SUPPLEMENT TO THE WELDING JOURNAL, MARCH 1976

Sponsored by the American Welding Society and the Welding Research Council ( "Ru l

Deoxidation Practice for Mild Steel Weld Metal

Quantitative metallography explains how weld composi­tion controls ductile and brittle fracture resistance

BY D. J. WIDGERY

ABSTRACT. The effects of Mn, Si, Ti and Al on inclusion contents, micro-structure, strength and toughness of as-deposited, single run C 0 2 welds made at 1.6 kJ /mm (40 kJ/ in) have been studied. A statistically planned series of experimental electrodes was used and toughness was assessed with the crack opening displacement (COD) test.

Resistance to cleavage and to fibrous fracture were differently af­fected by microstructure and com­position. Good cleavage resistance was achieved with acicular micro-structures, which manganese con­tents up to at least 1.6% in the de­posit helped to promote, and with low titanium levels to avoid dispersion harden ing. Good cleavage resis­tance was also achieved in a soft, homogeneous structure with a yield strength of 390 N/mm2 (56 ksi) but this might be difficult to guarantee in commercial practice.

Resistance to ductile fracture by microvoid coalescence was related to inclusion content, yield strength and strain hardening behavior of the weld. Low inclusion contents were favored by weld Mn/Si ratios of 4-5 but it may be more profitable to aim for soft microstructures if good tearing resis­tance under highly triaxial stress is required.

D. J. WIDGERY is associated with the Welding Institute, Research Laboratory, Abington Hall, Abington, Cambridge CBI 6AL, England.

Paper was presented at the AWS 56th Annual Meeting held in Cleveland, Ohio, April 21-25, 1975

Introduction The important influence which weld

deoxidation practice can have on the toughness of the deposit has long been recognized, but only in recent years have attempts been made to quanti fy the var ious mechanisms through which this influence is ex­erted (Refs. 1-4). A recent program at the Welding Institute set out to inves­tigate how deoxidation practice af­fected inclusion contents and micro-structure and thus how resistance to fracture in the ductile and brittle modes was control led. The results of this investigation are reported here.

Some oxidation almost invariably accompanies the welding of steel. In the slag shielded processes, the activ­ity of oxygen in the slag is not negligi­ble and perfect shielding from atmo­spheric oxidation difficult to achieve. The GMAW process generally relies on the use of a mildly oxidizing gas to allow some oxidation of the work-piece and stabilize the arc. Since ox­idation cannot be avoided, deoxida­tion procedures must be devised to minimize its adverse effects on the in­tegrity and performance of the weld and the optimization of these proce­dures was the object of this program. The solid wire COz (GMAW) process was chosen for this because it per­mits a moderate degree of oxidation, is readily mechanized and is uncom­plicated by the presence of a flux. Further incentives to the use of this process were its increasing commer­cial acceptance and the fact that many reactions in the weld pool are common to welds made with basic manual metal-arc electrodes or with

some submerged arc fluxes in which carbonates are incorporated.

Aims of Deoxidation

In C 0 2 welding, droplets transfer­ring between the wire and the pool pick up oxygen as

C 0 2 ± ; C O + [0 ] (D The dissolved oxygen can react with carbon in the steel to form CO,

[ C ] . + [0 ] Fe Fe

:CO (2)

which gives rise to porosity if its equi­l ibrium pressure p c o exceeds 1 atmo­sphere. To prevent this, it is neces­sary to limit the product [C] • [O] in the mass action equation

= constant. [C] . [O]

(3)

In practice, except in welding pure iron, the carbon content of the pool will always be such that the oxygen activity must be reduced by adding metallic deoxidants if porous welds are to be avoided (Ref. 5).

If the first aim of deoxidation is to prevent porosity, the second is to minimize the amount of oxygen re­maining in the deposit after sol id­ification. Since there is little evidence to contradict Sims' assertion (Ref. 6) that "for all practical purposes, the solubility of oxygen [in solid steel] is so low that it may be neglected," this objective may be restated as the min­imization of the amount of oxide inclu­sions in the weld.

In steelmaking practice, control of inclusion contents is generally achiev-

W E L D I N G R E S E A R C H S U P P L E M E N T ! 57-s

Table 1

Elec­trode

A B C D E

F G H J1 J2

K L M N 0

P Q R S T

U V

w X Y

z

— Chemical Analysis (wt.

C

0.05 0.09 0.10 0.09 0.09

0.07 0.11 0.11 0.08 0.08

0.08 0.06 0.08 0.09 0.08

0.08 0.09 0.08 0.09 0.07

0.07 0.08 0.10 0.07 0.09 0.07

Mn

1.14 1.11 1.63 2.15 1.10

1.42 1.10 1.53 1.51 1.42

1.95 0.98 1.01 1.02 1.92

0.98 1.98 1.62 2.27 0.63

1.42 1.47 1.46 1.51 1.41 1.38

%) of Experimental

Si

0.51 0.92 0.50 1.05 1.40

0.51 0.46 0.90 0.78 0.77

0.43 1.05 0.43 0.43 1.11

1.07 0.45 0.91 0.78 0.74

1.51 <0.01

0.76 0.79 0.78 0.70

Ti

0.19 0.22 0.36 0.14 0.46

0.25 0.44 0.40 0.19 0.14

0.09 0.08 0.23 0.09 0.23

0.24 0.20 0.08 0.14 0.15

0.14 0.17 0.32 0.01 0.15 0.16

Electrodes

Al

0.04 0.13 0.055 0.055 0.075

0.20 0.12 0.10 0.15 0.12

0.08 0.07 0.06 0.21 0.08

0.19 0.14 0.15 0.13 0.14

0.14 0.16 0.14 0.13 0.24 0.01

s 0.011 0.009 0.012 0.017 0.012

0.012 0.010 0.012 0.009 0.010

0.008 0.010 0.009 0.009 0.009

0.008 0.008 0.010 0.008 0.011

0.010 0.007 0.010 0 0 1 3 0.007 0.011

P

0.008 0.004 0.006 0.006 0.006

0.006 0.003 0.004 0.017 0.018

0.018 0.015 0.023 0.020 0.021

0.017 0.022 0.020 0.018 0.016

0.014 0.015 0.018 0.013 0.011 0.018

j \ ^ \/-—3mm mol md.

—1 5mm mot md.

Fig. 1 — Preparation of experimental welds: (a) plate grooved, (b) buttered and regrooved, (c) experimental weld com­pleted

ed by ensuring a low oxygen solubil­ity in the molten steel and relying on the removal of most of the deoxida­tion products before solidification. Much work has been devoted to eval­uating oxygen solubility in the liquid as a function of deoxidant additions and the oxygen content of the solid steel can closely approach the level thus predicted.

In welds, on the other hand, the final oxygen content is usually much

greater than the solubility in the melt: for example, the presence of 0.5% Mn with 0.5% Si reduces the oxygen sol­ubil i ty in mol ten i ron at 1600 C (2912 F) to -0 .007% (Ref. 7), where­as oxygen contents an order of mag­nitude higher may be encountered in welds with comparable or greater deoxidant contents (Ref. 2). Clearly, the short t ime available for the re­moval of deoxidation products from the pool means that the kinetics of coalescence and flotation have more influence on the final weld oxygen content than the exact level of oxygen solubility on the melt. This explains the limited success of thermodynam-ic arguments in predicting weld ox­ygen levels and shows the need to concentrate further studies on those properties of the deoxidation prod­ucts themselves which affect their rate of removal from the weld pool.

Since the elements used to deox­idize mild steel weld metals also con­stitute the principal, often the only alloying additions, a third aim of deoxidation practice must be to en­sure that the deoxidants which remain in solution after all the oxygen is chemically combined will promote the formation of desirable microstruc­tures and will not induce embritt le­ment through solid solution or dis­persion hardening. It is through such effects on the metal matrix, rather than in their str ict ly deoxid iz ing capacity, that the reactive elements in welding consumables exert most in­fluence on the cleavage resistance of the deposit, so these effects were of

particular concern in the present in­vestigation.

Experimental Program

The work was planned to study the effects of four deoxidants, Mn, Si, Ti and Al. The program centered on the measurement of fracture toughness, for which theoretical and practical considerations led to the adoption of the crack opening d isp lacement (COD) test (Ref. 8). Welding proce­dures were designed to provide COD testpieces with the notch in a single weld run. To provide the link be­tween the composit ion and mechan­ical properties of the deposit, quan­titative measurements of the micro-structure and inclusion contents were undertaken. The experimental plan was based on the use of multiple regression analysis to interpret the results.

Experimental Electrodes

A series of 26 vacuum melted experimental electrodes was pro­duced with varying amounts of Mn, Si, Ti and Al. The experimental design had to take account of the likelihood that the effects of the deoxidants would be nonlinear, as shown by Sa-kaki (Ref. 9): in general, small addi­tions are taken up in the reaction with oxygen while successive larger addi­t ions have predominant ly mic ro-structural effects. A central, rotat-able, composite design was adopted since these were developed as an economical means of investigation nonlinear behavior (Ref. 10) and have previously proved valuable in re­search on weld metals (Ref. 11). The electrode composit ions produced are shown in Table 1. Electrode J was duplicated by two manufacturers and 5 welds were made with the two filler metals as a check on reproducibil ity. The batch of electrodes constituting the first half-factorial did not conform well to the plan, but the experience gained in their production allowed the remainder to be made more accu­rately and the series as a whole was satisfactory in respect of range and orthogonality. Electrodes were drawn to 1.2 mm (3/64 in.) diam without copper coat ing. Two commerc ia l electrodes, one containing 0.48% Al, were also included for comparison.

Welding Procedure

One of the advantages of the COD testing technique is that since frac­ture initiation is controlled only by the material close to the notch tip, it is possible to sample predetermined re­gions of a single weld run rather than, as in the Charpy test, a composite of all the structures through which the fracture passes as it crosses the specimen. Previous work had strong­ly indicated that cleavage cracks ini-

58-S I M A R C H 1 9 7 6

tiated more readily in as-deposited regions than in reheated regions of multipass welds (Ref. 12), so a proce­dure was devised to produce COD specimens with the notch tip along the centerline of a single, as-depos­ited bead. A bead-in-groove spec­imen was found suitable for this pur­pose, but in order to achieve a wide range of deposit composit ions with­out changing the base material it was necessary to reduce the effective di lu­tion by means of a buttering tech­nique. A groove somewhat larger than that needed for the test bead was fi l l­ed with several runs of the experi­mental weld metal and the groove for the test bead was machined into this deposit, as shown in Fig. 1. By this means the dilution of the test bead with the mild steel base metal was re­duced below 10%.

The test bead was deposited at 30 V, 300 A, 330 mm/min (13 in./min), giving spray transfer with an arc energy of 1.6 kJ /mm (40 kJ/in.), and the C0 2 flow rate was 0.25 l/s (30 cu. f t /h). A preheat temperature of 100 C (212 F) was used throughout. To help in interpreting weld microstructures, thermal analysis was carried out on selected welds by inserting a ther­mocouple into the molten pool and using a derivative processor to plot T and dT/dt as the weld cooled.

Chemical Analysis

All welds were analyzed spectro-graphically using millings from the final weld run, remelted under argon. In addit ion, total nitrogen contents were measured using the conven­tional indophenol blue finish method and oxygen contents initially by the vacuum fusion method. When diff i­culties arose in reconciling these ox­ygen measurements with the results of inclusion volume fraction deter­minations, further oxygen analyses were carried out using fast neutron activation analysis (Ref. 13).

mm (0.005 in.) slitting wheel was ex­tended in fatigue to a total depth of about 6 mm (0.24 in.), so that an ap­proximately square ligament 22 X 22 mm (0.87 x 0.87 in.) remained under the notch. Testing was carried out in 3-point bending over a span of 127 mm (5 in.) and COD was monitored using a double cantilever beam clip gage mounted on knife edges at the specimen surface. Readings were converted to crack tip COD, <5, using the relationship

V„

& = -1 + a+z

r(W-a)

(4)

where Vg is the knife edge displace­ment, a the total crack depth, z the knife edge thickness, W the spec­imen width and r the rotational factor, taken here as 1/3 in accordance with the Bri t ish Standards Inst i tut ion's Draft for Development on COD testing (Ref. 14).

Tensile Testing

Longitudinal all-weld-metal tensile testpieces were extracted from each weld. These had a gage diameter of 6.41 mm (0.253 in.) and a parallel and gage length of 22.7 mm (0.894 in.). Testing was carried out at room tem­perature. In addition to measure­

ments of yield or 0.2% proof stress, tens i le s t r eng th , e l onga t i on and reduction of area, a simple computer program was used to evaluate true stress, a T , and true strain, e. The strain hardening exponent , n, as defined by the Holloman equation (Ref. 15)

cr = ke (5)

was determined as the slope of the log a vs log E plot over the range of strain between e = 0.025 and the maximum uniform strain.

Metallography

The complexity of weld metal microstructures confronts the inves­tigator with a choice of classifying them according to a simple system, permitting measurements of good reproducibil ity, or of devising a more complex system which may describe the transformation products more accurately in a qualitative sense at the

Fdtigue c / w * 15mm * « p '5mm &ep 3/ryn

X 3rmil\ S^r iy !± ,. 5mm.mot md.orPmm

Fig. 2 — Geometry ot COD specimen

W^b~' r s , ! , , <*-^~> -fT\ V1 S p s f 1

COD Testing

In addition to the practical advan­tage that the COD test gives in sam­pling well defined microstructural re­gions, its theoretical basis, which allows it to be used for the prediction of critical or allowable defect sizes in real fabrications, has led to its in­creased use in specifications for crit­ical applications such as offshore structures. A third benefit in using the test is that since it measures resis­tance only to fracture initiation and not to propagation, which may in­volve different mechanisms, metal­lurgical interpretation of the results tends to be simpler than when the Charpy test is used.

In the present work, specimens were produced as shown in Fig. 2 with the notch parallel to the welding direction. A slot produced with a 0.13

- ' . ' • > . . , IP •.4-//.^:

. i J U

JM : iv:v--'X;^,»-:~^:-•-•-••'«#'"' "Jf:W^ ,

Fig. 3 — Microstructure classification used in point counting, x200, reduced 18%

WELDING RESEARCH SUPPLEMENT! 59-s

Carbon replica

Fig. 4 — Mechanism ot inclusion extraction tor counting in the electron microscope, show­ing how inclusions of diameter x are extracted from a volume of depth x

expense of greater difficulty in achiev­ing adequate reproducibil ity in quan­titative work. For example, the former approach might rely on light micro­scopy while the latter would use the electron microscope. The decision was made hereto begin by using light microscopy so that large numbers of measurements could be made quick­ly and with minimal sampling errors. The measurements were backed by electron microscope studies on both replicas and thin foils, but it was con­cluded that quantitative data need not be sought from these if it was possi­b le to co r re la te the mechan i ca l proper t ies of the welds with the results of light metallography, as was found to be the case.

A Swift point counter* was used for quant i tat ive microst ructura l mea­surements. This moves the specimen under the microscope in a predeter­mined grid pattern, the investigator identifying the structure and oper­ating the appropriate counter at each grid point. Transverse weld sections were examined at a magnification of x200 and 100 microstructural iden­tifications were made at 0.05 mm (0.002 in.) intervals along each of 10 lines, parallel to the plate surface and 0.2 mm (0.008 in.) apart, starting 1 mm (0.04 in.) below the plate surface. Thus 1,000 counts were made in a 5 X 2 mm (0.2 X 0.08 in.) area in the center of the weld, corresponding to the notch pos i t ion in the COD specimens.

The s y s t e m of c l a s s i f i c a t i o n adopted is shown in Fig. 3, and repre­sents an attempt to reconcile the re­quirements of easy recognition with an accurate reflection of the trans­fo rmat ion process. To avoid the semantic confusion apparent in the literature on weld microstructures, the products were initially identified only by numbers: later, it was possible to ascribe to them places in the overall transformation scheme.

Inclusion Counting

Gloor et al showed in 1963 (Ref. 16) that the majority of weld metal inclu­sions are too small, typically <1 nm,

"James Swift and Son Limited, Basing-stoke, England.

to be resolved in the light micro­scope and most subsequent inves­tigators have relied on transmission or scanning electron microscopy (TEM or SEM) for inclusion measure­ments. However, while either method can provide a straightforward means of measuring the size distribution of inclusions on the surface of a frac­ture produced by microvoid coales­cence, only indirect methods are available to find the size and spatial distribution of inclusions in the matrix. These methods are based on two models of specimen preparation, the flat surface model and the deep etch model.

In materials polished flat for op­tical metallography the area fraction AA of any second phase on the sur­face may be equated to the volume fraction V v of the phase in the matrix. Against this ease of volume fraction measurement must be set the need to allow for sectioning effects in con­verting the observed distribution of inclusion section sizes to a volume size distribution. Procedures for this, based on a method proposed by Scheil (Ref. 17), assume, reasonably in the case of weld metals, that the in­clusions are spherical. Although these procedures are somewhat cumber­some and prone to cumulative errors, a more important objection to the use of the flat surface model is the diff i­culty of f inding an imaging mechan­ism which the model accurately represents. At first sight, it might appear that the condit ions of the model would be fulfilled by the use of polished specimens in the SEM and trials were undertaken with this in­strument.

As a "bench mark" weld, a manual metal-arc deposit made with rutile covered electrodes and containing 0.079% O (vacuum fusion determina­tion) and 0.022% S was selected. These figures could be converted to an estimated inclusion volume frac­tion with reasonable confidence by assuming that the deoxidation prod­ucts were the same as those an­alyzed by Gloor et al (Ref. 16) and that the sulfur was present as MnS. This led to an estimated inclusion volume fraction of 0.53% (Ref. 18). Since the Al203 and T i0 2 contents of the inclu­sions were expected to be low, any

errors in oxygen measurement arising from nonreduction of these oxides should also have been low.

When a mechanically polished sur­face of this weld was photographed in the SEM, the area fraction of inclu­sions was found to be 1.6 % or three times the calculated volume fraction. A similar result can be derived f rom data presented by Ruckdeschel (Ref. 19) for SEM inclusion measure­ments. Investigation showed that at the accelerating potential of 30 kV, the penetration of the electron beam was sufficient to image some inclu­sions which lay close beneath the pol­ished surface and to increase the ap­parent sectioning diameter of others which just broke the surface. With the instrument then available at the Weld­ing Institute, attempts to reduce beam penetration by lowering the accel­erating potential were frustrated by a simultaneous decline in resolution. Newer instruments may overcome this difficulty, as shown by an exam­ination of results recently produced by Hannerz and Lowery (Ref. 20), but in the meantime an alternative ap­proach has been developed.

A technique which has been used for particle size measurement in the TEM relies on etching or electropol-ishing the metal matrix so deeply that all sectioned inclusions are removed and those that remain are always seen at their true diameter. These may be extracted on replicas, Fig. 4. If it is assumed that all particles which touch the plane of polish are ex­tracted, then particles of any given diameter x are sampled f rom a volume of depth x and the number of particles of this diameter per unit area of replica, Ns(x), is related to the number per unit volume of the matrix, Nv(x), by

N„(x) = (1 /xN s (x ) (6)

This is the basis of the analysis pro­posed by Ashby and Ebeling (Ref. 21) and adopted in recent work at the Welding Institute (Refs. 3, 22). The use of an electropolishing technique found to give an acceptably flat sur­face, together with shadowed repli­cas to show nonextracted inclusions, al lows vo lume f ract ion measure­ments to be made without recourse to the doub le repl ica techn ique re­ported elsewhere (Ref. 23). The method readily yields all the impor­tant parameters describing the size and spatial distributions of the par­ticles using only simple computa­tions, and the calculated inclusion volume fractions are in acceptable agreement with those estimated from the weld chemistry. Although the method used by Gloor et al (Ref. 16) also produced credible results, this was of much greater analytical com­plexity, being based on a combina­tion of the deep etch and flat surface models.

60-s | M A R C H 1 9 7 6

In detail, the technique used in the present work was as follows: after rough polishing' on silicon carbide papers to 600 grade, electropolish-ing was carried out in a 7% perchlo­ric acid solution containing 2 I meth­anol, 60 ml ether, 25 ml distilled water and 150 ml perchloric acid. The pol­ishing potential was about 43 V, vary­ing by a few volts for opt imum results on some specimens, the current den­sity was about 10 m A / m m 2 and the polishing time 30 s. The specimen was then etched for 2 s in 2% nital. Carbon was deposited on the surface at an angle of 45 deg and scribed into 2.5 mm (0.1 in.) squares. Unwanted areas were lacquered off and the replicas were removed electrolyt-ically in 10% nital.

A typical micrograph of a replica prepared in this way is shown in Fig. 5. Following Stumpf and Sellars (Ref. 23), unshadowed inclusions, X, were considered to have floated into the replica f rom some other part of the spec imen and were not counted . Shadowing also revealed the sites of lost inclusions and allowed extrac­tion efficiencies to be calculated: these generally exceeded 95%. The magnification of the electron micro­scope was calibrated for each set of micrographs, which were enlarged to

an exact magnification of x5,000 on the print. Inclusion sizes were mea­sured with a specially made vernier device and recorded on punched tape. A simple computer program was then used to evaluate the fre­quency of the observed size distr ibu­tion, the frequency and percentage frequency of the volume size dis­tribution and the inclusion volume fraction, calculated after Ashby and Ebeling (Ref. 21) as

V v = ( T T / 6 ) S " X 3 N V (x) r)x =

> / 6 ) 2 " x 2 N s ( x ) 5 x (7)

Experimental Results

Thirty-one welds were made and their chemical analyses are given in Table 2. Sulfur and phosphorus con­tents were less than 0.02% in all the welds made with experimental elec­trodes. Carbon contents averaged 0.09% and showed a very low degree of correlation with the levels of metal­lic deoxidants, though in the single case of weld V a low carbon content of 0.03% was associated with a low total deoxidant content: this weld was too porous to be tensile tested or COD tested.

COD Tests

The behavior of a material in the COD test is most usefully character­ized by the critical value, 8C, of COD at which the specimen breaks. It is pos­s ib le , however , at t e m p e r a t u r e s above the ducti le/brit t le transition, for a stable ductile tear to propagate so far through the specimen that the load passes through a maximum while the specimen remains unbroken. In such cases, it is conventional to measure the COD at a point, 6m, where the load first attains its maximum value and to use this in plotting the transition curve, Fig. 6.

However, the use of Sm values in calculating critical defect sizes in real structures has no theoretical basis and moreover, in the present welds a stable crack had often propagated out of the weld metal and into the base metal before maximum load was attained. The results in such a case simply indicate that at the temper­ature concerned, cracks were un­likely to become unstable at low COD values, such as might be developed at very small defects under service loading.

The main concern of fabricators is to avoid materials which are sensitive to small defects, which for mild steel

Table 2 — Chemical Analysis (wt %) of Weld Deposits Made with Experimental and Two Commercial Electrodes

Weld Mn Ti Al -.(a) • • (b )

A B C D E

F G H J1 J2

J1R J2R J2RR K L

M N O P Q

R S T U V

w X Y

z Com. 1 Com. 2

(a) Vacuum

0.10 0.11 0.11 0.09 0.10

0.08 0.09 0.11 0.09 0.07

0.08 0.09 0.08 0.07 0.05

0.07 0.08 0.08 0.09 0.10

0.08 0.09 0.07 0.08 0.03

0.09 0.08 0.08 0.08 0.09 0.10

fusion (b) activation analysis

0.77 1.33 1.10 1.56 0.82

0.96 0.74 1.11 1.13 1.05

1.10 1.06 1.00 1.13 0.72

0.53 0.73 1.43 0.77 1.46

1.27 1.60 0.45 1.06 0.68

1.06 1.12 1.15 0.93 1.09 0.84

0.27 0.54 0.26 0.58 0.61

0.23 0.28 0.49 0.60 0.57

0.50 0.48 0.53 0.17 0.86

0.16 0.26 0.92 0.92 0.31

0.76 0.57 0.52 1.20

<0.01

0.58 0.58 0.66 0.45 0.58 0.28

0.04 0.04 0.07 0.04 0.12

0.06 0.10 0.11 0.07 0.05

0.06 0.05 0.05 0.02 0.03

0.04 0.03 0.08 0.09 0.08

0.03 0.04 0.05 0.05

<0.01

0.11 <0.01

0.06 0.03

<0.01 <0 .01

0.008 0.013 0.013 0.018 0.025

0.046 0.023 0.024 0.06 0.04

0.06 0.05 0.04 0.02 0.03

0.014 0.06 0.03 0.065 0.07

0.055 0.045 0.045 0.045

<0.005

0.05 0.04 0.12

<0.005 <0.005

0.19

0.017 0.017 0.012 0.013 0.014

0.013 0.011 0.013 0.010 0.012

0.007 0.011 0.013 0.008 0.007

0.009 0.007 0.009 0.011 0.008

0.008 0.008 0.011 0.012 0.008

0.010 0.010 0.010 0.010 0.011 0.013

0.018 0.018 0.012 0.015 0.012

0.013 0.014 0.010 0.016 0.016

0.018 0.015 0.017 0.012 0.007

0.015 0.011 0.013 0.014 0.013

0.013 0.017 0.014 0.013 0.014

0.014 0.020 0.015 0.014 0.021 0.015

0.010 0.010 0.010 0.011 0.007

0.012 0.009 0.008 0.004 0.012

0.009 0.007 0.007 0.010 0.003

0.004 0.006 0.010 0.008 0.006

0.005 0.014 0.009 0.009 0.004

0.007 0.009 0.005 0.016 0.006 0.011

0.055 0.055 0.048 0.055 0.047

0.067 0.058 0.057 0.060 0.056

0.054 0.063 0.063 0.064 0.063

0.048 0.053 0.045 0.048 0.042

0.047 0.040 0.052 0.043 0.080

0.049 0.041 0.046 0.029 0.034 0.036

— — — — —

0.088

— 0.094 0.097

— — — —

0.12

— 0.10 0.095 0.10

— — —

0.10 0.12 0.10

— — — — — — —

W E L D I N G R E S E A R C H S U P P L E M E N T ! 61-S

V

• •

• I

1 ) • I •

7 &

•

Table 3 — 0.1 mm COD Transition Temperatures of Experimental and Commercial Weld Metals <">

Weld To,C

A B C D E F G H J1 J2

- 3 9 - 5 1 - 2 8 - 9 4 - 1 - 3 8 + 16 - 1 9 - 7 6 - 4 9

Weld To, C Weld T„, C

J1R J2R J2RR K L Wl N 0 P Q

- 63 - 47 - 46 - 94 - 30 - 1 4 1 - 67 - 75 - 25 - 54

R S T U W X Y Z Com. 1 Com. 2

- 88 - 1 1 3 - 11 - 31 - 18

- 79 - 90 - 77 - 65 - 55

(a) Weld V was too porous to test

Fig. 5 — Transmission electron micro­graph of inclusions on a replica, as used in inclusion measurements. Unshadowed in­clusions, x, are not counted. x5,000, reduced 46%

0-7

0-6

05

1

0 3

0-2

or

• o

•

•

COO

COO

a at

fracture maximum

• • " • " " ^

toad

, . ^

- ^ - «

f*

•

I A

. -100 -SO -60 -AO -20 0 .1

remperattire. 'C

Fig. 6 — COD test results for weld J1

weld metals generally means that in­terest is concentrated at the lower end of the transition range. To compare this aspect of the behavior of the ex­perimental welds, a transition tem­perature Tc corresponding to 5c = 0.1 mm (0.004 in.) is tabulated for all the welds in Table 3. This may be re­garded as a measure of their cleav­age resistance.

The results of a regression anal­ysis of Tc on weld composit ion are shown in Table 4. These results may be presented graphically by omitting each term successively f rom the regression equation and using the de-completed equation to predict Tc. The error in prediction, ATC, is then re­lated to the level of the variable omitted. Thus Fig. 7 effectively shows the variation in Tc with Mn content, other elements being corrected to a constant level.

An unsatisfactory feature of Fig. 7 is the presence of a single point, indi­cating the greatest overestimate of Tc, but which by appearing at the low Mn end of the scale seems to reverse the general trend to lower transition tem­peratures with increasing Mn. This

point represents weld M. If similar plots are made to show the effects of Si, Ti and Al, the results for weld M again accord poorly with the rest of the data. For reasons which will appear below, it was later considered justifiable to repeat the regression analysis without including the result for weld M and the results of this are also shown in Table 4.

No very convenient measure of tearing resistance emerged from the COD tests. Examination of fractured COD specimens in the SEM showed that small amounts of ductile tearing preceded cleavage in all specimens broken in the transition range, the COD at the initiation of tearing, <5>, being typically of the order of 0.1 mm (0.004 in.), rising to 0.25 mm (0.010 in.) for weld M only. Precise defini­tion of Si, was hindered by the fact that in the early stages of tearing only a few isolated microvoids could be seen at high magnification, as compared with the fibrous thumbnail visible to the naked eye at higher values of <5C. Potential drop (Ref. 24) and stress wave emission (Ref. 25) techniques did not reliably detect the appear­ance of the first microvoids. Since h< is not a basis for practical design cr i ­teria, the ductile behavior of the ex­perimental welds was mainly inferred from tensile test results.

Tensile Tests

Tensile test results are given in Table 5. The range of yield or 0.2% proof strengths was f rom 388 to 582 N/mm2 and of tensile strengths from 486 to 734 N/mm2 . Regression anal­ysis showed that Mn, Si and Ti had a significant effect on both yield and tensile strengths, Table 6: weld M was included in this analysis. Carbon did not have a detectable effect in yield strength and its effect on tensile strength was not significant at the 95% probability level, but it seems possible that this reflects the small variation in -^thpr than the lack of real effect : . „ ^ C I . I -cients for Al were large and positive but not quite significant at the 95%

level unless results for the commer­cial electrodes were included. Aver­age RA values varied between 50 and 70%: individual results below 50% were always associated with a small pore. Tensile ductility was not d i ­rectly correlated with composit ion in any simple way.

Microstructure

Table 7 gives the results of classi­fying the weld microstructures ac­cording to the simple scheme of Fig. 3. Thin foil and replicas studied in the TEM and thermal analysis of the welds as they cooled allowed the structures seen to be rationalized as follows.

In low dilution situations, mild steel weld metals solidify as columnar 5 ferrite grains and austenite nucleates at the boundaries of these as they cool. If an austenite nucleus grows into only one of the adjacent 5 ferrite grains, the f> boundary will be pre­served as an austenite boundary: if the nucleus grows both ways, the boundary will disappear. In fact, the position of the original ferrite boun­daries can often be revealed by solute sensitive etches such as bromine-methanol and this enables the rela­tionship between the 5 ferrite and austenite structures to be inferred, since the austenite boundaries sub­sequently act as nucleation sites for pro-eutectoid ferrite. In the present work the austenite boundaries were found to be largely independent of the solidification boundaries, although the structure remained columnar.

Structure I in the classification scheme was pro-eutectoid ferrite, which in addition to forming thin veins at the austenite boundaries was also seen as polygonal grains within the columnar structure in weld metals of low hardenability. Structure II, mainly Widmanstatten ferrite side plates, often grew continuously from the massive grain boundary ferrite with no intermediate sub-boundaries. Any upper bainite which formed would have been included as structure II on the basis of its appearance in the light

62-s | M A R C H 1 9 7 6

microscope, but there was no ev­idence on replicas or thin foils exam­ined in the TEM of the sheaf structure of independently nucleated sub-units which has been cited as a dist in­guishing feature of upper bainite.

Structure III may be described as acicular ferrite and was the only one of the four whose incidence in the ex­perimental welds was negatively cor­related with transformation temper­ature. Weld S, containing 88% of acic­ular ferrite, transformed over the ap­proximate range 630-450 C (1170-840 F), while at the other extreme of hardenability weld V yielded only 8% of acicular ferrite after transforming in the range 710-545 C (1310-1010 F). The grain size of the acicular ferrite was represented by a mean linear in­tercept in the range 1-2 (tm, the finer structures corresponding to the lower transformation temperatures.

The fourth key on the point coun­ter was reserved for all structures which could not readily be assigned to the other three classes, but only one type of structure made an important contribution to this group in quantita­tive terms. This was a sub-pearlit ic aggregate of ferrite and carbides, rec­ognizable but not usually resolved in the light microscope. Electron micro­scopy showed the carbides to be al igned in rows, which often ap­peared parallel to a y-t-ct transforma­tion boundary, suggesting an inter­phase precipitation mechanism. The structure was found in welds of low hardenability and strength but no ad­verse ef fec ts on t o u g h n e s s was detected.

Other microstructural constituents, pearlite, retained austenite and mar­tensite, were also observed, though as these were not only low in total quantity, but also so finely dispersed that they were difficult to resolve and identify in the light microscope, they did not significantly influence the overall microstructural assessment. X-ray diffraction measurements of the retained austenite are included in Table 7. Since these were indepen­dent of the metallographic measure­ments the figures are not included in

Table 4 — Regression Coefficients of 0.1 mm COD Transition Temperature T, on Weld Metal Chemical Analysis

Regression coefficient (partial correlation coefficient)

Data set

A-U, W-Z

Constant

- 5 1 (V too porous)

A-Z, ex­cluding

A-Z, ex­cluding

M,V

M,V

Significance levels:'

Table 5 -

Weld

A B C D E

F G H J1 J2

J1R J2R J2RR K L

M N O P Q

R S T U W

X Y Z Com 1 Com 2

- 4

- 1

95%, "99%, *"99.9%

• Tensile Test Results

Yield or 0.2% proof stress,

N/mm 2

479 470 537 563 547

514 520 532 536 539

537 503 504 453 453

388 451 569 538 582

547 569 465 573 578

535 560 533 477 468

Mn

- 4 5 * ( -0.446) - 6 9 * * * ( -0.786) - 6 8 * * * ( -0.781)

Tensile str,

N/mm2

608 586 648 696 693

611 633 711 661 630

660 604 608 565 559

486 555 689 674 710

650 734 557 714 708

630 692 623 618 601

Si Ti

28 703*** (0.253) (0.632) 10 627*** (0.149) (0.784) not 630*** included (0.782)

Elonga­tion,

%

26 24 24 28 23

22 20 24 27 28

24 27 28 26 34

34 32 30 28 24

30 26 29 27 24

30 26 31 28 27

Al

- 3 0 1 ( -0.270) - 3 9 0 * * ( -0.535) - 3 7 0 * * ( -0.520)

RA,

%

60 60 64 56 50

58 60 57 57 58

60 60 60 55 69

59 63 61 60 58

57 52 65 54 59

59 56 66 55 58

Mult iple correlation coefficient

0.72

0.89

0.89

Strain hardening exponent,

n

0.16 0.15 0.12 0.14 0.12

0.12 0.12 0.13 0.11 0.13

0.13 0.16 0.15 0.13 0.16

0.17 0.14 0.13 0.13 0.09

0.13 0.13 0.14 0.14 0.10

0.15 0.12 0.14 0.13 0.13

Table 6 — Results of Regression Analysis of Weld Metal Yield and Tensile Strengths on Electrode and Weld Composition, for Strengths in N/mm2 and Composition in Wt. %<a)

Regression coefficient (partial correlation coefficient)

Dependent variable

<*Y

<7V

UTS

UTS

Source of analysis

Electrode

Weld

Electrode

Weld

Const.

347

331

400

358

C

not included 192 (0.07) not included 725 (0.26)

Mn

66*** (0.69) 93*** (0.66) 88*** (0.74) 111*** (0.71)

Si

45* (0.43) 56* (0.44) 62** (0.50) 78** (0.55)

Ti

1 2 7 " (0.49) 601* * (0.48) 215*** (0.64) 865*** (0.60)

Al

181 (0.34) 305 (0.27) 221 (0.36) 333 (0.28)

Multi correl. coeff.

0.77

0.85

0.88

0.90

(a) Significance levels: "95%; "99%; ""99.9%

WELDING RESEARCH SUPPLEMENT! 63-

the 100% sum of the latter results: electron microscopy showed the aus­tenite to be mainly present in the in­terstices of the acicular ferrite grains.

Inclusion Measurements

Table 8 lists the parameters de­scribing the inclusion distributions in selected experimental welds. Inclu­sions were initially measured in welds A-H, which formed a half-factorial section of the experimental plan. At­tempts were made to relate inclusion volume fraction to a series of com­positional factors, among them the Mn/Si ratio. This latter was included because a number of investigators (Refs. 26-28) have shown an appar­ent connection between Mn/Si ratio and weld inclusion contents, fo l ­lowing a suggestion by Sekiguchi (Ref. 29) that the fluidity of deoxida­tion products based on the Mn-Si -0 system, at temperatures just above the steel freezing point, depends crit­ically on this ratio. When it was found that the Mn/Si ratio, alone of the com­positional parameters tr ied, gave a useful correlation with the inclusion content of welds A-H, a second set of 8 welds with a wider range of Mn/Si

ratios was chosen for further mea­surements.

The results of all the inclusion mea­surements are plotted against the Mn/Si ratio of the electrodes in Fig. 8. Although the correlation for the com­plete set of results remained statis­tically significant at the 95% level, the regression line shown is that or ig­inally derived for welds A-H. Other authors (Ref. 28) have found that the relationship breaks down at high Mn/Si ratios, which may explain the high inclusion content of weld K. Rather than attempting to represent this by a second order regression equation, it was considered that the acceptable fit of the rest of the data warranted the use of the original equation. It is interesting to note that this accurately predicted the inclu­sion contents of weld L and T on the basis that both had a Mn/Si ratio of 0.9 although the absolute levels of Mn and Si were 50% higher in weld L; however, the prediction broke down for weld S which had the highest Mn level of all. Similar results were ob­tained plotting weld composit ion in place of electrode composit ion.

Although the present work pro­duced no grounds for disputing the

conventional view of the importance of the Mn/Si ratio in determining weld inclusion contents, energy-disper­sive x-ray analysis of inclusions in the SEM showed that most contained large amounts of Al and Ti in addition to Mn and Si. The results are thus not readily explained by the simple model of Sekiguchi, and the case for regard­ing the Mn/Si ratio as the main con­trolling influence on weld metal inclu­sions remains a phenomenological one.

Oxygen Analysis

The negligible solubility of oxygen in solid steel implies a close relation­ship between oxygen and inclusion contents in steel. The main reason for embarking on time consuming direct measurement of inclusions, how­ever, was the lack of general success, with certain exceptions (Ref. 20), in correlating weld oxygen contents with mechanical properties. One likely ex­planation for this was that not only the volume fraction, but also the size and spatial distribution of inclusions were important in determining fracture be­havior. The possibility of significant errors in oxygen analysis seemed

Table 7 — Microstructure of Welds as Determined by Point Counting, with Retained Austenite from X-ray Diffraction. Columns I-IV Sum to 100% and the Austenite is Dis­persed Through these Structures

Weld

A B C D E

F G H J1 J2

J1R J2R J2RR K L

M N O P Q

R S T U V

w X Y Z Com 1 Com 2

n.d.: not detected

I Proeutectoid

ferrite

37.2 26.5 21.4 16.7 27.3

26.3 36.4 23.7 22.4 24.2

21.7 29.1 24.3 18.9 59.4

45.5 25.2 15.7 28.9

2.1

20.8 8.9

41.6 28.5 29.6

10.4 23.2 11.9 24.0 21.4 27.2

II Lamellar

components

10.1 2.8 1.9 0.5 5.5

3.6 6.3 5.5 5.1 1.5

1.8 4.1 1.8

20.3 2.6

11.3 13.5

2.0 11.3 2.2

3.7 1.6

10.0 1.3 8.3

3.7 0.8 0.2 4.1

23.3 21.2

III Acicular ferrite

41.3 64.6 65.1 78.7 49.9

55.6 20.2 57.8 58.7 67.3

67.0 62.0 72.6 16.4 32.3

10.2 22.4 78.0 30.4 92.7

71.5 87.6 36.8 67.6

5.7

63.7 74.3 82.9 65.6 13.9

6.8

IV Other

structures

11.4 6.1

11.6 4.1

17.3

14.5 37.1 13.0 13.8 7.0

9.5 4.8 1.3

44.4 5.7

33.0 38.9

4.3 29.4

3.0

4.0 1.9

11.6 2.6

56.4

22.0 1.7 5.0 6.3

41.4 44.8

Retained austenite

— — — —

— — — — 2.3 2.3 2.3 1.0 1.0

n.d. 1.0 3.3 3.3 1.8

2.8 4.8 1.0 3.7 n.d.

1.7 1.8 1.4 2.2 2.6 1.9

Fig. 7 — Partial regression of 0.1 mm COD transition temperature on weld metal man­ganese content for welds A-Z

3 4 5 Mn/Si. wire

Fig. 8 — Variation of inclusion volume frac­tion, Vy, in the deposit with Mn/Si ratio of the wire

64-s I M A R C H 1 9 7 6

much more remote in view of the ex­tensive development of s tandard techniques.

In the attempt to relate metallo­graphically determined inclusion con­tents to weld composit ion, the degree of correlation with oxygen contents d e t e r m i n e d by a c o n v e n t i o n a l vacuum fusion method was very low. This at first caused little concern since the type of deoxidation product varied from weld to weld, but it was then realized that the density of the oxides involved is so closely linked to the equivalent weight of the deoxidants that a given weight of oxygen pro­duces quite similar volumes of MnO, Si02 , AI2O3, T i0 2 and intermediate products. When the estimated volume fraction of MnS was included in the calculations, the chemically deter­mined inclusion contents remained consistently lower than the metallo­graphic figures by up to 50%.

Towards the end of the work an op­portunity arose to check some of the oxygen contents by fast neutron ac­tivation analysis (Ref. 13). The amount of material remaining was not suffi­cient to do this for all welds, but the results of 10 analyses are shown in Table 2. Given that both the activa­tion analysis and the inclusion counts had a standard error of about 10%, there was now an acceptable agree­ment between the two sets of figures.

Discussion

The aims of weld deoxidation prac­tice were defined at the outset as the prevention of porosity, the minimiza­tion of weld inclusion contents and the provision, through the deoxidants re­maining in solution at the end of the deoxidation reaction, of adequately strong and tough weld microstruc­tures. The first of these requirements was the most easily satisfied, since only in the Si-free weld V was serious porosity encountered. Isolated pores were found in weld M, but the addi­tion of alloying elements required to give a yield strength of 450 N/mm2 (65 ksi) sufficed in all cases to produce sound welds. In practice, deoxidant additions somewhat higher than the minimum would be used to provide a margin of protection when welding in drafts or on dirty plate.

Weld Cleanliness

It has been concluded that the weld metal oxygen contents measured by the vacuum fusion method cannot at this stage be regarded as reliable indicators of inclusion contents. A tentative explanation for this is based on the supersaturation of CO neces­sary to generate bubbles on very small nuclei. The analytical method used requ i res the reduc t i on of a l u m i n a by c a r b o n at 1750 C (3180 F), which will only proceed at CO partial pressures of a few mi Hi-

Table 8 -- Parameters Describing the Size Distribution of Inclusions

Weld

A B C D E

F G H J1 K

L N S T U Y

Inclusion volume fraction.

% 0.49 0.61 0.32 0.55 0.75

0.43 0.41 0.50 0.41 0.51

0.69 0.48 0.65 0.69 0.63 0.57

Distribution and Spatial in Some Experimental Weld Metals

Est. std. error in

vol. f rac ,

% 0.055 0.06 0.03 0.06 0.07

0.03 0.05 0.06 0.03 0.03

0.06 0.04 0.04 0.07 0.08 0.08

No. of inclusions per mm3 , NUX108

1.36 0.96 0.54 0.98 1.74

0.45 0.52 1.03 0.29 0.65

0.50 0.36 0.64 0.45 0.40 0.40

Geometric mean

inclusion diam, nm

0.27 0.33 0.36 0.30 0.26

0.41 0.35 0.28 0.50 0.42

0.49 0.49 0.44 0.46 0.50 0.47

Avg. vol. nearest

neighbour spacing, A3, j im

1.1 1.2 1.5 1.2 1.0

1.6 1.5 1.2 1.8 1.4

1.5 1.7 1.4 1.6 1.6 1.6

meters of mercury. However, pres­sures much greater than this are re­quired to sustain bubbles of the size of weld metal inclusions: 1 atmo­sphere sustains a bubble of about 60 Mm (0.0024 in.) diam in molten steel (Ref. 30). If the inclusions do not coalesce into particles larger than this, it seems possible that some may not be reduced. Further investiga­tions into the discrepancy between different analytical methods are now in progress, but in the meantime both sets of oxygen figures have been treated with caution.

It n e v e r t h e l e s s s e e m s c lea r , whether we consider the chemical, physical or metallographic estimates of oxygen level, that this is indepen­dent, over a wide range, of the total amount of deoxidants present and their capacity to reduce oxygen activ­ity in the molten steel. Only in the porous weld V was there any indica­tion of a high oxygen level associated with a low total deoxidant level. The only compositional parameter which was significantly correlated with inclu­sion volume fraction was the Mn/Si ratio. This has been assumed in the past to influence the fluidity of the deoxidation products and thus their ability to coagulate with each other and with the top slag. A simple cal­culation (Ref. 18) shows that flotation according to Stokes' law is inad­equate to explain inclusion removal, a 1 itim inclusion moving only about 1.7 Mm under its own buoyancy in the 5 s or so for which the pool is molten, but the concept of inclusions in a tur­bulent pool impinging on and being held by the pool surface is attractive and invites a reappraisal, in future work, of the role of surface tension in addition to that of fluidity of the deox­idation products.

Fig. 9 — Variation of tensile fracture strain, e", with inclusion volume faction

Ductile Behavior

Given that it may be possible to use the Mn/Si ratio to control weld clean­liness, it becomes useful to know how this is related to the mechanical prop­erties of the deposit. It appeared from the inclusion counts on welds A-H that a useful correlation existed between tensile fracture strain e", calculated as In [100/(100-%RA)] and inclusion volume fraction Vv , as work on other materials has shown (31). However, when further welds were examined to see whether this corre lat ion was maintained over a wider range of properties, it became clear that the earlier view was oversimplif ied. The complete set of results is shown in Fig. 9 with the original regression equation indicated by a dashed line. The two welds which performed sig­nificantly better than predicted from the inclusion volume fraction alone, welds L and T, were characterized by

W E L D I N G R E S E A R C H S U P P L E M E N T ! 65-3

Fig. 10 — Partial regression of 0.1 mm transition temperature on weld metal manganese content, excluding weld M

low yield strength, tTY, and a high strain hardening exponent, n. The complete set of results suggested that both parameters played a part in determining fracture strains, though there was insufficient data to allow e* to be expressed quantitatively as a function of V v , "y and n.

Models of the microvoid coales­cence process predict the depen­dence of fracture strain hardening be­havior: if yield strength also has an effect, this implies that the void nucle­ation process is critical. The fact that this is so for mild steel weld metals is confirmed by the observation that not all the inclusions in the weld metal are capable of nucleating microvoids: in­deed, it has previously been shown (Ref. 32) that in the relatively soft (<TY = 460 N/mm2 = 67 ksi) "bench mark" weld metal used in developing the in­clusion counting technique, the size distribution of inclusions taking part in the fracture of a Charpy specimen at 0 C (32 F) peaked at 0.7 pm, whereas the volume dist r ibut ion peaked at 0.3 um and only 15% of the inclusions had a diameter of 0.7 um or more. Examination of fractured ten­sile specimens in the SEM confirmed this trend for the experimental welds and a similar "cutoff" effect has been reported by Hill and Passoja (Ref. 2).

It must be concluded that the prac­tical benefits of producing cleaner welds by suitable deoxidation prac­tice are probably not extensive for welds of relatively low strength. All the present welds achieved tensile elon­gations of more than 20%, which in the absence of stress concentrators should be sufficient guarantee against all failures other than those due to gross overloading. It is more inter­esting to consider the resistance to fracture initiating at a stress con­centration or defect in a structure and the simulation of this situation by means of a notched test. Both ex­perimental tests (Ref. 33) and the­oretical models (Ref. 34) have demon­strated the strong interaction be-

150

V* 50

^ 40

30

20

10

0

•w

-20

•30

- 4 0

-

O

1

o

o oo

o o

1

---

o i

O 0 0

<?% __

o o

0

I I .

0 01 02 03 Ot, 05 06 07 06 09 10 11 12 Si. V.

Fig. 11 — Partial regression of 0.1 mm transition temperature on weld metal silicon content, excluding weld M

tween strain hardening behavior and triaxiality of stress. As the triaxiality increases, the rate of strain hard­ening plays an increasing part in determining the fracture strain and the possibility of effecting useful im­provements by producing cleaner materials is correspondingly reduced.

In the welds studied here, the vol ­ume fraction of inclusions was reduc­ed by raising the Mn/Si ratio. Lower­ing the Si level, as far as was consis­tent with the prevention of porosity, carried no penalty in terms of tearing resistance but the addition of Mn re­fined the structure of the deposit and regression analysis showed that this raised its yield strength and lowered its strain hardening exponent, n. Be­cause the tensile fracture strain is rel­atively weakly dependent on n, some improvement can result f rom the in­creased cleanliness as Mn is added, but if the triaxiality of stress is in­creased to the point where unstable fracture is likely to be of practical con­cern, the improvement in cleanliness is outweighed by the fall in n. Thus there was no indication, in the pres­ent tests, that the COD at tearing ini­tiation increased as the inclusion con­tent fell over the range studied, and only in the case of weld M, with a very low strength and high n value, was <5i significantly greater than the average. Of course, where it is possible to im­prove c lean l iness w i thou t ma jo r microstructural changes, for exam­ple by reducing sulfur contents or the oxygen content of the shielding gas, more unequivocal benefits may be obtained (Refs. 20, 22). In multipass welds also, the effects of weld clean­liness may be less masked by micro-structural variations.

Cleavage Resistance

In order to explain the variations in cleavage resistance among the ex­perimental welds, the strengthening mechanisms at work were first inves­tigated. The good combination of

strength and toughness possible in mild steel welds is known to result at least in part from the development of fine, acicular microstructures and the first hypothesis tested was that the proportion of acicular ferrite (Struc­ture III) present directly controlled the yield strength. A significant correla­tion was found, but the correlation coefficient at 0.78 left room for im­provement and other strengthening mechanisms were considered.

While the strengthening effects of Mn and possibly Si may be supposed to result at least partly f rom changes in the nature of the austenite decom­position products, that of Ti as shown in Table 6 is so large in relation to any effect of this element on hardenabil­ity either detected here or reported elsewhere (Ref. 35) that only a dis­persion hardening effect seems able to explain it. This leads to the idea that it might be possible to represent the mechanisms controll ing weld metal strength by two factors, one related to the nature and morphology of the t ransformat ion products and the other to matrix strengthening by solid solution and dispersion hardening. The simplest way of doing this would be to take, for example, the propor­tion of acicular ferrite as the first fac­tor and to ignore all elements but t i ­tanium in the second. This hypoth­esis was tested by regression anal­ysis and led to the equation

n Y = 390 + 1.69 (% acicular ferrite)

+ 644 Ti /g)

in which strength is in N/mm2 , both terms are significant at the 99.9% level and the correlation coefficient is 0.90.

Applying the same approach to the prediction of COD transition temper­ature, simple regression of Tc on the various microstructural components did not produce significant results but with the inclusion of titanium the equation giving the best correlation was

T c = - 1 3 2 + 1.24 (% pro-eutectoid ferrite) + 807 Ti (9)

with a correlation coefficient of 0.70. However, inspection of the data show­ed that the behavior of weld M was out of line with the general trend.

In most of the experimental welds, the soft pro-eutectoid ferrite formed narrow veins separating blocks of stronger acicular ferrite. When such structures are tested, strain is con­centrated in the pro-eutectoid ferrite and it is here that cleavage cracks first initiate, as was shown by metallo­graphic examination of fracture pro­files. Thus the embritt l ing effect of the pro-eutectoid ferrite results, not from any inherent brittleness of this com­ponent, but from its "mortar- in-br ick-work" morphology.

66-s I M A R C H 1 9 7 6

When the propor t ion of softer phases reaches 80 or 90%, as in weld M, this analogy breaks down and strain concentration may increase as the proportion of acicular ferrite, now the minority phase, increases. Edel-son and Baldwin (Ref. 31) have re­marked on the similarities in the effect, as stress concentrators, of hard and soft second phases.

In interpreting the present experi­mental results, two approaches are possible. If it is allowed that weld M differs qualitatively f rom the other welds, data for this weld may legit­imately be excluded from the regres­sion analyses. Alternatively, second order terms may be incorporated in the regression equations so that these fit the whole data set including weld M, whose behavior is only anom­alous in the sense that it is not pre­dicted by the first order equations. In view of the fact that the coefficients of the second order terms could not be determined with good precision, while all the data except that for weld M could be satisfactorily fitted to first order equat ions, the former ap­proach was adopted. The data set without weld M now gave the regres­sion equation

Tc = - 1 4 2 + 1.87 (% pro-eutectoid ferrite) + 771 Ti / 1 0 )

whose correlation coefficient at 0.87 was little short of that obtained using the full chemical analysis.

Consumable Design

A rational basis for weld deoxida­tion pract ice has now begun to emerge. This acknowledges that different factors control fracture in the ductile and brittle modes and that in formulating welding consumables, priority must be assigned to perfor­mance in one or other mode. If the ductile behavior is important, for example Charpy upper shelf level, there may be some advantage in us­ing a Mn/Si ratio approaching 3-4 in the electrode or 4-5 in the deposit to limit inclusion contents. It is more profitable to achieve this by limiting the Si content, provided porosity is avoided, than by raising the Mn con­tent: the latter increases strength and lowers the rate of work hardening, both of which can offset the benefits of improved cleanliness.

If, as is more usually the case, cleavage resistance is the more im­portant criterion, the effects of the deoxidants are those shown in Figs. 10-13. These have been derived in the same way as Fig. 7, but without using the results from weld M. For the re­maining welds, that is, all those with yield strengths above 450 N/mm2 (65 ksi), manganese had a uni formly beneficial effect within the range studied, that is, from 0.45 to 1.6% in the deposit. This is achieved through

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'

o

•

o s£

1

o

-

--

0 001 002 003 001. 005 006 007 0 08 009 0 10 On 012 r,. •/.

Fig. 12 — Partial regression of 0.1 mm transition temperature on weld metal ti­tanium content, excluding weld M

replacement of pro-eutectoid ferrite by fine, interlocking acicular ferrite in the weld structure and is accompa­nied by an increase in yield strength of some 66 N/mm2 (10 ksi) per 1%Mn in the deposit.

No statistically significant effect of silicon could be detected over the range from 0.17 to 1.2% in the de­posit. Titanium increased the transi­tion temperature uniformly, within the limits of experimental accuracy, as its level in the deposit rose f rom 0 to 0.12%: it should be noted, however, that these were single run welds and some of the repor ted benefic ial effects of Ti in multipass welds could result from a refinement of the re­heated structure. There is also scope for further investigation of Ti levels in the range 0.01-0.015% which recent work has shown to be important, (Ref. 36). Aluminum lowered the transi­tion temperature over the range of 0-0.12% in the deposit, though this result was rather heavily weighted by a single point at 0.12%, and the 0.19% Al in the second commercial weld did not produce any detectable effect on cleavage resistance: fur ther work would be needed to conf irm this ev­idence of non-linearity in the effect of Al.

Over the range of composit ion cov­ered by the experimental weld metals with strengths above 450 N/mm2 (65 ksi) the above results are embodied in the equation

-1 - 6 8 Mn 370 Al

+ 630 Ti

(11)

which is shown graphically in Fig 14. If a low strength weld metal is

acceptable, with yield and tensile strengths below about 400 and 500 N/mm2 (60 and 75 ksi) respectively, an alternative design philosophy is suggested by the excellent tough­ness of weld M. This would involve re­ducing the amount of acicular ferrite in the weld to very low levels and pro­ducing a homogeneous, soft struc-

0 001 002 003 004 005 0 06007 006 009 0 10 0 ti 0 12

Fig. 13 — Partial regression of 0.1 mm transition temperature on weld metal alu­minum content, excluding weld M

-120-HO-100-90-80 -70-60 -50-40 -30-20 -10 0 10 Predicted TC.'C

Fig. 14 — Comparison of observed 0.1 mm COD transition temperature Tc with value predicted from the regression equation Tc = -1-68 Mn ' 630 Ti-370 Al. The results for weld M and the commercial welds 1 and 2 were not used in deriving the re­gression equation

ture. Some additional deoxidation by non-hardening elements such as Al would have to be provided since weld M was barely adequately deoxidized and conta ined occasional pores: careful consideration would also have to be given to the possible effects of dilution in preventing achievement of the desired structure.

Conclusions

1. Deoxidation practice affects both weld inclusion contents and microstructure and so can influence resistance to both microvoid coales­cence and cleavage fracture.

2. In as-deposited single pass C0 2

welds wi th s t reng ths above 450 N/mm2 (65 ksi), good cleavage resis­tance is promoted by a fine acicular ferrite structure and an absence of d ispers ion hardening effects. For welds made at 1.6 kJ /mm (60 kJ/in.) the 0.1 mm (0.004 in.) COD transition temperature was related to composi-

WELDING RESEARCH SUPPLEMENT! 67-S

t i on a n d m i c r o s t r u c t u r e by

Tc = - 1 - 6 8 M n + 630 Ti

- 379 Al Tc = - 1 4 2 + 1 . 8 7 (% p r o ­

e u t e c t o i d fe r r i te )

+ 771 Ti w h e r e Tc is m e a s u r e d in d e g r e e s C.

3. T h e s e r e l a t i o n s h i p s d i d no t a p p ­ly t o a w e l d of 3 9 0 N / m m 2 (56 ks i ) y i e ld s t r e n g t h , w h i c h d e r i v e d g o o d t o u g h n e s s f r o m a re la t i ve ly so f t s t r u c ­t u r e w i th on ly 1 0 % of ac i cu l a r f e r r i t e .

4 . I nc lus ion c o n t e n t s in C 0 2 w e l d s a re not re la ted to t he to ta l a m o u n t of d e o x i d a n t s p r e s e n t un less t h e s e a re at s u c h low leve ls tha t p o r o s i t y is i m ­m i n e n t . T h e r e a p p e a r e d to b e a re la ­t i o n s h i p b e t w e e n i n c l u s i o n c o n t e n t a n d M n / S i ra t io in t h e e x p e r i m e n t a l w e l d s , t he c l eanes t w e l d s b e i n g p r o ­d u c e d w i t h M n / S i ra t i os of 3 -4 in t h e e l e c t r o d e or 4 -5 in t he d e p o s i t .

5. The r e l a t i o n s h i p b e t w e e n w e l d c l e a n l i n e s s a n d r e s i s t a n c e to f i b r o u s f r a c t u r e is not s i m p l e bu t d e p e n d s o n t h e s t r e n g t h a n d s t ra in h a r d e n i n g b e ­hav io r of t h e ma t r i x . A s t h e t r iax ia l i t y of s t ress i nc reases , m a t r i x p r o p e r t i e s b e c o m e i nc reas ing l y i m p o r t a n t a n d h i gh y ie ld s t r e n g t h s a n d low s t ra in h a r d e n i n g e x p o n e n t s r e d u c e t e a r i n g res i s tance .

Acknowledgments

The author is indebted to Dr. R. Week, Director General of the Welding Institute, for permission to present this paper. His thanks are also due to his colleagues and to Dr. J. F. Knott for many helpful discus­sions and to Mr. B. G. Ginn and all those who helped with the experimental work.

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