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Transverse cracks present on the surface of continuously cast semis prior to roll-

ing often cause undesirable imperfections on rolled product. Such cracks can occur on all types of strand cast product: billets, blooms, slabs, rounds and beam blanks. Transverse cracks range in severity from those that are short, thin and difficult to find on a scaled sur-face, to those that are long, gross and readily apparent. A typical fine transverse crack found on the broadface of a slab after cold machine scarfing is shown in Figure 1. Although not a rule, transverse cracks often occur at or near as-cast corners.

Microcracking, which is invisible on a scaled surface, is a more subtle as-cast surface prob-lem and is often associated with transverse cracking. Microcracks are most evident after a sample of as-cast surface is deep-etched in hot hydrochloric acid. Upon acid etching, the problem appears as an irregular network of fine intergranular cracks,2–3 which has also been called craze cracking, network cracking, hairline cracking, shatter cracking, hot-short-ness cracking, checking and spider-web crack-ing. However, from this point forward, the term crazing is used when referring to the condition. Other methods used to reveal crazing have included dye-penetrant testing, magnaflux test-ing, and light scarfing or grinding. Crazing may be widespread or localized. Localized crazing is often found at the base of oscillation marks or in other surface depressions.3 More wide-spread crazing, along with some fine transverse cracks, is shown in Figure 2. The combination of crazing and transverse cracking can result in irregular surface seams on hot rolled product. An example of such seams on heavy-gauge plate is shown in Figure 3.

Although semis cooled to ambient tem-perature can be inspected and conditioned, the necessity for removing surface cracks of any type is obviously an undesirable expense. Moreover, when strand casting is coupled with subsequent hot rolling (e.g., hot charging or

hot direct rolling), inspection and/or condi-tioning is precluded. Therefore, elimination of transverse cracking and crazing on strand surfaces would go a long way toward improv-ing caster profitability.

Review of Surface CrackingTransverse Cracking — Numerous studies of transverse cracking have been reported in the literature. From these reports, several charac-teristics of the problem seem almost universal, namely:

• There is a greater propensity for crack-ing in microalloyed steels, e.g., high-strength low-alloy (HSLA) steels con-taining Nb, V, Al and N.

• Cracks are usually more severe on the top (inside radius) face of the strand, the face that undergoes tension during straightening (with strands having an arc).

• Transverse cracks occur most often at the base of oscillation marks and/or other surface depressions.

• The cracks are invariably intergranular and follow along boundaries of excep-tionally large prior-austenite grains.

Hot Ductility — The finding that transverse cracks are usually worse on the strand face that undergoes tensile strain during straightening (unbending) has prompted a plethora of lab-oratory studies on the hot ductility of steels in

Strand Surface Cracks — The Role of Abnormally Large Prior-austenite Grains

Excessive grain growth is a probable prerequisite for

transverse cracking during unbending of continuously

cast strands. Crack formation when grain diameters

are greater than 1 mm is reviewed, and experiments

to evaluate the problem are described.

Authors

Rian Dippenaar (left), professor of casting and steelmaking, University of Wollongong, Wollongong, Australia (rian@uow.edu.au); Suk-Chun Moon (center), casting manager — Minimill Dept., POSCO Gwangyang Works, Jeonnam, South Korea (msc9241@posco.co.kr); and Edward S. Szekeres (right), principal consul-tant, Casting Consultants Inc., Rochester, N.Y. (casting@fes-net.com)

106 ✦ Iron & Steel Technology

the temperature range of 600–1,100°C. These tests usually involve pulling apart an appropri-ately heated cylindrical specimen using equip-ment such as a Gleeble. At strain rates similar to those in a caster, most steels exhibit a duc-tility “dip” or “trough,” with a minimum in ductility somewhere between 750 and 850°C (Figure 4). Generally, the temperature at which the ductility minimum occurs coincides with the temperature at which weaker ferrite starts to form on austenite grain boundaries. However, it is found that certain elements, particularly Al, N, S, Nb and V, can lower the minimum ductility and/or extend the weakness to higher temperatures. This sug-gests that ductility loss can be exacerbated by

precipitates forming on the austenite grain boundaries prior to ferrite formation.

Because of the ductility trough, casters are usually operated according to one of two strand surface temperature scenarios. With the “higher temperature approach,” the surface temperature at the straightener is kept above the ductility trough, generally above 900°C. With the “lower temperature approach,” the aim is to keep the surface temperature below the trough, e.g., below 650°C. In either case, temperatures within the range of the ductility trough are avoided. Some success in decreas-ing transverse cracking has been reported with both programs. However, it is usually impractical to get surface temperature at the

Transverse cracks on the top surface of a machine-scarfed slab of C-Mn-Al-Nb plate steel.1

Figure 1

Localized crazing and some transverse cracks on the top corner region of a bloom of 1018-grade steel. Etched in hot HCl.2

Figure 2

Irregular surface imperfections on heavy-gauge hot rolled plate due to transverse cracks and crazing on the as-cast slab. Unetched.2

Figure 3

Typical hot ductility curve for C-Mn-Al steel, sensitized at 1,300°C, cooled to test temperature, strained at 3 x 10–3/second. Redrawn from Ref. 4.

Figure 4

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straightener below the ductility trough with strands having a liquid core in that zone. Therefore, newer casters tend to keep strand surface temperature at straightening above the ductility trough by casting faster and minimiz-ing secondary spray-water cooling. In a recent plant study, Tsai et al. found that an HSLA peritectic grade containing Nb had a higher percentage of transverse corner cracks.5 They obtained a decrease in cracking with a combi-nation of (1) casting “hot,” i.e., higher speeds and less secondary cooling; (2) maintaining a uniform strand surface temperature, e.g., avoiding “cold” corners; and (3) keeping Al, N, S and Cu at lower levels.

Primarily because it reveals a ductility dip, hot ductility tensile testing remains a popular laboratory tool. However, some authors4,6 have correctly cautioned that laboratory speci-mens stretched to failure with 5–100% reduc-tion-in-area have a strain history far different than that experienced by the strand surface during straightening. One major complica-tion with straining laboratory specimens to failure is dynamic recrystallization of austenite that inherently occurs at temperatures above about 950°C. In commercial casters, dynam-ic recrystallization is very unlikely because the maximum surface strain during strand straightening is typically around 1% or less (although surface strain in the range of 1–1.7% does occur with some caster designs).

To make laboratory hot ductility tests more commercially relevant, the gauge length of a hot ductility test specimen is sometimes melted and resolidified in situ. This technique allows the testing of “virgin” austenite and avoids the fact that reheating a test specimen from ambi-ent temperature does not actually simulate solid steel in the strand shell, which has not yet transformed fully to alpha ferrite. But even with these modifications, hot ductility testing

has not provided a complete understanding of the transverse cracking problem. There have been industrial cases in which a “good” strand surface temperature at the straightener yielded cracks, while a “bad” temperature yielded none. Moreover, one could question, “If transverse cracks are caused simply by a ductility loss, why do cracks occur at the base of oscillation marks or similar depressions? Is the surface temperature that much different a short distance away from the base of a mark? And why, at a similar temperature, do some marks and/or depressions have cracks and others do not?” Obviously, conditions other than ductility loss alone are contributing to the transverse cracking mechanism.

Crazing — Crazing, as shown in Figure 2, is the acid enhancement of fine microcracks that also follow the boundaries of abnormally large prior-austenite grains. Metallographically, unetched specimens of crazing have revealed that microcrack depth is usually 1–2 mm. However, at some locations, cracks may extend beyond 6 mm. Although shallow crazing can be eliminated by scaling in a reheat furnace, deeper cracks are not removed and can cause surface tearing in the first pass of roughing (Figure 5). With continued rolling, such sur-face tears elongate and become unacceptable, irregular, arrow-shaped imperfections on the surface of hot rolled flat product, such as shown in Figure 3. Crazing on as-cast billets can become parallel seams on final bar prod-uct (Figure 6).

Star Cracks and Cu Infiltration — Star cracking is a highly localized form of crazing (Figure 7). However, in almost all cases, copper can be found within the microcracks. Often, the infiltrated copper is from the strand abrading the mold liner or some other Cu-bearing caster

Crazing on the top surface of a slab opening up after about 10% reduc-tion in the first roughing pass. Etched in hot HCl.2

Figure 5

Surface of a 32-mm-diameter cold drawn bar of 1018-grade steel exhibiting seams due to craz-ing on the as-cast billet. Etched in hot HCl.2

Figure 6

108 ✦ Iron & Steel Technology

fixture, such as mold-exit grids. Preventing the hot strand surface from contacting copper, e.g., by the use of Cr and/or Ni-plated mold liners, significantly lowers the incidence of star cracking. However, with increased tonnage, liner plating eventually wears away, especially at the mold bottom and in the corners, and the potential for strand-to-copper contact returns.

The microcracks associated with widespread crazing can also be caused by a Cu-rich phase. However, the source is generally the buildup of copper beads in the surface scale, and the nature of the scale significantly affects the problem. The subject of Cu buildup in scale has been covered elsewhere and need not be repeated here.7–9 Suffice it to say that when-ever liquid copper contacts the strand surface in the early stages of shell development, at sur-face temperatures above about 1,050°C, there is the propensity for copper to infiltrate the boundaries of abnormally large grains.

The Influence of Prior-austenite Grain Size on Cracking and Crazing — The one characteris-tic common to transverse cracking, crazing and star cracking is that, in all cases, the cracks follow the boundaries of unusually large prior-austenite grains. As early as 1974, Schmidt and Josefsson showed evidence that transverse cracks occur only in the presence of abnor-mally large prior-austenite grains.10 Others have made a similar claim.11 More recently, Alvarez et al., in a study of transverse corner cracks in billets, found that cracks occurred only in surface regions having abnormally large prior-austenite grains, which had a film of blocky ferrite along the boundaries.12 In regions without cracks, grain size was normal, and the ferrite was fine and distributed uni-formly within the grains. Similarly, Tsai et al. found that fine transverse cracks followed the soft ferrite film that outlined prior-austenite grains having a diameter of about 1 mm.5

Although the association between transverse cracks and large prior-austenite grains has long been noted, it has not been adequately emphasized that an abnormally large grain condition is the key factor and a mandatory prerequisite for transverse cracking, as well as for crazing and star cracking. Without such large grains, microcracks and/or transverse cracks would not form to the extent they do, and other casting factors, such as surface temperature at the strand straightener, would be far less critical. Abnormally large prior-aus-tenite grains at the as-cast surface have been referred to simply as “blown grains.”2,13 To quantify the term, a blown grain is defined herein as a prior-austenite grain about 1 mm or greater in diameter when measured either on the as-cast surface plane or on a plane nor-mal to the as-cast surface. A big step toward alleviating the transverse cracking, crazing and star cracking problems could be made with a better understanding of the blown grain phenomenon.

An obvious question arises: What causes blown grains on the as-cast surface? The full answer to this question is yet to be found. It has been argued that austenite grains tend to grow at the base of oscillation marks, and at the base of deeper surface depressions caused by mold level variations and/or turbulence, because of a lack of local contact between the strand shell and the mold wall, which creates a “hot spot.”4,7 In support of this argument, minimizing the depth of oscillation marks does help decrease cracking, in some cases. However, the “hot spot” argument does not explain the fact that not all strand surface depressions have blown grains and not all flat surface regions of the strand are free of blown grains. Also, it is known that blown grains develop after the surface has completely solid-ified. Metallographic evidence of this fact was shown recently by Tsai et al.5 Cracks along blown grain boundaries were not interden-dritic. Instead, the cracks, as well as the blown grain boundaries, cut the trough preexisting dendrite arms.

Duplicating Blown Grains — To study how fast grains grow in virgin austenite, Maehara et al. machined cylindrical specimens from rolled laboratory heats having different levels of carbon.11 For the test, the gauge length of the specimens was remelted in situ and continuously cooled from 1,580°C at a fixed rate of 16.8°C/minute (relief of strain was not indicated). The results showed that aus-tenite grain size increased very rapidly in the temperature range of 1,450–1,350°C (Figure 8). Furthermore, in near-peritectic steel with 0.16% C, grains started growing earlier and attained a relatively larger size. Because the

Star cracks on the surface of a slab of linepipe steel, as revealed by a light scarfing pass. Unetched.2

Figure 7

July 2007 ✦ 109

surface temperature of most strands is below 1,300°C at mold exit, the results suggest that blown grains must develop while the surface region is still in the mold.

The results also suggest that 1,300°C, which is often used to “sensitize” (solution-treat) laboratory hot ductility specimens, is prob-ably not hot enough to create the very large austenite grains that are a prerequisite for cracking in commercial casting. The diameter of blown grains associated with cracks on strand surfaces is roughly an order of mag-nitude larger than grain diameters normally reported in solution-treating hot ductility test specimens.

Nevertheless, many researchers continue to test specimens with a prior-austenite grain size far smaller than the sizes known to be associ-ated with the aforementioned surface crack-ing problems. It is hard to see how laboratory testing of specimens that do not have blown grains provides data very meaningful to the surface cracking problems in industry.

Laboratory InvestigationsPurpose — The hot ductility of steel is known to decrease as grain size increases. However, most previous hot ductility studies have not tested specimens having anywhere near the abnormally large prior-austenite grain size found to be associated with various forms of strand surface cracking. Therefore, a labora-tory investigation was made to determine the effect of blown grains on hot ductility. To isolate the role of austenite grain size alone, hot ductility specimens were made from plain Fe-C steels. The aim was to avoid confound-ing factors such the presence of low-melting-point impurities, e.g., iron sulfides, and/or the grain boundary precipitation of carbo-nitrides, which are known to influence the relationship of grain size and hot ductility in microalloyed steels.4,6

Materials and Experimental Work — The chemical composition of three Fe-C alloys prepared for this investigation is shown in Table 1. Steel B was made to approximate a

peritectic composition. The melts were pre-pared in a vacuum induction furnace with car-bon as the only alloying element deliberately added. Aluminum was added as a deoxidizer. An attempt was made to keep tramp elements and impurities at low levels. However, ingress of a small amount of nitrogen (15–16 ppm) could not be avoided. (Because of the possibil-ity that some precipitates may have formed on grain boundaries during testing, the faces of some specimens fractured after cooling in liq-uid nitrogen were examined.) Approximately 37 kg of each melt was cast as an ingot measur-ing about 100 x 170 x 280 mm.

The solidified ingots were reheated to 1,200°C and then hot rolled to 15-mm-thick plate in a laboratory-size rolling mill. Cylindrical Gleeble specimens, 10 mm in diameter and 115 mm long, were machined from the plates. The axis of the specimens was kept parallel to the rolling direction. Hot duc-tility tensile tests were made using a Gleeble 3500 thermomechanical simulator, and the tests were run in a vacuum of approximately 10–3 atmospheres. Ductility was assessed by the reduction-in-area (RA) method.

Figure 8

Rate of austenite grain growth in in-situ melted specimens cooled from 1,580°C to test temperature at a rate of 16.8°C/minute. Redrawing of data from Ref. 11.

Chemical Composition of Plain Carbon Steels Investigated (wt %)

Steel ID C Mn P S Si Ni, Cr, Mo, Cu, Sn Ti, V Nb Altot N

A 0.05 <0.01 0.002 0.002 <0.005 <0.002 <0.003 <0.001 0.016 0.0015

B 0.18 <0.01 0.002 0.002 <0.005 <0.002 <0.003 <0.001 0.034 0.0016

C 0.45 <0.01 0.002 0.002 <0.005 <0.002 <0.003 <0.001 0.025 0.0016

Table 1

110 ✦ Iron & Steel Technology

Solution Treatments — The cycle for solu-tion-treating specimens is shown in Figure 9a. Each specimen was heated at a rate of 10°C/second to a solution temperature and held for 10 minutes. To vary grain size, three different solution temperatures were used: 1,100°C, 1,200°C and 1,350°C. Also, extra specimens of Steel B were treated at 1,280°C. Following solution treatment, each speci-men was cooled at a rate of 200°C/minute to a test temperature between 1,100 and 600°C. After holding for 1 minute at the test temperature, each specimen was pulled to fracture at a strain rate of 7.5 x 10–4/second, a rate similar to that obtained in commercial casting.

In-situ Melting — To simulate as-cast condi-tions, some Gleeble specimens were melted in situ (direct cast) and then cooled to the test temperature (Figure 9b). Collapse of the mol-ten zone during melting was prevented by use of a quartz tube around the middle portion of the specimen. Also, a compressive deforma-tion of 7% was applied to the specimen during solidification to avoid formation of shrinkage cavities. To determine the effect of cooling rate on hot ductility, two different cooling rates were used: 100°C/minute to simulate conventional casting and 200°C/minute to simulate thin-slab casting.

Grain Size Evaluation — Prior-austenite grain size was determined metallographically on specimens not pulled to failure. These speci-mens were subjected to the above heat treat-ments (without applying compression), but then water-quenched after cooling to a tem-perature within the two-phase region (830–700°C, depending on carbon content). This allowed ferrite to delineate austenite grain boundaries.

Results and DiscussionAustenite Grain Growth — The appearance of prior-austenite grains on the cross-section of thermally treated Gleeble specimens of the three steels is shown in Figure 10. Grain size increased nearly linearly with solution treat-ment temperature (STT) (Figure 11a). Grain size was greater than 1 mm even at an STT of 1,200°C. At an STT of 1,350°C, the average grain size was about 4 mm in all three steels (although a maximum grain size of almost 6 mm occurred with Steel A). Interestingly, the grain sizes obtained in this study using plain carbon steel are significantly greater than the 0.2- to 1-mm-size generally reported in the lit-erature by others testing Nb-containing steel or plain C-Mn steel under similar thermal con-ditions.11,14 Possibly, blown grains do not form as readily in solution-treated Gleeble speci-mens of microalloyed steels because elements and/or compounds pin grain boundaries and prevent growth at a high STT. Obviously, more study is needed to better understand austenite grain growth in Gleeble specimens.

With direct cast (in-situ melted) specimens, the slower cooling rate after solidification produced somewhat larger grains (Figure 11b). At a slower cooling rate, more time is spent in the high-temperature region; hence the grain size is larger. Steel A had a larger grain size than Steel C, which may be due to delta-ferrite grains growing less impeded in the single delta-ferrite phase on cooling, after solidification. There is an argument that prior-austenite grain size is largely influenced by delta-ferrite grain size.15

A comparison of average grain size versus carbon content for direct-cast specimens and specimens solution-treated at 1,200°C and 1,350°C is shown in Figure 12. In the direct-cast specimens, the largest grains occurred in

Schematic of thermomechanical cycles for hot ductility tests under (a) solution treatment and (b) direct casting (in-situ melting).

Figure 9

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(a) 1,100°C

Appearance of prior-austenite grains on the cross-section of 10-mm-diameter Gleeble specimens. Solution treatment temperatures: (a) 1,100°C, (b) 1,200°C and (c) 1,350°C. Cooling rate under direct casting condition: (d) 100°C/minute and (e) 200°C/minute. Sections etched in nital or a picric-acid-based solution. For clarity, grain boundaries were photo-enhanced in photographs (d) and (e) of Steel A and photograph (e) of Steel B.

Figure 10

Condition SteelA,0.05%C SteelB,0.18%C SteelC,0.45%C

Solutiontreatment

(b) 1,200°C

(c) 1,350°C

Direct casting

(d) 100°C min–1

(e) 200°C min–1

112 ✦ Iron & Steel Technology

the near-peritectic Steel B (0.18% C). This is consistent with the findings of Rezaeian et al.6 and Maehara et al.11 Grain growth is a sharp function of temperature, and upon cooling from liquid, austenite forms at a higher tem-perature in steel of peritectic composition. Based on the Fe-C phase diagram, austenite in Steel B formed at a temperature almost 100°C higher than in either Steel A or C. On the other hand, with solution treatment, grain growth in Steel B lagged, relatively. At the 1,350°C STT, a larger grain size was obtained with Steels A and C than with Steel B. With Steel B, a 1,350°C STT was required to cre-ate an average grain size equivalent to that obtained by direct casting.

Hot Ductility of Solution-treated Specimens — The hot ductility of specimens solution-treated at 1,100°C, 1,200°C and 1,350°C is shown graphically in Figure 13. The grain size obtained with each STT is shown on the graphs. Also included are the Ae1 tempera-ture (the equilibrium eutectoid transforma-tion temperature) and the Ae3 temperature (the equilibrium austenite/ferrite transforma-

tion-start temperature), which were calculated using MTDATA software.

At an STT of 1,100°C and 1,200°C, Steel A had significantly better ductility than either Steel B or Steel C. At an STT of 1,350°C, minimum ductility of all three steels was about equal (Figure 13d). The graphs in Figure 13 show that, as grain size increased with higher STT, the ductility trough (defined as RA val-ues below 40%) for all three steels deepened and extended toward higher temperatures, well into the fully austenitic region above the Ae3 temperature. Nothing was found in the lit-erature to indicate that deformation-induced ferrite can form beyond the Ae3 temperature. Therefore, the reason for the extension of low ductility to temperatures beyond the Ae3 temperature remains uncertain. It may be due to grain boundary sliding, but others have sug-gested that this failure mechanism is favored only in steel containing more than about 0.3% carbon.15 Another possibility for the widen-ing of the ductility trough is that nitrides precipitated on austenite grain boundaries. However, the nitrogen content of the steels investigated was quite low, about 16 ppm. Furthermore, examination of the fracture faces of specimens broken after cooling in liquid nitrogen revealed neither nitride nor sulfide precipitates. Therefore, the low ductil-ity of the austenite may be a result solely of the larger austenite grain size obtained. Grain size did not have a major effect on ductility at test temperatures below about 800°C.

The relationship between minimum ductil-ity and the reciprocal of prior-austenite grain size (1/D) obtained at the three solution treatment temperatures is shown in Figure 14. Although grain size and minimum duc-tility was similar in all three steels at an STT of 1,350°C, the ductility of Steel A improved more significantly as STT (hence, grain size) decreased. Thus, the minimum ductility for Steel A was a much stronger function of grain size than for either Steel B or Steel C. These data suggest that if extremely large blown

Prior-austenite grain size as a function of (a) solution treatment temperature and (b) cooling rate in direct casting.

Figure 11

ba

Average grain size versus %C for direct-cast and solution-treated specimens.

Figure 12

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grains can be avoided in commercial casting, then 0.05% C steel may be inherently less sus-ceptible to transverse cracking than steels with higher carbon content.

Hot Ductility of Direct-cast Specimens — The hot ductility of direct-cast specimens cooled at 100 and 200°C/minute is shown in the two graphs in Figure 15. The most striking result is the relatively good ductility exhibited by Steel A as compared to the higher-carbon steels. At the slower cooling rate, this low-carbon steel showed only a minimal loss in ductility.

Steel B had the largest grain size at both cooling rates. As discussed earlier, the large grains in Steel B result from higher tempera-ture in the austenite region as compared to that of the other two steels. The loss of ductil-ity seems to be directly related to an increase in grain size, and at a test temperature of 800°C, Steel B had the lowest RA value. Also, the austenite grains in Steel B were much more elongated perpendicular to the surface than grains in the lower- and higher-carbon alloys. This elongated grain shape is clearly evident in Figure 10.

The reason for the formation of elongated grains in the near-peritectic Steel B is unclear. It may be related to the larger temperature gradient during cooling of specimens of this

composition, or to the occurrence of the peritectic reaction followed by the peritectic transformation to austenite.11 Tsai et al.5 found blown grains elongated inward at least 6 mm from a slab surface. The inward elonga-tion of blown grains has practical significance because the effective grain size influencing intergranular fracture is the length of the elongated grains rather than the average grain diameter on the surface plane.11 With a

Hot ductility curves at 1,100°C, 1,200°C and 1,350°C STT for (a) Steel A, (b) Steel B, (c) Steel C and (d) all three steels at a 1,350°C STT alone.

Figure 13

ba

c d

Minimum RA value versus reciprocal of austenite grain size (1/D) for solution-treated specimens.

Figure 14

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greater effective grain size, the depth to which surface cracks can propagate is greater.

Although Steel C had the smallest grain size, it was still relatively brittle, possibly due to grain boundary sliding being dominant at this carbon level. Crowther and Mintz showed that increasing the carbon content of steel to above 0.3% (with a grain size of about 300 µ) caused intergranular failure of austenite to occur by grain boundary sliding.16 This result-ed in a very wide ductility trough. Increasing the carbon content was found to increase the activation energy for dynamic recrystal-lization, and hence to encourage more grain boundary sliding and linkage of cracks.

Ductility Comparison: Solution-treated Versus Direct-cast — One might expect that an as-cast structure is inferior to a reheated structure with respect to high-temperature mechanical properties due to differences in microstructure. However, this study showed that the ductility of direct-cast specimens of low-carbon steel (Steel A) is far superior to that of specimens solution-treated at tempera-tures above 1,100°C. Even with Steel B, the hot ductility of as-cast structure at test tem-peratures above 800°C was better than that of solution-treated specimens, although there were only minor differences in grain size. Tsai et al. found very few transverse corner cracks on LCAK grades (the highest percentage of corner cracks was found on a peritectic grade with Nb).5 Therefore, it appears that testing of solution-treated Gleeble specimens of low-carbon steel may have less relevance to the cracking problem in commercial casting than does hot ductility testing using the direct cast-ing (in-situ melting) method.

Summary and ConclusionsIt has long been recognized that transverse cracking is associated with large prior-austen-ite grains. What has not been fully appreci-ated is that the key factor responsible for transverse cracking, crazing and star crack-

ing is the presence of such abnormally large prior-austenite grains in local regions on the strand surface. On the as-cast surface plane, such grains are often greater than 1 mm in diameter (blown grains). It is not clear why some austenite grains grow to the extent they do, but the condition develops after solidifica-tion of the strand surface is complete. Reports in the literature have indicated that austenite grain size can increase rapidly at temperatures above 1,350°C, which means that blown grains must be forming while the surface region of a strand is still in the mold, probably within a distance of 200–300 mm below the meniscus.

To better understand the influence of blown grains on mechanical properties at elevated temperatures, the present study tested the hot ductility of specimens from three plain Fe-C alloys. Prior-austenite grains greater than 1 mm diameter were readily produced at solu-tion treatment temperatures of 1,200°C and greater. At an STT of 1,350°C, a grain size of 4 mm or larger was obtained. In addition, direct casting (melting in situ) and cooling to test temperature at a rate of either 100 or 200°C/minute resulted in grains greater than 2 mm in diameter. In testing a 0.05% C alloy, the hot ductility of direct-cast specimens was found to be superior to that of solution-treated specimens. This result indicates that testing of direct-cast specimens gives results more con-sistent with cracking experience in industry. Also, the direct casting technique produced elongated blown grains in specimens of a near-peritectic alloy (0.18% C). Elongated grains increase the effective grain size that influences cracking, and such elongated grains may par-tially explain the observed higher cracking propensity of peritectic steels.

A scenario for surface cracking based on blown grain formation is schematically sum-marized in Figure 16.

Stage I represents the normal grain size of new steel just solidified on the mold wall. The grain diameter at the surface is probably less than 500 µ.

Hot ductility under direct casting conditions (a) cooled at 100°C/minute and (b) cooled at 200°C/minute.

Figure 15

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Stage II represents austenite grains that have grown to a size several times larger than the original solidification grain size. The sur-face temperature is probably above 1,350°C. Strain energy can also drive grain boundary motion, but this factor was not studied in the present work. Future studies may clarify the role of strain in the development of blown grains (strain levels similar to that of mold friction). Regardless of how they form, the boundaries of blown grains become subject to weakening mechanisms and eventual crack-ing. The cracks are definitely intergranular, but are not necessarily interdendritic.

Stage III allows that solid-state sulfides may precipitate on blown grain boundaries and aid in microcrack formation. Furthermore, if liquid copper (or a Cu-rich alloy) is present at the scale/matrix interface, it can infiltrate the blown grain boundaries, weaken them, and allow microcracks to initiate and/or grow. However, in surface areas where the diameter of the austenite grains is normal, e.g., < 500 µ, the same volume fraction of precipitates and/or liquid copper does no harm.

Stage IV indicates further weakening of blown grain boundaries by precipitation of nitrides — e.g., AlN, Nb(C,N) or V(C,N) — at lower temperatures. Subsequently, or simul-taneously, the nucleation and growth of pro-eutectoid ferrite in a film-like fashion weaken the grain boundaries and cause a significant loss of ductility. Crack propagation is easier along a continuous ferrite film, as opposed to a discontinuous one.

Stage V represents the top side of the strand (inside radius), which experiences tension during straightening. Thus, any microcracks that are aligned with decorated blown grain boundaries are easily extended. Most micro-cracks are present prior to straightening, but new ones may develop. If the strand surface temperature at the straightener is above the Ae3, this stage may actually precede Stage IV.

References1. McPherson, N.A., and Mercer, R.E., “Continuous

Casting of Slabs at BSC Ravenscraig Works,” Ironmaking and Steelmaking, 1980 (4), pp. 167–179.

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Figure 16

Schematic illustration of events in the formation of surface cracks related to blown grains on the strand surface.

This paper was presented at AISTech 2006 — The Iron & Steel Technology Conference and Exposition, Cleveland, Ohio, and published in the AISTech 2006 Proceedings.