Heat-treatment Processing of Austenitic Manganese Steels

Seluk Kuyucak, Renata Zavadil, Val Gertsman

CANMET Materials Technology Laboratory, Ottawa, Ontario, Canada

ABSTRACT

Hadfields austenitic manganese steels are best described as retained austenites. Steels heat-sensitized at 400 800C range, such as slowly cooled castings, de-velop intergranular embrittlement caused by hypereutectoid carbide precipitation. A solution annealing and water quenching heat treatment restores the mechanical properties. The present paper is an overview of research results describing the quantitative effect of grain boundary constituents on impact toughness and effect of section size on heat treatment, and a method of measuring of quench time using quench water temperatures.

Keywords: manganese steel castings, heat-treatment, embrittlement

INTRODUCTION

Hadfields steels with a basic composition of 1.2% C, 13% Mn are metastable austenites. This basic composition has remained the same since its discovery by Robert Hadfield in 1882. Both carbon and manganese help stabilize austenite against martensite formation. The martensite start temperature Ms is typically below the temperature of liquid nitrogen (196C). The embrittling carbides present in the as-cast structure are removed by solution annealing above 1000C and quenching. The resulting retained austenite structure with a large amount of carbon in solution gives these steels their high impact toughness and high strain-hardening rate and capacity, which give them their excellent impact-abra-sion resistance.

This overview of our recent studies on manganese steels [1-10] was funded by the American Foundry Society and the North American metal casting industry. It encompasses the nature of embrittling phases and their quantitative relationship to impact toughness, the behaviour of carbides and phosphide eutectic, structural soundness and microporosity distribution, macro- and micro- segregation, and a method of measuring quench time as a process control tool. The overview fo-

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cuses on results, the specific details of procedures and materials are found in individual research papers.

BEHAVIOUR OF CARBIDES

The embrittling intergranular carbides in manganese steels form during slow cooling or reheating through the 400 800C range. They are removed by solu-tion annealing above 1000C followed by rapid quenching. The kinetics of car-bide formation follow the typical C-curve of an isothermal transformation diag-ram, with the fastest growth (carbide nose) occurring at 600 650C. Essentially, two types of carbides form: the thin carbides form very rapidly and appear as grain boundary delineations in the etched structure at 200-1000 (Fig. 1a-d). Later, thick carbides nucleate on thin carbides and grow along the grain bounda-ries. The thin carbides do not embrittle the steel significantly, whereas the thick carbides embrittle severely. The latter are distinguished from the thin carbides by having a resolvable cementite interior with a clear austenite / carbide interphase boundary on both sides of the cementite film. The thin carbides remain as deline-ations even at high magnifications, their cementite is not resolved at 1000.

In regular manganese steels (ASTM A 128, Grade B), the thin carbide delinea-tions are less than 0.2 m thick, and the thick carbides have a starting thickness of 0.5 1.5 m, and appear as a step where they meet the thin carbides (arrows in Fig. 1). In Cr-bearing manganese steels (Grade C), intermediate thickened car-bides are also observed (Fig. 1e,f). These are effectively similar to thin carbides. They appear as thick diffuse delineations and the cementite is not resolved by light microscope.

Figures 2a-d show extracted replica images of grain boundary carbides by trans-mission electron microscopy (TEM). The thickened carbides in Grade C steel are seen to have a straight interface with one austenite grain and a serrated interface with the other. Although, their total width in the thickened areas have reached 3 4 m, their centers are 1.0 1.5 m thick when their finger-like growth is exclu-ded.

Figures 2e-g show thin foil TEM micrographs of the three carbides. The thin car-bide delineations are seen to be discreet particles. They nucleate on a grain boun-dary, but growth usually takes place into only one austenite grain. The thicken-ed carbides in Grade C steel also consist of many discreet particles but with grea-ter protrusions into austenite matrix. There are indications of stress on the lattice particularly as the particles get larger (Fig. 2f). The thick carbide films, on the other hand, are continuous (Fig. 2g).

Figure 3 shows energy dispersive X-ray analysis (EDX) line scans across a thin and a thick carbide in a Grade C steel using the TEMs nanoprobe. The thin car-bides Mn and Cr contents are close to those of austenite matrix, 13% and 2.5%, respectively. The thick carbide has significantly higher concentrations of Mn and Cr in solid solution (18% and 6%, respectively) and the diffusion of these ele-ments into the carbide causes a depleted layer in the surrounding matrix. The lite-

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rature on phase diagrams shows that in a 13% Mn, 1.2% C steel, the Mn content of the cementite that forms at 600C varies between 17% Mn (Benz 1973, in [11]) and 32% Mn (Koch and Keller 1964, in [11]). The 18% measured composition of the thick carbide in Fig. 3 falls within this range. This supports the hypothesis that the thick carbides are the stable phase, and the thin carbides are metastable.

The thin carbides appear as delineations under light microscope, primarily be-cause they are comprised of many small, discreet particles. The thickened car-bides in Grade C steel are essentially the same, with greater protrusions into auste-nite matrix. Crystallographically, all carbides index to an M3C cementite struc-ture but their chemical compositions are significantly different. It is proposed that the following mechanism operates for carbide precipitation during quenching:

The thin carbides are less stable but have a good lattice match with the austenite matrix. This gives rise to a low energy interface and a low nuc-leation barrier. The thick carbides, on the other hand, are the more stable constituents, but have a higher energy interface and a higher nucleation barrier.

Discrete, metastable thin carbides whose Mn and Cr contents are similar to those of the matrix nucleate rapidly at the grain boundaries. They are comprised of many discreet particles, which give rise to the observed deli-neations in the etched microstructure. These carbides have good cohesion with austenite because of their small size and low energy, possibly cohe-rent, interface. They often show preferred growth into one austenite grain. In Grade C steel, this growth is more pronounced, resulting in the appea-rance of thickened carbides (Figs. 1f, 2c,d,f).

Soon after they form, the thin carbide growth slows down and stops be-cause of the increased lattice mismatch and stress at the interface, and the driving force being low, is not able to overcome the increased interfacial energy. Eventually, a thick carbide with a larger driving force and an in-coherent interface forms.

The thick carbide thus formed grows rapidly in the lateral direction in a continuous film, along the grain boundary, incorporating any thin carbides along its way. Growth in the normal direction slows after a micron or so thickness because of the required diffusion of Mn and Cr from the matrix and of Si from the carbide. The thick carbides embrittle the steel either from their poor cohesion with austenite, in which case cracks propagate along the carbide / austenite interface, or because of their continuous na-ture and the cracks propagate through the brittle carbide.

QUANTITATIVE RELATIONSHIP BETWEEN EMBRITTLING PHASES AND IMPACT TOUGHNESS

The degree of embrittlement in austenitic manganese steels depends on the degree of grain boundary coverage by the embrittling phases. In this respect, intergran-

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ular carbides play a key role but other embrittling phases such as phosphide eutec-tic, non-metallic inclusions and microporosity also are involved. It is important to note that the abundance of an embrittling phase is not as important as its distribu-tion. As low as 5 ppm bismuth can embrittle high copper alloys [12] and the pre-sence of 60 ppm aluminum nitride as grain boundary films is sufficient to em-brittle steels [13]. These constituents, even in very thin layers, cause a loss of co-hesion with the matrix. Oxides have the least cohesion with a metal matrix, but typically their distribution is in the form of random discrete particles, hence they are not as harmful to toughness. Phosphorous is a greater concern in manganese steels it tends to segregate at grain boundaries, liquefies during solution annealing and forms an embrittling phosphide eutectic film. Sulphur in low carbon steels can form Type II grain boundary iron-sulphide eutectic during solidification [14]; but in manganese steels, tends to form globular manganese sulphides which are not particularly detrimental.

The very heterogeneous distribution of the embrittling phases in austenitic manga-nese steels makes it difficult to predict their impact toughness from metallogra-phic observations. A sufficiently large area must be characterized and averaged to achieve adequate accuracy. This was done for three steels, regular manganese steel (Grade B3: 1.2% C, 13% Mn), high C steel (Grade B4: 1.3% C, 13% Mn), and Cr-bearing steel (Grade C: 1.2% C, 13% Mn, 2% Cr). Charpy specimens from these steels were tested and their ends were prepared for metallography. The samples were examined at 200 magnification to easily distinguish the thin, thickened and the thick carbides from each other. Although each field presented a 0.1 mm2 area, approximately 1000 fields were scanned to qualitatively establish the grain boundary coverage by a given constituent. For rapid characterization, albeit with some loss of precision, a procedure based on a multi-field comparison was employed. The fraction of grain boundary covered by a certain constituent was mentally recorded while scanning the sample by moving the microscope stage. The overall average was recorded and the procedure was repeated for each constituent. A trained eye can accomplish this within 10% precision in 20 mi-nutes each sample. As each constituent embrittles Hadfields steel differently, each is assigned a weight. For example, we know that as-cast material, comp-letely covered with thick carbides, has toughness less than 90% of its optimum. Therefore, the thick carbides are given a weight of 0.9. Phosphide eutectic is about as detrimental as the thick carbides, also is weighted at 0.9. On the other hand, a heat-treated steel displaying only the thin carbides is very close to its optimum toughness; therefore, the thin carbides are weighted at 0.1. Thickened carbides are considered somewhere between and are weighted at 0.5. Microporo-sity, of course, creates complete decohesion, and is weighted at 1.0. Therefore, according to this scheme, a 90% coverage by the thin carbides is equivalent to 10% coverage by thick carbides.

Figure 4 shows the room temperature CVN impact toughness plotted against the weighted grain boundary coverage in the three steels mentioned [5]. A good cor-relation was obtained regardless of the steel type and processing. The results can be summarized as follows:

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Weighted g.b. Coverage Toughness CVN, J < 20 % Excellent >100

20 - 40% Good >50 > 50% Poor < 20

Impact toughness decreased at a rate of 3.5 J per percent weighted grain boundary coverage, from a nominal value of 180-200 J.

MACRO- AND MICRO- SEGREGATION

Macro-segregation

The kinetics of carbide precipitation in manganese steels increase sharply with carbon content, so that steels having more than 1.3% C are difficult to produce. Any positive macrosegregation therefore would compound this problem, as this steels poor thermal diffusivity and now higher carbon content would make it very difficult to heat treat the section centers. However, chemical analyses of top-poured and side-poured blocks showed negligible or inverse macrosegregation. Carbon was richer around the perimeter of the casting and leaner towards the center [4,7]. This was true for most of the other elements, phosphorous in partic-ular showed a clear inverse macrosegregation (Fig. 5). The problem of heat trea-ting the thick-sectioned austenitic manganese steels is therefore strictly a heat transfer problem.

Micro-segregation and its Effect on Phosphide Eutectic

Steels having more than 0.06% P or solution annealed above 1150C can be sus-ceptible to phosphide eutectic embrittlement, where a phosphide-carbide eutectic liquefies at the grain boundaries, gradually spreads along the grain boundary during the course of heat treatment and remains in place on quenching [3]. The as-cast structures contain this eutectic at the centers of carbide colonies as glob-ular constituents with a mottled appearance [2,4]. In the as-cast steels, the cooling castings do not spend as much time at high temperatures as in heat treatment, the eutectic does not have time to spread over the grain boundaries. Phosphide eutec-tic embrittlement is reversible. When castings are heat treated at the correct solu-tion annealing temperature, it re-dissolves into the austenite matrix. In severely embrittled steels, however, a trail of Kirkendall voids and cavities are left behind and the toughness is not completely restored [3].

Most solutes in manganese steels, notably Mn and Cr, and to some extent P, mic-rosegregate towards grain boundaries and intercellular regions. The degree of microsegregation increases with section size [7]. This makes the thick sections, especially near the casting surface, where the phosphorous content and the tem-peratures are higher, more susceptible to phosphide eutectic formation.

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MICROPOROSITY

Austenitic manganese steels have a long freezing range, and microporosity is a common feature in these castings, regardless of feeding. Microporosity occurs in clusters as random pockets, and its occurrence does not correlate with any thermal criteria such as Niyamas [8]. Figure 6 shows the typical distributions of porosity in 1-in. dia. sections in side-poured and well-fed ([Niy]>0.5)1, 4-8 in. thick, 15 in. wide, 20 in. long wedge-block castings in regular manganese steel. The random distribution of microporosity in clusters suggests heterogeneous nucleation. Each cluster is interconnected through intercellular regions. Microporosity can nucleate when the pressure drop ahead of the solidification front caused by the need to feed the solidification shrinkage leads to the supersaturation by the gaseous solutes. Usually, a large supersaturation and a favorable site (a non-wetting surface such as an oxide inclusion) are required to nucleate a pore [15]. When a pore nuc-leates, it quickly grows to fill the interdendritic channels by the exsolution of gas until the liquid is no longer supersaturated. This process may repeat itself at other nucleation sites, leading to the observed microporosity clusters in a 2-D section.

Gaseous solute content, oxide inclusion content and distribution, and the solidifi-cation rate (pouring temperature, molding media, use of chills, and section size) primarily affect the formation of microporosity. Among these, the gaseous solute content and solidification rate are probably the more important as oxide inclusions are a common feature in aluminum-killed steels.

Microporosity forms preferentially in interdendritic areas (Fig. 6e,f); therefore, its grain boundary fraction is much greater than its volume fraction. Indeed, the grain boundary fraction of microporosity as an embrittling constituent in this steel was estimated to be 5-10% [5]. Since the embrittling phases decrease Charpy im-pact toughness at a rate of 3.5 J per percent coverage, the contribution of micropo-rosity to embrittlement is expected to be a 15-30 J decrease in impact toughness from an expected maximum of 180-200 J. Microporosity is a secondary but a sig-nificant concern for loss of toughness next to intergranular carbides.

The microporosity of 1-in. dia. core sections varied from 0.04 to 1.5% [8]. The average porosity in the two blocks was 0.37%. As this is not large enough to compensate for the 3% solidification shrinkage, risers are required to feed the cas-tings.

QUENCH TIME MEASUREMENT

The quenching operation is critical in the production and heat treatment of auste-nitic manganese steels. The quench speed depends on the load distribution on a pallet, initial quench water temperature, section size and bath agitation. Among

1 [Niy] is Niyamas criterion value (G/T/t) in (C s) mm-1, where G is the local thermal gradient and T/t is the cooling rate.

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these, the load distribution is the least effectively controlled. If the water permea-tion through the load is poor, then the effective section size of the load becomes much larger, and with the thermal diffusivity of the steel being poor, the quen-ching operation can slow down considerably. The other process variable that is of interest is the initial quench water temperature. Although each metal caster has its own specifications for maximum initial and final water temperatures, in hot sum-mer months and with large loads, these can become difficult to attain. Typically, specifications apply to worst case scenarios (large section sizes). It is desirable to know if a certain load can be quenched (or has been quenched) even though the initial water temperature is (or has been) above the specified limit.

The steel chemistry plays an important role in determining the required quenching speed. Increasing carbon content increases the rate of carbide precipitation, and steels having more that 1.3% C are difficult to make. Chromium also increases the carbide precipitation rate by increasing the carbide stability and causes car-bides to start forming at higher temperatures.

The measurement of quench water temperatures provides a cost-effective means to evaluate a quenching operation. As castings cool, the quench water warms; hence, there is a one-to-one temperature relationship between the two. Quench water temperatures and heat capacity data of the steel, can be used to derive the average temperature of the castings at a given time during the quench, starting from their furnace (solution annealing) temperature. A useful, single-valued de-termination that lends itself to comparison would be the time to extract 90% of the heat from the load. This is schematically shown in Fig. 7 for a laboratory quench [3]. The total temperature rise in the quench water corresponds to the total heat taken from the castings. Therefore, 90% of that rise corresponds to 90% heat ex-traction. From the construction in the graphic, 90% heat extraction (t90%) in this quench took 4.2 minutes. In regular manganese steels having 1.2% C, satisfactory toughness (>100 J) is obtained if t90% is kept below 10 mins [3]. In higher carbon manganese steels (>1.3% C), the required t90% decreases to less than 5 minutes.

Figure 8 shows an industrial quench where there is continuous water circulation from a cooling tower [6]. The cooling rate of the quench tank (T ) attributable to the cooling water is given by:

)(

)(

CWW

CWWWW

TTMmT

TTCmTCM

(1)

where m is the cooling water flow rate, MW is the mass of quench water, TCW is the cooling water temperature as it enters the tank, and T is the tank temperature. The cooling water flow rate is usually known, but if not, it can be estimated from the data. At the end of a quench, when the castings are sufficiently cool, their contribution to heating the tank could be neglected and the tank temperature is solely affected by the cooling water. In Figure 8, 30 minutes after the quench, the tank cooling rate was 0.136C/min., the tank temperature was 23.4C, the cooling water temperature could be taken to be the tank temperature prior to the quench,

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20C. The 171213 ft. quench tank contained 75,100 kg water. Using this in-formation in Eqn. 1 yields a cooling water flow rate of 3000 kg/min. This figure is probably more reliable than a nominal design value and could be used to double check for any blockages in the system. The water temperatures must be adjusted by adding the effect of cooling water:

iiCWiW

i tTTMmT )( ,

(2)

The adjusted water temperature curve thus obtained represents the water tempe-rature in the absence of any cooling. The quench times can now be found by a similar construction as in Fig. 7.

The precision of the quench time analysis can be improved by taking temperatures in more than one location in the tank and averaging the results, the cooling water temperature as it enters the tank, and the ambient temperature. The data can be further processed to obtain cooling rates that show transitions from vapor blanket to nucleate boiling, and lastly convective cooling, and to conduct quench factor analysis using information from an isothermal transformation diagram [9].

CONCLUSIONS

Heat treatment processing of austenitic manganese steels has been studied with the following results:

Intergranular embrittlement by hypereutectoid carbide precipitation. The detri-mental thick carbides have been distinguished from the relatively harmless thin carbides. A model has been proposed for the different types of carbide for-mation. A quantitative relationship has been established between impact tough-ness and grain boundary coverage by the embrittling phases.

Chemical segregation. Macrosegregation was either absent or negative (concen-trations were higher around the perimeter of the castings than in the center). As carbon and chromium are not any richer in the center, the thick section problem in quenching the manganese steels is strictly a heat transfer problem. Microsegrega-tion increases with section size, and together with inverse macrosegregation, com-pounds the phosphide eutectic embrittlement around the casting perimeter.

Microporosity and structural soundness. Microporosity was observed as ran-domly and heterogeneously distributed colonies in castings. Although its volume fraction averaged 0.37%, it had a preferred location at intergranular and inter-cellular regions, where its grain boundary fraction was 5-10%. Its embrittling effect was secondary to carbides, decreasing the impact toughness by an estimated 15-30 J.

Quench time measurement. A cost-effective procedure was developed making use of quench water temperatures, to quantitatively assess the quenching opera-tion.

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ACKNOWLEDGEMENTS

The authors would like to thank Mark Charest for TEM replica study and Cathy Bibby for sample preparation.

REFERENCES

1) KUYUCAK S, Newcombe P and Zavadil R On the heat treatment of Had-field's austenitic manganese steels, Part I: step-down temperature and resid-uals. AFS Trans., 108, Paper No. 00-154, (2000).

2) KUYUCAK, S and Zavadil, R On the heat treatment of Hadfield's auste-nitic manganese steels - Part II: metallographic studies. AFS Trans., 108, Paper No. 00-126, (2000).

3) KUYUCAK S, Zavadil R and Newcombe On the heat treatment of Had-fields austenitic manganese steels, Part III: heat transfer model, macrosegre-gation and phosphide eutectic. AFS Trans., 109, Paper No. 01-117, (2001).

4) KUYUCAK S and Zavadil R On the heat treatment processing of austenitic manganese steels. Proc. 21st Heat Treat Conf., 5-8 Nov. 2001, ASM Interna-tional, Paper No. 16.

5) KUYUCAK S and Zavadil R On the heat-treatment of Hadfields austenitic manganese steels, Part IV: microstructure vs. impact toughness relationship. AFS Trans., 110, Paper No. 02-116, (2002).

6) KUYUCAK S and Newcombe P On the heat-treatment of Hadfields auste-nitic manganese steels; Part V: measuring quench time in industrial castings. AFS Trans., 111, Paper No. 03-102, (2003).

7) KUYUCAK S and Zavadil R On the heat-treatment of Hadfields austenitic manganese steels; Part VI: impact toughness, microstructure, macro- and micro- segregation in large wedge-block castings. AFS Trans., 111, Paper No. 03-133, (2003).

8) KUYUCAK S, Zavadil R and Ouellet G On the heat-treatment of Had-fields austenitic manganese steels; Part VII: casting soundness and micro-porosity in large wedge-block castings. AFS Trans., 111, Paper No. 03-134, (2003).

9) KUYUCAK S, Newcombe P, Bruno P, Grozdanich R and Looney G Mea-surement of quench time as a process control tool. Proc. Technical and Ope-rating Mtg., Nov. 2003, Steel Founders Society of America.

10) KUYUCAK S, Gertsman V Y and Zavadil R On the heat-treatment of Had-fields austenitic manganese steels; Part VIII: studies on microcharacteriza-tion. AFS Trans., 112, Paper No. 04-129, (2004).

11) RAYNOR G V and Rivlin V G in Phase Equilibrium in Iron Ternary Al-loys, The Institute of Metals, London, (1988), pp. 172-173.

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12) SADAYAPPAN M, Zavadil R and Sahoo M Effect of impurity elements on the mechanical properties of aluminum bronze alloy C95800. AFS Trans., 107, pp. 329-336, (1999).

13) WOODFINE B C and Quarrell A G . Effect of Al and N on the occurrence of intergranular fracture in steel castings. J Iron and Steel Inst., 195, (1960), pp. 409-414.

14) EEGHEM V J and DeSy A Side effects of cast steel deoxidation. AFS Trans., 72, (1964), pp. 142-148.

15) CAMPBELL J in Castings, Butterworth-Heinemann, (1991), p. 203. 16) PIWONKA T S, Kuyucak S and Davis K Shrinkage-related porosity in

steel castings. AFS Trans., 110, Paper No. 02-113, (2002).

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Fig. 1. Intergranular, hypereutectoid carbides in austenitic manganese steels [2,5]. (1a,b) regular manganese steel. Arrows show transition from thin to thick car-bides where a double austenite / carbide phase boundary on the carbide film be-comes visible. (1c,d) Further examples from regular manganese steel: carbide-free and thin carbide grain boundaries (1c), thick carbide laterally growing into a thin carbide (1d). (1e,f) Cr-bearing manganese steel. The thickened carbides in (1e) do not show a clear austenite / carbide phase boundary, whereas the thick carbides in (1f) do. Etched in equal parts of water, conc. HCl and conc. HNO3.

a b

c d

e f

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Fig. 2. Transmission electron microscopy images of carbides in Cr-bearing man-ganese steel [10]. (2a-d) carbon replica images, (2e-g) thin foil images. (2a) thin carbide, (2b) thick carbide, (2c,d) thickened carbide, (2e) thin carbide resolved into discreet particles (arrows indicate growth direction), (2f) thickened carbides (the extended tips in the growth direction are caused by dislocations), (2g) thick carbide showing as a continuous film.

f

g

e

c d

a b

13

0

2

4

6

8

10

12

14

-0.2 -0.1 0 0.1 0.2

Distance (m)

Mn

(wt.%

)

0

1

2

3

4

5

6

7

Cr,

C (w

t.%)

Mn

CrC

0

4

8

12

16

20

-3 -2 -1 0 1 2 3

Distance (m)

Mn

(wt.%

)

0

2

4

6

8

10

Cr,

C (w

t.%)

Mn

Cr

C

Thickness Composition (wt. pct.) (m) C Mn Cr Fe Grade C steel 1.2 12.4 1.9 83 Thin carbide 0.07 4.5 (6.8) 13 2.5 78 Thick carbide 0.8 6.0 (6.8) 18 6 69

Note: carbon content measured by the TEM nanoprobe is lower than the expected amount from the stoichiometry of cementite (given in parentheses). This is almost always the case for EDX analysis when a light element is present along with hea-vier elements in the compound. Iron is from balance.

Fig. 3. EDX profiles across thin (3a) and thick (3b) carbides in Grade C steel [10].

a

b

14

0

50

100

150

200

0 20 40 60 80 100Weighted Grain Boundary Coverage, pct.

Cha

rpy

Toug

hnes

s, J

Regular manganese steel, Grade B3High C manganese steel, Grade B4Cr-bearing manganese steel, Grade C

Fig. 4. Correlation between weighted grain boundary coverage by the embrittling constituents and impact toughness [5].

0.02

0.04

0.06

0.08

-3.5 -2.5 -1.5 -0.5 0.5 1.5 2.5 3.5Distance from Center, in.

Pct.

P

5 in.6 in.

7 in.

Block 1Block 2

5 in. 6 in. 7 in.

Drag Cope

Fig. 5. Macrosegregation behaviour of phosphorous in side-poured 4-8 in. thick, 15 in. wide, 20 in. long wedge-blocks (1.2% C, 13.7% Mn, 0.6% Cr, 0.3% Mo, 0.050% P) [7]. Samples were taken from 5-in., 6-in. and 7-in. sections.

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Fig. 6. Observed microporosity in sections taken from 1-in. dia. core samples, cut from wedge-block castings (6a-d). (6a) B1-7-4-6 (key: 1st block, core taken from 7-in. section, 4 in. from left, sectioned at 6 in. from drag). (6b) B2-7-6-7, (6c) B1-6-2-2, (6d) B2-7-4-2, from [8]. (6e) Intergranular and intercellular microporosity in a 1-in. thick block casting [16]. (6f) An extreme example of intergranular porosity [5]. (6a-e) as-polished, (6f) etched in equal parts of water, conc. HCl and conc. HNO3.

a b

c d

e f

16

Fig. 7. Construction to determine quench time from quench water temperatures in a laboratory quench tank [3].

16

20

24

28

32

36

40

0 5 10 15 20 25 30Quenching time, minutes

Que

nch

wat

er te

mpe

ratu

re,

C

Actual Temp.

Block 2, Quench Severity H = 3.0 per in.

Sivyer HT.xls

Cooling rate:0.25C/min at 32C

Adjusted Temp.

T 90% = 35.5C

t 90% = 6.3 min

T i = 22C

T max = 37C

Fig. 8. Quench time determination in an industrial quench tank with cooling water recirculating from a cooling tower [6].

0

5

10

15

20

25

30

0 2 4 6 8 10Quench Time, minutes F9034-HT.xls

T

t 90% = 4.2 min.

T90%

Water Temperature, C