Inclusion Phases and the Nucleation of Acicular Ferrite in Submerged Arc Welds in High Strength Low Alloy Steels

J.M. DOWLING, J.M. CORBETT, and H. W. KERR

Series of submerged arc welds of HSLA steel made with three different fluxes and metallic additions of Ti, Mo, and Cr have been examined to study the inclusions and their role in the nucleation of acicular ferrite. Inclusion phases and compositions have been analyzed by electron diffraction and X-ray microanalysis. These analyses have shown that the inclusions contained many different com- pounds, the proportions of each depending upon both the flux and metallic additions. Six inclusion phases have been identified: galaxite (A1203 �9 MnO), a titanium-rich compound (probably TiO), a copper sulfide, a manganese sulfide, a silica, and an aluminum-rich phase. No correlation was found between the amount of acicular ferrite in the weld metal and either average inclusion composition or individual inclusion phases. No epitaxial relationships between inclusions and adjacent ferrite grains could be identified. It has been concluded that inclusions nucleate acicular ferrite by acting as inert substrates according to the classical theory of heterogeneous nucleation. Because most inclusions are multi-phase and are touched by several ferrite grains, it has also been concluded that each inclusion can nucleate several ferrite grains, due to local regions of high surface energy on the inclusion.

I. INTRODUCTION

FOR good toughness over a range of temperatures, sub- merged arc welds (SAW) in high strength low alloy (HSLA) steels should have a high proportion of acicular ferrite. However, the factors which promote acicular ferrite for- mation are not fully understood.

During the decomposition of austenite, 1'2 the first a-ferrite often forms as coarse proeutectoid ferrite at the austenite boundaries, from which side plates develop. As transformation proceeds, fine grained acicular ferrite may be nucleated at suitable intragranular sites. At still lower temperatures, any remaining austenite transforms to bainite or martensite. The relative proportions of these trans- formation products are governed by their nucleation and growth rates, but are now thought to be influenced by the inclusions in the weld deposit, 3'4'5 which provide nucleation sites for acicular ferrite within the prior austenite grains. However, it should be emphasized that transformation of austenite to acicular ferrite will occur only if conditions in the weld pool are favorable. The effects of weld com- position and cooling rate must be such that the acicular ferrite transformation is promoted. If these conditions are not met, the austenite will transform to bainite or coarse ferrite depending on the cooling rate. 3

The fraction of acicular ferrite depends to some extent upon the weld metal oxygen content. It has been shown that too few oxide inclusions allow a bainitic structure while too many promote early proeutectoid ferrite nucleation at the expense of the acicular ferrite. 2 More recently, the uncom- bined weld pool oxygen content has been proposed as the critical factor:3 if this oxygen level lies within a certain range prior to and during solidification, the oxide inclusions formed will promote acicular ferrite formation. Alterna- tively, the oxide inclusions may determine the trans-

J. M. DOWLING, Research Associate, Department of Mechanical Engi- neering, J.M. CORBETT, Professor, Guelph-Waterloo Program for Graduate Work in Physics, Waterloo Campus, and H.W. KERR, Pro- fessor, Department of Mechanical Engineering, are with The University of Waterloo, Waterloo, ON N2L 3G1, Canada.

Manuscript submitted September 9, 1985.

METALLURGICAL TRANSACTIONS A

formation product by controlling the grain size of the prior austenite, large grain sizes being associated with bainitic structures. ~'6

Other reports have proposed that acicular ferrite is nucle- ated by certain inclusion compositions and phases. Com- pounds such as titanium oxide, TiO, 7 boron nitride BN precipitated on rare earth metal oxysulfides, s aluminum-rich inclusions, 9 and titanium nitride, TiN, ~~ have all been sug- gested as nucleant phases. Inclusions rich in manganese and inclusions covered with a skin of sulfide are reported to be ineffective nucleants. 9A2

Some authors consider that inclusion phase or com- position is less important than low mismatch between inclu- sion phases and a-ferdte. ~ Others have proposed that high surface energy of inclusions is the main requirement for ferrite nucleation. The inclusions then act as inert substrates and reduce the energy barrier for nucleation. 4 The nuclea- tion barrier may also be reduced by strain and dislocation fields which arise from differences in contraction around the inclusions during weld cooling. 6

The present paper is part of a larger investigation on microstructure-toughness relationships in submerged arc welds. 13'14,15 The nature of inclusions has been studied in relation to the nucleation of acicular ferrite. Preliminary studies showed that inclusions contained more than one compound and sometimes had a surface layer rich in copper and sulfur. 5 These studies have been extended to include other weld series made with different fluxes and metallic additions. The weld inclusions have been examined to iden- tify their constituent phases and to study their role in the nucleation of acicular ferrite. The results of these in- vestigations have been interpreted in terms of the various proposed nucleation mechanisms and phases.

II. EXPERIMENTAL DETAILS

A. Welding Procedures

Three series of welds were made with different fluxes. A full description of the procedures has been reported else- where. 13'14A5 The compositions of the fluxes, base plates,

VOLUME 17A, SEPTEMBER 1986-- 1611

and welding wires for each series are given in Tables I and II. Table III gives the compositions of the various welds which are discussed in this paper.

The W welds were made with a flux rich in alumina, and titanium and molybdenum were added to the welding grooves to produce a series of welds each with a different level of titanium (Table III). The heat input for this series was approximately 4.9 kJ mm -1 on plates 18 mm thick which gave a time for cooling from 800 ~ to 500 ~ of about 42 seconds.

The B and K welds were made with two different fluxes with chromium added to some welds in both series. For these two series, the heat input was 3.2 kJ mm -~, the plate thickness 13.7 mm, and cooling times of about 60 seconds were recorded for a fall in temperature from 800 ~ to 500 ~

B. Quantitative Metallography

To estimate the proportions of different microstructural phases (proeutectoid ferrite, acicular ferrite, bainite, etc.)

1000 points for each weld were examined, using an optical microscope, in the vicinity of the weld centerline after etch- ing in 2 pct nital. ~3'15

C. Transmission Electron Microscopy

For thin foils, sections of thickness 0.1 mm or less were prepared by cutting the welds on diamond slitting wheels followed by chemical polishing alternated with light grind- ing. Mechanically punched discs, 3 mm in diameter, were polished to perforation at - 2 0 ~ in a water-based solution of 5 pct perchloric acid (70 to 72 wt pct grade) and 50 pct glacial acetic acid.

Carbon extraction replicas were made from polished weld sections which had been etched in 2 pct nital.

Elements in inclusions and average compositions were analyzed using energy dispersive X-ray spectroscopy. Inclu- sions on extraction replicas were analyzed in a Philips EM200 transmission electron microscope (TEM), operated at 100 kV, and fitted with a Kevex energy dispersive spec- trometer. The condensed beam diameter was approximately

Table I. Estimated Flux Compositions (WtPct)

Flux CaO CaF2 MgO* MnO SiO2 A1203 TiO2 FeO ZrO2 B203 BI §

W 1 8 14 8 17 34 10 5 4 0 0.7 B 20 10 10 0 10 20 30 0 0 0.4 1.1 K 37 5 10 5 25 18 0 0 0 0 1.6

*includes small amounts of K~O, Na20 § = Basicity Index

CaO + CaF2 + MgO + K20 + Na20 + �89 + FeO)

SiO2 + �89 + TiO2 + ZrO2)

Table II. Compositions of Base Plates and Welding Wires

Composition (Wt Pct)

C Mn P S Si Cu Ni Cr Mo V Nb A1

Plates W 0.05 1.88 0.002 0.009 0.28 0.032 0.243 0.019 0.387 <0.005 0.057 0.074

*K + B 0.040 1.71 0.019 0.002 0.26 0.014 0.016 0.037 0.330 0.013 0.050 0.023

Wires W 0.041 1.02 0.006 0.011 0.19 0.079 0.025 0.031 0.006 <0.005 <0.005 <0.005

K + B 0.064 1.65 0.002 0.006 0.22 0.055 0.264 0.028 0.455 <0.005 <0.005 <0.005

*Weld K6: Base plate contained 0.25 wt pct Ni but other elements as above.

Table III. Weld Metal Chemical Analyses (Wt Pet)

Weld C Mn P S Si Cu Ni Cr Mo V Nb A1 Ti B* O* N*

W2 0.039 1.69 0.003 0.007 0.34 0.08 0.170 0.028 0.374 <0.005 0.038 0.010 <0.005 - - 454 98 W25 0.038 1.69 0.004 0.009 0.31 0.08 0.160 0.025 0.463 <0.005 0.034 0.008 0.018 - - 433 110 W16 0.040 1.71 0.002 0.008 0.34 0.08 0.165 0.028 0.373 <0.005 0.034 <0.005 0.033 - - 303 133 WI1 0.037 1.80 0.004 0.007 0.52 0.08 0.170 0.028 0.355 <0.005 0.039 0.010 0.200 - - 523 104 W12 0.040 1.77 0.002 0.007 0.48 0.08 0.171 0.028 0.468 <0.005 0.039 0.012 0.230 - - 662 102 B1 0.051 1.38 0.013 0.005 0.26 0.06 0.084 0.035 0.344 0.016 0.031 0.009 0.020 39 357 62 B4 0.047 1.37 0.012 0.005 0.26 0.06 0.I00 0.370 0.351 0.019 0.034 0.011 0.029 41 407 66 K8 0.051 1.79 0.012 0.005 0.29 0.05 0.092 0.032 0.340 0.012 0.032 0.008 <0.004 7 377 62 K6 0.047 1.73 0.016 0.004 0.34 0.06 0.267 0.425 0.373 0.009 0.034 0.010 <0.005 - - 394 72

- not analyzed *ppm

1612--VOLUME 17A, SEPTEMBER 1986 METALLURGICAL TRANSACTIONS A

1/xm. The relative proportions of the elements detected in the inclusions were calculated using the Cliff-Lorimer thin film correction method. 16 Inclusion sizes were also mea- sured from extraction replicas.

The constituent phases of inclusions in both foils and replicas were analyzed and identified from electron diffrac- tion patterns and X-ray microanalysis. Fine probe analysis (1 or 9 nm) was performed using a VG-HB5 scanning trans- mission electron microscope (STEM), and a Philips EM 400T was used for intermediate probe sizes of 20 to 40 nm. Additional diffraction data were obtained using a Philips EM200 and a Philips EM300.

To study orientation relationships between inclusion phases and adjacent ferrite grains, diffraction patterns were obtained from thin foils examined in a Philips EM300. Grain arrangement around the inclusions was also studied using thin foils.

I I I . RESULTS

A. Weld Metal Compositions and Microstructure

The compositions of the welds are given in Table III and the proportions of the different ferrite morphologies are presented in Table IV. Except in W11 and W12, which were fully bainitic, the welds discussed in this paper contained a high proportion of acicular ferrite ranging from about 80 pet in the W series to 100 pet in the B series.

In the W series, where titanium and molybdenum were added to the weld groove, there was a slight increase in the fraction of acicular ferrite with molybdenum. Titanium, on the other hand, produced a small increase in acicular ferrite content up to additions of about 0.04 wt pet. Titanium addi- tions beyond this increased the bainite fraction, a fully ba- initic structure being observed at 0.23 wt pet titanium. ~3 Increasing the titanium content from about 0.04 to 0.23 wt pet lowers the transformation temperature of austenite from about 590 to 600 ~ (for acicular ferrite) to about 550 ~ ~3 Therefore in W l l and W12, acicular ferrite does not form because the hardenability effect of titanium favors the bainitic transformation. Although the oxygen and nitrogen contents vary from weld to weld (Table III), these variations do not account for the bainitic structures observed in Wl l and W12, since welds W2 and W25 have similar oxygen and nitrogen contents (but very little titanium) and contain about 90 pet acicular ferrite (Table IV). Thus it has been concluded that the bainitic structures of Wl l and W12

result from an increase in hardenability produced by high titanium contents.

Weld microstructures cot~taining 100 pet acicular ferrite were observed in all of the B welds which were made with a flux containing boron oxide and a high proportion of titanium oxide.15 The presence of boron oxide in the flux transferred small amounts of boron (40 ppm) to the weld deposit. Boron can influence the microstructures of steels in many ways, although the mechanisms are the subject of much debate. The addition of boron suppresses the nuclea- tion of ferrite at austenite grain boundaries, apparently by the segregation of boron to the grain boundaries.IS This probably explains why no proeutectoid ferrite was observed in the B welds. Boron is also known to refine the ferrite grain size in coarse-grained austenite, apparently due to intragranular nucleation on such phases as boron nitride and/or iron boro-carbides. 8'17'19 This aspect will be discussed in Section B-1. Although up to 0.5 wt pet chromium was added to the B series, the element had no noticeable effect on microstructure. 15

Up to 0.73 wt pet chromium was added to the K series, and in these welds the fraction of acicular ferrite remained constant at about 90 pet. 15 However, proeutectoid ferrite decreased with chromium content and was replaced with bainite at additions of 0.4 wt pet and above. 15

B. The Nature of the Inclusions

1. Constituent elements and morphology Table V contains a summary of the average com-

positions, in atomic pet, for inclusion sections in the W, B, and K series. Inclusions in W12 and W25 were not ana- lyzed because the weld metal compositions (Table III) were similar to those of welds W 11 and W16, respectively. Com- positions are presented in terms of the three elements (aluminum, titanium, and manganese) which were common to all inclusions. In the W series, provided no titanium was added to the weld groove, the inclusions were rich in alumi- num, but the titanium content increased with titanium addi- tion. The B inclusions were rich in titanium while man- ganese was the major constituent in the K inclusions.

Small amounts of copper and sulfur were detected in many analyses but the peak counts were too low to give reliable quantitative data. Silicon was detected in the K and B weld inclusions but not in the W welds. Although molyb- denum was added to the W series and chromium to the K and B series, neither of these elements was detected in the

Table IV. Summary of Quantitative Micro-Structural Analyses

Average Wt Pet Wt Pet Wt Pet Pet Proeutectoid Pet Acicular Pet Bainitic Inclusion

Weld Ti Cr Mo Ferrite Ferrite Laths Size (/xm)

W2 <0.005 0.028 0.374 4 84 7 0.5 W25 0.018 0.028 0.463 3 88 5 - - Wl6 0.033 0.028 0. 373 6 82 9 0.6 W11 0.200 0.028 0.355 0 0 100 0.5 W12 0.230 0.028 0.468 0 0 100 - - B 1 0.020 0.035 0.344 0 100 0 0.4 B4 0.029 0.370 0.351 0 100 0 0.4 K8 <0.005 0.032 0.032 11 89 0 0.4 K6 <0.005 0.425 0.373 5 92 3 0.4

METALLURGICAL TRANSACTIONS A VOLUME 17A, SEPTEMBER 1986-- 1613

Table V. Average Inclusion Compositions in Atomic Percent

Weld A1 Ti Mn Other Elements

W2 57 +-- 13 2 - 2 29 --- 8 sulfur silicon

W16 49 -+ 18 38 --- 21 14 + 7 sulfur

Wll 18 --- 17 76 -+ 17 3 --- 5 sulfur B1 33--- 4 56--- 6 7-+ 4 sulfur

copper silicon: 3 --- 2

B4 29 - 12 40 --- 14 11 -+ 6 sulfur copper silicon: 16 -+ 20

K8 12 -+ 15 7 + 3 71 -+ 20 sulfur copper silicon: 6 - 7

K6 18--- 8 4+- 1 52+- 17 sulfur copper silicon: 23 --+ 11

inclusions. Data for oxygen, carbon, nitrogen, and boron are not presented because elements with atomic number less than nine were not detectable with the energy dispersive spectrometer.

The shape of the inclusions varied with the flux and groove addition, but the average inclusion size was about 0 .5/~m in all welds (Table IV). In the W series, W 11 and WI2 (with high titanium additions) contained angular inclu- sions (Figure 1). Fewer angular inclusions were seen in W16 and W25 where the inclusions tended to be irregular in shape with some angular portions (Figure 1). W2 inclusions (no titanium addition) were more rounded, and angular in- clusions were rare.

No angular inclusions were seen in B and K welds. In these welds, the inclusions were of a uniform rounded shape, similar to those in W2 (Figure 2).

2. Phases in inclusions X-ray microanalysis in STEM showed that the com-

position varied from area to area within the inclusions of the K and B welds. These observations were similar to earlier STEM data obtained for inclusions in W welds, s

Figures 3 and 4 show inclusions in welds K6, B4, and W25 together with the qualitative compositional analyses. In the central, thick areas of B and K inclusions, a mixture of manganese, aluminum, titanium, and silicon was usually observed (Figure 3). Elsewhere, there were titanium-rich areas and also surface patches of copper plus sulfur, both of which have been detected on W inclusions (Figures 3 and 4). The core of rounded W inclusions consisted of aluminum, manganese, and titanium (Figure 4(a)), but angular inclusions contained mostly titanium (Figure 4(b)). No continuous single-phased skin was observed on any W, K, or B inclusion.

Analysis of diffraction patterns obtained in both TEM and STEM confirmed the existence of several phases within a single inclusion. Five compounds have been identified so far: a titanium-rich phase, an aluminum-manganese phase, copper sulfide, manganese sulfide, and silica. The titanium-rich phase had an fcc structure with a lattice parameter of 0.44 nm. This lattice parameter was consis- tent with TiO (a0 = 0.418 nm), TiN(a0 = 0.425 nm), or

(b) Fig. 1--Extraction replicas of W welds showing the effect of titanium additions on inclusion shape. (a) Weld Wll , 0.200 wt pct Ti. (b) Weld W16, 0.033 wt pct Ti.

1614--VOLUME 17A, SEPTEMBER 1986 METALLURGICAL TRANSACTIONS A

Fig. 2--Extraction replica of weld B4 with rounded inclusions.

TiC (a0 = 0.418 nm) z~ within the limits of accuracy of elec- tron diffraction analysis. Because the inclusions were too thick for electron energy loss spectroscopy (EELS), the identity of the anion was not determined.

The aluminum-manganese phase had a spinel structure with a lattice parameter of 0.83 nm and was ascribed to an aluminate compound, A1203 �9 MnO, which is also known as galaxite. 2~

Manganese-sulfur rich areas corresponded to the e~-MnS compound which had an fcc structure and lattice parameter 0.52 nm. z~ Diffraction patterns from TEM and STEM have indicated that the copper-sulfur rich patches have a cubic structure with a lattice parameter in the range 0.55 to

0.59 nm. This value is consistent with the structures of CuxS, where 1.8 -< x -< 2, and CuS2. 2~

A silica phase has been identified from diffraction pat- terns as /3-cristobalite which has a spinel structure and a lattice parameter of 0.71 nm.20

In addition to these five compounds, an aluminum-rich phase has been detected in W12 and W25 inclusions but the compound has not yet been identified. Diffraction pattern analysis has indicated that the phase is not one of the many aluminas listed in the Powder Diffraction File. 2~

Table VI gives the observed frequency of the six com- pounds in inclusions from the three weld series. Where an inclusion contained several discrete patches of a single com- pound, for example copper sulfide in K inclusions, this has been reported as one observation. The relative abundances of the different phases in the inclusions showed trends which were similar to the overall inclusion compositions (Table V). In the W inclusions, both the Ti-phase and the spinel phase, A1203 �9 MnO, were observed. The Ti-phase became more common with increasing titanium in the weld metal, the inclusions in the high titanium welds (W11 ,W12) containing virtually no spinel phase. For the B and K series, both the A1203 �9 MnO and Ti-phases were observed, the spinel phase being more common in the K inclusions where the titanium content was low. The aluminum-rich phase has been observed only in W25 and W12, while the silica phase, fl-cristobalite, has been found in B 1 and K6.

The copper and manganese sulfides have been observed in B, K, and W welds. Although the sulfide phases were easily detected in STEM by X-ray analysis or selected area diffraction, in conventional TEM studies it was rare to obtain diffraction data because of the small size of the sulfide patches.

Diffraction studies in TEM have yielded many patterns which the authors have not been able to solve or match with known compounds in the Powder Diffraction File.

(a) (b) Fig. 3--Scanning transmission electronmicrographs of multi-phase inclusions in B and K welds. The marked regions are rich in A: Al, Mn, Ti, Si; B: Ti; C: Cu, S. (a) Inclusion in B4. (b) Inclusion in K6.

METALLURGICAL TRANSACTIONS A VOLUME 17A, SEPTEMBER 1986-- 1615

(b)

(a)

Fig. 4--Scanning transmission electronmicrographs of multiphase inclusions in weld W25. The marked regions are rich in A: A1, Mn, Ti; B: Ti; C: Cu, S. (a) Rounded inclusion. (b) Angular inclusion.

Table VI. Identified Compounds in Inclusions (Diffraction Data)

Number of Inclusions with:

SiO2 Ti Rich Phase A1203" MnO CuxS a-MnS Weld 0.71 nm 0.44 nm 0.83 nm 0.57 nm 0.52 nm

W2 - - 1 9 - - - - W25 + - - 7 8 - - - - W16 - - 6 5 1 - - W l l - - 6 1 - - 1 W12 + - - 4 - - - - - -

B1 2 3 5 2 1 B4 - - 4 4 - - 2

K8 - - 2 9 2 1 K6 2 5 11 4 - -

§ of Al-rich phase; W12:8 of Al-rich phase

3. Orientation relationships between inclusions and fer- rite grains

Using thin foils, the number of ferrite grains touching individual inclusions in welds W2, W16, B1, B4, K6, and K8 were counted. Welds W l l and W12 had bainitic struc- tures and were excluded. Table VII gives the number of inclusions surveyed in the three weld series and the per- centage touching one ferrite grain, two grains, and three

or more grains. For all the welds, the proportion of inclu- sions associated with just one grain was 30 pct or less while three or more grains emanated from most inclusions. One example of this multiple grain arrangement is shown in Figure 5.

Thin foils were also used to look for epitaxial re- lationships between inclusion phases and contiguous ferrite grains. A few diffraction patterns had reciprocal lattice di-

Table VII. The Number of Ferrite Grains Touching Inclusions

Weld

Total Number One Grain per Two Grains per Three or More Inclusions Inclusion Inclusion Grains per Inclusion

(--) (Pct) (Pct) (Pct)

W2 + W16 47 21 26 53 B1 + B4 98 13 28 59 K6 + K8 210 25 31 44

1616--VOLUME 17A, SEPTEMBER 1986 METALLURGICAL TRANSACTIONS A

Fig. 5 - - Arrangement of ferrite grains around inclusions in a thin foil from weld B4.

(a)

rections which were nearly parallel but the zone axes were not parallel (Figure 6). Nearly parallel lattice directions in other patterns involved high indices. No epitaxial re- lationships were found between ferrite grains and specific inclusion phases including angular inclusions with flat crys- tallographic surfaces (Figure 7). Numerous ferrite-inclusion patterns have been examined, and as yet the authors have not found convincing or consistent evidence for epitaxial relationships in W and B welds. Epitaxial relationships in K welds have not been studied.

IV. DISCUSSION

A. Composition and Phases in Inclusions

1. Origin of inclusions in steels In steel castings, three types of inclusions occur: primary

inclusions, secondary inclusions, and exogenous inclusions. Primary inclusions are formed before solidification and sec- ondary inclusions form in the supersaturated melt in the interdendritic regions while exogenous inclusions originate from outside the melt. 23 The roles of metal deoxidants and the molten slag-metal interface in the formation of in- clusions are relatively well understood for molten steels using thermodynamic data 24 and/or quaternary and ternary phase diagrams. 25 The subse 4uent behavior of the inclusions (coalescence, redistribution, and assimilation in the slag) can be elucidated. 23

In submerged arc welds, the reactions are more complex, and slag/metal equilibrium is not necessarily reached be- cause melting and solidification are rapid. 26 There is also some debate on slag/metal interaction and the release of oxygen and metal ions into the molten weld p o o | . 26'27'28

Consequently, it is difficult to predict the origins and phases of weld inclusions. However, the presence of the observed inclusion phases in W, B, and K welds can be understood in a qualitative way using data from deoxidation processes in molten steel ingots.

2. Composition of inclusions The overall inclusion compositions (Table V) varied pri-

marily with the flux, provided no strong metallic deoxidant

(b)

(c) Fig. 6 - - T h i n foil micrographs showing lack of epitaxial relationships between an inclusion and adjacent ferrite grain in weld B4. (a) Bright field. (b) Selected area diffraction pattern from central inclusion and surrounding ferrite grain. (c) Schematic of diffraction pattern. Ferrite grain (0) : (111) zone axis. Inclusion (e): (110) zone axis, spinel.

METALLURGICAL TRANSACTIONS A VOLUME 17A, SEPTEMBER 1986-- 1617

(a)

(b)

(c)

Fig. 7 - - L a c k of epitaxial relationships between an angular inclusion and adjacent ferrite grain in a thin foil of weld W2. (a) Bright field. (b) Selected area diffraction pattern from inclusion and adjacent ferrite grain. (c) Sche- matic of diffraction pattern. Ferrite: Q, inclusion: e.

was added. In the W and B series, the dominant metallic ion was the same in both the inclusions and the fluxes: alumi- num in the W welds and titanium in the B welds. Since the base plates and wires were very low in aluminum and ti- tanium, the W and B fluxes were the main sources of these elements. For similar reasons, it has also been concluded that in the B welds the flux was the source of the second major element, aluminum. In W2, which was titanium free, the second major element was manganese. Although there was 1 to 2 wt pct of manganese in the wires and base plate, the W flux was also influential here because these inclusions contained more manganese than the B inclusions formed from a manganese-free flux.

In the K inclusions, the major element was manganese followed by aluminum and silicon. This was somewhat sur- prising considering that the manganese content in the flux was less than that in the W series. It appears that the com- bined influences of the increased Mn content of the wire plus the presence of less alumina and more silica in the flux promoted deoxidation of manganese.

In the presence of strong metallic deoxidants, the inclu- sion composition was no longer determined by the flux alone. In the W series, increasing amounts of metallic ti- tanium in the weld resulted in higher titanium levels in the inclusions. Titanium is a very powerful deoxidant, and me- tallic additions of the element form inclusions in seconds in steels. 24

Although metallic chromium was added to the B and K welds, no chromium was present in the inclusions. Using the criteria of the standard free energy of oxide formation, chromium is a relatively weak deoxidizer compared with either aluminum or silicon combined with manganese. 29 It appears that the existing aluminum, manganese, and silicon in the molten welds were adequate for efficient deoxidation in the B and K welds, despite the addition of the metallic chromium. However, addition of chromium to B and K welds apparently promoted silicate formation. In welds B4 and K6 with about 0.4 wt pct added chromium, the inclusions contained about 20 pct silicon (compared to only about 5 pct silicon in B 1 and K8 with no chromium addition).

Calcium was not found in any inclusion, despite the pres- ence of up to 37 wt pct calcium oxide in the fluxes. Other authors 6,~2,30,3~ have not reported calcium in weld inclusions. Absence of calcium suggests that either the calcium does not dissolve, or calcium inclusions form early in the molten weld and are large enough to float out and be removed by slag entrapment.

3. Phases in inclusions Despite the variation in flux formulation, the titanium-

rich phase and the aluminum-manganese oxide (A1203" MnO) were observed in inclusions in all the welds. How- ever, the relative abundance of these two phases in the inclusions varied from weld to weld and closely followed the overall inclusion compositions (Tables V and VI).

In the W welds, at high levels of titanium addition, the inclusions were dominated by a titanium-rich phase (Table VI). The titanium-rich phase has been ascribed to the fcc compounds TiO, TiC, or TiN, all of which have similar lattice parameters. Considering the deoxidation powers of titanium, TiO would be a logical product. TiN or Ti(O, N) or Ti(O, C) may also be present because oxygen, nitrogen, and carbon are all soluble in the fcc lattice TiX. 29 The

1618--VOLUME 17A, SEPTEMBER 1986 METALLURGICAL TRANSACTIONS A

Ti-rich compound has also been observed in B and K welds and in W2 where no metallic titanium was added. In the absence of information on the anion, the Ti-phase will be referred to as TiX in this text.

The other phase common to all weld series, A1203 �9 MnO, is the stable deoxidation product for molten iron. McLean 32 has shown that when manganese and aluminum are pres- ent in solution, the formation of alumina occurs only at low oxygen concentrations. For molten iron containing 1.5 wt pct manganese, this oxygen content must be less than 0.005 wt pct. Since the manganese content of all the welds was about 1.7 wt pct and the total oxygen content about 0.04 wt pct, much higher than the critical value given by McLean, galaxite would be the expected deoxidation product. An aluminum-rich phase has been observed in W 12 and W25, but the electron diffraction data do not fit any of the alumina phases in the Powder Diffraction Fi le] ~

Sulfides of manganese and copper have been observed on some inclusions, et-MnS is common in steels unless treated with rare earth metals or calcium. 33 The calcium present in the fluxes either did not dissolve or did not remain in the metal. The origin of the copper sulfide is more difficult to explain. Although other elements can substitute for man- ganese in its sulfide, copper is not one of the group. 29'33 The authors are certain that the copper was not an artifact in the X-ray microanalysis. Copper-sulfur rich areas have been detected in STEM using extraction replicas on nickel grids, and analysis of the carbon film showed the absence of cop- per and sulfur. Patches of copper and sulfur on weld inclu- sions have been seen by other authors, ~2'3~ and the cubic compound Cul.sS(1) (a0 = 0.57 nm) has been reported in the matrix of mild steel welds. 34

In the welds B 1 and K6, diffraction patterns correspond- ing to /3-cristobalite (SiO2) were occasionally obtained. Silicon-rich areas have also been found in B4 and K6 in STEM microanalyses. Silica phases can originate from both deoxidation reactions and exogenous sources, 29 but ex- ogenous silica is associated with a quartz structure. Thus the silica in B and K inclusions was assumed to be indige- nous although in cast steels the lower temperature form, a-cristobalite, is reported to always appear. 29

As mentioned previously, there were diffraction patterns which were not solved, and additional phases may exist. Certainly examination of the overall inclusion compositions in Table V indicates that some phases are yet to be identi- fied. For inclusion compositions in W2, the atomic ratio of aluminum to manganese is in agreement with the stoi- chiometry of the spinel, A1203 ' MnO. Thus this inclusion composition may be accounted for by the phases already detected: A1203 �9 MnO and TiX with CuSx and MnS. W11 inclusions, which are rich in titanium with traces of alumi- num and manganese, can be accounted for in terms of TiX, A1203 �9 MnO, and the aluminum-rich phase. The latter and A1203 �9 MnO can also account for the composition of the W16 inclusions. However, silicon is present in B and K inclusions, and the latter contain more manganese than alu- minum (Mn : A1 = 3:1). Therefore, B and K inclusions may contain compounds apart from those given in Table V.

If it is assumed that titanium exists only as TiX, the ternary phase diagram for A1203-MnO-SiO2 provides an indication of other phases which may occur in the inclu- sions. In Figure 8, the overall inclusion compositions have been plotted in terms of wt pct A1203, MnO, and SiO2. The

3 MnO. M,203 . 3SiO z (Spessartite) ~ ,~ t

( RM:oOd ~i i01% ) ~ 1

2 MnO.SiO z /'"~'~l/ _ ( T e p h r o ~

/ I I I I MnO

SiO 2

2MnO.2A~ 2 0 3 . 5 S i O 2 (Mn -Cordierite)

nMO: _A ~An~:t'h2i tSe i)2

3 AP,203 . 2 Si02 (Mullite)

Wl6 Wll A~z03 A~zO3.MnO (Golaxite)

Fig. 8--Ternary diagram of coexisting phases m annealed mixes for the system AI203-MnO-SiO2. Composition data for W, B, and K welds are also given on the diagram. Coexisting phases for triangles A, B, C, D are taken from Reference 35. A: A1203" MnO; A1203; 3MnO' A1203" 3SiO~. B: A1203 �9 MnO �9 2SIO2; A1203; 3MnO �9 A1203 �9 3SIO2. C: A1203 �9 MnO; 2 M n O . S i O 2 ; MnO. D: A1203" MnO; 2 M n O . S i O 2 ; 3MnO. A1203 �9 3SIO2.

compositions and formulae of known compounds have been given on the diagram. Also included are the data from Rait and Pinder ~5 who determined the high temperature pseudo- equilibrium phases for this ternary system.

The compositions of the W and B inclusions lie in the same triangle (A) in Figure 8. The predicted coexisting phases are A1203 (corundum), A1203 ' MnO (galaxite), and 3MnO �9 A1203 �9 3SIO2 (spessartite). W2, WI6, and W11 lie on the pseudo-binary A1203-MnO with W2 coinciding with the composition for Al203 �9 MnO, and W11 and W16 being nearer to the A1203 corner. As already discussed, these three inclusion compositions are adequately explained by the ob- served phases given in Table VI.

B inclusions also contain aluminum in excess of the 2 : 1 A1 : Mn ratio required for galaxite. According to Figure 8, this excess will be taken up mainly as corundum but some aluminum, together with silicon and manganese, will form spessartite. B4 lies on the boundary between two triangles, A and B, in Figure 8, and thus B4 inclusions may contain an additional compound MnO �9 A1203 �9 2SIO2 (manganese- anorthite). Neither the latter nor spessartite has been de- tected in B inclusions.

The K inclusions lie in triangle C at the MnO corner of the phase diagram. Within this area, the coexisting com- pounds are MnO (manganite), 2MnO �9 SiO: (tephroite), and A1203 �9 MnO (galaxite). K6, straddling the boundary be- tween C and D, may also contain spessartite. Apart from the galaxite phase, none of these phases has been observed in K inclusions.

Because equilibrium is not reached in molten welds, the equilibrium phases in Figure 8 may not form. In addition, many of the weld inclusions contained titanium and there- fore an A1203-TiO-MnO or A1203-TiO-MnO-SiO2 diagram might be more appropriate. Despite extensive searches in the literature, no such phase diagrams have been found.

METALLURGICAL TRANSACTIONS A VOLUME 17A, SEPTEMBER 1986--1619

The compositions and compositional range of the inclu- sions may also be partly explained by the solid solubility ranges of the compounds given in Figure 8. For example, galaxite (A1203 �9 MnO) is known to exist with an excess of aluminum or manganese. 29'36 When manganese ions substi- tute for aluminum, the composition ranges from A1203 �9 MnO to A1Mn204 while retaining the spinel structure. 36 Because the lattice parameter changes from 0.83 nm to only 0.84 nm, 2~ electron diffraction patterns from the two phases will be indistinguishable. Yet a third phase, Mn304, also has a cubic spinel structure with lattice parameter 0.84 nm. 2~ Dekker and Rieck 36 have shown that some spinels with Mn:A1 ratios greater than 2:1 transform to a tetragonal structure at low temperatures. The lattice parameters depend upon the Mn:A1 ratio but for Mn22Alo804 a0 = 0.82 nm and Co = 0.88 nm. Two A1-Mn rich areas in K6 inclusions have yielded STEM electron diffraction patterns with zone axes (010) which were consistent with a tetragonal lattice and a0 = 0.82 nm and Co = 0.85 nm.

It has been observed that most inclusions contained sev- eral phases which form at different stages of solidification. The spinel (A1203 �9 MnO) and TiO are formed in the melt as solids (primary inclusions) while sulfides separate out from the interdendritic liquid at a later stage in the solidification of steels. 33 Many inclusions in the W series had the form of an angular particle attached to a round one (Figure 1). This morphology might result from coalescence of primary de- oxidation phases, i. e., A1203 �9 MnO and TiX during solidi- fication or nucleation of one phase on another. 29 The morphologies of B and K inclusions were rounded without angular protuberances, suggesting that they were formed as liquid droplets, and then crystallized within solid metal.

The melting points of both observed and suggested inclu- sion compounds are given in Table VIII. Compounds in W inclusions, galaxite and TiX, have high melting points com- pared with the solidification temperatures of mild steels. The phases detected in B and K welds also have high melting points, but other possible phases such as tephroite and spessartite have melting points of 1345 ~ and 1195 ~ respectively. Since A1203 �9 MnO is found to some extent in all B and K inclusions, the low melting point phases may form around it. The low melting point phases may have an amorphous structure arising from a decreased ionic mobility. So far, close examination of diffraction patterns has failed to reveal any diffuse rings.

B. Nucleation of Acicular Ferrite

1. Nucleant phases As discussed in the introduction, several nucleant phases

and mechanisms have been proposed for acicular ferrite. The observed inclusion phases and microstructures in the W, B, and K welds will be discussed firstly in relation to nucle- ating phases and secondly in terms of suggested nucleation processes.

A titanium-rich phase, TiX, has been observed in all the weld inclusions, being more prevalent where the inclusions contained substantial amounts of titanium (Tables V and VII). With the exception of Wl l and W12, all the welds contained a high proportion of acicular ferrite, some 80 to 100 pct. This supports the suggestion that TiX is a nucle- ating phase for acicular ferrite. 7 However, the inclusions in W11 and W12 contained mainly the TiX phase and yet both

Table VIH. Melting Temperatures of Selected Oxides and Sulfides

Melting Temperature Compound (~ Reference

A1203 2050 [37] p. 21 (Corundum) 3A1203" SiO2 1850 [37] p. 116 (Mullite) TiO approximately 1800 [38] p. 9 SiO2 1723 [29] p. 17 (/3-Cristobalite) A1203"MnO approximately 1650 [37] p. 21 (Galaxite) MnS 1620 [38] p. 528 Mn304 1560 [37] p. 21 2MnO" SiO2 1345 [29] p. 48 (Tephroite) 3MnO" A1203" 3SIO2 1195 [29] p. 52 (Spessartite) Cul.96S approximately 1125 [38] p. 515 Cul.sS

welds had a 100 pct bainitic microstructure. As mentioned in Section III-A, titanium additions of about 0.2 pct lower the transformation temperature in these welds, resulting in a bainitic structure instead of acicular fenite. Apparently, in Wl l and W12 the nucleation and growth of bainite occurs before intragranular nucleation of acicular ferrite on TiX or any other nucleant. Even in low titanium welds, with high acicular ferrite contents, the nucleant powers of TiX are open to debate. In Figure 9(a) the overall titanium contents in inclusions have been plotted against the pct of acicular ferrite in the welds. Inspection of the graph indicates little correlation between titanium in inclusions and acicular fer- rite in the metal matrix. Therefore, it is concluded that the presence of the Ti-rich phase, be it TiO, Ti(O, C), Ti(O, N), or TiN, is not sufficient for the nucleation of acicular ferrite if hardenability effects due to alloying promote lower tem- perature transformations. Furthermore, the suggestion of TiN or TiO as nucleants was prompted by possible low mismatch relationships between these compounds and fer- rite. So far the authors have seen no evidence of such re- lationships in B and W welds.

The other phase common to all the welds was the spinel A1203 �9 MnO. This compound has not been identified as a weld inclusion phase or nucleant by other authors. Bhatti et al. 9 have observed that welds with a high A1: Mn ratio in the inclusions contained a high proportion of acicular fer- rite. The AI:Mn ratio in W, B, and K welds has been plotted against pct acicular ferrite in Figure 9(b). There is no correlation between AI:Mn ratio and acicular ferrite. Consequently, it is concluded that a high AI:Mn ratio in inclusions is also not necessary for the formation of a microstructure with a high percentage of acicular ferrite. However, the AI:Mn ratio gives no indication of the amount of A1203 �9 MnO in an inclusion. A1203 �9 MnO was observed commonly in B and K inclusions and W welds with low titanium addition but rarely in Wl l and W12 with bainitic microstructures. Thus A1203 �9 MnO may play a role in the nucleation of acicular ferrite.

It has been proposed that surface layers on inclusions may be important in the nucleation of acicular ferrite.l X-ray

1620--VOLUME 17A, SEPTEMBER 1986 METALLURGICAL TRANSACTIONS A

~- I00 z w I-. z 80 0 r

60.

u , 4 - - . ,R / ' ~

z 0

20 _J

~ 0 I00

Z

K6 j [ j

90 ACICULAR

t W L6

K8 W2

85 80 FERRITE IN WELD

(a)

Wll

(%)

z

Z

e -

. . |

0 I--

0 I--

_

q BI 5

4 -

D-B4

2 -

I -

%

OWl6

eW2

Wl

K6 �9 K8

I l e I I 4 , _ _ 95 90 85 80 V 0

ACICULAR FERRITE IN WELD (%)

(b) Fig. 9--Percentage acicular ferrite in welds as a function of inclusion composition. (a) No correlation with titanium in inclusions. (b) No cor- relation with AI : Mn ratio in inclusions.

microanalysis and diffraction in STEM has shown that cop- per sulfide and manganese sulfide exist as small surface patches on inclusions in B and K welds (Figure 3). In earlier studies, copper-sulfur rich regions were found on W inclu- sions. 5 Because copper-sulfur patches were observed on in- clusions in W12, it must be concluded that the presence of copper sulfide is not sufficient for the development of a fully acicular microstructure. Indeed, there is some evidence from other authors that inclusions covered with a layer of manganese sulfide are ineffective sites for nucleation. 12

Boron nitride has also been suggested as a nucle- ant phase. 8'~9 So far the presence of this phase has not been confirmed in any of the examined welds, even those (B 1, B4) where boron is known to be present. However, this does not imply that boron nitride is absent from the inclu- sions, merely that it has not been detectable with the research facilities at the authors' disposal. Thus the possi- bility remains that, in the B welds, the highly acicular microstructures result because boron improves the harden- ability and provides, as boron nitride, intragranular nuclea- tion s i tes , s'lT-19

Oxides have also been proposed as nucleating phases. 2'3 Both A1203 �9 MnO and the TiX phase have been found to

METALLURGICAL TRANSACTIONS A

some extent in all welds regardless of transformed micro- structure. Hence, the presence of oxides is not a sufficient condition for the formation of acicular ferrite. This empha- sizes the importance of considering all the metallurgical parameters of weld metal transformation. Phases, such as A1203 �9 MnO, may promote acicular ferrite nucleation but only when other parameters such as hardenability, austenite grain size, and cooling rate are favorable.

2. Nucleation mechanisms Three mechanisms for ferrite nucleation at inclusions

have been put forward. They are all based on the classical theory of heterogeneous nucleation: nucleation by low mis- match interfaces between nucleant and nuclei, 7 nucleation at a high energy inert substrate, 4'3~ and nucleation in a region of high strain energy. 6

When nucleation proceeds by creation of low mismatch interfaces, there will be an epitaxial relationship between the nucleant and nuclei. Thus a diffraction pattern from an in- clusion and contiguous ferrite grain should show parallel lattice directions. No convincing crystallographic rela- tionships have been seen in the diffraction studies of thin foils. In many cases, the inclusions lacked well-defined crystallographic faces and thus epitaxial relationships were probably not to be expected. However, even when inclu- sions did possess fiat faces, no consistent relationships were observed. Therefore, the authors have concluded that nu- cleation by creation of low mismatch interfaces does not occur in the welds studied.

Acicular ferrite may nucleate in the strain field around inclusions. 4 To investigate this, data are required on the thermal expansion coefficients for all phases in the inclu- sions. Considering the multiphase nature of the inclusions, calculation of the volume change on cooling will be diffi- cult. The computation is further complicated because of uncertainty about the origin and formation of the inclusions. Some phases, such as TiO and A1203 �9 MnO, are formed as solids in the melt. Others have low melting points and may still exist as liquids after solidification of the metal matrix. Yet another problem arises if compounds undergo a phase transformation during cooling. Certainly, the lattice struc- tures of copper sulfides change with decreasing temperature as do some of the aluminum-manganese spinels. 21'35 In view of the lack of data and uncertainties about the formation of inclusion phases, the possibility of nucleation at strain en- ergy fields will not be discussed further.

Because no epitaxial relationships have been observed, it appears that inclusions nucleate ferrite by acting as inert substrates. An inclusion will then be a favorable nucleation site if a high energy interface exists between the inclusion and the prior austenite, thus lowering the energy barrier for nucleation. In the welds studied, the inclusions presented several phases at their surfaces. Presumably, each of these phases will have a different surface energy. Local regions with high surface energies may then promote heterogeneous nucleation of ferrite. Using the general assumption that sur- face energy increases with temperature of melting, then high melting point phases such as TiO, A1203" MnO, SiO2, MnS, and A1203 (Table VIII) would be expected to have high surface energies. Other phases such as copper sulfide and the manganese silicates, all with low melting points, would then be less efficient nucleants.

It has been observed that most inclusions contained sev- eral phases and therefore each inclusion may present several

VOLUME 17A. SEPTEMBER 1986--1621

suitable high energy nucleation sites. If this is so, then several grains of ferrite will be associated with one inclu- sion. Some support for this proposed multiple nucleation comes from the observations of ferrite grains around inclu- sions (Table VII). Since most inclusions were touched by two or more grains with many touching up to six grains, it has been concluded that multiple nucleation of ferrite occurs at inclusions. Multiple nucleation at inclusions has also been observed by Ricks et al. 4 However, they also propose that within a given weld, acicular ferrite can nu- cleate by an additional mechanism: sympathetic nucleation on Widmanst~itten ferrite, itself inclusion nucleated. 4 Little evidence for this mechanism has been observed in the welds studied.

Despite the existence of high melting point phases, such as TiO and TiC, in Wl l and W12 inclusions, the micro- structures of these welds were fully bainitic. Therefore, the weld chemistries were not appropriate for the decomposition of austenite to ferrite for the cooling rates used. The bainitic structures in these high titanium welds can be explained in terms of increased hardenability, 13 but it has been suggested that the formation of acicular ferrite requires a critical oxygen level just prior to the solidification of the weld. Terashima and Hart 3 have proposed that at this level the resulting oxide inclusions are of the right size and type to nucleate acicular ferrite. If too much deoxidant, such as titanium, is added to the weld, large inclusions are formed at high temperatures and the oxygen level is too low at solidification to form the required inclusions. When insuf- ficient deoxidant is present, they propose that the oxygen level is too high at solidification and the oxides are too small to act as nucleants. However, this hypothesis is apparently not applicable to the W series. W2, with no added titanium, has an average inclusion size of 0.6/zm and a high propor- tion of acicular ferrite whereas WI1, with a high titanium addition, is bainitic with an inclusion size of only 0.5/zm (Table IV).

From the above considerations, it is apparent that inclu- sions play a complex role in the transformation of weld metal. There is evidence that they nucleate acicular ferrite provided thermal and chemical conditions in the weld pool are suitable. Therefore, it is not surprising that nucleation of acicular ferrite is independent of specific phases in inclu- sions. Provided conditions in the weld pool favor the trans- formation, ferrite will nucleate at inclusions if these lower the energy barrier for heterogeneous nucleation. Because inclusions contain several phases, there will usually be at least one high energy site on the surface. However, inspec- tion of grain arrangements around inclusions indicated that two or more grains emanated from most inclusions. Thus it has been concluded that multi-nucleation of ferrite takes place on inclusion surfaces because these possess several sites with high surface energy.

V. CONCLUSIONS

1. Inclusion compositions varied with the flux provided no strong deoxidants were added. Metallic titanium addi- tions produced inclusions rich in titanium, but there was no chromium in inclusions when this metal was added.

2. The inclusions contained many phases. These and their proportions depended upon the flux and metallic addi-

.

.

.

tions. Six phases have been identified: A1203" MnO (galaxite), a titanium-rich compound, probably TiO or Ti(O, N) or Ti(O, C), an aluminum-rich phase, a-MnS, a copper sulfide, and SiO2 (a-cristobalite). Inclusion composition and electron diffraction studies have indi- cated the existence of other inclusion phases. No correlation between pct acicular ferrite in weld micro- structure and inclusion phase or composition was ob- served. There were no epitaxial relationships between ferrite grains and inclusions. Most inclusions nucleated several grains of acicular fer- rite. This was inferred from the observation that most inclusions touched two or more ferrite grains. Inclusions nucleate acicular ferrite by acting as inert sub- strates according to the classical theory of heterogeneous nucleation. Because inclusions are multiphase, there is a high probability that many high surface energy regions exist on an inclusion. Thus multi-nucleation of ferrite grains takes place at most inclusions.

ACKNOWLEDGMENTS

This work was supported by a grant from the Natural Sciences and Engineering Research Council of Canada, which is gratefully acknowledged. The authors would like to thank Dr. B. Robertson and Mr. F. Pearson at McMaster University for operating the VG HB5 STEM, and one of us (J. M. C.) is grateful to the Facility for High Resolution Electron Microscopy, Arizona State University, for provid- ing access to the Philips 400T TEM. The Steel Company of Canada provided welding materials and facilities and carded out the chemical analyses. Dr. E. S. Kayali examined the W welds in the STEM.

REFERENCES 1. R.C. Cochrane: IIW Doc. IX-1248-82, 1982. 2. D.J. Abson and R.E. Dolby: Welding Institute Research Bulletin,

1978, pp, 202-06. 3. H. Terashima and P. H.M. Hart: International Conference on The

Effects of Residual, Impurity and Micro-Alloying Elements on Weld- ability and Weld Properties, The Welding Institute, London, England, 1983, Paper 27.

4. R.A. Ricks, P.R. Howell, and G. S. Barfitte: J. Mater. Sci., 1982, vol. 17, pp. 732-40.

5. E.S. Kayali, J.M. Corbett, and H. W. Kerr: J. Mater. Sci. Letters, 1983, vol. 2, pp. 123-28.

6. P.L. Harrison and R.A. Farrar: J. Mater. Sci., 1981, vol. 16, pp. 2218-26.

7. N. Mori, H. Homma, S. Okita, and M. Wakabayashi: IIW Doc. IX-1196-81, 1981.

8. T. Funakoshi, T. Tanaka, S. Ueda, M. Ishikawa, N. Koshizuka, and K. Kobayashi: Transactions I. S. 1. J., 1977, vol. 17, pp. 419-27.

9. A.R. Bhatti, M.E. Saggese, D.N. Hawkins, J.A. Whiteman, and M.S. Golding: Welding Journal, 1984, vol. 63, pp. 224s-30s.

10. Y. Ito and M. Nakanishi: IIW Doc. XII A-113-75, 1975. 11. T.N. North, H. B. Bell, A. Koukabi, and I. Craig: Welding Journal,

1979, vol. 58, pp. 343s-54s. 12. G.M. Evans: IIW Doc. II-A-640-85, 1985. 13. A.J. Pacey, S. Kayali, and H.W. Kerr: Can. Metall. Q., 1982,

vol. 21, pp. 309-18. 14. E.S. Kayali, A.J. Pacey, and H.W. Kerr: Can. Metall. Q., 1984,

vol. 23, pp. 227-36. 15. G. Snieder and H.W. Kerr: Can. Metall. Q., 1984, vol. 23,

pp. 315-25. 16. G. Cliff and G.W. Lorimer: J. Microscopy, 1975, vol. 103,

pp. 203-07.

1622--VOLUME 17A, SEPTEMBER 1986 METALLURGICAL TRANSACTIONS A

17. J. Tsuboi and H. Terashima: IIW Doc. IX-1246-82, 1982. 18. J.E. Morral and T. B. Cameron: Boron in Steel, Proceedings of the

International Symposium on Boron Steels, TMS-AIME, Warrendale, PA, 1980, pp. 19-32.

19. S. Ueda, M. Ishikawa, and N. Ohashi: Boron in Steel, Proceedings of the International Symposium on Boron Steels, TMS-AIME, Warrendale, PA, 1980, pp. 181-98.

20. Powder Diffraction File, J. C. P. D. S., International Centre for Dif- fraction Data, 1980.

21. S. Djude: Acta Chem. Scand., 1958, vol. 12, pp. 1415-26. 22. R.A. Munson: Inorganic Chemistry, 1966, vol. 5, pp. 1296-97. 23. M.C. Flemings: Int. Metals Rev., 1977, vol. 22, pp. 201-07. 24. W.O. Philbrook: Int. Metals Rev., 1977, vol. 22, pp. 187-201. 25. R.K. Iyengar and W.O. Philbrook: Metall. Trans., 1973, vol. 4,

pp. 2189-93. 26. T.H. North, H. B. Bell, A. Nowicki, and I. Craig: Welding Journal,

1978, vol. 57, pp. 64s-75s. 27. T.W. Eagar: Welding Journal, 1978, vol. 57, pp. 76s-80s. 28. J.H. Palm: Welding Journal, 1972, vol. 51, pp. 358s-60s.

29. R. Kiessling and N. Lange: Non-Metallic Inclusions in Steel, 2nd ed., The Metals Society, London, England, 1978.

30. G.S. Barritte, R. A. Ricks, and P. R. Howell: Conference on Quan- titative Microanalysis with High Spatial Resolution, The Metals Soci- ety, London, England, 1981.

31. S.G. Court and G. Pollard: J. Mater. Sci. Letters, 1985, vol. 4, pp. 427-30.

32. A. McLean: J. Metals, 1968, vol. 20, pp. 96-100. 33. A.R. Palmer: Report MRP/PMRL 84-37(TR), Physical Metallurgy

Research Laboratories, Canada Centre for Mineral and Energy Tech- nology, Ottawa, ON, Canada, May 1984.

34. J.E. Harbottle and S. B. Fisher: Nature, 1982, vol. 299, pp. 139-40. 35. J.R. Rait and H. W. Pinder: J. 1. S. 1., 1946, vol. 154, pp. 371p-98p. 36. E.H.L.J. Dekker and G.D. Rieck: Z. Anorg. Allg. Chem., 1975,

vol. 415, pp. 69-80. 37. Phase Diagrams for Ceramists, 1st ed., The American Ceramic Soci-

ety, Columbus, OH, 1981, vol. IV. 38. Phase Diagrams for Ceramists, 1969 Supplement, 1st ed., The Ameri-

can Ceramic Society, Columbus, OH, 1969, vol. II.

METALLURGICAL TRANSACTIONS A VOLUME 17A, SEPTEMBER 1986-- 1623