Kirkwood PR. Developments in the weldability and toughness of steels for offshore structures. In: Conference Proceedings: Microalloyed HSLA steels, Chicago, IL, September 1988:175–188. [7] Naylor JP. The influence of the lath morphology on the yield stress and transition temperature of martensitic–bainitic steels. Metallurgical Transactions, 1979;10A:861–873.
Weld HAZ embrittlement of Nb containing C–Mn steels Roelof J. Hattingh a,*, Gert Pienaar b b
a Itec R&D, Iscor Ltd, Roger Dyason Road, Pretoria, South Africa Materials Science and Metallurgical Engineering Department, University of Pretoria, Pretoria, South Africa
Received 12 June 1998; accepted 22 June 1998
Abstract The adverse influence of Nb on weld HAZ properties is still an active issue of discussion between construction companies and steel manufacturers. Some controversy exists in the literature concerning the influence of Nb on HAZ properties under certain conditions, and this investigation was subsequently performed over a range of C and Nb compositions, typical of the steels concerned, and for three different single cycle heat inputs ranging from 1.5 to 6 kJ/mm. Simulated thermal cycles were employed, using a Gleeble 1500 thermomechanical simulator, followed by CVN testing. It is shown that Nb additions can have a detrimental or beneficial effect at low heat inputs, depending on the C level, but a severe detrimental effect of Nb on HAZ toughness is observed at high heat inputs and C levels. q 1998 Elsevier Science Ltd. All rights reserved Keywords: CGHAZ embrittlement; Nb; C–Mn steels; Simulated HAZ; CVN toughness; FATT 50; High heat input welding
1. Introduction Normalised carbon–manganese steels are still being used for the manufacturing of large constructions such as containment tanks for liquids and gases. These tanks can be several storeys high, with very large diameters. They are manufactured by welding the steel plates together in a staggered brick-like pattern along the circumference of the tank. Plates thicker than 50 mm are used for larger vessels, depending on the specific size and application. The welding of these plates constitutes a large fraction of the total cost of manufacturing due to the thickness of the weld, and the total length of the welds. Extremely high heat input welding processes are used for this reason, as a higher deposition rate can be obtained this way. Processes such as electro-slag and electro-gas welding are used with heat inputs as high as 35 kJ/mm. Due to the safety hazard of such massive storage tanks which can hold millions of liters of liquid or gas, the construction companies must ensure that material with sufficient strength and toughness is used, and that the construction as a whole has sufficient mechanical properties to prevent rupture. This implies that the weld metal and base metal heat affected zones of the welds must have sufficient strength and toughness. * Corresponding author. Tel.: +27-12-307-3335; Fax: +27-12-325-6769
The steel manufacturer, however, is mainly concerned with the mechanical properties of the steel, and would like to consistently obtain the minimum specified properties. Obtaining the required through-thickness mechanical properties for thick plate, normalised carbon–manganese steels can become borderline. The steel manufacturer would like to add micro-alloying elements, such as Nb, to consistently obtain the minimum specified properties to ensure a high yield. Addition of Nb up to a certain level, depending on the application, is allowed by certain construction companies, but the steel manufacturer would like to exceed these specified limits to obtain the specified strength and toughness properties more easily. The typical steels concerned are BS 1501 grades 223 and 225, BS 4360 grade 50D, and ASTM A516 grades 60, 65 and 70. This investigation was initiated as a result of the above mentioned conflicting interests between construction companies and steelmakers. The construction companies indicated that the limit on Nb additions is due to references in the literature concerning the adverse influence of Nb on weld metal and HAZ toughness. There are, however, conflicting reports in the literature on the influence of Nb, but in consideration of the catastrophic consequences associated with the failure of large storage tanks, construction companies prefer a more conservative approach. It is the objective of this investigation to fundamentally quantify the
0308-0161/98/$ - see front matter q 1998 Elsevier Science Ltd. All rights reserved PII: S0 30 8 -0 1 61 ( 98 ) 00 0 66 - 0
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Fig. 1. Schematic diagram of the experimental procedure.
influence of Nb on HAZ properties during high heat input welding, in order to assist resolution of the issue.
2. Experimental procedure The experimental procedure is schematically represented in Fig. 1. Each of the procedure steps is subsequently discussed in detail, including motivation for its inclusion in the procedure.
2.1. Laboratory casts To isolate the influence of Nb, and to fundamentally study its influence, as many of the variables as possible must be excluded or kept constant. The C content has a major influence on the behaviour of Nb and must also be accommodated. All other alloying elements are kept constant between tight tolerances. The C contents are typical of the relevant steels, while the Nb contents have been chosen inside as well as outside acceptable limits dictated by construction
Table 1 Chemical compositions of laboratory casts with the aim values in brackets [wt%] Cast no.
C
Nb
Mn (1.2)
Si (0.3)
P (0.005 max)
S (0.005 max)
Al (0.03)
N (0.007)
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.071 (0.06) 0.057 (0.06) 0.062 (0.06) 0.124 (0.12) 0.133 (0.12) 0.120 (0.12) 0.20 (0.19) 0.187 (0.19) 0.192 (0.19)
0.004 (0) 0.033 (0.03) 0.055 (0.06) 0.003 (0) 0.032 (0.03) 0.059 (0.06) 0.003 (0) 0.032 (0.03) 0.065 (0.06)
1.27 1.14 1.24 1.18 1.19 1.28 1.19 1.18 1.16
0.26 0.33 0.28 0.32 0.34 0.28 0.34 0.33 0.33
0.017 0.021 0.017 0.020 0.018 0.008 0.020 0.020 0.021
0.009 0.010 0.010 0.010 0.010 0.003 0.013 0.011 0.010
0.033 0.029 0.033 0.032 0.035 0.021 0.035 0.037 0.019
0.0069 0.0062 0.0079 0.0053 0.0052 0.0066 0.0050 0.0046 0.0065
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companies and material specifications. The aim and obtained chemical compositions are given in Table 1. 2.2. Reheating Initially it was thought to use the same reheating temperature as used in practice for Nb microalloyed steels, namely 12508C. However, to prevent excessive austenite grain growth in the non-Nb containing steels, it was decided to use a slightly lower reheat temperature of 12008C as most of the Nb carbides in the steels investigated should be dissolved at this temperature [1]. The soaking time at the reheating temperature is determined by the Nb(CN) precipitate size in the billets. In the case of laboratory casts these precipitates will be relatively small due to the fast cooling rate. The precipitate size is not expected to be larger than obtained in practice during continuous casting. It has been shown during a previous investigation [2] that a soaking time of 1 h is sufficient to dissolve all the Nb(CN) precipitates in continuously cast BS 4360 grade 50B containing 0.028%Nb and 0.16%C. 2.3. Rolling Rolling instead of forging down to the final thickness of 11 mm is preferred in order to obtain similar directional properties as in practice and to obtain controlled deformation, resulting in strain-induced Nb(CN) precipitation, resembling practical conditions. A total strain of 1.5 is obtained by rolling 50 mm billets, forged from 20 kg laboratory casts, down to 11 mm, thereby simulating practical rolling of a plate with a final thickness of 50 mm. The reduction in temperature per pass allows only four passes to finish at 9008C, which is the recommended finishing temperature for these steels [2]. 2.4. Normalising The material was normalised after rolling to equalise properties and because 50 mm plate material is usually normalised in practice. Thick plate HSLA material is normalised to obtain the specified through-thickness properties. 2.5. Reference specimens Nine CVN specimens and two tensile specimens were removed from each cast to serve as reference specimens of the mechanical properties of the base material. Hardness measurements were carried out on these specimens.
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influence of Nb on the toughness properties obtained in specific microstructures generated by a weld thermal cycle. As many variables as possible had to be excluded, of which the welding process has the most. The purpose therefore, was not the exact replication of practical welding processes or practical welding conditions. (2) Weld thermal simulation is an inexpensive, simple and rapid test method to investigate fundamental phenomena in a large number of specimens. (3) CVN impact testing is a valid indication of impact toughness in simulated specimens due to the much larger homogenous HAZ microstructure obtained. (4) A Gleeble simulator was available. The simulation route was chosen to obtain the most adverse HAZ toughness properties. This is obtained by a single thermal cycle with a peak temperature of 13508C, representing the coarse-grained HAZ (CGHAZ) at a position of 0.4 mm from the fusion line. The thermal cycles were varied to represent welding of a 50 mm thick plate at a low, medium and high heat input of 1.5, 3 and 6 kJ/mm. Analytical Rosenthal equations were used with refinements as suggested by Ion et al. [3] to predict thermal cycles as a function of plate thickness, preheat and heat input. 2.7. Mechanical testing Impact transition curves were constructed from CVN data and hardness measurements were made. 2.8. TEM analysis and optical metallography The microstructures were analyzed by optical metallography, and TEM analyses were performed in order to investigate precipitation behaviour and microstructures where they could not easily be characterised by optical methods. 2.9. Mechanisms The TEM, optical metallography and mechanical testing results were used to identify embrittling mechanisms for each of the chemical compositions and thermal routes. 2.10. Operating window From the identified mechanisms and data generated, an operating window was constructed. This window presents data on the HAZ toughness behaviour of Nb-containing steels in terms of the Nb and C content, heat input and position in the HAZ (0.4 mm from the fusion line). 3. Results and discussion
2.6. Thermal simulation
3.1. Base material properties
The thermal simulation route, using 11 mm 3 11 mm CVN blank specimens, was chosen for the following reasons. (1) This is a fundamental investigation into the
3.1.1. Chemical composition The obtained chemical compositions are listed in Table 1. LC signifies low C, MC medium C and HC high C.
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3.1.2. Microstructural analysis A definite grain refinement effect is noticed with an increase in Nb level. The microstructures consist of ferrite and pearlite, with an increase in pearlite fraction with increased C content, as expected. A TEM analysis was performed on material from selected compositions for two reasons, namely to explain some unexpected mechanical property behaviour, and to determine the size and distribution of Nb(CN) precipitates to confirm proper dissolution during reheating before rolling. For a 0.19%C–0.06%Nb steel, practically all the Nb carbides and nitrides will be in solution at 12008C [1]. It was on this basis that the reheat temperature before rolling was chosen as 12008C, at which temperature excessive grain growth in the lower alloyed steels is also not expected to occur. The combined effect of C and nitrogen was not taken into account at that stage, and it is expected that Nb(CN) precipitates will be more stable at a particular temperature than either NbC or NbN [1]. Some undissolved precipitates are subsequently expected in the higher alloyed steels, and the TEM analysis was performed to confirm this. TEM micrographs of the 0.12%C–0.06%Nb and 0.19%C–0.03%Nb rolled and normalised base material are shown in Fig. 2. Some relatively large undissolved Nb rich precipitates are indeed observed, which could have an effect on lowering the strength as observed in Table 3 (MC3). A network of extremely fine Nb-rich precipitates is also observed in the matrix, which is an indication that enough Nb and C were taken into solution. The properties obtained conform to the specified properties for these steels. For instance, the specified strength and toughness values for a normalised 50 mm BS4360 grade 50D steel plate are 490– 620 MPa tensile strength, 340 MPa minimum yield strength, and a CVN impact toughness of 27 J at ¹208C. The composition of this steel compares well with that of the 0.19%C–0.03%Nb steel, which delivered similar or better properties in the as-rolled and normalised condition, namely a tensile strength of 546 MPa, a yield strength of 367 MPa, and CVN impact toughness in excess of 60 J at ¹208C. The precipitation of aluminium nitride also plays a significant role in the behaviour of Nb micro-alloyed steels. It has been shown that the ductile-to-brittle transition temperatures of Nb-containing steels are further reduced with the addition of aluminium. An optimum aluminium addition, which determines the Al/N ratio for a particular N content in the steel, is required for maximum advantage [4]. The effect of aluminium on ductile-to-brittle transformation temperature is due to grain refinement following precipitation, and the removal of free nitrogen from solution, which is the most severe interstitial solid solution hardening element. Numerous AlN precipitates were observed during the TEM analysis of the base material. The ferrite grain size and pearlite fraction of the base materials were measured by image analysis, and are shown in Table 2. A general grain refining effect is obtained
Fig. 2. TEM micrographs showing large Nb(CN) precipitates, but also a fine distribution of Nb(CN) precipitates.
with an increase in Nb and C levels, and an increase in pearlite fraction is obtained with an increase in C level, as expected. 3.2. Mechanical properties The base material mechanical properties are listed in Table 3.
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Cast composition
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.06C–0Nb 0.06C–0.03Nb 0.06C–0.06Nb 0.12C–0Nb 0.12C–0.03Nb 0.12C–0.06Nb 0.19C–0Nb 0.19C–0.03Nb 0.19C–0.06Nb
Ferrite grain size [mm] 13.1 10.5 9.4 11.5 9.1 5.2 8.5 7.8 6.7
d ¹1/2 [mm ¹1/2]
ASTM no. 10 10 11 10 11 12 11 11 12
8.7 9.7 10.3 9.3 10.5 13.9 10.9 11.4 12.3
Pearlite fraction [%] 16 7.5 10.6 19 28 18 23 24 36
exception of the 0.12%C casts, where the increase in Nb content had a significant increase in hardness. The FATT 50 was plotted as a function of hardness with the purpose of correcting the toughness for hardness, but a poor relationship was obtained. The trend does indicate a general increase in FATT 50 with increased hardness. However, this is not necessarily always the case, as lower FATT 50 values are obtained with increased Nb content, resulting in increased hardness, for some of the lower C steels. A decrease in grain size with increased Nb content is obtained due to more efficient pinning of grain boundaries by Nb(CN). More Nb at a particular C content, will result in more stable precipitates and more efficient pinning. The exceptional and poor toughness behaviour of the 0.06%C–0.03%Nb material cannot be explained at this stage. The cast was repeated, followed by the exact reheating, rolling, cooling and normalising route that were used as previously. Very much the same mechanical properties were obtained as with the original material. Further work is required in this regard and the matter will not be pursued further here.
3.2.1. CVN toughness properties The influence of different strengthening mechanisms on toughness and yield strength is thoroughly discussed by Gladman [1] and Pickering [5], of which a graphical representation is shown in Fig. 3. It can be seen that increased pearlite fraction has no influence on yield strength, but has an adverse influence on toughness. Refined grain size and the addition of manganese and aluminium are the only options which contribute positively to strength, while having a positive influence on toughness as well. The positive influence of aluminium on toughness is due to the removal of nitrogen from solution, while both elements have a grain refining effect. The observed properties will be analysed with reference to this model. The rolled and normalised base material CVN toughness properties, as a function of C and Nb content, are shown in Figs 4 and 5. FATT 50 and CVN upper shelf energy behaviour are shown. Three-dimensional plots show the overall trends in properties as a function of all the parameters very clearly. Additional points were added inbetween the existing data points, and were obtained from two-dimensional second-order regression curves. It follows that the influence of C on toughness properties is dominant over that of Nb, with one exception at 0.06%C– 0.03%Nb. As shown in Table 3, a significant increase in hardness is obtained with the increase in C content, but only a slight increase is obtained with the increase in Nb content, with the
3.3. Simulated CGHAZ properties 3.3.1. Metallography The optical micrographs are attached in Appendix A. The following aspects must be kept in mind during interpretation of these microstructures, and in the discussion of the
Table 3 Mechanical properties of the rolled and normalised base material Tensile properties
Toughness properties
Cast no.
Cast composition
Hardness [HV 30]
Rm [MPa]
Re [MPa]
El [%]
RA [%]
FATT 50 [8C]
US [J] a
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.06C–0Nb 0.06C–0.03Nb 0.06C–0.06Nb 0.12C–0Nb 0.12C–0.03Nb 0.12C–0.06Nb 0.19C–0Nb 0.19C–0.03Nb 0.19C–0.06Nb
124 129 (133) 132 134 145 151 153 153 160
438 444 (447) 457 475 510 497 543 546 556
305 318 (325) 330 313 367 347 350 367 380
40 37 (35) 38 37 35 36 34 34 32
60 53 (55) 57 47 48 52 44 45 43
¹87 ¹40 (¹53) ¹100 ¹28 ¹40 ¹56 ¹24 ¹34 ¹25
266 121 (132) 283 109 97 128 91 90 63
Note: Values in bracketts are those of the repeated LC2 cast. a US–CVN upper shelf energy.
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Fig. 3. The influence of strengthening mechanisms on yield strength and impact transition temperature [1,5].
mechanical properties of the simulated CGHAZ at the different weld heat inputs. There are primarily two aspects that determine the microstructure of a steel, namely the characteristics of the thermal profile during reheating and cooling, and the particular chemical composition. The parameters during the thermal cycle, such as the heating rate, maximum temperature, time at maximum temperature and cooling rate, determines
Fig. 4. Base material FATT 50 as a function of C and Nb content.
the behaviour of precipitates, phase transformations and grain growth or refinement at a particular chemical composition. A faster heating rate will increase transformation temperatures and precipitate dissolution temperatures with a subsequent influence on the grain growth behaviour. Time at some maximum temperature will determine the precipitate dissolution behaviour and once again the subsequent
Fig. 5. Base material CVN upper shelf energy as a function of C and Nb content.
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grain growth obtained. The cooling rate will determine the precipitation behaviour and transformation behaviour, as observed on the CCT diagram for a particular chemical composition and grain size, cooled from a given reheat temperature. Faster cooling rates will depress transformation temperatures, resulting in lower transformation temperature products being formed. Precipitation reactions will also be suppressed. Materials of different chemical composition will also deliver different microstructural compositions when subjected to the same thermal cycle. Several mechanisms are responsible for this behaviour, namely the depression of transformation temperatures by virtually all alloying elements in solid solution, differences in precipitation behaviour, the influence of alloying elements in solid solution on transformation and other kinetics, resulting in changes in hardenability and diffusion-dependent phenomena. Some alloying elements, such as C and N, have more significant influences on the transformation behaviour and the resultant microstructure than others such as Nb. The relative effect and quantity of a particular alloying element must therefore be taken into account. Less precipitates, which could pin grain boundaries and prevent austenite grain growth, are expected at low C levels. Furthermore, plain cementite precipitates are less stable at high temperatures than Nb(CN) precipitates, and an increase in the Nb level at a particular C level is expected to prevent austenite grain growth more efficiently. A microstructural refinement is therefore expected with Nb additions. However, during fast cooling rates, insufficient time may be available for complete precipitation to occur, and some alloying and interstitial elements may remain in solid solution. Slower cooling rates obtained during high heat input welding are expected to result in more complete
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precipitation. The relative influences of precipitation strengthening, elements in solid solution, the specific microstructure and the resultant grain size creates a complex situation which is rather difficult to quantify. At low alloying element levels, martensite and bainite start temperatures can be relatively high, and a larger fraction of these phases can form than at higher alloying element contents. Due to the higher transformation temperature, autotempering can take place, which will deliver structures with superior mechanical properties. Large austenite grain sizes will increase the hardenability of the material, resulting in a larger fraction of lower temperature transformation products. An increased final grain size and higher bainite fraction will have an adverse effect on toughness. A significant increase in alloying elements, and C in particular, will depress transformation temperatures and increase hardenability. This will result in low temperature transformation products such as bainite and martensite being formed at low temperatures, without autotempering being possible, during fast cooling rates. These microstructures will have high strength, but poor toughness properties. A decreased cooling rate and longer times at peak temperature, as obtained during higher heat input welding, will result in increased austenite grain size. The transformation temperatures are then depressed less and more complete precipitation is obtained during cooling. Relatively coarse, higher transformation temperature products are expected at higher heat inputs, with accompanying poor strength and toughness. A qualitative representation of the typical microstructures formed as a function of transformation temperature and C content, as well as the relative toughness properties of these microstructures, are shown in Fig. 6. Lower bainite is often generally classified as having better toughness properties than upper bainite. This,
Fig. 6. The toughness properties of microstructures formed at varying C levels and transformation temperatures [6].
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however, is only true for lower C steels. In the vicinity of 0.2%C, lower bainite exhibits poorer toughness than upper bainite. The other important observation is the formation of martensite at higher C levels which occurs at lower transformation temperatures without significant autotempering. A relatively high FATT 50 is subsequently obtained, as compared to autotempered martensite forming at higher transformation temperatures in low C steels. 3.3.2. Mechanical properties The strengthening mechanism obtained in pure iron is described by the Hall–Petch equation as follows [1,5]: jy ¼ ji þ ky d ¹ 1=2
(1)
where: j y ¼ lower yield stress (MPa); j i ¼ Peirls friction stress (MPa); k y ¼ constant; and d ¼ grain diameter (mm). The Hall–Petch equation has been extended for steels to include terms for dislocation density strengthening, solid solution strengthening and precipitation strengthening, but the exact relationship of these to each other (additivity) is still under discussion [1]. A similar equation exists for ductile-to-brittle transition temperature of pure iron (DBTT) [5] DBTT (8C) ¼ ln B ¹ ln C ¹ ky d ¹ 1=2
(2)
where: B, C and k y are all constants and d ¼ grain diameter (mm). This equation has also been extended for steels to include the effects of substitutional and interstitial solid solution, dislocation strengthening and precipitation strengthening [1]. This is schematically represented in Fig. 3. It must, however, be kept in mind that more subtle influences could be responsible for changes in the strength-totoughness ratio. Precipitation, for instance, removes alloying elements out of solid solution, and the effect of precipitation on strength and toughness must then be weighed up against that of reducing elements in solid solution. The influence of the particular microstructure is described in terms of grain size, precipitation and dislocation density. Lower transformation temperature products such as martensite have significantly higher dislocation densities than that of undeformed polygonal ferrite. Other deviations from the relationships above are concerned with the definition of grain size in non-polygonal microstructures such as acicular ferrite, bainite and martensite. It has been shown by Naylor [7] that the lath diameter and lath packet sizes become the important parameters to describe mechanical property behaviour, in particular toughness, and that prior austenite grain size is not the determining factor in lath type microstructures as was always accepted. The morphology and size of secondary phases such as carbides and other precipitates has also been shown to influence fracture behaviour. The FATT 50 was plotted as a function of HV 30 hardness in order to use the relationship to correct the toughness for
Fig. 7. FATT 50 as a function of C and Nb content for the 1.5 kJ/mm weld HAZ.
hardness (strength). Relatively poor correlations exist for the 1.5 and 3 kJ/mm heat inputs, but it is acceptable for the 6 kJ/mm heat input. The gradients obtained differ widely between the different heat inputs, where it is 0.19 for the 1.5 kJ/mm, 0.064 for the 3 kJ/mm, and 0.65 for the 6 kJ/mm simulated weld heat affected zones. This is understandable in the light of the widely varying microstructures obtained from one situation to another. It was with these reasons in mind that the decision was made not to correct the FATT 50 values for hardness. The individual hardness values will be kept in mind during the interpretation of the results. Graphs of FATT 50 and DFATT 50 as a function of C and Nb content for the various simulated weld heat inputs are
Fig. 8. DFATT 50 as a function of C and Nb content for the 1.5 kJ/mm weld HAZ.
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Fig. 9. FATT 50 as a function of C and Nb content for the 3 kJ/mm weld HAZ.
Fig. 11. FATT 50 as a function of C and Nb content for the 6 kJ/mm weld HAZ.
shown in Figs 7–12. During the interpretation of the change in property graphs, the shape of the base material curve (Fig. 4) must be kept in mind. The FATT 50 behaviour of the 0.06%C–0.03%Nb discussed earlier is referred to in particular. The high FATT 50 of this material in the rolled and normalised condition is clearly reflected in the DFATT 50 graphs of the weld HAZ material, which might create the impression that this material has superior resistance to toughness deterioration during welding. In fact, it is only a reflection of poor base material toughness properties. The above mentioned aspects have been considered during the interpretation of the results obtained for each condition.
3.4. Mechanical properties of the simulated 1.5 kJ/mm weld HAZ material
Fig. 10. DFATT 50 as a function of C and Nb content for the 3 kJ/mm weld HAZ.
Fig. 12. DFATT 50 as a function of C and Nb content for the 6 kJ/mm weld HAZ.
The mechanical properties of the simulated weld HAZ at a heat input of 1.5 kJ/mm are listed in Table 4. Taking into account the various factors which determine the specific microstructure, namely composition, thermal cycle and austenite grain size, the following attempt is made at describing the formation of these microstructures, and to explain their influence on the obtained mechanical properties. The factors described in the Hall–Petch equations for strengthening and DBTT will be taken into account, namely precipitation, solid solution, dislocation
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Table 4 Mechanical properties of the 1.5 kJ/mm simulated weld HAZ material Toughness properties Cast no.
Cast composition
Hardness [HV 30]
DHardness [HV 30]
FATT 50 [8C]
DFATT 50 [8C]
US [J]
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.06C–0Nb 0.06C–0.03Nb 0.06C–0.06Nb 0.12C–0Nb 0.12C–0.03Nb 0.12C–0.06Nb 0.19C–0Nb 0.19C–0.03Nb 0.19C–0.06Nb
233 199 215 225 310 246 433 410 445
109 70 83 91 165 95 280 257 285
8 3 13 20 ¹19 ¹26 62 49 30
95 43 113 48 21 30 86 83 55
120 107 115 37 33 68 29 29 40
DUS [J] ¹146 ¹14 ¹168 ¹72 ¹64 ¹60 ¹62 ¹61 ¹23
density (specific microstructure) and grain size. The toughness properties are shown in Figs 7 and 8.
3.5. Mechanical properties of the simulated 3 kJ/mm weld HAZ material
3.4.1. Interpretation of properties obtained during the 1.5kJ/mm weld haz simulation The increase in C from 0.06% to 0.19% dominates the behaviour of the material. Insufficient pinning of grain boundaries at low C and Nb levels are obtained. Higher bainite–ferrite type microstructures are obtained at low C content with effective grain sizes smaller than that of the base material, but with a higher dislocation density and more alloying and interstitial elements in solid solution due to the relatively fast cooling rate. The best strength and toughness properties are obtained at the intermediate C level with a Nb addition of 0.03% due to the formation of autotempered martensite. Grain refinement due to increased precipitate stability and an associated decrease in elements in solid solution with a Nb addition of 0.06% delivers a softer material with the lowest FATT 50. Untempered bainite and martensite are obtained at the high C level, associated with a high hardness of 410HV 30 to 445HV 30, and poor toughness properties. A gradual and consistent increase in hardness is observed with the increase in C content, while inconsistent hardness behaviour is obtained with the increase in Nb at the various C levels.
The mechanical properties of the simulated weld HAZ at a heat input of 3 kJ/mm are listed in Table 5. Figs 9 and 10 show the toughness properties. 3.5.1. Interpretation of properties obtained during the 3 kJ/ mm weld HAZ simulation The increase in C from 0.06% to 0.19% once again dominates the behaviour of the material. The overall behaviour of the material at a heat input of 3 kJ/mm is very much the same as that obtained for the 1.5 kJ/mm heat input, except at the high C range where autotempered martensite and bainite are obtained with better toughness properties and a lower hardness. Ferrite–bainite type microstructures are obtained at lower C levels. The microstructures are coarser than that obtained at 1.5 kJ/ mm due to more time at high temperatures, which means that more complete dissolution of precipitates will occur, leading to poorer grain boundary pinning. However, more complete precipitation will be obtained during cooling due to the slower cooling rate. The best strength and toughness properties are obtained at the intermediate C level with a Nb addition of 0.06% due to the formation of autotempered martensite and the refinement of the
Table 5 Mechanical properties of the 3 kJ/mm simulated weld HAZ material Toughness properties Cast no.
Cast composition
Hardness [HV 30]
DHardness [HV 30]
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.06C–0Nb 0.06C–0.03Nb 0.06C–0.06Nb 0.12C–0Nb 0.12C–0.03Nb 0.12C–0.06Nb 0.19C–0Nb 0.19C–0.03Nb 0.19C–0.06Nb
175 191 231 273 341 364 366 365 372
51 62 99 139 196 213 213 212 212
FATT 50 [8C] 9 17 6 19 8 ¹3 35 30 44
DFATT 50 [8C]
US [J]
96 57 106 47 48 53 59 64 69
110 74 116 66 43 81 41 38 33
DUS [J] ¹156 ¹47 ¹167 ¹43 ¹54 ¹47 ¹50 ¹52 ¹30
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Cast composition
Hardness [HV 30]
DHardness [HV 30]
FATT 50 [8C]
DFATT 50 [8C]
US [J]
LC1 LC2 LC3 MC1 MC2 MC3 HC1 HC2 HC3
0.06C–0Nb 0.06C–0.03Nb 0.06C–0.06Nb 0.12C–0Nb 0.12C–0.03Nb 0.12C–0.06Nb 0.19C–0Nb 0.19C–0.03Nb 0.19C–0.06Nb
195 195 206 223 232 302 308 285 326
71 66 74 89 87 151 155 132 166
¹14 1 1 46 42 68 53 83 83
73 41 101 74 82 124 77 117 108
278 128 178 54 88 98 60 32 43
structure due to the grain boundary pinning effect of more stable precipitates. 3.6. Mechanical properties of the simulated 6 kJ/mm weld HAZ material The mechanical properties of the simulated weld HAZ at a heat input of 6 kJ/mm are listed in Table 6. Figs 11 and 12 show the toughness properties. 3.6.1. Interpretation of properties obtained during the 6 kJ/ mm weld HAZ simulation The toughness behaviour over the whole C range differs significantly from that obtained for heat inputs of 1.5 and 3 kJ/mm. Contrary to expectation, the lowest FATT 50 values for all of the conditions investigated are obtained at the low C range, with a rather sharp increase to high FATT 50 values at intermediate and high C levels. The highest overall FATT 50 values in the vicinity of 70–808C are obtained at intermediate-C and high-Nb levels, and at high C and intermediate and high Nb levels. This is a clear indication that Nb has an adverse influence on toughness properties at high heat inputs, typically above 3 kJ/mm. It is further shown that C has a very significant influence on the general toughness behaviour of steels, and that an increase in C, associated with Nb additions has a significantly adverse effect on HAZ toughness of high heat input welds. The long times at high temperatures and the slow cooling rates obtained at a simulated weld heat input of 6 kJ/mm result in excessive grain growth, leading to the coarse microstructures observed. The slow cooling rate results in more complete precipitation, particularly at higher C and Nb levels. Coarse precipitates have a significant adverse effect on toughness. Larger prior austenite grain sizes mean higher hardenability, and at high C and Nb levels, the hardenability is high enough to deliver a mainly lower bainitic microstructure with poor toughness properties as shown in Fig. 6. Due to the lath-like microstructures obtained in the simulated heat affected zones, the findings of Naylor [7] are more applicable in this case, where a parameter describing lath packet geometry (size) is used to describe structure
DUS [J] 12 7 ¹105 ¹55 ¹9 ¹30 ¹31 ¹58 ¹20
property relationships. The lath packet size and geometry are, however, functions of prior austenite grain size, meaning that the same tendency of prior austenite grain size on toughness will also hold true for the lath packets. The TEM analysis showed that solute enriched areas transform to brittle untempered martensite at high C and Nb levels, which has a further adverse effect on toughness. The extremely poor toughness properties at high C and Nb levels are due to a combination of these microstructural effects, namely a coarse prior austenite grain size, a combination of high-C lower bainite and solute enriched untempered martensite areas with poor toughness properties, and coarse precipitates.
4. General discussion With reference to the three-dimensional graphs, it is clear that the toughness behaviour of the 1.5 and 3 kJ/mm simulated weld HAZ material performed in much the same way, except at high C levels where the low heat input simulation resulted in the formation of brittle untempered martensite. The best toughness properties are in both cases obtained at an intermediate C level of 0.12%C and at intermediate to high Nb levels. The toughness behaviour of the 6 kJ/mm weld HAZ simulation is clearly different to that of the lower heat input simulations. The best toughness properties are obtained at low C levels with a significant decrease in toughness at intermediate and high C levels. Nb additions at the higher C levels have a further significant adverse effect on toughness at a heat input of 6 kJ/mm, to deliver the poorest toughness for all the conditions investigated. These findings are summarised in Fig. 13 for each of the simulated welding conditions. The optimum operating windows for each condition is evident from these contour plots. From Fig. 13 it can be seen that the best toughness properties in the simulated HAZ at a heat input of 1.5 kJ/mm are obtained in the Nb region of 0.03–0.06%, and at a C level of 0.12%. In this case, minimum Nb and C levels are required
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Fig. 13. FATT 50 contour plots (8C) as a function of C and Nb content for weld HAZ simulation at 1.5, 3 and 6 kJ/mm.
for optimum toughness. It can further be seen that the tolerance for C increases with increased Nb additions above a C level of 0.12%. The positive gradients of the contours mean that the same FATT 50 can be obtained at higher C levels, with the addition of Nb. Almost the same behaviour is observed for the simulated HAZ at a heat input of 3 kJ/mm, with the best toughness properties at a Nb level of 0.06% and a C level of 0.12%. The lowest FATT 50 obtained in this area is, however, higher than that obtained at a heat input of 1.5 kJ/mm, and the tolerance for C decreases instead of increases at Nb levels above 0.03%. A completely different picture emerges at a simulated heat input of 6 kJ/mm. The best toughness properties are obtained at low C and Nb levels. The negative gradients
of the contour lines further indicate less tolerance towards C with increased levels of Nb. This intolerance is more severe at higher C levels, and extremely high FATT 50 values of the order of 808C are obtained at high C and Nb levels. It is shown that the behaviour of C–Mn steels with Nb additions can vary quite significantly depending on the particular chemical composition and weld thermal cycle. The adverse influence of Nb at high heat inputs has been confirmed, which supports the conservative attitudes of the construction companies. However, it is also shown that Nb additions can have a beneficial effect on HAZ properties at lower heat inputs over plain C–Mn steels. An operating window in terms of composition and heat input is established. It must, however, be verified by actual welding
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tests, as weld thermal simulations do not necessarily deliver the same results as obtained during actual welds.
5. Conclusions The adverse influence of Nb on HAZ toughness at high heat inputs is confirmed. The highest FATT 50 values of 838C for all the investigated conditions are obtained at a simulated weld HAZ heat input of 6 kJ/mm at intermediate to high C levels (0.12%C up to 0.19%C), and intermediate to high Nb levels (0.03%Nb up to 0.06%Nb). This confirms the fears of construction companies concerning the adverse influence of Nb on high heat input weld HAZ toughness properties, as reported in the literature. A combination of a few microstructurally observed phenomena is given as explanation for the observed poor toughness behaviour in this area, namely: (1) an overall coarser prior austenite grain size at a simulated heat input of 6 kJ/mm as compared to the other simulated heat inputs. The extended Hall–Petch [1,5] relationship for DBTT shows a negative effect of grain size on toughness (DBTT). Due to the lath-like microstructures obtained in the simulated heat affected zones, the findings of Naylor [7] are more applicable in this case, where a parameter describing lath packet geometry (size) is used to describe structure property relationships. The lath packet size and geometry is, however, a function of prior austenite grain size, meaning that the same tendency of prior austenite grain size on toughness will also hold true for the lath packet geometry, and (2) the larger prior austenite grain size increases hardenability, and although the cooling rates at a simulated weld thermal cycle of 6 kJ/mm are lower than at the other heat inputs, hardenability is sufficient to cause transformation to coarse upper and lower bainite and pockets of martensite at the 0.19%C level with additions of Nb. The slower cooling rate at a simulated heat input of 6 kJ/mm also causes more complete precipitation and coarsening of
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carbides. These coarse carbides, combined with brittle upper bainite and untempered martensite, are considered to be the main contributors towards the poor toughness properties. The C content is shown to determine the toughness properties in general and tends to be particularly detrimental to HAZ toughness at higher C levels (0.19%C) in combination with Nb. Nb does not have a significant influence on HAZ toughness at low C levels (0.06%C) during high heat input thermal cycling. Good toughness properties can be obtained at intermediate C levels of 0.12%C with intermediate to high Nb additions at lower heat inputs of 1.5 and 3 kJ/mm. High C levels (0.19%C) combined with a low heat input result in the formation of untempered brittle martensite and lower bainite with poor toughness properties. An operating window is established in terms of chemical composition and heat input, which shows the influence of these variables on HAZ toughness and strength properties. The strengthening and embrittling mechanisms has been identified to better understand the behaviour of Nb in C– Mn steels during weld thermal cycling. The operating window needs to be verified by actual weld runs.
Acknowledgements This work was financially sponsored by the Product Technology Department, Iscor Vanderbijlpark Works. Their contribution is gratefully acknowledged.
Appendix A Simulated weld HAZ microstructures at different heat inputs Figs A1–A3
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Fig. A1. 1.5kJ/mm Weld simulation (micrographs 5003 magnification)
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Fig. A2. 3kJ/mm Weld simulation (micrographs 5003 magnification)
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Fig. A3. 6kJ/mm Weld simulation (micrographs 5003 magnification)
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