Dynamic Fracture Toughness of High Strength Cast Steels

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1 Paper pdf, Page 1 of 17 Dynamic Fracture Toughness of High Strength Cast Steels Copyright 2012 American Foundry Society L.N. Bartlett, A. Dash, D.C. Van Aken, V.L. Richards, K.D. Peaslee Missouri University of Science and Engineering, Rolla, Missouri ABSTRACT The dynamic fracture toughness of Cr and Mo steels with nickel contents of 0, 1.56, and 5.5 wt.% was evaluated and compared to a lightweight steel of composition Fe-30.40%Mn-8.83%Al-1.07%Si-0.90%C-0.53%Mo. Each steel was heat treated to a Rockwell C-scale hardness range of 36 to 38. The 4130, 4325, and HY130 steels were quench-hardened and tempered. The lightweight steel was solution treated, water quenched and age hardened. Of the alloys tested, the lightweight steel, the 4325 steel and the Al-killed and Ca-treated HY130 steel had similar dynamic fracture toughness values of 153, 153 and 165 kj/m 2, respectively. The 4130 steel had a much lower toughness of 94 kj/m. 2 The lightweight Fe-Mn-Al-C alloy performed better at Rockwell C32, producing the highest measured dynamic fracture toughness of 376 kj/m 2. Toughness of the Cr and Mo steels was strongly dependent on deoxidation practice. Alloys treated with ferro-titanium showed a reduction in toughness, which was attributed to TiN particles and in one case eutectic Type II sulfides. Addition of misch metal to an aluminum and ferro-titanium treated HY130 steel eliminated the Type II sulfides and increased the dynamic fracture toughness from 58 to 88 kj/m 2. HY130 obtained the highest toughness (165 kj/m 2 ) when aluminum deoxidation was followed by calcium treatment. INTRODUCTION The fracture toughness of high strength steel can be characterized by either a stress intensity or a J-integral approach to fracture mechanics. For materials with low toughness where cleavage fracture dominates, there is little plastic deformation around the crack tip and linear elastic fracture mechanics (LEFM) dominates. 1 Toughness is then evaluated as a critical stress intensity factor, K Ic. For sufficiently ductile materials, such as austenitic Fe-Mn-Al-C steels, failure is governed by the flow properties around the crack tip and the LEFM approach is no longer valid. 1 Therefore, for materials exhibiting crack tip blunting, an elastic-plastic fracture mechanics (EPFM) approach is required and the fracture behavior is described by the path independent J integral, which is equivalent to the energy release rate in elasticplastic materials. 1 The fracture toughness of ductile materials, J Ic, is defined as the critical value of J near the onset of stable crack growth. 2 In both of the above determinations of fracture toughness, the material is assumed to be under quasistatic loading conditions of less than 2 MPa m/s. 3 A material s resistance to fracture is often dependent on the loading rate; therefore, the static J Ic may not be representative of material behavior at high loading rates. Dynamic fracture toughness (DFT) under high loading rates is often difficult to obtain because crack extension during impact loading is difficult to measure. Instrumented Charpy impact tests provide a reproducible way of measuring the time dependency of force and crack displacement at elevated loading rates and therefore, provide a means of measuring DFT. Brittle fracture or Type I failure for linear elastic materials is characterized by crack initiation at the maximum load, which is followed by unstable crack propagation to failure. Failure of elastic-plastic materials is characterized by Type II, Type III, or Type IV behavior. Type II failure occurs when there is enough plasticity around the crack tip to allow for a small amount of stable crack extension before fracture at the peak load. For structural applications subject to shock loading, Type III or Type IV failure of the steel is desired. In Type III failure, there is significant stable crack extension after the peak load followed by unstable fracture. Type IV fracture is characterized by stable crack growth and the material fails by ductile tearing only, which is the desirable behavior for military armor. Schindler 4 has proposed a method of determining toughness from instrumented Charpy impact tests that is based on the crack tip opening displacement (CTOD) and crack tip opening angle (CTOA) models of crack nucleation and growth. This results in an algebraic expression for the dynamic J-R curve from which J 1d can be evaluated in an analogous way to the determination of J Ic. 2 This method is a single specimen approach with experimental inputs of the peak load (P max ), the energy up to peak load (E max ) and the total facture energy (E tot ), which are easily determined from the instrumented Charpy results. A complete description of the test procedure and methods for the analysis of the data may be obtained from a reading of Schindler 4, ASTM E and ASTM E FRACTURE OF HIGH STRENGTH STEELS The ability of a cast high strength steel to resist fracture is a function of many different metallurgical factors, including the inherent matrix toughness, the segregation of impurities that have limited solubility and the composition, morphology, and distribution of second phase particles. For steels of similar strength and microstructure, fracture toughness at elevated loading

2 Paper pdf, Page 2 of 17 rates is a strong function of steel cleanliness. Deoxidation practice can affect the toughness because it affects the inclusion type, size, morphology and distribution; and, therefore, can affect the fracture behavior. The use of strong deoxidizers such as Al and Ti can promote the formation of eutectic Type II sulfides, which form in interdendritic regions and reduce both tensile ductility and notch toughness. 5,6 Titanium is often added after Al deoxidation to reduce nitrogen related gas porosity, but this practice can promote Type II sulfides. Type I and Type III sulfides are less detrimental to toughness because they form at higher temperatures and are less likely to form eutectic cells along dendritic boundaries during solidification. Type I sulfides are globular and have the least effect on toughness while Type III sulfides are often faceted and these sharp crystallographic features can lead to void nucleation at small strains. Additions of misch metal have been shown to promote globular Type I sulfide formation near the liquidus as well as convert Type II and Type III sulfides, which have formed by deoxidation with Al or Ti into Type I sulfides. 6 Additions of Ti and Al are chemically active in the melt and can also form nitrides during solidification and subsequent cooling, which can also reduce toughness. Adding excess Al in amounts more than necessary for deoxidation has been shown to produce grain boundary precipitation of AlN, which can lead to rock candy fracture. 6 Quench-hardened and tempered (Q and T) Cr and Mo steels are used extensively for structural applications. Some of the more important variables that influence the capability of Q and T martensite to suppress brittle fracture during impact loading are the fineness of the microstructure and the ability of screw dislocations to cross slip at high strain rates. 5, 7 There is a rich metallurgical history associated with the addition of Ni to improve notch toughness. It has been shown that additions of Ni promotes cross slip of screw dislocations, produces solid solution softening at low temperature and lowers the ductile to brittle transition temperature in high strength steels. 5 HY130 steels are high strength Cr and Mo steels that have been modified with additions of Ni that typically range between 2 and 5.3 wt.% Ni. 5 All compositions in the following text are in weight percent. Yield strength values for Q and T HY130 steels generally range from 897 to 980 MPa and tensile strengths greater than 1047 MPa with elongations as high as 19% have been reported in wrought HY130. 5, 8 Dynamic fracture toughness of Ni modified Cr and Mo steel has been shown to increase with loading rate and toughness values greater than 250 kj/m 2 have been reported in wrought HY80 steel. 9 Austenitic steels in the Fe-Mn-Al-C system have gained much interest as a lightweight alternative to traditional high strength, cast, and Q and T steels. Depending on composition, the addition of Al produces a 12 to 18% reduction in density below that of traditional steel without a sacrifice in mechanical properties Typical alloy compositions contain from 20-30% Mn, 5-11% Al, and % C and are age-hardenable by the coherent precipitation of nano-sized κ-carbide, (Fe,Mn) 3 AlC, which precipitates homogeneously in the austenitic 11, matrix. The mechanical properties of Fe-Mn-Al-C alloys depend on the composition and the degree of age hardening. In a recent study, Van Aken et al. reported a yield strength of 873 MPa, an ultimate tensile strength of 953 MPa and an elongation to failure greater than 20% for a cast Fe-30.40%Mn-8.83%Al-1.07%Si-0.90%C-0.53%Mo alloy aged to a Rockwell C-scale (HRC) hardness of An austenitic matrix and high work hardening rates, which are characteristic of Fe-Mn-Al-C steels, contribute to high energy absorption rates under impact loading. However, the toughness of Fe-Mn-Al-C alloys is a strong function of P content and P contents greater than 0.01% have been shown to promote brittle cleavage fracture. 18 Impact toughness has also been shown to be dependent on inclusion type and density. Schulte et al. have shown that an increasing amount of AlN precipitation reduced the breaking energy at -40C (-40F) by almost 50% in a Fe-30.30%Mn-8.76%Al-1.02%Si-0.94%C-0.38%Mo alloy. 19 However, with a clean melt practice, using high purity charge materials, and keeping P levels below 0.006%, dynamic fracture toughness values greater than 400 kj/m 2 have recently been reported in age hardened alloys. 20 In the study reported here, the Schindler 4 method was used to determine the dynamic fracture toughness (DFT) of three different Ni-modified Cr and Mo Q and T steels. The role of deoxidation practice was also examined and the results are compared with toughness values obtained from a lightweight Fe-Mn-Al-C alloy aged to an equivalent hardness. DESIGN OF EXPERIMENTS Five different heats of HY130 were prepared to study the effects of different deoxidation practices on the same nominal composition. Three HY130 heats were prepared at Missouri S and T (HY130 heats A, B and C) and two heats were prepared at two different commercial foundries (HY130 heats D and E). Deoxidation practice for each alloy is listed in Table 1. The furnace charge for HY130 heats A, B and C consisted of high purity induction iron, De-sulco graphite, ferro-silicon (75%), ferro-molybdenum (60%), ferro-vanadium, electrolytic manganese, electrolytic nickel and electrolytic chromium. HY130 heats A and B were prepared under Ar cover in 200 lb and 100 lb coreless induction furnaces, respectively, and poured into bonded silica sand Y-block molds. HY130 heat C was prepared in a 20 lb vacuum induction furnace and cast into an investment Y-block mold. HY130 heats D and E were also prepared in induction furnaces at commercial foundries.

3 Paper pdf, Page 3 of 17 Table 1. Deoxidation Practices of The Alloys Used in the Current Study Alloy/Heat Quantity/Furnace Atmosphere Filter Deoxidation practice HY130 - A 140 lb melt in coreless induction furnace Ar gas cover No filters used 0.07% Al added to pouring stream in ladle followed by 0.05% Ca added as Ca-Si wire HY130 - B HY130 - C HY130 - D HY130 - E Fe-Mn-Al-C 100 lb in coreless induction furnace Ar gas cover 64 ppi filter used Same as HY130 - A Melt was poured into mold directly. 20 lb in vacuum Deoxidation same as induction furnace Vacuum of 1 torr No filters used HY130 - A 1000 lb in coreless induction furnace Inert atmosphere by SPAL (surface protective argon liquid) No filters used 0.08% Al added in furnace, followed by 0.12% ferro-titanium, and 0.08% misch metal. All additions were added within 20 s in the furnace. 200 lb in coreless induction furnace No protective atmosphere No filters used Pre-alloyed ingot was remelted without additional deoxidation by Al, and ferro-titanium was added prior to pouring (0.11% Ti final). coreless induction furnace (size not reported) Ar gas cover No filters used Al addition (0.007% final) coreless induction furnace (size not Al addition (0.008% final) followed by Ca-Si (0.008% Ca final) and ferro-titanium (0.044% Ti final) reported) Ar gas cover No filters used 200 lb in coreless induction furnace Ar gas cover No filters used No deoxidation required The HY130 castings were double normalized for one hour at 941 C (1726F) and one hour at 891C (1636F), austenitized at 841C (1546F) for one hour, and then oil quenched. The castings were tempered at 611C (1132F) for one hour to produce similar microstructures and hardness values. The 4325, 4130, and the lightweight Fe-Mn-Al-C steels were cast in commercial foundries. The 4325 alloy was deoxidized with Al and cast into bonded silica sand tensile bar molds. As cast 4325 bars were normalized for one hour at 870C (1598F), austenitized at 875C (1607F) for one hour, water quenched, and tempered for 30 min at 611C 1132F). The 4130 alloy was deoxidized with Al followed by Ca addition in the form of Ca-Si, treated with ferro-titanium additions and cast into bonded silica sand molds. Specimens were rough cut from the casting and austenitized at 870C 1598F) for one hour, water quenched, and tempered for one hour at 480C (896F). The Fe-Mn-Al-C steel heat was prepared using high purity induction iron, high purity aluminum, ferro-silicon, ferro-molybdenum, graphite and electrolytic manganese. The charge was melted under argon cover, triple calcium treated and argon stirred prior to casting into bonded silica sand plate molds, which were coated with a zircon wash. Because of the high Al content in Fe-Mn-Al-C steels, no additional deoxidation practice was required. Specimens from the Fe-Mn-Al-C plates were solution treated for two hours at 1050C 1922F) water quenched and age hardened from 13 to 60 hr at 530C (986F). Tensile test specimens were prepared from the 4325, HY130 and Fe-Mn-Al-C steels per ASTM E8 and ASTM E8M specifications. All tensile specimens were tested on a MTS servo-hydraulic load frame. DFT specimens were machined from the center of the castings as standard rectangular Charpy bars, i.e. 10 mm x 10 mm x 55 mm. Additional specimens from HY130 heat C were hot isostatically pressed to determine the effects of porosity on toughness. A sharp 0.25 mm wide notch with a depth of 2.0 to 2.5 mm was machined into one side of each test bar at midspan using a diamond saw (Cr and Mo steels) or by wire EDM (Fe-Mn-Al-C steel). The notched specimens were fatigue pre-cracked in 3-pt bending (R = 0.1) to a total initial crack length (a 0 ) of between 3.5

4 Paper pdf, Page 4 of 17 and 5.5 mm as per ASTM E1820. Finished DFT bars were fractured at room temperature on a Tinius Olson Charpy impact machine outfitted with an MPM instrumented striker. The load versus displacement data was used to estimate the dynamic fracture toughness (J 1d ) using the single specimen technique developed by Schindler. 4 The force displacement data required smoothing because of ringing oscillations in the results caused by shear wave propagation and reflection within the specimen during impact. A customized function for smoothing of the data that takes into account the wavelength of the oscillations was used as outlined in Kalthoff and Gregor. 21 Chemical analysis of the castings was determined by using arc spectroscopy for the Cr and Mo steels and by inductively coupled plasma spectrometry (ICP) after sample dissolution in perchloric acid for the Fe-Mn-Al-C steel. Additionally, the oxygen and nitrogen content in the castings was measured using a LECO O-N analyzer. Fractography was performed on the ends of broken test specimens using a Hitachi S570 scanning electron microscope (SEM). Inclusion analysis was determined using an ASPEX PICA 1020 SEM with automated feature analysis. Reported uncertainties were calculated as sample standard deviation for a sample size greater than three. RESULTS CHEMICAL ANALYSIS Chemical analysis for each steel is listed in Table 2. The oxygen and nitrogen content of the five different HY130 heats is given in Fig. 1. The nitrogen content of the HY130 steels varied between 40 and 190 ppm. As expected, the vacuum melted HY130 heat C showed the least amount of total N content. HY130 heat D proved to have the lowest O content. The higher oxygen content in HY130 heat C is attributed to the presence of moisture in the ceramic investment mold, which was fired but not pre-heated prior to casting. The low oxygen content of HY130 heat D was attributed to the surface protective argon liquid (SPAL) practice used during melting. The lack of protective atmosphere during melting of HY130 heat E resulted in the highest total N and O contents of 190 and 350 ppm. INCLUSION ANALYSIS The inclusion density by type for steels without Ti treatment is shown in Fig. 2. All of the inclusion analyses for the different steels were determined for a particle or pore major diameter of 1 to 80 µm. Cr and Mo steels without ferro-titanium additions consisted mainly of calcium aluminate type inclusions and sulfides containing Ca and Mn. HY130 heat A, which was prepared using Ar cover and cast without a filter, had the largest population of aluminate and sulfide inclusions. Analysis of the Fe-Mn-Al-C steel showed that inclusions were mainly complexes of AlN and MnS. In many cases, MnS was Table 2. Chemical Analysis of Different Steel Castings Given in Weight Percent Alloy/Heat C Si Mn P S Cr Mo Ni Al Ti V Fe HY130 / A < bal HY130 / B < bal HY130 / C <.0012 < bal HY130 / D < > bal HY130 / E <.0012 < bal bal <.0012 < <.002 bal Fe-Mn-Al-C bal

5 Paper pdf, Page 5 of 17 Fig. 1. The oxygen and nitrogen contents of the HY130 heats show that the highest amount of both oxygen and nitrogen were obtained for HY130 heat E which was melted without protective atmosphere. Fig. 2. The inclusion density by type for the alloys without Ti addition shows that a large percentage of the inclusions in the HY130 and 4325 alloys were calcium aluminate and (Mn, Ca)S inclusions. Inclusions in the Fe-Mn-Al-C steel were mostly complex AlN and MnS. found to precipitate around pre-existing AlN inclusions. The Fe-Mn-Al-C steel also contained the most porosity. Of the HY130 heats analyzed, heat C was found to contain the most porosity with a volume fraction of The inclusion density by type for Cr and Mo steels with ferro-titanium treatment is shown in Fig. 3. Most of the inclusions in all three Ti-treated steels consisted of TiN and MnS with a small percentage of aluminate inclusions. In many cases, MnS was found to have precipitated on pre-existing TiN inclusions. HY130 Fig. 3. The inclusion density by type for the alloys with ferro-titanium addition shows that the largest percentage of the inclusions in the Ti treated HY130 and 4130 alloys were TiN and MnS. The HY130 heat E steel, which was melted without protective atmosphere, had more than twice the amount of inclusions as the other steels. heat E, which was prepared without protective atmosphere, had more than twice the amount of TiN and. MnS inclusions as the 4130 and the SPAL processed HY130, heat D. The total average inclusion density was determined for each steel and the results are given in Fig. 4a. HY130 heat A, which was Al and Ca treated and cast without a filter, and HY130 heat E, which was Ti-treated and prepared without a protective atmosphere, had the greatest inclusion densities at 267 and 247 particles/mm, 2 respectively. In fact, both HY130 heats A and E contained more than twice the amount of inclusions as the other steels. The 4325 steel had the least amount of inclusions at 46 particles/mm. 2 The HY130 heats D and C had the largest average inclusion size of greater than 3 μm (Fig. 4b). This is an interesting result since HY130 heats D and C were produced under SPAL and a vacuum, respectively. The smallest (< 2.5 μm) inclusions were found in the HY130 heat A and Fe-Mn-Al-C steels. It should be noted that direct correlation between nitrogen and oxygen contents (Fig. 1) with inclusion measurements (Fig. 4) is not possible since inclusions with diameters less than 1 µm are not included in the later. The total average volume fraction of inclusions was determined for the different steels and the results are given in Fig. 4c. The reported inclusion volume fraction was obtained directly from the inclusion area fraction coverage. HY130 heat E, which was Ti treated and melted without protective cover, had the largest volume fraction of inclusions with a value of The 4325 was shown to have the least volume fraction of inclusions at The volume fraction of inclusions was statistically the.

6 Paper pdf, Page 6 of 17 (a) (b) (c) (d) Fig. 4. (a) The total inclusion density as a function of alloy type and heat number shows that HY130 heat A, which was prepared under Ar cover and cast without a filter, and HY130 heat E, which was Ti treated and prepared without protective atmosphere, have more than twice the amount of inclusions as the other steels. (b) HY130 heat D, which was Ti treated and melted under SPAL, had the largest average inclusion size while HY130 heat A showed the smallest average inclusion size. (c) HY130 heat E had the largest volume fraction of inclusions and the 4325 showed least amount. (d) The average inclusion spacing was the highest for the 4325 and the lowest for Al and Ca treated HY130 heat A, which was cast without a filter. same for the other alloys regardless of deoxidation practice or chemistry. Average inclusion spacings for the different steel castings are shown in Fig. 4d and are reported as a combination of both non-metallic inclusions as well as voids and pores. Average inclusion spacing was determined directly from the automated feature analysis data, where the closest neighbor distance was determined from the Cartesian coordinates of each measured particle. The 4325 steel had the largest spacing with an average distance of 63 µm and HY130 heats A and E had the smallest inclusion spacing with values of 29 and 30 µm, respectively METALLOGRAPHY All of the Cr and Mo steels had similar microstructures of tempered martensite with an average martensite block width of 3.5 to 6 μm. Optical micrographs of HY130 heats B and D are shown in Figs.5a and b. Both show similar microstructures and fineness of the tempered martensite. Large TiN inclusions, greater than 5 µm in diameter, were observed in the microstructure of the HY130 heat D specimen (Fig. 5b). The 4325 and 4130 steel are shown in Figs. 5c and 5d and are similar to the microstructures of the HY130 steels. The microstructures of the Fe-Mn-Al-C specimens are shown in Figs. 5e and 5f. The Fe-Mn-Al-C specimen was aged for 13 hr at 530C (986F) and shows a fully austenitic matrix with carbide and or intermetallic phases that have precipitated on grain boundaries. The Fe-Mn-Al-C alloy that was aged for 60 hr at 530C (986F) also shows precipitation on grain boundaries as well as approximately 1% of ferrite as shown in Fig. 5f.

7 Paper pdf, Page 7 of 17 TiN (a) (b) (c) (d) Ferrite (e) (f) Fig. 5. Optical micrographs of the (a) HY130 heat B, (b) HY130 heat D, (c) 4325, and (d) 4130 steel specimens show similar microstructures of tempered lath martensite. (c) A single TiN is shown in the HY130 heat 4 specimen. The optical micrographs of the Fe-Mn-Al-C alloy (e) aged at 530 C for 13 hr and (f) aged at 530C for 60 hr have been etched to reveal the dendritic pattern of the austenitic matrix. Small amounts of ferrite (< 1%) were present in the 60 hr aged specimen in (e). Small angle grain boundaries are highlighted by precipitation of carbides or intermetallic compounds in the Fe-Mn-Al-C alloys. All specimens were etched with 2% nital followed by etching with Klemms s reagent.

8 Paper pdf, Page 8 of 17 TENSILE PROPERTIES Tensile properties were evaluated from select heats. The results are listed in Table 3. Yield and ultimate strengths of the HY130 specimens were shown to be insensitive to deoxidation practice and both the yield strength (YS) and ultimate tensile strength (UTS) were statistically the same at 970 and 1180 MPa, respectively. Total elongation to failure (% e f ) and percent reduction in area (% RA) was lowest in HY130 heat E, which was Ti treated and produced eutectic Type II sulfides. Yield and ultimate tensile strengths of the 4325 alloy were lower than the HY130 specimens at 957 and 1094 MPa, respectively. Elongation to failure was nearly the same between the Al and Ca deoxidized HY130 heat A castings and the 4325 castings at 11%. However, the HY130 heat A specimen showed the greatest reduction in area at 43% compared to the 4325 steel with a reduction in area of 25%. Tensile properties of the lightweight and austenitic Fe- Mn-Al-C steel are also shown in Table 3 for two different aged conditions. Increasing the aging time from 26 hr to 60 hr at 530C (986F) increased the hardness from 32 to 38 HRC, increased the yield strength from 728 to 873 MPa and increased the ultimate tensile strength from 873 to 953 MPa. However, the elongation decreased from 28 to 20% with increased aging. At equivalent hardness values, ie. 36 to 38 HRC, the Fe-Mn-Al-C steel showed the greatest elongation to failure. The large cast grain structure of the Fe-Mn-Al-C steel produced an irregularshaped necked region that was difficult to accurately evaluate for reduction of area measurements and was thus not reported. given in Fig. 6 as a function of crack displacement. The Fe-Mn-Al-C steel specimens that were aged for 13 hr at 530C (986F) to a hardness of 32 HRC had the greatest total energy at fracture as shown in Fig. 6e. Upon further aging of the Fe-Mn-Al-C steel to a hardness of 38 HRC, the absorbed energy of fracture decreased (Fig. 6f) and the alloy displayed similar fracture energies as the Al and Ca treated HY130 heat B (Fig. 6a) and the Al killed 4325 steel as shown in Fig.6c. The dynamic fracture toughness values of the respective steels are given in Table 4. Of the steels heat treated to similar hardness values of 36 to 38 HRC, the HY130 heat B obtained the highest average toughness of 165 kj/m 2. The 4325 and the Fe-Mn-Al-C steel (aged for 60 hr at 530C [986F)) had similar toughness values of 153 kj/m 2. The Fe-Mn-Al-C steel achieved a much higher toughness of 376 kj/m 2 when aged to a hardness of 32 HRC (13 hr at 530C [986F]). Of the Al killed and Ca treated, HY130 steels, heat C proved to have the lowest toughness of only 114 kj/m. 2 Hot isostatic pressing (HIP) of HY130 Heat C specimens resulted in a 20% reduction in the amount of porosity, a 46% reduction in the size of the remaining voids and recovery of toughness to an average value of 162 kj/m. 2 The HY130 and 4130 steels, which were Ti-treated, showed a severe decrease in toughness as compared to steels deoxidized with Al only or by a combination of Al and Ca. The HY130 heat E steel, which was Ti treated and contained eutectic Type II sulfides had the lowest toughness of only 59 kj/m 2. DYNAMIC FRACTURE TOUGHNESS Select load versus displacement curves from the instrumented Charpy impact tests are given in Fig. 6 for the various steels. Force data was smoothed using a sixth order polynomial fit to eliminate the specimen ringing and to calculate the dynamic fracture toughness. The HY130 heat B, 4325, and both of the Fe-Mn-Al-C steel specimens show Type IV behavior with stable crack extension to failure. The 4130 steel exhibited Type III behavior with stable crack propagation followed by a drop in the load versus displacement curve, which is indicative of unstable crack propagation, subsequent crack arrest and continued stable crack growth (Fig. 6). The HY130 heat E showed brittle Type I behavior with unstable crack propagation and failure of the specimen soon after reaching the peak load (Fig. 6b). The total energy of the fracture process for the representative specimens is also FRACTOGRAPHY The fracture surface of the HY130 heat B specimen is shown in Fig. 7 a. Failure was mostly ductile in nature with microvoid initiation and coalescence around globular calcium aluminate and (Ca, Mn)S inclusions. Additions of ferro-titanium to the HY130 steels produced large, > 7 µm, TiN inclusions that are shown fractured on the surface of the broken specimen from HY130 heat D (Fig. 7b). In many cases, voids were shown to nucleate from the fractured TiN inclusions. In one instance, a quasi-cleavage crack initiated from the fractured TiN particle as shown in Fig. 7b. The lack of sustained fracture by cleavage would indicate that the failure was not hydrogen related Table 3. Tensile Properties of the Various High Strength Steel Alloys Alloy/Heat Heat Treatment YS, MPa UTS, MPa % e f % RA Hardness, HRC HY130 - A 2 Q&T 970 ± ± ± 2 43 ± 9 36 ± 1 HY130 E 2 Q&T 975 ± ± ± ± 1 36 ± Q&T ± 1 Fe-Mn-Al-C 1 Aged C 728 ± ± ± ± 1 Fe-Mn-Al-C 1 Aged C 873 ± ± ± ± 1 1 Elongation measured in a 25 mm gage, 2 Elongation measured in a 30 mm gage, 3 Elongation measured in a 50 mm gage

9 Paper pdf, Page 9 of 17 (a) (b) (c) (d) (e) (f) Fig. 6. The force vs. displacement curves obtained from the instrumented Charpy tests for the (a) HY130 heat B,(b) HY130 heat E, (c) 4325, (d) 4130, (e) 13 hr aged Fe-Mn-Al-C steel and (f) 60 hr aged Fe-Mn-Al-C steel show Type IV behavior for the HY130 heat B, 4325, and both aged Fe-Mn-Al-C steel specimens. The HY130 heat E specimen in (b) shows brittle Type I behavior and failure soon after peak load. The 4130 specimen in (d) exhibits type III behavior. The 13 hr aged Fe-Mn-Al-C specimen had the greatest total fracture energy but the lowest peak load.

10 Paper pdf, Page 10 of 17 Table 4. Dynamic Fracture Toughness of the High Strength Steels Alloy/Heat Deoxidation practice Hardness (HRC) Avg. J Id (kj/m 2 ) Inclusion density, #/mm 2 HY130 heat A Al + Ca 37.7 ± ± HY130 heat B Al + Ca 36.9 ± ± HY130 heat C Al + Ca 36.2 ± ± 8 HY130 heat C, HIP Al + Ca HY130 heat D Al + Ti + RE 37.7 ±.8 88 ± 9 78 HY130 heat E Ti 37.4 ± 1 59 ± Al only 37 ± Al + Ca + Ti 38 ± Fe-Mn-Al-C aged C Fe-Mn-Al-C aged C Al 32 ± ± 69 Al 38 ± ± The HY130 heat D alloy also contained a large amount of globular MnS inclusions. These globular MnS inclusions failed by decohesion of the matrix resulting in ductile tearing (Fig. 7b). In the HY130 heat E specimen, fracture was also shown to initiate at cracked TiN inclusions. However, a great deal of the fracture surface showed failure facilitated by the presence of large, > 50 µm, eutectic cells of Type II MnS inclusions on prior dendritic boundaries (Fig. 7c). The fracture surface of one of the 4325 specimens is shown in Fig. 7d and shows ductile fracture around globular MnS inclusions. Fracture of the 4130 steel specimen revealed the presence of a high density of cracked TiN particles that were closely spaced to one another as shown in Fig. 7e. No evidence of embrittlement from tempering was observed in the 4130 steel despite the unorthodox tempering temperature used. As in the Ti treated HY130 specimens, fracture was observed to have initiated in the TiN particles. The fracture surface of the austenitic Fe-Mn-Al-C steel specimen, which was aged to a hardness of 38 HRC, is shown in Fig. 7f. Failure of the Fe-Mn-Al-C specimen was ductile in nature with microvoid nucleation around globular MnS inclusions or AlN particles. DISCUSSION TENSILE PROPERTIES Yield and tensile strengths of the Cr and Mo steels were relatively uniform whereas deoxidation practice had a significant influence on ductility. Total elongation to failure decreased by 50% and percent reduction in area decreased by almost 66% in HY130 heat E when compared to HY130 heat A. The total inclusion density between these two steels was largely the same at 267 particles/mm 2 for the HY130 heat A and 247 particles/mm 2 for the HY130 heat E. However, additions of ferro-titanium in the HY130 heat E produced large, greater than 5 µm, TiN inclusions and eutectic Type II sulfides that contributed to reduced ductility. Strength and elongation values for the Fe-Mn-Al-C steel were in good agreement with values reported by Van Aken et al. for similar compositions of alloys aged to equivalent hardness levels. 18 DYNAMIC FRACTURE TOUGHNESS Fracture in the high strength steels was dependent on the nature of second phase particles within the microstructure. During fracture, void nucleation can occur at the particle/matrix interface by decohesion or can be created at brittle particles that fracture during deformation. Failure then progresses by the growth of voids due to further plastic strain and then linkage of the voids leading to final fracture. 1 In the presence of a sharp crack or a notch, crack progression is often dependent on inclusions and second phase particles that fracture or nucleate voids within the crack tip process zone. 22 In the case of the Ti-treated materials it is apparent that under dynamic loading the fracture strength of the TiN is less than the cohesive strength of the particle and matrix interface. In contrast, the calcium aluminate particles fail by decohesion, which indicates that the fracture strength of the calcium aluminate particle is greater than the cohesive interface strength.

11 Paper pdf, Page 11 of 17 (a) (b) (c) (d) (e) (f) Fig. 7. (a) The fracture surface of a specimen from HY130 heat B shows ductile failure with microvoid nucleation and coalescence of voids around globular calcium aluminate and (Ca, Mn)S inclusions. (b) The surface of the HY130 heat D specimen shows void nucleation initiated at fractured TiN inclusions. A singular and isolated area of quasicleavage fracture which initiated from the corner of a coarse (>7 µm) TiN particle is highlighted by the arrow in (b). (c) Fracture in the HY130 heat E specimen was initiated both by fractured TiN and by the presence of Type II MnS stringers in interdendritic regions. (d) The 4325 specimen shows ductile failure around globular MnS. (e) The fracture surface of the 4130 alloy shows fracture through a high density of TiN particles which are closely spaced. (f) Failure in the Fe-Mn-Al-C steel was mostly ductile in nature with microvoid nucleation around MnS and AlN particles.

12 Paper pdf, Page 12 of 17 The dynamic fracture toughness of the high strength steels in the current investigation was a strong function of inclusion content, morphology and spacing. Heats deoxidized with Al or a combination of Al and Ca contained mostly globular calcium aluminate and MnS inclusions that facilitated void rupture at the particle interface as shown in Figs. 7a and 7d with Type IV elastic-plastic behavior (Figs. 6a and 6c) and a higher toughness than heats treated with ferro-titanium. Toughness of the Al and Ca treated HY130 and 4325 castings varied between 114 and 165kJ/m 2, which is much higher than reported literature values of 92 kj/m 2 for a wrought 4340 steel. 23 It was determined that the uncharacteristically low toughness of HY130 heat C, which was vacuum induction melted, was due to the presence of a high volume fraction, , -4 of porosity. Hot isostatic pressing was effective at reducing the volume fraction of porosity from to and the average pore/void size from 2.7 to 1.4 µm which resulted in a restoration of fracture toughness to 162 kj/m 2, which was statistically equivalent to the toughness value of 165 kj/m 2 that was observed from HY130 heat B. It should be noted that all of the cast steels had microporosity on the order of 1.0x Ferro-titanium treatment produced large and widespread precipitation of TiN in HY130 and 4130, which contributed to Type I, failure in HY130 heat E and Type III failure in the 4130 steel with an accompanying loss in toughness (Figs. 6b and 6e). In addition to the steels in this study, coarse TiN precipitation has been shown by several researchers to lead to brittle fracture and reduced toughness in HSLA (high strength low alloy) steels as determined by instrumented Charpy testing The deleterious effects of TiN during fracture was also supported by the fractography that revealed large void formation from fractured TiN (Fig. 7b). Fracture toughness was lowest in the HY130 Heat E steel and this was attributed to a large concentration of TiN particles as well as Type II MnS inclusions, which produced unstable crack growth along the eutectic cells (Fig. 7c). It is well known that without rare earth additions, the toughness of Cr and Mo steels is a strong function of S content because of the tendency to form eutectic Type II MnS inclusions in the vicinity of grain boundaries. 5, 27 Strong deoxidizers such as Ti can magnify this effect. 5, 6 Addition of misch metal to the ferro-titanium treated HY130 heat D produced globular MnS inclusions and improved the toughness from 59 to 88 kj/m. 2 However, the presence of TiN appears to adversely affect void nucleation by particle cracking rather than interface void nucleation and as a result produce larger initial voids with commensurate loss in toughness. No evidence of hydrogen embrittlement could be discerned for the Ti-treated steels. In fact, in a study by Kuslitskii et al. 28 on a plain carbon steel with different types of nonmetallic inclusions, H-embrittlement was found to be mainly a function of surface area of the inclusion/metal interface rather than chemistry. Elongated silicate and alumina inclusions trapped more H and contributed to a reduction in notched tensile strength. TiN particles were the least efficient at trapping H and the TiN containing steels showed the least susceptibility to H-embrittlement. 28 High work hardening rates associated with high Mn-C austenitic steels have been shown to produce excellent Charpy V-notch (CVN) toughness. In the solution treated condition and at room temperature, Hale and Baker recorded a CVN breaking energy of 206 J for a wrought Fe-30Mn-8Al-1C alloy. 29 Precipitation of κ-carbide during aging decreases notch toughness and recent studies of the current Fe-Mn-Al-C steel gave an average value of 115 J for the room temperature breaking energy of a casting aged in the hardness range of 29 to 32 HRC. 20 The dynamic fracture toughness of the lightweight Fe-Mn-Al-C steel was reported to exceed that of the Q and T Cr and Mo steels. 20 Fe-Mn-Al-C specimens that were aged for 13 hr at 530 C (986F) to a hardness of 32 HRC obtained much higher toughness, 376 kj/m 2, as compared to the Q and T Cr and Mo steels. However, with an increase in aging to a hardness of 38 HRC, the Fe-Mn-Al-C steel attained statistically equivalent DFT values, 153 kj/m 2, as the Al and Ca treated HY130 and 4325 steels. Inclusions present in the Fe-Mn-Al-C steel were multiphase inclusions composed mainly of AlN and globular MnS. In most cases, the faceted AlN particles were coated with globular MnO or MnS that prevented early void nucleation at the AlN corners or possible particle fracture. An attempt was made to model the effects of inclusion density, size, volume fraction and spacing on the dynamic fracture toughness of the high strength steels studied here. The energy of fracture in high strength steels is due in part by the size of the plastic zone and the type and distribution of second phase inclusions. 22 The plastic zone size increases with applied stress and is controlled by the yield strength and fracture toughness. Crack propagation occurs when the plastic zone size equals the spacing of the second phase inclusions. Large plastic zones in front of the crack tip increase the chance of fracture taking place by direct linkup of voids produced by fibrous rupture of the matrix. As the yield strength of the material decreases, the size of the plastic zone increases and this correlates with an increase in toughness. 22 When the inclusion spacing in the matrix is large, the interaction between the plastic zones of the individual particles is limited and this will permit large strains to occur before void coalescence. Low energy ductile fracture can occur when there is a high density of fractured particles with small spacing between them. A number of researchers have attempted to model the effect of second phase particles on ductile rupture of both ferrous and non-ferrous alloys. McClintock 30 and Beremin 31 assumed that fracture occurred when the volume fraction of cavities reached a critical value. Broek 32 studied the role of second phase particles in

13 Paper pdf, Page 13 of 17 initiating microvoids in a variety of Al alloys and determined that larger particles with greater inclusion spacing produced considerable void growth before failure while smaller and closer spaced particles nucleated smaller voids, which grew quickly to failure. Broek 32 therefore suggested that the fracture strain was inversely proportional to some function of the volume fraction of particles or voids and that toughness would be expected to increase with decreasing particle density. Observations by Deimel et al. on the relationship of inclusion density to the fracture toughness of high Cr and Mo pressure vessel steel showed that fracture toughness was inversely proportional to the total inclusion density. 33 A plot of inclusion density versus dynamic fracture toughness for the steels in the current study is given in Fig.8. regular array of spherical particles with L 0 = D 0 (π/6) 1/3 V f. -1/3 The fracture toughness was shown to fit the following model. K IC ~ V f -1/6 [2(π/6) 1/3 YE D 0 ] 1/2 Equation 1 Where Y is the tensile yield strength and E is Young s modulus. Prior researchers 36, 37 have also shown that the fracture toughness of high strength steel varied linearly with An attempt was made to fit the above model which relates the quasi-static fracture toughness to inclusion size and distribution to the dynamic fracture toughness of the steels studied here. The relationship between inclusion spacing and the volume fraction of inclusions is given in Fig. 9 and shows an increase in the spacing between inclusions with decreasing volume fraction. The slope of the regression line in Fig. 9 gives a value of -0.37, which is in reasonable agreement with the 34, 35 value of given by Hahn et al. For application of this model to the current study, K Id must be converted to an appropriate J Id value by use of the following 1, 33 equation: K Id = (EJ Id ) 1/2 Equation 2 Substituting Equation 2 into Equation 1 yields the following relationship between J Id and the volume fraction of inclusions. J Id = V f -1/3 [2 (π/6) 1/3 YD 0 ] Equation 3 Fig. 8. Inclusion density was found to have an inverse relationship with toughness. Two distinct trends are observed that suggest that void nucleation by particle cracking (Ti treated steels) is inherently worse than void nucleation (Ca and Al treated steels) at the particle and matrix interface. The fracture toughness of both the Cr and Mo steels, as well as, the lightweight Fe-Mn-Al-C steel show a decrease in toughness with increasing inclusion density (Fig. 8). The HY130 heat B steel showed a 10 kj/m 2 increase in toughness over the 4325 steel even though the HY130 steel had more than twice the total inclusion density as the 4325, suggesting a beneficial effect on toughness from the additional Ni content. However, Ni was shown to have no noticeable effect on toughness in the Ti treated steels and the 4130 steel obtained the highest toughness among the Ti treated steels at 94 kj/m 2. This suggests that in the Ti treated steels, the favorable effect of Ni at toughening the matrix is outweighed by the deleterious effect of TiN particle fracture. Hahn et al. modeled the effect of inclusion diameter (D 0 ), spacing (L 0 ), and volume fraction (V f ) on the plane strain fracture toughness, K IC, of aluminum alloys. 34, 35 The model by Hahn et al. uses an ideal case consisting of a Fig. 9. Inclusion spacing decreased as the volume fraction of inclusions increased. The slope of the line gives a value of which is comparable to the 34, 35. expected value of Figure 10 shows the dependence of dynamic fracture toughness on the volume fraction of inclusions for the Cr and Mo steels as well as the Fe-Mn-Al-C alloy. A linear relationship between V f -1/3 and DFT is hard to claim for the steels that were deoxidized with Al or a combination of both Al and Ca. In fact, DFT was relatively insensitive

14 Paper pdf, Page 14 of 17 to the volume fraction of inclusions for all of the Al and Ca treated steels. The steels that were treated with ferrotitanium, however, show that DFT may vary linearly with V f, -1/3 but additional testing would be required to verify the trend (Fig. 10). At the same nominal volume fraction of inclusions, the Ca and Al treated steels obtained much higher toughness than the steels treated with Ti. Again, this may be related to particle cracking in the Ti treated steel versus interface void nucleation for Al and Ca treated steel. with Ti. This decrease in toughness can again be attributed to the nucleation of large voids resulting from TiN particle fracture (Fig. 7b). Conversely, in the steels which were deoxidized with Al or a combination of Al and Ca, voids were nucleated by decohesion of the matrix from the mostly globular calcium aluminate or MnS inclusions which did not crack during dynamic loading. Fig. 10. DFT was observed to be a linear function of V f -1/3 in the Ti treated steels. At equivalent volume fractions of inclusions, the Al and Al + Ca treated steels showed much higher toughness than the steels which were treated with Ti. Fig. 11. Both the Al and Al + Ca treated castings show a linear increase in toughness with increasing inclusion spacing, except for the 4325 casting which did not fit the expected trend. The castings that were treated with ferro-titanium also showed a linear trend but at lower toughness. From the model by Hahn et al., fracture toughness (K Ic ) is a parabolic function of the inclusion spacing as given in the following relationship. 34 K IC ~ [2EYL 0 ] 1/2 Equation 4 Combining Equations 2 and 4 gives a linear relationship between J Id and the average inclusion spacing. J Id ~ 2YL 0 Equation 5 Figure 11 shows a general trend between increasing particle spacing (L 0 ) and increasing DFT for all of the Al and Al plus Ca treated castings and possibly the Ti treated steels albeit lower. According to Equation 3, the relationship between D 0 and J Id should also be linear. A plot of average inclusion diameter versus DFT for the high strength steels is shown in Fig. 12. Toughness was found to increase linearly with D 0 for the Al and Ca treated steels. However, while there is a general increase in toughness with increasing particle size in the Ti treated heats, a direct and linear relationship is difficult to claim. At an equivalent average inclusion diameter, D 0, the Al and Ca deoxidized steels obtained almost twice the toughness of the steels that were treated Fig. 12. Although toughness increased linearly with increasing average inclusion diameter for the Al and Al + Ca deoxidized castings, a linear relationship is less obvious in the ferro-titanium treated castings. At a similar average inclusion size, the steels which were treated with ferro-titanium showed up to a 100 kj/m 2 loss in toughness when compared with steels which were deoxidized with Al or a combination of both Al and Ca.

15 Paper pdf, Page 15 of 17 In sufficiently ductile materials, fracture is the combination of void nucleation (by particle/matrix decohesion or particle cracking), void growth and ultimate coalescence to failure. The true fracture strain is then a function of the strain to nucleate voids, ε n, and the strain to grow them to failure, ε g. ε fracture = ε n + ε g Equation 6 In general, ε n increases if particles are small and round, but decreases when particles are faceted, brittle and clustered. The Brown-Embury 38 equation relates the strain associated with void growth to the volume fraction of inclusions as seen in the following: [( ) ( ) ] Equation 7 True fracture strain as determined from tensile testing was used along with the Brown-Embury equation to determine the strain to nucleate, ε n, for the Al and Ti treated steels. However, all of the calculated values of ε n were negative, suggesting that the strain to fracture was almost entirely a function of the strain for void growth for the steels in the present study. Deoxidation with Ti produced a large density TiN particles that cracked during deformation producing large voids, which decreased the distance between major voids and reduced the toughness. A similar observation for the toughness was observed between HY130 heat C HIP and the as cast, vacuum melted HY130 heat C. Hot isostatic pressing was effective at reducing the volume fraction of porosity in heat C from to and reducing the average pore or void size by 46%, from 2.7 to 1.4 µm. Upon healing the large voids by HIP, the toughness increased by 42% going from 114 kj/m 2 to 162 kj/m 2. CONCLUSIONS The dynamic fracture toughness of quenched and tempered Cr and Mo steels with Ni contents of 0, 1.56 and 5.5 wt.% was evaluated with regard to deoxidation practice and compared with the toughness values obtained for a lightweight Fe-30.40%Mn-8.83%Al-1.07%Si- 0.90%C-0.53%Mo steel aged to an equivalent hardness. The highest toughness, 165 kj/m 2, was obtained for the 5.5 wt.% HY130, which was Al killed, Ca treated and cast utilizing a 64 ppi filter. A high amount of porosity drastically decreased the toughness of the HY130 heat C to 114 kj/m 2. However, hot istostatically pressing of specimens from HY130 heat C was effective at closing large voids and restoring toughness to 162 kj/m 2. Statistically speaking, the 4325 and the lightweight Fe-Mn-Al-C achieved equivalent dynamic fracture toughness with values of 153 kj/m 2. At a lower hardness of 32 HRC, the Fe-Mn-Al-C alloy obtained a much higher toughness of 376 kj/m. 2 The dynamic fracture toughness of the high strength steels in the current investigation was a strong function of the density, type, morphology and distribution of the nonmetallic inclusions. Cr and Mo steels that were treated with Ti showed as much as a 100 kj/m 2 reduction in toughness in comparison with the Al and Ca treated steels. The reduction in toughness observed in the ferro-titanium treated steels was attributed to void nucleation at coarse TiN particles, which fractured during dynamic loading as well intergranular fracture facilitated by eutectic Type II MnS inclusions. Addition of rare earths in the form of misch metal increased the toughness of the ferro-titanium heats to 88 kj/m 2 by converting Type II MnS inclusions to globular MnS. Deoxidation with Al and Ca produced mainly globular calcium aluminate and Type I MnS inclusions. which promoted ductile failure by void nucleation resulting from decohesion of the particle from the matrix. Increasing the inclusion spacing and average inclusion diameter generally increased the toughness, while increasing the density and volume fraction of inclusions decreased the toughness. At equivalent inclusion diameter and distribution, the Al and Ca deoxidized heats had much higher toughness than the heats treated with ferro-titanium. Addition of Ni from 1.6% in the 4325 steel to 5.4% in the HY130 resulted in a 10 kj/m 2 increase in toughness in a steel with three times the inclusion density. Therefore, for maximum toughness in cast steels with higher inclusion contents, the Ni contents in Q and T Cr and Mo steel should be kept at levels greater than 5%. However, the benefit of additional Ni content was not apparent for the steels treated with Ti. Therefore, the use of Ti in Cr and Mo steels is strongly discouraged if high toughness is required. However, if Ti treatment is necessary for control of nitrogen, it should be accompanied by additions of rare earth elements in the form of misch metal to eliminate the formation of eutectic Type II sulfides. The best practice is to use Al and Ca deoxidation in heats prepared under a protective atmosphere and cast using an appropriate filter. ACKNOWLEDGMENTS This work was supported in part by Army Research Laboratory under contracts from Battelle Memorial Institute (contract W911NF-07-D-0001) and Benet Laboratories (contract W15QKN ). Laura Bartlett was also supported by a U.S. Department of Education GAANN fellowship under contract P200A The following foundries are gratefully acknowledged for providing test materials: Waukesha Foundry, Inc., Nova Precision Casting Corporation and Precision Castings, Inc. The authors also gratefully acknowledge Bodycote Inc. for HIP of selected test bars as well as Mr. Brandon Ensor, Mr. Tyler Preall, and Mr. Joseph Brookshire for help with specimen preparation.

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