Module 21 Welding Metallurgy of steel.pdf

5J Welding Metallurgy of Steels Module 21 PART OF THE GCIL CERTIFICATE PROGRAMS A FOR INDUSTRIAL LEARNING Training today for the needs of the future.  Gooderham Centre for Industrial Learning...

5J Welding Metallurgy of Steels Module 21 PART OF THE GCIL CERTIFICATE PROGRAMS A FOR INDUSTRIAL LEARNING Training today for the needs of the future.  Gooderham Centre for Industrial Learning J Welding Metallurgy of Steels Module 21 £: j Copyright © 2005 by The CWB Group All rights reserved. Although due care has been taken in the preparation of this module neither the Gooderham Centre nor any contributing author can accept any liability arising from the use or misuse of any information contained herein or for any errors that may be contained in the module. Information is presented for educational purposes and should not be used for design, material selection, procedure selection or similar purposes without independent verification. Where reference to other documents, such as codes and standards, is made readers are encouraged to consult the original sources in detail. -j SI i 7250 West Credit Avenue, Mississauga, ON L5N5N1 Tel: 905-542-2176, Fax: 905-542-1837 www.gciltraining.com [CWB] t? Gooderham Centre for Industrial Learning o' MODULE 21 WELDING METALLURGY OF STEELS CONTENTS Introduction and Objectives........ 1 Steel Metallurgy.......................... 2 Crystal structure..................... 2 Iron-iron carbon phase diagram 3 HAZ grain size...................... 11 Isothermal diagrams............... 16 CCT diagrams........................ 19 Hardenability...................... 22 Properties of Steels...................... 30 Effect of carbon...................... 30 Effect of thickness.................. 32 Effect of work hardening........35 Heat Affected Zone...................... 37 Weld Metal.................................. 46 Structure................................ 46 Weld metal composition........ 48 GMAW electrodes................. 50 Fluxes.................................... 50 Properties of welds................. 56 Summary..................................... 62 Additional Resources.................. 63 Guides and Exercises................... 64 i  ! *.JJ Gooderham Centre for Industrial Learning J MODULE 21 WELDING METALLURGY OF The increasing sophistication of steel property requirements— L. especially fracture toughness—has resulted in a new generation of L steels made by complex processing routes. These present important L. challenges to the welding engineer who must select welding consum­ ables and develop procedures giving specific properties and which avoid cracking problems. This task is made all the more difficult by i: > the growing sophistication of steel metallurgical design and demands L. a thorough knowledge of underlying metallurgical principles. This is the subject of this module. L- Here you will leam about the structure of steels and how they L. obtain their strength and toughness. You will use various diagrams to determine the effects of heat treatment on the structure of steel. We will discuss what factors determine the properties of welds—both weld metal and the heat affected zone—and finally we show you how to determine a welding procedure that will avoid cracking problems. We do not discuss steel making or ways of classifying and specifying steel. These topics are covered in Module 8 and we recommend you study that Module in addition to this one. Objectives After successfully completing this module you will be able to: Describe the ways steels derive their strength and toughness Describe the effects of heat on steel properties Explain how the welding procedure affects the weld metal and heat affected zone properties Determine procedures necessary to avoid various cracking r4 problems 1  tJ Gooderham Centre for Industrial Learning ZA -J STEEL METALLURGY 'J Steels form the largest group of commercially important alloys. "J There are several reasons for this: The great abundance of iron in the earth's crust The relative ease of extraction and low cost The wide range of properties that can be achieved as a result IH of the solid state phase transformation Other unique properties such as magnetism za Steels are alloys of iron and less than 2% carbon plus a wide range za rJ of other elements. Some of these are added deliberately to impart special properties and others are impurities not completely removed ZA during the steel making process. Elements may be present in solid solution or combined as intermetallic compounds with iron, carbon or other elements. Some elements, namely carbon, nitrogen, boron and ZA “J hydrogen, form interstitial solutions with iron whereas others such as manganese and silicon form substitutional solutions. Beyond the limit “J of solubility these elements may also form intermetallic compounds -J with iron or other elements. ZA -J Iron has the special property of existing in different crystallogra­ Crystal structure phic forms in the solid state. Below 910°C the structure is body- ZA centred cubic (bcc) but between 910°C and 1390°C it assumes a face- centred cubic structure (fee). Above 1390°C up to the melting point ZA at 1534°C the structure reverts back to the body-centred cubic form. These are known as allotropic forms of iron. The face-centred cubic ZA form is a close-packed structure being more dense than the body- centred cubic form. Consequently iron will actually contract as it is heated above 910°C when the structure transformation takes place. Fig. 1 illustrates this. ZA ZA ZA Figure 1. Heating a bar of iron causes it to expand but when transformation occurs from the bcc structure to the Expansion bcc liquid ZA fee structure at 910°C a small contraction takes place. The fee crystal structure is tcc close-packed and occupies less volume bcc than the bcc structure. 910 1390 1534 Temperature °C 2 ZA ^ Gooderham Centre for Industrial Learning 700 r 600 - Maximum solubility o{ 500 - carbon in a iron is O O 0.02 wt % at 723°C P 400 - CD cl 300 E Extremely low solubility o of carbon in a iron at 200 room temperature 100 - 0 1 1 1 J 0.005 0.01 0.015 0.02 weight % carbon Figure 2. Solubility of carbon in a (bcc) iron as a function of temperature. c* us? The solubility of carbon in the bcc form of iron is very small, the ulity of carbon maximum solubility being only about 0.02 weight per cent at V 723°C. As Fig. 2 shows there is negligible solubility of carbon in iron at room temperature (less than 0.0001 wt %). Since steels nearly always have more carbon than this, the excess carbon is not in solution but rather present as the intermetallic compound iron- carbide Fe3C known as cementite. In contrast the face-centred cubic form of iron dissolves up to 2.0 % carbon, well in excess of the usual carbon content of steels. A steel can therefore be heated to a temperature at which the structure changes to fee and all carbon goes into solution. The way in which the carbon is obliged to redistribute itself upon cooling back below the transformation is the origin of the wide range of properties achievable in steels. Iron-Iron carbide Fundamental to a study of steels is an understanding of the iron- phase diagram carbon phase diagram. The diagram we will discuss is actually the metastable iron-iron carbide system. The true stable form of carbon in iron is graphite, but except for cast irons this only occurs after prolonged heating. Since in steels carbon is normally present as iron -> carbide, it is this system that we consider. Fig. 3 shows the iron-iron carbide system up to 6 wt % carbon. We will now discuss several important features of this diagram. 3 VLJ  Gooderham Centre for Industrial Learning ZJ ZJ t zj. zz 1500 delta ferrite tj -I 1400 liquid ‘ ' ‘I 1300 liquid + -J 1200 austenite -J austenite -J 1100 ZJ ZA 1000 -J austenite + cementite ZJ 900 ZJ 800 ZJ Temp , ZJ, °C 700 zJ 600 ferrite zA 500 ferrite + cementite ZJ 400 300 ZJ ZJ 200 ZJ 100 ZJ ZJ 0 1 ZJ 1.0 2.0 3.0 4.0 5.0 6.0 ZJ Weight percent of carbon ZA Figure 3. The iron-iron carbide phase diagram. ZJ ZJ ZJ zJ 4 ZJ ZJ  Gooderham Centre for Industrial Learning Z Jr 1600 1400 Si 1200 mS' || Austenite y phase ||| | carbon dissolved Sis O 1000 '||SS| in fee iron IP Figure 4. The austenite region !§: of the iron-iron carbide diagram. 3 G 800 Austenite dissolves up to about i 2% carbon. E e 600 400 200 1 1 I I 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 Weight % carbon A eutectic is formed at 4.3 % carbon. Liquid of this composition on solidification at 1147°C transforms to a mixture of two phases (austenite + cementite). This region is important when discussing cast irons but is not relevant to steels. The region in which iron is face-centred cubic, identified in Austenite v I4 Fig. 4, dissolves up to 2% carbon. This phase is termed austenite or gamma phase. With no carbon present it begins at 910°C on heating but with 0.8% carbon it starts at 723°C. When a steel is heated into the austenite region all carbon and most other com­ pounds dissolve to form a single phase. m The region shown in Fig. 5 where carbon is dissolved in bcc rite: a iron is very narrow, extending to only 0.02% carbon at 723°C. This phase is termed ferrite or alpha phase. Although the carbon content of ferrite is very low other elements may dissolve appreciably in it so ferrite cannot be considered as "pure iron". 1600 1400 1200 o 1000 © 2 800 Figure 5. The ferrite region of the I iron-iron carbide diagram. E 600 © Ferrite a phase carbon dissolved in 400 bcc iron. Maximum solubility only 0.02% 5 Jr 200 1 at 723°C X 1 1 JL X X 1 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 Weight % carbon 5 I ri zj Gooderham Centre for Industrial Learning «j zj zj 0.8% C steel cooled from the austenite region zJ ZJ zJ Figure 6. The eutectoid part of the iron-iron carbide 723°C ZJ phase diagram. a> ZJ =3 CC ferrite \ cementite ZJ 03 peariite Fe3c ZJ Q. E 03 a eutectoid mixture ot ferrite and cementite zJ ZJ ZJ 0.4 0.8 1.2 J. J ZJ Carbon (wt%) 2J Pearii At 0.8% carbon and 723°C a eutectoid is formed. This is similar to a eutectic but occurs completely in the solid. As Fig.6 illustrates a.P W steel of eutectoid composition will, when cooled from the austenite 2J region, transform at 723°C to a mixture of two phases, ferrite and ZJ cementite (Fe3C). This eutectoid mixture is called peariite. Fig. 7 shows an example. Note that peariite is not a single phase itself but ZJ rather a constituent made of two phases. Peariite is but one of many ZJ constituents given different names but made of the same two phases ZJ ferrite and iron carbide. Cementite itself is extremely hard—about ZJ 1150 Hv—but when mixed with the soft ferrite layers in peariite the ZJ average hardness of peariite is considerably less. ZJ i*r* ZJ /- '2 xi'-V- AW tJ-v-V Figure 7. A typical example of ZJ peariite showing the two phases sSi ferrite and cementite (iron carbide) in layers. Light areas are ferrite ZJ and dark areas are cementite. ZJ ZJ ZJ 6 ZJ ZJ J* J* ~j* Gooderham Centre for Industrial Learning z J- The region at the top left portion of the diagram enlarged in PerltectfC Fig. 8 is where iron reverts back to the bcc structure. Here again the solubility for carbon is low, only 0.1 wt % at 1493°C. This phase is known as delta ferrite. The part of the diagram at 0.16% carbon having the appearance of an inverted eutectoid is called aperitectic. At this point a two phase mixture of delta ferrite and liquid transforms on cooling to a single phase solid of austenite. This portion of the diagram will not be discussed in detail, but the general shape of the phase diagram in this region should be recognized since it has been invoked to explain various hot cracking phenomena in welding. 1600 LIQUID L+5 1500 Figure 8. Peritectic region of o the iron-iron carbide diagram. 03 L+T 44 "5 Q3 q. E 1400 Delta ferrite 8 phase I 03 AUSTENITE Y 1 1 1 i J 0.1 0.2 0.3 0.4 0.5 0.6 0.7 Carbon (wt%) The part of the diagram that is of most interest to a study of steels is the region containing the eutectoid in the bottom left comer of the iron-iron carbide diagram and we will discuss this in a little more detail. A steel with 0.8 wt% carbon, it will be recalled, transforms on cooling through 723°C to the two phase eutectoid constituent Pearlite growth mm pearlite. In pearlite the two phases fenite and cementite are mixed closely together in fine layers. As the fenite contains very little carbon while the cementite has 6.7 %, carbon atoms must diffuse to the growing cementite plates as Fig. 9 shows. The distance they can ^9 diffuse—and hence the spacing of the plates—depends on how fast the pearlite is growing. A fast growth rate means less time for diffusion and a finer pearlite. 7 TV J I :i Gooderham Centre for Industrial Learning :i zj Direction of pearlite growth Fast growth rate Cementite plate Ferrite plate Austenite Austenite Carbon atoms must With a fast growth rate diffuse to cementite there is less time for plates away from the ferrite carbon diffusion and atoms cannot travel far. Plates 'I must therefore be closer together giving a fine pearlite -J ZJ Figure 9. The growth of pearlite is diffusion controlled. The morphology therefore depends on the -J growth rate which in turn depends on the cooling rate of the steel. ZJ ZJ Proeutect* If the steel has less than 0.8% carbon (termed hypo-eutectoid) ferrite ferrite will be formed first from the austenite. The example in Fig. 10 shows a steel of 0.4% carbon. This ferrite is called proeutectoid 'J ferrite. As transformation continues and the temperature drops the Zl remaining austenite becomes richer in carbon. At 723°C the steel comprises ferrite and remaining austenite with 0.8% carbon. The U. -J latter then transforms to pearlite so the final structure of the steel is a mixture of pro-eutectoid ferrite and pearlite. The amounts of pro- ZJ ZJ ©J 1 ZJ 816 ZJ ZJ 723 4 o Steel of 0.4% carbon cools from austenite region O © On reaching about 816°C some ferrite starts to form ZJ ® Remaining austenite becomes richer in carbon ZJ -J i -*- t. 1. '* hi rJ ft£&?s Figure 17. Photomicrograph of martensite. St&t.'.Srf t ’.-fi ii» 4 rJ ‘S>? !< rJ *- *?A - ZJ W > f*. V rJ. J-.-J ZJ -J.4 ZJ ' 5* f- ; I ZJ zj This constituent is known as martensite. Under the microscope zj Martensrte martensite has the appearance (Fig. 17) of a mass of needles. Marten­ site can be a very hard and brittle constituent when it contains ZJ appreciable carbon. The hardness depends almost exclusively on the - ZJ carbon content with other elements having litde effect. Fig. 18 shows zj the relationship between martensite hardness and the carbon content. zJ zJ zj ZJ looor zj ZJ 800 - Alloy v!v ZJ Figure 18. Graph showing the martensites § Retained ZJ effect of carbon content on the hardness of martensite. Other X > 600 - austenite ZJ alloy elements have only a small to CO zj \ effect on martensite hardness. a> P 400 - >:j 8S Plain carbon martensite ZJ ns X 200 ZJ ZJ 1 J. 1 1 ZJ 0 0.2 0.4 0.6 0.8 1.0 ZJ Weight % carbon ZJ ZJ ZJ ZJ 14 ^J ZJ ja Gooderham Centre for Industrial Learning > HA2 HARDNESS (Hv) 475 450 300 275 ;. *! Steel composition C = 0.36% Mn = 0.77% Si =0.24% Figure 19. Examples of four welds of different heat input. The small welds cool rapidly enough to form martensite in the HAZ with a hardness exceeding 450Hv. Illustrated in Fig. 19 are a number of welds in a carbon steel that have cooled at various rates. In each weld the heat affected zone transforms to a microstructure dependent on the cooling rate of that weld. For the small, rapidly cooled welds martensite is formed. For the large, slowly cooled welds the HAZ structure is pearlite. The hardness of the HAZ is much higher in those welds in which martensite is present. Intermediate between a rapid quench that produces martensite and a slow cool producing pearlite other constituents may form, particularly in alloy steels. The most important of these is bainite. :::::: Bainite is still a two phase mixture of ferrite and iron carbide but unlike the cementite plates in pearlite the carbide in bainite is spherical (Fig. 20). Bainite formed above 300°C contains relatively coarse particles of the Fe3C form of iron carbide (cementite) and is termed upper bainite. When formed below 300°C bainite has a much finer structure with the carbides tending to form striations across the ferrite laths. This is lower bainite. The carbides in lower bainite are Fei4 C ( £ carbide). - i; ~r": ^^ “ -s Z. ,'x.^ v * V. s J- v m X; ;< N. Figure 20. Photomicrograph of >c '-Ov.... mm5-v. bainite. The growth rate is too rapid v for carbides to form as plates of cementite as in pearlite, and instead the carbides form as particles. ;-v - i&Zyi W504 : V 3 ‘ Sd* t m. S. '-5*^ / :> g-’-r:-', v.'.-T :u. 15 r Gooderham Centre for Industrial Learning U zj Transformation Since the iron-iron carbide diagram is only valid for very slow @SSSS! diaqrams cooling rates, alternate diagrams for determining the constituents SUM present in a more rapidly cooled steel have been developed. There are two types covering the case of continuous cooling and the case of isothermal transformation where the steel sample is held at a constant temperature until transformation is complete. We will consider the latter type first 910°C Figure 21. Procedure for determining quench ZJ time-temperature-transformation 723°C 3 behaviour. Steel is heated to the austenitic P region, quenched to a specific temperature a> a. E hold ZJ then held at that temperature until transformation is complete. a> zj time zj % carbon ZJ ZJ Imagine a sample of steel heated until fully austenitic then quenched to some temperature below the equilibrium transformation ZJ temperature (Fig. 21). If we hold the steel at this temperature we find zj MNilil mm there is a delay before transformation begins and a further elapse of. zj : time while transformation takes place. The delay depends on the temperature at which the steel is held and we can plot this information zJ on a diagram of temperature against time for a given steel composi­ zj tion. ZJ ZJ An example of such a time-temperature-transformation (TTT) diagram for a carbon steel is shown in Fig. 22. Note that at high ZJ ZJ 800 ZJ Pearlite start ZJ 700 Pearlite finish ZJ ZJ O 600 P ZJ 3 co is 500 Bainite Bainite finish ZJ CD CL start zj i Ms Martensite start zj 400 M Martensite finish Figure 22. Typical TTT or isothermal ZJ f 300 transformation diagram. ZJ ZJ 1 1 I 102 1 1 ZJ 104 105 1 10 103 ZJ Time (seconds) ^J 16 ZJ  Gooderham Centre for Industrial Learning : 4, \ \ 2) (3 \ \ &--k Specimen quickly cooled from austenite region then held at about 650°C 4 © Proeutectoid ferrite begins to form © a: 20 - 1 1 X i J 0.2 0.4 0.6 0.8 1.0 1.2 1.4 Weight % carbon At 0.8% the steel is 100% pearlite and the maximum strength is obtained. Any increases of carbon above 0.8% serve only to produce pro-eutectoid cementite with no further increase in strength. Plain carbon steels of about 0.8% carbon are often specified where maxi­ mum strength is desired, for example rails and piano wire. ss Increasing carbon content, however, reduces the ductility (as on ductility shown in Fig. 38) and this often limits the applications for plain carbon steels. The impact toughness is also lowered and the effect of carbon content on the chaipy transition curve is illustrated in Fig. 39. The strength of a plain carbon steel depends to a certain extent on the coarseness of the pearlite, a very coarse pearlite giving a strength about 20 ksi (137 MPa) less than a fine pearlite. Very fine pearlite can be achieved in an isothermal transformation by selecting a tempera­ ture level with the nose of the TTT diagram where transformation occurs in the shortest time. In the wire industry this heat treatment is known as patenting and is used to impart very high strength and good ductility on the wire enabling it to undergo the drawing operation. 200 0.11%C g; iso 0.20% C Figure 39. Effect of carbon , CL 50 0.80% C TO s: O ^ j* x x 0 j -100 -50 0 50 100 150 200 Temperature °C 31 Kk 1  A Gooderham Centre for Industrial Learning TJ A A A Table 2. Maximum specified carbon levels for various thicknesses of ASTM A516. A Thickness range Carbon content (max) for Grade 70 A 1 in. (25 mm) and under 0.31 -J Over 1 to 2 in. (25 to 50 mm),incl. Over 2 to 4 in. (50 to 100 mm), incl. 0.33 A Over 4 to 8 in. (100 to 200 mm), inci. 0.35 0.35 A Over 8 in. (200 mm) 0.35 A A Under continuous cooling conditions the coarseness of pearlite A ::::: ! ess depends on cooling rate. A thick plate cooling slower produces a coarser structure with lower strength. This is reflected in the specifi­ A m. cations for carbon steels which may have lower specified strength A requirements for thick sections or may allow higher carbon levels to A. compensate. An example for A516 is shown in Table 2. This has important implications for welding since not only does a weld on a A. thick plate cool faster than one on a thin plate, but the thick plate is ZJ likely to have a higher carbon content. These two factors greatly A increase the risk of cracking in thick plates. A. The coarsest pearlite and ferrite structure is achieved in a carbon steel after a full anneal, i.e. heating to fully austenitize at (900°-950°C) A. followed by a very slow cool in a furnace. A full anneal is used when maximum softening is required but the coarse microstructure has very TJ low toughness. Although the ductility in a tensile test is improved, the TJ impact toughness is poor and the fully annealed condition is rarely desirable in a steel for structural use. A full anneal might be specified A prior to machining, particularly when this is followed by a further heat A. treatment A Normalizing The toughness of carbon structural and pressure vessel steels can A be improved by normalizing to reduce the grain size (Fig. 40). This A treatment involves austenitizing at about 900°C followed by an air A A A Conventional A Normalized carbon manganese steel as-rolled Figure 40. The toughness of A. O" Li- carbon steels is improved by normalizing. The diagram A shows a typical distribution of NDT temperatures for as-rolled A i and normalized steels. A -40 -20 0 20 A Nil ductility temperature °C A. 32 A J Gooderham Centre for Industrial Learning r^ cool. The temperature is the lowest practical temperature at which the steel becomes fully austenitized, and the holding time need only be Normalizing sufficient to ensure complete austenitization has occurred. Prolonged temperature heating in fact may cause grain growth which is undesirable. A holding X':-:,. time of 1/2 hour at temperature for carbon steels up to 75 mm is typical although various codes and standards may specify differently. Some of the older steel specifications allow a hot forming opera­ tion carried out by the fabricator to count as the normalize. If the fabricator chooses this option there are two important things to note: First, the steel may be supplied in the as-rolled (un-normalized) state, but the mill test certificates will report tests on coupons normalized in the laboratory. Second, the temperature for hot forming must not significantly exceed the normalizing temperature otherwise grain growth and toughness deterioration could occur. In this regard fine grain steels are more tolerant to higher temperatures than those with coarse grain. Cooling after To prevent coarse pearlite forming, normalizing is followed by an normalizing air cool rather than a furnace cool. An inadequate cooling rate could result in toughness deterioration. If several components are normal­ -* ized together they must be properly stacked to allow good air circula­ tion between them. With thick sections a forced air cool using fans may be specified, but care must be taken to avoid uneven cooling thatresults in distortion. Alternate methods The negative effects of carbon on toughness and weldability have of strengthening led steel makers to employ other strengthening mechanisms. The most recent and important for structural steels are precipitation hardening and grain refinement The first generation of high-strength low-alloy structural steels used carbo-nitride forming elements vanadium and niobium (columbium) as precipitation hardening agents. These elements had the advantage they could be added to semi-killed steels because of their limited tendency to form oxides. The increased strength due to precipitation hardening allows a reduction in carbon content so im­ proving weldability and toughness. Typical of these steels are ASTM A441 and A572. More recent steels have employed grain refinement in addition to precipitation hardening to produce high strength coupled with excel­ lent toughness. Fine grain is produced by thermo-mechanical process­ ing in which the reduction of plate thickness by rolling is integrated -* 33 BL.J  3 ZJ Gooderham Centre for Industrial Learning ZJ zj zj w rJ ZJ ZJ ZJ ZJ Austenite grains deformed during hot rolling ZJ Original austenite grains If rolling is finished at high temperature ZJ grains recrystallize and grow again ZJ ZJ ZJ ZJ ZJ If rolling is finished at lower temperature tendency to regrow is reduced ZJ ZJ ZJ If rolling is finished at low temperatures and recrystallization is retarded by ZJ niobium small grain size will be retained ZJ Figure 41. Using lower finishing temperatures during hot rolling and micro-alloy elements that retard recrystallization can result in steels of very fine grain size. ZJ ZJ with temperature. During hot rolling the austenite grains breakup, but ZJ at high temperatures they recrystallize and grow again (Fig. 41). If ZJ Controlled rolling rolling is performed at lower temperatures (controlled rolling) the tendency to regrow austenite grains is less. The recrystallization of ZJ the austenite grains is further retarded by the presence of microalloy ZJ elements—particularly niobium—and the austenite range is extended ZJ to lower temperatures by alloys such as manganese. ZJ A steel maker may use controlled rolling for the lighter sections of plate and normalizing for thick sections when excellent toughness is required. Several other devices are being developed to further ZJ improve properties among them being on-line accelerated cooling ZJ and direct quenching. zA 34 Jt J j Gooderham Centre for Industrial Learning J0 * Since heating a steel above 400°C causes recovery and recrystal­ lization to begin, any process involving temperatures above this removes the effects of work hardening. Thus castings, hot forged, hot rolled, or any heat treated product will not gain any strength from Effect of work work hardening. But any cold finished product or one subject to hardening subsequent cold working such as bending or forming will have a strength component that results from work hardening. In Fig. 42 we show the increase in strength due to cold deformation while Fig. 43 reveals the accompanying reduction in ductility. Work hardening as a strengthening method is particularly valuable in the production of wires. The wire drawing process itself provides a greater degree of reduction than is achieved with any other method. High strength achieved in low carbon steel wires is particularly valuable in making gas metal arc welding electrodes. GMAW wires are usually pushed by a wire feeder along a conduit to the welding gun, and a high column strength is desirable to prevent kinking. Previous editions of AWS A5.18 "Specification for Carbon S teel Filler Metals for Gas Shielded J* 1000 05 s 800 05 £ 600 CO CD CO o' 1 i 1 ! J ZJ 400 300 200 100 0 ZJ Temperature of weid during cooling °C Figure 53. Diagram showing the fraction of hydrogen remaining in an example weld as it cools. U ZJ When preheat is used very little hydrogen remains by the time the weld cools to room temperature. ZJ ZJ 44 ZJ ZJ Gooderham Centre for Industrial Learning The preheat temperature required depends on the susceptibility of cm formu the HAZ to hydrogen cracking, and much work has been done to find compositional formulae to indicate this. One widely used formula due to Ito and Bessyo is Pcm = C + Si + Mn + Cu + Ni + Cr + Mo + V + 5B 30 20 20 60 20 15 10 You should note that this relates directly to cracking susceptibility in steels of less than about 0.2% carbon and should be distinguished from the hardenability carbon equivalent given earlier. The level of preheat also depends on the original hydrogen level in the weld metal. These two factors—the plate composition and the hydrogen—can be combined in an index of susceptibility equal to 12Pcm + log H where H is the hydrogen determined by the standard IIW method. This index is used in CSA W59-Appendix P and AWS D 1.1-Appen­ dix XI as part of an alternate guide for estimating preheat levels. s relief cracki During thermal stress relief the residual elastic strains are con­ verted to plastic strains by a drop in the material yield point and creep. If these strains are not spread uniformly through the metal but are concentrated in local regions, they may be sufficient to cause crack­ ing. This phenomenon, known as stress relief or re-heat cracking, has been observed in the heat affected zones of welds in particular steels. Cracking is more likely in highly restrained, very heavy sections and where there is a stress concentration such as an unfused root or the toe of a large fillet weld. Precipitation of carbides of secondary hardening elements during stress relief is the primary cause of cracking. Precipitation occurs preferentially within the grain causing the grain interior to strengthen relative to the grain boundary. This shifts the strain onto the grain boundary where cracks may form. Studies of stress relief cracking reveal that composition has a dominant effect, and empirical relations have been established to predict whether a steel is sensitive to this type of defect One formula suggested in the research literature is: Psr= Cr + Cu + 2Mo + 10V + 7Nb + Ti -2 When Pa > 0 stress relief cracks may occur. This formula does not workfor low carbon (1.5%), these steels being resistant to stress relief cracking. 45  cr Gooderham Centre for Industrial Learning zr ' cr WELD METAL er Since weld metal starts off as a small casting, cools rapidly and does not undergo heavy deformation it is not suprising that the properties of weld metal can be quite different from those of the parent plate and heat affected zone. Consequently the optimum composi­ tions of weld metals are rarely exactly the same as the steels they are joining. For example, a thick plate of pressure vessel steel of 0.25% carbon may be joined with a weld metal containing only 0.1% carbon. On the other hand, when alloy steels are selected for specific proper­ ties such as corrosion resistance or high temperature strength, the weld metal will likely have a similar composition to the base metal. The microstructure of weld metal is influenced first by the solidification structure that exists after the metal freezes. U sually long austenite grains extend from the fusion boundary towards the centre of the weld (Fig. 54). They are epitaxial with the grains in the HAZ meaning they grow directly from them. This is an important observa­ tion since the grain size of the weld metal can be influenced by the grain size of the HAZ. In most steels the HAZ grain size at the fusion boundary depends only on the heat input and not on the steel composition so the weld metal grain size is also independent of the steel composition. As we saw earlier, however, some modem titanum steels have very small HAZ grain size, and this is reflected in a smaller weld metal grain size. m i iv Figure 54. Weld metal structure Sp- Wr, m , r- :/> I'ZSC < u showing austenite grains extending i* Z- r. from the fusion boundary where they are epitaxial with the HAZ grains. te2 Wh 4» CV »; mT22 S’-; a, SwSils'*/ WjS. if € W- m- u -... - - ^ V Sypy v * > *>; 46 J j Gooderham Centre for Industrial Learning j t* j X j ' t- * Jyi*-" 0 ’* : ::« 1 j I - “ > * # Figure 55. Proeutectoid ferrite V , 'Vv Vg-A. ; 'veins' formed at the austenite grs.;.V grain boundaries in a C-Mn-Si i weld metal. C5 ‘iT ?r< K\ *g£ % - V' « --V i.- ". J&S In a typical carbon manganese weld metal, austenite grains transform to a mixture of ferrite and carbide. Ferrite may form first x! __ at the grain boundaries of the austenite and give rise to die familiar :&! 'veins' commonly observed in weld metals and visible in Fig. 55. Ferrite may also form within the grain if nucleation sites are available. Small oxide particles provide excellent sites for nucleation and a weld metal having more than about 250 ppm of oxygen present has enough S* oxide particles to nucleate many tiny grains of ferrite within the austenite grain. This intragranular ferrite, shown in Fig. 56, is termed acicular ferrite and is associated with weld metals having good toughness. Ferrite may also be present in the form of spikes or laths Jr growing into the grain from the edge. This form, sometimes termed side plate ferrite, has been associated with poor toughness. After ferrite is formed the remaining austenite becomes richer in carbon and can transform to a number of different structures. The carbon may therefore be present as a carbide, in small islands of martensite or in retained austenite. **r- ~r. 'rr-.. / *.k. Y Figure 56. Acicular ferrite formed within the austenite grain of weld r metal. The very small grain size of ~v\ this intragranular ferrite contributes to good toughness. -r 't r> -5* f rj, -j&F '4k * \ f,\r' y- Ox- F * : -k& F A ffiSijGsefer 47 F Gooderham Centre for Industrial Learning ZJ *i zj The final composition of the weld metal is determined by three mtj factors: ZJ Weld metal ZJ The filler metal composition ZJ composition Dilution from the parent metal Chemical reactions ZJ