Introduction

High strength and toughness alongwith good weldability are required for ad-vanced structural steels. The history ofthese steels, specifically the high-yield(HY) and high-strength low-alloy (HSLA)series, and their use for naval combatantships has been reviewed by Czyryca (Ref.1). Recent research conducted at North-western University (Refs. 2–8) has devel-oped a next-generation alloy that improvesupon the properties of the HSLA steels.NUCu-140 is a high-strength low-carbon(HSLC) steel that exhibits excellent me-chanical properties and a simple chemicalcomposition. NUCu-140 has a ferritic ma-trix microstructure with nanoscale Cu-richprecipitates as the primary strengtheningagent. In comparison to HSLA-100, use ofNUCu-140 can reduce costs in three ways:reduction and elimination of expensive al-loying elements, reduction in total mate-rial used due to increased strength levels,and savings in production via simple pro-cessing techniques (Ref. 9).

Lis et al. (Ref. 10) investigated the in-fluence of Cu concentration in HSLA-100,and observed that the addition of copperfrom 1.6 to 2.0 wt-% enhances the fracture

toughness (the standard specification forCu in HSLA is 1.15 to 1.75 wt-%). This isattributed to the nature of the copper pre-cipitates and their role in austenite forma-tion. Although the mechanism was notfully defined, HSLA-100 in the overagedcondition results in incoherent ε-Cu pre-cipitates, and these precipitates allow foreasier formation of new stable austenitethat can provide a strong barrier to growthof cleavage cracks in the ferrite matrix.Thus, an increase in copper concentrationcan lead to an increased amount of newaustenite that can be formed, which inturn, provides resistance to cleavage crackpropagation. An enhanced resistance tocleavage crack failure will result in highertoughness. Conversely, copper precipitatesthat are coherent with the matrix restrictplastic deformation and limit the forma-tion and growth of microvoids, which leadsto decreased toughness (Ref. 11).

More recent research has been con-

ducted on NUCu-140 in order to deter-mine the fracture properties of the basemetal. It was found that ductile fractureoccurs over a range of temperatures andstrain rates as the NUCu-140 base mate-rial exhibits a failure mode of microvoidcoalescence in all cases, with the dimplesmore homogeneously dispersed at higherstrain rates (Ref. 12). Some limited workhas been conducted to investigate the me-chanical properties of the HAZ and hasshown that there is local softening in theHAZ following welding. The local soften-ing was attributed to coarsening and dis-solution of copper precipitates since thegrain size of the ferritic matrix throughmost of the HAZ remained unchanged(Refs. 13, 14).

More recently, Farren et al. performeddetailed microstructural characterizationand modeling studies to determine theevolution of the copper precipitates dur-ing weld thermal cycles with four differentpeak temperatures corresponding to thefour regions of the HAZ: subcritical(675°C), intercritical (800°C), fully recrys-tallized (900°C), and coarse grain (1350°C)(Ref. 14). This research confirmed that thelocally softened HAZ occurred because ofcoarsening and dissolution of copper pre-cipitates, and fracture occurred in this lo-cally softened region during tensile testing.The evolution of the copper precipitatesduring weld thermal cycles was deter-mined via local-electrode atom-probe(LEAP) tomography and kinetic model-ing. The subcritical region of the HAZshowed partial dissolution of the copperprecipitates, and all other HAZ regionsshowed complete dissolution on heating,followed by varying amounts of reprecipi-tation on cooling.

Although the changes in precipitatemorphology and resultant hardness andstrength in the HAZ have been explored,the fracture toughness behavior of the dif-ferent HAZ regions for this alloy has notyet been investigated in detail. Thus, theobjective of this research is to understandhow microstructural and precipitate evo-lution affects the fracture behavior in var-ious regions of the HAZ for NUCu-140steel.

Fracture Toughness of Simulated Heat-Affected Zones in NUCu-140 Steel

A detailed investigation was conducted on the relationship between copper precipitate evolution and HAZ toughness

BY B. M. LEISTER AND J. N. DUPONT

KEYWORDS

Fracture ToughnessHeat-Affected ZoneHigh-Strength SteelLow-Carbon SteelCopper PrecipitatesB. M. LEISTER (bml204@lehigh.edu) and J. N.

DUPONT are with Department of Materials Sci-ence and Engineering, Lehigh University, Bethle-hem, Pa.

ABSTRACT

The fracture toughness of simulated heat-affected zones (HAZ) in NUCu-140steel was investigated. The subcritical, intercritical, and fully recrystallized HAZ re-gions exhibited a reduction in hardness and increases in toughness relative to the basemetal. The base metal and these HAZ regions all exhibited a fracture mode of mi-crovoid coalescence. The coarse-grain HAZ exhibited a slight increase in hardnessand a mixed fracture mode of cleavage/microvoid coalescence, but the toughness wascomparable to that of the base metal. The impact toughness was higher in the longi-tudinal-transverse (LT) orientation compared to the transverse-longitudinal (TL) ori-entation. The increased toughness of the subcritical, intercritical, and fullyrecrystallized regions was attributed to softening associated with changes in the cop-per precipitates during welding. Differences in impact toughness observed betweenthe LT and TL orientations were due to anisotropy from the rolling procedure, andthe anisotropy persisted throughout the HAZ samples.

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Procedure

The chemical composition of theNUCu-140 plates used for this research issummarized in Table 1. The alloy wasgiven an initial homogenization heat treat-ment at 1150°C for 3 h. The alloy was thenhot rolled at 950°C to a final thickness of12.7 mm and air cooled. Next, the plateswere solutionized at 900°C for 1 h, waterquenched, and then aged at 550°C for 2 hand air cooled. The plates used for testing

had a yield strength of 807 MPa and a ten-sile strength of 848 MPa.

HAZ thermal simulations were con-ducted on a Gleeble 3500 series thermo-mechanical simulator. The thermal cycleswere derived from calculations made withthe Sandia optimization and analysis androutine (SOAR) (Ref. 15) representativeof a 750 J mm–1 heat input, and were con-trolled using the Gleeble QuikSim™ soft-ware package. The specific peaktemperatures were chosen to represent

the subcritical (675°C), intercritical(800°C), fully recrystallized (900°C), andcoarse-grain (1350°C) heat-affected zoneregions. The complete thermal cycles as-sociated with each region are shown in

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Fig. 1 — SOAR and Gleeble profiles used for NUCu-140-simulated HAZ samples. A — Subcritical(675°C); B — intercritical (800°C); C — fully recrystallized (900°C); D — coarse-grain HAZ (1350°C).

Fig. 2 — Schematic showing the labeling code forCharpy impact specimens and orientation of im-ages in Fig. 4.

Fig. 3 — Dimensions of fracture toughness barused for testing. All dimensions are in mm, andthe sample is 8 mm thick.

Fig. 4 — Light optical photomicrographs showingbanding in cross-rolled NUCu-140 base metal. A —Longitudinal direction; B — transverse direction.Etched with modified Winsteard’s reagent.

Fig. 5 — Light optical photomicrographs of NUCu-140. A — Base metal; B — subcritical; C — inter-critical; D — fully recrystallized; E — coarse-grainHAZ microstructures. Etched with 3% nital.

A

A

A

B

B

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C D

C D

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Fig. 1, which show both the calculatedthermal cycles from SOAR and the lin-earized thermal cycles used for the Glee-ble simulations. These peak temperatureswere based on recent dilatometry resultsin which the Ac1 and Ac3 temperatureswere found to be ≈750° and 850°C, re-spectively (Ref. 14). Samples examinedunder a light optical microscope were pre-pared using standard metallographic tech-niques and etched using either ModifiedWinsteard’s reagent or 3% nital. Micro-hardness was performed on a LECOM400-FT tester using a Vickers indenter,300-g load, 15 s dwell time, and measure-ments were done using CAMS-Win™ pro-gram. Grain size measurements wereconducted according to ASTM E-112(Ref. 16) using the three circle method,and five fields were measured for eachsample. Charpy impact testing was con-ducted at room temperature for both thelongitudinal-transverse (LT) and trans-verse-longitudinal (TL) orientations [asspecified by Hertzberg (Ref. 17) andshown in Fig. 2] according to ASTM A370(Ref. 18) and E23 (Ref. 19) using a 700 ft-lb impact tester. The longitudinal direc-tion for these samples was taken to be inthe direction of the longest dimension ofthe plate, corresponding to the majorrolling direction. Fracture toughness test-ing using the J-integral was performed ac-cording to ASTME-1820 (Ref. 20), with asingle edge notched bend (SENB) sampleof size 8 × 16 × 80 mm (Fig. 3), with a ma-chined notch depth of 4 mm. These frac-ture toughness samples were taken fromthe TL orientation within the plate. Aver-age and standard deviation fracture tough-ness values were calculated from five testsfor each material condition.

Precracking was conducted on a servo-hydraulic Instron frame, with a clip gaugeattached to the sample and interfaced withthe control unit. The precrack was inducedat a starting load of 13.3 kN and frequencyof 10 Hz until the a/W ratio was 0.5, wherea is the total length of the crack and the ma-chined notch and W is the height of thespecimen (16 mm for these specimens, asshown in Fig. 3). Fracture toughness testing

was performed in displacement control onan Instron 5567 load frame using a resist-ance curve test method with a three-pointbend span of 65 mm. Each cycle utilized acrosshead displacement load of 0.075 mmfollowed by an unload displacement of0.025 mm, both at a rate of 0.25 mm min–1.Fracture surfaces were observed on an FEIXL-30 scanning electron microscope(SEM) using an accelerating voltage of 15kV. Energy-dispersive spectroscopy (EDS)analysis and imaging of particles observedon the fracture surfaces was performed on aHitachi 4300SE/N SEM using an accelerat-ing voltage of 15 kV.

Results

As shown in Fig. 4, the microstructureof the base metal shows banding in boththe longitudinal and transverse directions.Figure 2 is a schematic showing the sampleorientation within the rolled plate for the

images provided in Fig. 4. The banding ismore apparent in the longitudinal direc-tion. The microstructures associated withthe base metal and each HAZ thermalcycle can be seen in Fig. 5. As shown inFig. 5A–D, the base metal, subcritical, in-tercritical, and fully recrystallized HAZeach exhibit equiaxed ferrite grains of sim-ilar size. The matrix microstructure of thecoarse-grain HAZ (Fig. 5E), exhibits amore complex microstructure with regionsof acicular ferrite, bainite, and possiblymartensite.

Figure 6 shows the results of the grainsize measurements for the base metal andthe four simulated HAZ regions. There isno significant difference between the basemetal and subcritical, intercritical, andfully recrystallized HAZ, but there is alarge increase in grain size for the coarse-grain HAZ. The grain size measurementsof the coarse-grain HAZ samples repre-sent the prior austenite grain size, whereas

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Table 1— Chemical Composition of NUCu-

140 (wt-%)

Element NUCu-140Al 0.65C 0.04Cu 1.35Fe BalMn 0.47Nb 0.07Ni 2.75P 0.009S 0.002Si 0.47

Fig. 6 — Grain size as a function of peak temper-ature for NUCu-140-simulated HAZ samples.

Fig. 7 — Vickers microhardness as a function of peaktemperature for NUCu-140 simulated HAZ samples.

Fig. 8 — Charpy impact toughness as a function of peak temperature for the simulated NUCu HAZ sam-ples. A — TL orientation; B — LT orientation.

Fig. 9 — Crack growth resistance curves (J-Rcurves) for NUCu-140-simulated HAZ samples.

Fig. 10 — Average fracture toughness (JQ) as a func-tion of peak temperature for NUCu-140-simulatedHAZ samples.

A B

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measurements for the other samples rep-resent equiaxed ferrite. The Vickers mi-crohardness, shown in Fig. 7, shows a

decrease from the base metal (295 HV) tothe minimum hardness of 228 HV for thefully recrystallized HAZ, followed by anincrease to the peak hardness of 305 HVin the coarse-grain HAZ.

Charpy impact toughness results forboth the TL and LT orientations are shownin Fig. 8A and B, respectively. Both orien-

tations display an increase in impact tough-ness from the base metal up to the fully re-crystallized HAZ, followed by a slightdecrease in toughness for the coarse-grainHAZ. There are significant differences intoughness between the two orientations,with the LT orientation exhibiting a highertoughness compared to the TL orientation.

The results of the fracture toughnesstests are shown in Figs. 9 and 10. Figure 9shows representative J-R curves for each ofthe thermal cycles and the base metal plot-ted as data points along with the ASTM-re-quired exclusion lines (Ref. 20). The resultsof all of the fracture toughness tests areplotted in Fig. 10 as averages with errorbars. These results show that the base metalhas the lowest fracture toughness (JQ) of224 kJ m–2 increasing up to the maximum of425 kJ m–2 in the fully recrystallized HAZ,followed by a decrease to 255 kJ m–2 in thecoarse-grain HAZ. According to ASTMstandard E1820, these JQ values meet thecriteria to be classified as JIC.

The fracture surfaces of the samples arepresented in Figs. 11 and 12. Figure 11shows low-magnification SEM images ofthe entire fracture surface of each sample aswell as “splitting” that was observed normalto the direction of the primary crack growthplane. A schematic of the splitting and pri-mary crack growth plane are shown in Fig.13. The higher magnification images seen inFig. 12 show a failure mode of microvoid co-alescence for the base metal, subcritical, in-tercritical, and fully recrystallized HAZ(Fig. 12 A–D, respectively). The coarse-grain HAZ samples (Fig. 12E) displayed a

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Fig. 11 — Low-magnification SEM images of fracture surfaces. A — NUCu-140 base metal; B — sub-critical; C — intercritical; D — fully recrystallized; E — coarse-grain HAZ. Note that the fracture sur-face above the white line was induced following testing.

Fig. 12 — SEM images showing fracture surfacesof NUCu-140-simulated HAZ samples showingductile failure via microvoid coalescence. A —Base metal; B — subcritical; C — intercritical; D— fully recrystallized HAZ. Areas of microvoidcoalescence and cleavage fracture are present inE, the coarse-grain HAZ.

Fig. 13 — Schematic showing relative orientationof primary crack growth plane and splits.

A

A

B

B

C

C

E

D

D E

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mixed mode of microvoid coalescence andcleavage fracture.

Discussion

The grain size measurements (Fig. 6)show the differences between the basemetal, subcritical, intercritical, and fullyrecrystallized HAZ are within experimen-tal error. Therefore, the changes in hard-ness (Fig. 7) and fracture toughness (Fig.

10) properties for these regions of theHAZ can be attributed to changes in thecopper precipitates that occur during thethermal cycle.

Previous work by Farren et al. (Ref. 14)has shown that the copper precipitates un-dergo considerable change in the HAZduring exposure to the weld thermal cycle.Figure 14 shows local electrode atomprobe (LEAP) data for a NUCu-140 welddescribing the average radius, numberdensity, and phase fraction of copper pre-cipitates as a function of thermal cyclepeak temperature. These results were ac-quired from the HAZ of a gas metal arcweld, and the peak temperatures fromeach HAZ region were estimated usingthe SOAR routine. Note that there is aconsiderable decrease in the precipitateradius and phase fraction with increasingpeak temperature. Kinetic modeling re-sults have indicated these changes occurby a combination of coarsening and disso-lution. For the subcritical HAZ, partialdissolution occurs upon heating, followedby a small amount of reprecipitation dur-ing cooling. For the intercritical, fully re-crystallized, and coarse-grain HAZ, fulldissolution occurs followed by varyingamounts of reprecipitation during cooling.

The increase in fracture toughnessfrom the base metal through the fully re-crystallized HAZ is due to the decrease inphase fraction and radii of copper precip-itates within these regions that lead tosoftening and increased plasticity. Evi-dence for this is observed by crack tipblunting that occurs prior to the onset ofstable crack extension. As an example, acomparison of the J-R curves of the sub-critical and fully recrystallized HAZ re-gions is seen in Fig. 15 where the fullyrecrystallized HAZ sample follows theblunting line for higher crack extensions in

comparison to the subcritical HAZ sam-ple. This indicates that more energy wasabsorbed by the fully recrystallized HAZsample due to the plastic deformation oc-curring in the form of crack tip blunting.There is a similar drop in phase fraction ofcopper precipitates for the coarse-grainHAZ sample, but its fracture toughness islower than that of the other HAZ regions.The reduction in toughness in this regioncan primarily be attributed to the coarsergrain size. The change in matrix mi-crostructure may also contribute to thechange in fracture toughness. However, itis important to note that the toughness ofthis region is still comparable to that of theunaffected base metal.

The splitting observed normal to the di-rection of the primary crack growth planewas examined in the base metal samplewhere a larger number of smaller splits wereobserved (Fig. 11A) and in the fully recrys-tallized sample where the ridges were muchlarger and deeper — Fig. 11D. Figures 16Aand 17A show the line of linked microvoidsahead of the splits for the base metal andfully recrystallized HAZ samples, respec-tively. These voids were commonly ob-served to develop ahead of the splits andappear to form by particle/matrix decohe-sion. Figure 16C shows an EDS spectra cor-responding to the particle (denoted by anarrow) shown in Fig. 16B for the base metalsample, and Fig. 17B and C detail the sameinformation for the fully recrystallized HAZsample. The EDS spectra show significantpeaks for Fe and Al and smaller peaks fromNi and Cu. Further work is required inorder to determine the exact nature of theseparticles.

The difference in Charpy impact tough-ness between the TL and LT orientation iscommonly observed (Ref. 17) and can be at-tributed to anisotropy due to rolling. The

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Fig. 15 — Close-up view of crack tip blunting regionof NUCu-140 JR curves for the intercritical (675°C)and fully recrystallized (900°C) HAZ samples.

Fig. 14 — LEAP tomography data from a NUCu-140 gas metal arc weld showing radius, numberdensity, and phase fraction of copper precipitates asa function of thermal cycle peak temperature.

Fig. 16 — A and B — SEM photomicrographsshowing a line of microcracks and particles aheadof a split in base metal sample; C — EDS spectrumof particle at the edge of a microcrack.

Fig. 17 — A and B — SEM photomicrographs show-ing a row of microcracks and particles ahead of a splitin fully recrystallized HAZ sample; C — EDS spec-trum of particle at the edge of a microcrack.

A B

C

A B

C

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crack plane in the TL samples is orientedalong the rolling direction where inclusionsbecome elongated during rolling. Crackgrowth along this direction is relatively easy.In contrast, the crack plane of the LT sam-ples is perpendicular to the rolling direction,so the toughness is higher. It is interestingto note that the anisotropy persists through-out the HAZ samples, indicating that thenature of the inclusions and their influenceon toughness is not significantly affected bythe weld thermal cycle. Also note that thefracture toughness results shown in Fig. 10were derived from samples in the TL orien-tation (due to material limitations). Frac-ture toughness results from the LTorientation would be expected to producesimilar trends but higher toughness values.Thus, the results shown in Fig. 10 can beconsidered conservative.

Conclusions

The fracture behavior of simulatedheat-affected zones (HAZ) in NUCu-140was studied via fracture toughness testing,impact testing, scanning electron, and lightoptical microscopy. The following conclu-sions were drawn from this research.

1) The increase in fracture toughnessfrom the base metal through the fully re-crystallized HAZ is attributed to a de-crease in the phase fraction and radii ofcopper precipitates caused by coarseningand dissolution of the particles during thethermal cycles. This results in softeningand increased energy absorption due toplastic deformation.

2) The toughness in the coarse-grainHAZ is slightly lower than that of theother HAZ regions. This is primarily at-tributed to the coarser grain size, and thechange in matrix microstructure may alsocontribute to the reduced fracture tough-ness. However, the toughness of this re-gion is still comparable to that of theunaffected base metal.

3) Ductile fracture occurs by microvoidcoalescence in the base metal, subcritical,intercritical, and fully recrystallized HAZregions. The coarse-grain HAZ exhibited

a mixed fracture mode of cleavage and mi-crovoid coalescence.

4) The difference in Charpy impacttoughness observed between the TL and LTorientation is caused by anisotropy due torolling. The anisotropy persists throughoutthe HAZ samples, indicating that the na-ture of the inclusions and their influence ontoughness is not significantly affected by theweld thermal cycle.

Acknowledgments

The authors would like to acknowledgethe financial support provided by the Of-fice of Naval Research under grantsN00014-07-1-0331 and N00014-09-1-0361.

References

1. Czyryca, E. J. 1993. Advances in highstrength steel technology for naval hull con-struction. Key Engineering Materials (Switzer-land). pp. 84, 85, 491–520.

2. Isheim, D., and Seidman, D. N. 2004.Nanoscale studies of segregation at coherentheterophase interfaces in α-Fe based systems.Surface and Interface Analysis 36: 569–574.

3. Isheim, D., Gagliano, M. S., Fine, M. E.,and Seidman, D. N. 2006. Interfacial segrega-tion at Cu-rich precipitates in a high-strengthlow-carbon steel studied on a sub-nanometerscale. Acta Materialia. 54: 841–849.

4. Kolli, R. P., and Seidman, D. N. 2007.Comparison of compositional and morphologi-cal atom-probe tomography analyses for a mul-ticomponent Fe-Cu steel. Microscopy andMicroanalysis 13: 272–284.

5. Gagliano, M. S., and Fine, M. E. 2004.Characterization of the nucleation and growthbehavior of copper precipitates in low-carbonsteels. Metallurgical and Materials TransactionsA. Physical Metallurgy and Materials Science.35A: 2323–2329.

6. Isheim, D., Kolli, R. P., Fine, M. E., andSeidman, D. N. 2006. An atom-probe tomo-graphic study of the temporal evolution of thenanostructure of Fe-Cu based high-strengthlow-carbon steels. Scripta Materialia 55: 35–40.

7. Kolli, R. P., and Seidman, D. N. 2008. Thetemporal evolution of the decomposition of aconcentrated multicomponent Fe-Cu-basedsteel. Acta Materialia 56: 2073–2088.

8. Kolli, R. P., Wojes, R. M., Zaucha, S., andSeidman, D. N. 2008. A subnanoscale study of

the nucleation, growth, and coarsening kineticsof Cu-rich precipitates in a multicomponent Fe-Cu based steel. International Journal of Materi-als Research (Zeitschrift fur Metallkunde). 99:513–527.

9. Fine, M. E., Ramanathan, R., Vaynman,S., and Bhat, S. P. 1993. Mechanical propertiesand microstructure of a weldable high perform-ance low carbon steel containing copper. Inter-national Symposium on Low-Carbon Steels forthe 90s, pp. 18–21.

10. Lis, A. K., Lis, J., and Jeziorski, L. 1997.Advanced ultra-low carbon bainitic steels withhigh toughness. Journal of Materials ProcessingTechnology (Netherlands) 64: 255–266.

11. Das, S. K., Sivaprasad, S., Das, S., Chat-terjee, S., and Tarafder, S. 2006. The effect ofvariation of microstructure on fracture me-chanics parameters of HSLA-100 steel. Materi-als Science and Engineering A, 431: 68–79.

12. Vaynman, S., Fine, M. E., Lee, S., andEspinosa, H. D. 2006. Effect of strain rate andtemperature on mechanical properties and frac-ture mode of high strength precipitation hard-ened ferritic steels. Scripta Materialia 55:351–354.

13. Vaynman, S., et al. 2008. High-strengthlow-carbon ferritic steel containing Cu-Fe-Ni-Al-Mn precipitates. Metallurgical and MaterialsTransactions A, 39: 363–373.

14. Farren, J. D., et al. 2011. Microstructuralevolution and mechanical properties of fusionwelds in an iron-copper based multi-componentsteel. Science and Techonology of Welding andJoining.

15. Sandia National Laboratories. 2004.Software optimizes parameters for automatedwelding. Advanced Materials & Processes 162:18, 19.

16. American Society for Testing and Mate-rials. 2010. ASTM E112-10, Standard Test Meth-ods for Determining Average Grain Size. ASTMBook of Standards.

17. Hertzberg, R. W. 1995. Deformation andFracture Mechanics of Engineering Materials.

18. American Society for Testing and Mate-rials. 2011. ASTM A370-11, Standard Test Meth-ods and Definitions for Mechanical Testing ofSteel Products. ASTM Book of Standards.

19. American Society for Testing and Mate-rials. 2007. ASTM E23-07ae1, Standard TestMethods for Notched Bar Impact Testing ofMetallic Materials. ASTM Book of Standards.

20. American Society for Testing and Mate-rials (2009) ASTM E1820-09e1, Standard TestMethod for Measurement of Fracture Toughness.ASTM Book of Standards.

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