Characterization of microstructural strengthening in the ...

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Acta Materialia xxx (2010) xxx–xxx

Characterization of microstructural strengthening in theheat-affected zone of a blast-resistant naval steel

Xinghua Yu a, Jeremy L. Caron a,b, S.S. Babu a,b,*, John C. Lippold a,b,Dieter Isheim c,d, David N. Seidman c,d

a Department of Materials Science and Engineering, The Ohio State University, Columbus, OH 43221, USAb Welding and Joining Metallurgy Group, Welding Engineering Program, 1248 Arthur E. Adams Drive, The Ohio State University, Columbus, OH

43221, USAc Department of Materials Science and Engineering, Northwestern University, 2220 North Campus Drive, Evanston, IL 60208, USA

d Northwestern University Center for Atom-Probe Tomography, 2220 North Campus Drive, Evanston, IL 60208, USA

Received 14 June 2010; accepted 17 June 2010

Abstract

The influence of simulated heat-affected zone thermal cycles on the microstructural evolution in a blast-resistant naval steel was inves-tigated by dilatometry, microhardness testing, optical microscopy, electron backscatter diffraction and atom-probe tomography (APT)techniques. Coarsening of Cu precipitates were observed in the subcritical and intercritical heat-affected zones, with partial dissolution inthe latter. A small number density of Cu precipitates and high Cu concentration in the matrix of the fine-grained heat-affected zone indi-cates the onset of Cu precipitate dissolution. Cu clustering in the coarse-grained heat-affected zone indicated the potential initiation of Cureprecipitation during cooling. Segregation of Cu was also characterized by APT. The hardening and softening observed in the heat-affected zone regions was rationalized using available strengthening models.� 2010 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Steels; Atom-probe field-ion microscopy (APFIM); Electron backscatter diffraction (EBSD); Precipitation strengthening; Grain boundarysegregation

1. Introduction

The requirements for blast-resistant steels for US Navyhull and deck applications are rigorous: specifically, highimpact fracture toughness (Charpy toughness >115 J(85 ft-lbs) at �64.4 �C (�84 �F)), high-strength (yieldstrength: 1030–1240 MPa (150–180 ksi)) combined withgood formability (strain at fracture: 15%), weldabilityand resistance to hydrogen-induced cracking. A new steel,BlastAlloy 160 (BA-160), was developed at NorthwesternUniversity to meet these requirements [1,2], using a combi-nation of multiscale materials modeling and detailed

1359-6454/$36.00 � 2010 Acta Materialia Inc. Published by Elsevier Ltd. All

doi:10.1016/j.actamat.2010.06.031

* Corresponding author. Tel.: +1 614 292 2440; fax: +1 614 292 6842.E-mail addresses: [email protected] (X. Yu), [email protected] (S.S.

Babu).

Please cite this article in press as: Yu X et al. Characterization of micrtant naval steel. Acta Mater (2010), doi:10.1016/j.actamat.2010.06.03

microstructural characterization. BA-160 is a low-carbonmartensitic/bainite steel strengthened primarily by nano-scale (2–5 nm radius) Cu-rich precipitates and M2C (whereM = Cr, Mo, and V) carbides. The minimum yield strengthof 1103 MPa (160 ksi) was achieved by employing calcu-lated heat treatment conditions. In addition to strengthen-ing by nanoscale precipitates, a maximum Charpy impacttoughness of 176 J (130 ft-lb) at room temperature wasachieved through the suggested precipitation of Ni-stabi-lized austenite precipitates within the matrix. These precip-itates improve the toughness through dispersed phasetransformation toughening. In order to employ these steelsin shipbuilding applications, it is important to understandthe fundamentals of their weldability.

The BA-160 steel has been assumed to be weldablewithout susceptibility to hydrogen-induced cracking based

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ostructural strengthening in the heat-affected zone of a blast-resis-1

http://dx.doi.org/10.1016/j.actamat.2010.06.031

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2 X. Yu et al. / Acta Materialia xxx (2010) xxx–xxx

on a low-carbon concentration (<0.05 wt.%). Thisassumption is based on the tendency for hydrogenembrittlement in the heat-affected-zone (HAZ), whichrequires a hard microstructure (such as martensite witha high carbon content), hydrogen concentration and a tri-axial state of stress [3]. In addition to hydrogen-inducedcracking, the weldability of steel must consider solidifica-tion cracking in the weld metal (WM) region, liquationcracking in the HAZ and WM region, ductility dip crack-ing, reheat cracking, along with the final mechanicalproperties of the various weld regions. The focus of thecurrent research is to develop a scientific basis for theweldability evaluation of BA-160 using a Gleeble� ther-momechanical simulator, optical microscopy (OM),microhardness testing, electron backscatter diffraction(EBSD) and atom-probe tomography (APT). Theobserved microstructure and strengthening behavior arerationalized using computational thermodynamic andkinetic models.

As noted, BA-160 steel is strengthened by nanoscale Cuprecipitates within the banite and martensitic matrix [1,2].The weldability of this alloy is related to stability of theseprecipitates during continuous heating (growth, coarseningor dissolution) and cooling (reprecipitation) conditions ofwelding. Therefore, a brief overview of Cu precipitate sta-bility in Fe–Cu alloy systems is reviewed to provide a basisfor the current work.

2. Cu precipitates in the a-Fe matrix

The Fe–Cu phase diagram exhibits only limited solu-bility of Cu within the a-Fe phase (ferrite) region. How-ever, by austenitizing at temperatures above 900 �C,more than 3.6 wt.% Cu can be dissolved in solution.Quenching this Cu-enriched c-Fe phase (austenite) toroom temperature results in a ferritic phase supersatu-rated in Cu. During subsequent aging at temperaturesin the range of 450–650 �C, Cu precipitates from thesupersaturated a-Fe phase as spherical, metastable,body-centered-cubic (bcc) precipitates [4,5]. Goodmanet al. [6,7], using an atom-probe field-ion microscope(APFIM), detected bcc Cu precipitates with a composi-tion of approximately 50 at.% Fe and 50 at.% Cu for pre-cipitates up to 2.5 nm diameter in a binary Fe–1.4 wt.%Cu alloy. More recent APT investigations exhibit generalagreement with Goodman et al. [2,8] with 50–90 at.% Cuin the core of precipitates. Kozeschnik [9] evaluated theCu precipitate concentration in a-Fe based on thermody-namic data. Employing MatCalc�, it was found that atmole fractions of Cu above 0.3–0.5, bcc Cu clustersbecome thermodynamically stable. For this condition,the “most likely” nucleus Cu content is 92 at.% for Fe–0.75Cu and 37 at.% for Fe–3Cu.

As the bcc Cu precipitates grow to a critical diameterbetween 4 and 12 nm, bcc Cu will transform martensiticallyto a twinned 9R structure [4]. The precipitates are observa-ble by high-resolution transmission electron microscopy


(HRTEM) and they exhibit a characteristic herringbonecontrast pattern for both the bcc and 9R structure. Thebcc and 9R structures can be differentiated by the herring-bone angle. The continued growth of spherical 9R precip-itates eventually leads to the formation of face-centered-cubic (fcc) e-Cu phase with a rod-like shape. Many studiesindicate that the orientation relationship between the fccCu precipitates and the Fe matrix is of the Kurdjumov–Sachs (KS) type [4,10,11].

While copper precipitation from a-Fe during isothermalaging has been studied extensively, limited research hasbeen devoted to the temporal evolution of Cu precipitatesduring non-isothermal conditions. Kimura and Takaki[11] reported that a dispersion of fcc e-Cu precipitatesoccur for both air- and furnace-cooling in an Fe–4 wt.%Cu alloy. For welding conditions, coarsening of Cu precip-itates along martensite lath boundaries were observed byTEM in the HAZ of HSLA-100 steel [12]. Additionally,Cu precipitates were shown to dissolve in re-austenitizedregions of the HAZ. Some Cu precipitates in the fine-grained HAZ (FGHAZ) were assumed to be undissolvedwith no reprecipitation of Cu observed in the coarse-grained HAZ (CGHAZ) [12].

Most characterization of Cu precipitates, specifically theAPFIM experiments, has been performed for ferriticmicrostructures. In the current study, the focus is on Cuprecipitation in martensitic matrix microstructure. Froma thermodynamic standpoint, precipitation phenomena inmartensite and ferrite are similar. Since martensite has ahigh dislocation density (approximately 1014 or 1015 m�2)and small lath dimensions (typically 0.3 � 2.8 � 100 lm3),both the nucleation and growth of Cu precipitates are mod-ified [14]. For example, simulation shows Cu precipitates atdislocations and boundaries grow at the expense of Cu pre-cipitates in martensite after holding at 550 �C for morethan 1000 s [13].

It is agreed that the maximum strengthening of Fe–Custeels is achieved for bcc Cu precipitates that are coherentwith the a-Fe matrix. Strengthening of low-carbon Fe–Custeels is commonly described by modulus strengtheningbased on a model developed by Russell and Brown [14].The difference in the shear modulus of the bcc Cu precipi-tates and the a-Fe matrix results in a difference in disloca-tion energy between these phases, and this differenceimpedes dislocation movement. Since fcc Cu precipitateshave a KS orientation relationship with the a-Fe matrix,there are many slip planes, which results in easy dislocationglide with a negligible hardening effect [15].

Since the dimensions and morphology of Cu precipitatesplay an important role in strengthening BA-160, it is cru-cial to understand the evolution of these precipitates inthe HAZ regions. The goals of the current study are tostudy the morphology, composition and changes in Cu pre-cipitates in the HAZ of BA-160, using APT. Strengtheningand softening in different regions of the HAZ was studiedby evaluating strengthening due to precipitate and mar-tensite high misorientation boundaries.



Table 2The heat treatment procedure for BA-160 steel.

Step Temperature, �C Duration Post-step procedure

1. Austenitization 900.0 1 h Water quench2. Liquid nitrogen

hold�196.0 30 min Air warm to room

temperature3. Tempering 550.0 30 min Water quench4. Tempering 450.0 5 h Air cool to room

temperature

X. Yu et al. / Acta Materialia xxx (2010) xxx–xxx 3

3. Experimental procedures

3.1. Chemical composition and heat treatment of BA-160

The BA-160 experimental material was received in theform of 34.925 mm (1.375 in.) diameter bar stock, providedby QuesTek Innovations LLC, Evanston, IL. The mea-sured chemical composition of BA-160 is given in Table 1.The details of the heat treatment procedure for BA-160 areoutlined in Table 2. Test samples were machined using wireelectrodischarge machining (EDM). The microhardness,following the heat treatment procedure, was verified to bein accordance with the expected 41 HRC value associatedwith the fully hardened condition. The microstructure fol-lowing heat treatment consists of a mixed bainitic/martens-itic matrix with nanoscale Cu precipitates and M2Cprecipitates [1].

3.2. HAZ simulation procedure

To investigate the HAZ microstructural evolution, ther-mal cycle simulations representative of the various HAZregions were performed. For steels, there are four distinctHAZ regions: (1) the subcritical HAZ (SCHAZ), whereno detectable transformation to austenite occurs since thepeak temperature is below the Ac1; (2) the intercriticalHAZ (ICHAZ), where partial transformation to austeniteoccurs; (3) the fine-grained HAZ (FGHAZ), where fulltransformation to austenite occurs at a peak temperatureslightly above the Ac3 temperature; and (4) the coarse-grained HAZ (CGHAZ), where full transformation to aus-tenite occurs at a peak temperature significantly above theAc3 temperature.

The thermal cycle simulations were performed with aGleeble�3800 thermal–mechanical simulator using solidcylindrical samples 6.35 mm (0.25 in.) in diameter and101.6 mm (4.0 in.) long. The sample temperature was con-trolled with a Type K thermocouple, which was wire per-cussion welded at the midsection of the sample. Thesamples were heated to the peak temperature at a linearrate of 100 �C s�1 and free-cooled at a rate simulating ahigh heat-input weld (Dt8/5 = 45 s, see Fig. 1). The simula-tions were conducted with the test chamber at a pressure of1.3 � 10�4 Pa (10�6 torr) to limit surface oxidation andthermocouple detachment. Phase transformation tempera-tures were determined using a dilatometer.

3.3. Microstructural characterization

The microhardness of the simulated HAZ regions wasdetermined utilizing a Leco M-400-H1 hardness-testing

Table 1Chemical composition of the BA-160 experimental material.

Element C Si Mn Cu

wt.% 0.059 0.015 0.001 3.39at.% 0.277 0.030 0.001 3.005


machine employing a 981 N (1 kgf) load, in accordancewith ASTM Designation E-384-08 [15]. Specimens usedfor the microhardness testing and microstructure charac-terization were taken from a single test sample to ensureconsistent results.

To characterize the original microstructure of the steeland the microstructural evolution in the HAZ, a PhilipsESEM FEG-30 scanning electron microscope equippedwith an EBSD camera was used; this microscope has anaccelerating voltage of 20 kV and a spot diameter of5 nm, with a scanning step size of 0.1 lm. EBSD maps wereanalyzed using OIMe Analysis Software.

Coupons (0.3 � 6 � 6 mm3) were cut from the center ofGleeble� specimens utilizing a LECO� VC-50 precisiondiamond saw. APT tip blanks, 0.3 � 0.3 � 6 mm3, werecut from the coupons and electropolished using a two-stepmethod [16]. Initial polishing was performed with a solu-tion of 10 vol.% perchloric acid in acetic acid at 10–25Vdc at room temperature. This was followed by a manuallycontrolled pulsed final-polishing step using a solution of2 vol.% perchloric acid in butoxyethanol at 10–25 Vdc atroom temperature, producing a tip with a radius <50 nm.

The APT data was collected utilizing the local-electrodeatom-probe (LEAP) at the Northwestern University Cen-ter for Atom-Probe Tomography (NUCAPT), Evanston,IL. The data was acquired at a specimen temperature inthe range of 75–85 K under ultrahigh vacuum (UHV) con-ditions of approximately 1.0 � 10�8 Pa (7.5 � 10�11 torr).Short-duration laser pulses (1 nJ pulse-1) are used to inducefield evaporation at a pulse repetition rate of 5 � 105 Hz.The steady voltage on the specimens is up to 12 kV. Thenthe acquired atomic position data are calibrated and recon-structed using the Imago Visualization and Analysis(IVAS�) program.

4. Results and discussion

4.1. Microstructural evolution and kinetics

Plots of temperature vs. time for the HAZ simulationsare provided in Fig. 1. The approximate t8/5 time (time to

Ni Cr Mo Ti Fe

6.8 1.9 0.61 0.016 Balance6.527 2.058 0.358 0.019 Balance



Fig. 1. Experimental measured thermal profile for the HAZ simulations.


cool from 800 to 500 �C) for the simulations was 45 s. Thethermal arrest that occurs around 350 �C on-cooling isindicative of the martensite phase transformation. Theabsence of a thermal arrest for the simulated SCHAZ sam-ple is noted, confirming the 650 �C peak temperature asbeing below the Ac1 temperature, where no transformationto austenite occurs on heating.

The dilatation curves for the simulated HAZ samplesare presented in Fig. 2. The Ac1 and Ac3 temperatures forBA-160, determined from the dilatation data, are660 ± 10 and 810 ± 10 �C, respectively. The on-coolingphase transformation temperatures are summarized in

Fig. 2. Dilatometry data for the simulated HAZs: (a) volume fraction of atransformation strains as a function of temperature; (c) volume fraction of mtransformation strains as a function of temperature.


Table 3. Martensite was found to be the only austenitictransformation product for each of the samples heated toa peak temperature above the Ac1 temperature. Some vari-ations were found for the martensite start (Ms) tempera-tures. The simulated CGHAZ exhibits the highest Ms

temperature, which is also in good agreement with the360 ± 8.4 �C value reported by Saha and Olson [1]. TheFGHAZ simulations exhibited the lowest measured Ms

temperature (Table 3). The non-linear slope of the dilatom-etry curve in the ferrite region suggests the presence ofsome retained austenite in the room temperaturemicrostructure.

A plot of Vickers microhardness vs. peak temperaturefor the simulated HAZ samples is presented in Fig. 3. Alsoshown is a dashed line indicating the base metal microhard-ness value of 402 HV. The microhardness profile exhibits amaximum value in the ICHAZ, with the lowest microhard-ness occurring in the CGHAZ sample. A slight decrease inmicrohardness was experienced for both the SCHAZ andFGHAZ samples. Previous study of Cu-bearing martens-itic naval steel shows only hardening in HAZ due to theformation of untempered martensite [17]. The microhard-ness variation in HAZ of BA-160 is unexpected.

4.2. Morphology and crystallography of the martensitic

matrix

The current picture of lath martensite is that the mar-tensite inside one prior austenite grain can be divided intopackets (groups of laths with the same habit plane), with

ustenite as a function of temperature during heating; (b) correspondingartensite as a function of temperature on-cooling; and (d) corresponding



Table 3The determined on-cooling phase transformation tem-peratures for the HAZ simulations.

Sample Ms. (�C) 0.99 M�f (�C)

CGHAZ 356 201FGHAZ 329 196ICHAZ 343 199

Fig. 3. Measured Vickers microhardness as a function of peak temper-ature for the simulated HAZ samples exhibits anomalous hardening insamples heated to 750 �C.


each packet being further subdivided into blocks (group oflaths of the same orientation or same variant) [18]. Sincethe packet and block size of lath martensite plays an impor-tant role with respect to strength and toughness [19–21],the current investigation of the martensitic matrix aims tounderstand the martensite morphology and crystallogra-phy by utilizing EBSD.

EBSD maps for all simulated HAZ samples are dis-played in Fig. 4. Large-angle grain boundaries (GBs) withmisorientations >15� are indicated by solid lines. The sam-ples exhibit a typical lath martensitic microstructure for allheat treatments. Since no detectable phase transformationoccurrs for the SCHAZ sample, no significant martensiteand prior austenite grain morphology change is observedcompared to the reference base metal specimen. Fine mar-tensite packets with diameters of approximately 1 lm arefound in the ICHAZ. The fine packet size results fromsmall austenite grains formed during heating which didnot have sufficient time to grow, due to the short timeabove the Ac1 temperature. Also visible in the ICHAZare some larger martensite packets (indicated by arrowsin Fig. 4c) which suggests they did not transform to austen-ite. Fine and coarse prior austenite grains are clearlyrevealed in Fig. 4c and d. In the base metal and simulatedHAZ samples with a peak temperature 6900 �C, the mar-tensite packet size is close to the prior austenite grain diam-eter, which implies there is only one packet inside eachprior austenite grain. There are several packets, however,


in one prior austenite grain in the CGHAZ: two packets,designated P1 and P2, and two blocks, designated B1 andB2, are shown in Fig. 4e.

4.3. Cu precipitate characterization

Since the BA-160 base metal in the heat-treated state hasbeen studied extensively [1,2,22], no attempt was made tocharacterize the Cu precipitates in this material. In theheat-treated condition, the number density (Nv) and aver-age radius (hRi) of the Cu precipitates, and Cu concentra-tion of the matrix ðCa

CuÞ are 4.20 � 1023 m�3, 2.4 nm,0.23 at.%, respectively [24]. Hereafter, we refer to these val-ues when discussing the BA-160 base metal condition.

Cu atom maps for different simulated HAZ conditionsare displayed in Fig. 5a, c, e, and g. The regions of highCu concentration are clearly discernible in Fig. 5a and c.The presence of high Cu concentration regions in theSCHAZ and ICHAZ indicates that the Cu precipitatesare not fully dissolved. In Fig. 5e and g, no Cu-enrichedzones are observable as the Cu atoms are almost homoge-neously distributed in the FGHAZ and CGHAZ samples.The absence of Cu precipitates implies their dissolutionin the FGHAZ and CGHAZ.

The Nv,hRi and composition of the Cu precipitates weredetermined by the envelope method [16]. The followingparameters were used: maximum Cu atom separa-tion = 0.6 nm; minimum number of Cu atoms in a clus-ter = 30; grid resolution = 0.12 nm. Precipitatecharacteristics including radius of gyration, Guinier radius,center-of-mass of the precipitates and number of atoms inthe precipitates, were determined for the Cu precipitates.The Nv is calculated using the following equation [16]:

Nv ¼N p1nX

; ð1Þ

where Np and n are the number of particles and total num-ber of atoms detected in the volume, respectively, X is theaverage atomic volume, 1.2 � 10�29 m�3 for bcc Cu, and Bis the detection efficiency of the multichannel plate (MCP),which is taken to be 0.5. The 3-D reconstructions for Cuatoms in the precipitates are displayed in Fig. 5b, d, f,and h for all conditions investigated. Even though Cu-con-centrated regions are not observed in the Cu atom maps forthe FGHAZ and CGHAZ regions, some small precipitateswith hRi = 1.6 nm are detected by the envelope method. Itis realized that the size of the Cu precipitates is dependenton the parameters chosen in the cluster searching algorithmand that the precipitates detected in the FGHAZ andCGHAZ regions could represent artifacts produced in thecurrent analysis. However, the major focus is related tothe evolution of the Cu precipitates in the HAZ, withemphasis placed on comparing the characteristics of theCu precipitates for the HAZ regions and base metal. Theabsolute properties of the Cu precipitates are of secondaryimportance. The summary of the results for the Cu precip-itates is given in Table 4. The concentration of the matrix



Fig. 4. Measured EBSD images from (a) base metal; (b) SCHAZ; (c) ICHAZ; (d) FGHAZ; and (e) CGHAZ show different martensite sub-structures.


and the precipitates, as determined by the envelope meth-od, is given in Tables 5 and 6. The best way to describethe composition of Cu precipitates is to determine the corecomposition of the precipitate by a proxigram method [23].


The interface cut-off with the envelope method is deter-mined by the maximum Cu atom separation; the cut-off oc-curs a few atomic layers into the precipitates. Since thereare thousands of atoms in the Cu precipitates in the



Fig. 5. Three-dimensional LEAP reconstructions of simulated HAZ regions: Cu atom maps are shown in (a) SCHAZ; (c) ICHAZ; (e) FGHAZ; (g)CGHAZ. Corresponding reconstruction of copper atoms in precipitates, as detected by the envelope method, are shown in (b) SCHAZ; (d) ICHAZ; (f)FGHAZ; and (h) CGHAZ.


Please cite this article in press as: Yu X et al. Characterization of microstructural strengthening in the heat-affected zone of a blast-resis-tant naval steel. Acta Mater (2010), doi:10.1016/j.actamat.2010.06.031


Fig. 5 (continued)

Table 5Composition in matrix (at.%).

C Cr Fe Ni Mo Cu V

SCHAZ 0.16 2.03 89.76 6.50 0.47 0.84 0.01ICHAZ 0.11 1.79 86.80 7.90 0.98 2.00 0.01FGHAZ 0.07 2.10 88.66 6.05 0.61 2.43 0.00CGHAZ 0.11 1.94 87.84 6.73 0.36 2.86 0.00

Table 6Composition in Cu precipitates (at.%).

C Cr Fe Ni Mo Cu V

650 0.12 1.34 62.62 6.38 0.33 29.04 0.00750 0.10 1.10 59.48 6.74 0.73 31.48 0.00900 0.00 0.00 11.67 0.00 0.00 88.33 0.00

1300 0.04 0.43 17.22 1.45 0.12 80.66 0.00


SCHAZ and ICHAZ regions, the effect of including atomsat the interface is negligible. On the other hand, since theCu precipitates in the FGHAZ and CGHAZ are muchsmaller, the 8 at.% Cu iso-concentration surface cannotbe generated and a proxigram method cannot be used forcomposition determination. As a result, the present studyemploys the envelope method for determining Cu precipi-tate composition.

For the Cu precipitates detected in the SCHAZ, a largerhRi and smaller Nv compared to the base metal indicatesthat the precipitates are undergoing coarsening. The matrixCu concentration of 0.84 at.% for the SCHAZ is shown tobe significantly greater than the 0.23 at.% for the basemetal [24]. This indicates that when the steel is heated upto the 650 �C peak temperature of the SCHAZ simulation,Cu coarsening occurs first. As the temperature increases,the solubility of Cu in the martensitic matrix increasesand Cu precipitates tend to dissolve, which increases theCu concentration of the matrix.

In the ICHAZ, the number density of the Cu precipi-tates increases slightly compared to the SCHAZ. Becausethe area analyzed by atom probe is relatively small(�103–105 nm3), the small difference in number density ofCu precipitates between ICHAZ and SCHAZ could beconsidered to be not significant. However, based on thesmaller average Cu precipitates radius (2.3 nm) and highCu concentration in the matrix in the ICHAZ, we can rea-

Table 4Radius and number density of Cu precipitates for the different HAZ regions.

As-received SCHAZ

Radius (nm) 2.4 ± 1.2 3.3 ± 1.4Number density (1023 m�3) 4.2 ± 2.2 1.83 ± 0.5


sonably conclude that partial dissolution of the Cu precip-itates occurs upon heating to 750 �C. In the ICHAZ, onlypart of the microstructure is transformed to austenite uponheating. Based on the determined Ac1 and Ac3 temperaturesand dilatometry analysis, the volume fraction of theaustenitized region at 750 �C is approximately 0.90. Also,since austenite has a much larger Cu solubility than ferrite(see Fig. 9), it is assumed that the atom probe data for theICHAZ is from an austenitized region. In the untrans-formed regions, Cu concentration in matrix is expected to

ICHAZ FGHAZ CGHAZ

2.3 ± 1.1 1.6 ± 0.4 1.7 ± 0.52.26 ± 0.7 0.32 ± 0.2 2.02 ± 1.1




be much lower due to the low solubility of Cu in ferrite.Further selected-area investigations of Cu precipitates inthe untransformed regions may prove useful.

A significant decrease in Nv (0.32 � 1023 m�3) and hRi(1.6 nm) was observed in the FGHAZ. The matrix Cu con-centration (2.43 at.%) is very close to the Cu concentration(2.48 at.%) of the bulk material analyzed by atom probe.The low Nv and small hRi indicate that Cu precipitatesare almost fully dissolved upon heating to a peak tempera-ture of 900 �C. The small precipitates detected could beconsidered either non-dissolved from the base metal orre-precipitated during cooling.

For the CGHAZ, a significantly higher Nv

(2.0 � 1023 m�3) and similar hRi (1.7 nm) compared to

Fig. 6. (a) Three-dimensional LEAP reconstruction of regions containing a boCr composition profiles across boundary A; (c) composition profiles across bouICHAZ sample.


the FGHAZ suggests reprecipitation of Cu from the a-Fematrix during cooling. The optimum hRi for strengthening,2–5 nm, is not achieved because of the short time spent inthe Cu precipitation aging temperature range [2].

4.4. Evidence of Cu segregation

In addition to the detection of Cu precipitates, evidenceof Cu segregation at suspected grain boundaries was cap-tured by the APT experiments. Segregation phenomenafor the ICHAZ and CGHAZ regions were observed inthe current study. Fig. 6a shows Cu and Fe atom mapsnear boundaries for the simulated ICHAZ. The regionsof increased Cu concentration are indicated by arrows,

undary for a ICHAZ sample displaying Cu and Fe atoms; (b) Cu, Ni, andndary B; and (d) 8 at.% Cu iso-concentration surface near the boundary in



Fig. 6 (continued)


with boundary A and boundary B as denoted in the figure.The Cu atom map for the ICHAZ reveals that Cu precip-itates near the boundaries are elongated (marked by boldarrows in Fig. 6a) due to segregation.

In refining the analysis of boundary segregation, twocylinders were placed in the reconstructed volume so asto intersect selected boundaries as shown in Fig. 6a. Theradius of each cylinder is 6 nm and the length 24 nm. Com-position profiles along the cylinders are plotted in Fig. 6band c for boundaries A and B, respectively. Compositionprofiles were obtained by generating a proxigram alongthe main axis of the selected cylinder with a selected binsize of 0.2 nm [23]. Cu segregation and Fe depletion wererevealed at boundaries A and B. As expected, the Cu con-centration away from the boundary is approximately


2 at.%, which is comparable to the matrix, with increasingconcentration found towards the boundary. The maximumCu concentration found at boundaries A and B was 17 ± 4and 8 ± 1 at.%, giving coefficients of segregation of 8.5 and4, respectively. It should be noticed that compared withboundary A, boundary B has lower Cu concentration buthigh Cu and Fe atom density in the reconstruction map(Fig. 6a). The high Cu and Fe atom densities at boundaryB could be attributed to the trajectory aberrations effect,which would result in increased atomic density in thereconstruction [25,26].

Composition profiles of Cu and Fe are shown in Fig. 6b.Cu is shown to be enriched on the boundary and depletednear the boundary. Boundaries A and B could either bemartensite lath boundaries or prior austenite grain bound-




aries. Since a relatively small volume of material was ana-lyzed, the boundary type cannot be determined.

Fig. 7 shows the segregation of Cu atoms at boundariesin a CGHAZ sample. A cylindrical region of interest, witha radius of 6 nm and a length of 24 nm, was selected to gen-erate concentration profiles across the boundary. The cor-responding composition profile is plotted in Fig. 7b. Themaximum concentration at the peak for is 20 ± 6 at.%.Compared with the ICHAZ region, more severe segrega-tion of Cu was revealed in the CGHAZ, indicating thatthe segregation of Cu is promoted by higher temperatures.In addition, the segregation of Ni was also observed atboundary B with a maximum concentration of 17 ± 6 at.%.

Fig. 7. (a) Three-dimensional LEAP reconstruction of regions containing marteprofiles across boundary C.


The Cu segregation result implies that as the peak tem-perature of the HAZ increases, the Cu concentration at theboundaries also increases. Since an increase in Cu concen-tration decreases the melting temperature of the Fe–Cu sys-tem, the segregation of Cu may be a concern for HAZliquation cracking. Susceptibility to HAZ liquation crack-ing for BA-160 is currently being evaluated.

4.5. Modeling of strengthening mechanisms

As stated previously, unexpected hardening and soften-ing was observed in the HAZ of BA-160. To interpret thehardening and softening phenomena, the current study

nsite boundaries in a CGHAZ sample and (b) Cu, Ni, and Cr composition




applies available models to evaluate the microstructuralstrengthening in the HAZ regions. Emphasis has beenput on the effects of lath martensite morphology and thesize of Cu precipitates on strengthening.

4.5.1. Martensite matrix strengtheningFor martensitic strengthening, the yield stress of the lath

martensite is calculated by the summation of the hardeningmechanisms [27]:

rYS ¼ r0 þ rs þ rp þþrg þ rp; ð2Þwhere r0 is the friction stress to move dislocations for pureFe, rs is the yield strength increment due to solid solutionhardening, rq is the strengthening term as a function of dis-location density, rg is the grain boundary strengthening,and rp is the precipitate strengthening. In all four condi-tions, Cu is the major element in the matrix. Accordingto Ref. [28], the yield strength increment due to 1 wt.%Cu in solid solution is less than 50 MPa. In addition, thedislocation density is function of Ms. As shown in the dila-tometry results in Fig. 2, the variation of Ms is small. Onthe other hand, the strengthening contribution on yieldstrength from Cu precipitates and martensite block andpacket boundaries could be several hundred MPa [1,28].As a result, in the present study, only precipitation andgrain boundary strengthening are considered in the overallstrengthening.

The packet size and block width in the different regionsof the HAZ were measured from the EBSD maps using themean linear intercept method. Most researchers take thecontributions from block width and packet size intoaccount separately when considering grain boundarystrengthening [21,29]. As two important parameters whichdescribe the morphology of lath martensite, packet size andblock width should be considered together for grainboundary strengthening. Since the block boundaries insidea packet act as obstacles to dislocation motion, slip planelength can be described to a certain extent by block widthand packet size. Assuming that the block is rectangular andthe length of the block is equivalent to the packet size, theaverage slip plane length, M, can be calculated from theblock width, db, and packet size, dp [21]:

M ¼ 1

p=2

Z a cosðdb=dpÞ

0

db

cos hdhþ

Z p=2

a cosðdb=dpÞdpdh

" #: ð3Þ

From the calculated average slip plane length, thestrengthening introduced by block boundaries can be cal-culated from the Hall–Petch equation:

Table 7Packet size, block width, slip plane length, and rg for different HAZregions.

SCHAZ ICHAZ FGHAZ CGHAZ

Packet size, dp (lm) 9.4 1.58 3.36 56.0Block width, db (lm) 6.26 1.48 2.19 12.4Slip plane length, M (lm) 8.20 1.57 2.90 2.53rg (MPa) 127 290 213 72.3


rg ¼ KyM�1=2; ð4Þwhere Ky is 0.363 MPa m1/2 [21]. The calculated grainboundary strengthening contribution is summarized inTable 7.

4.5.2. Cu precipitation strengthening

Due to the complicated evolution of the Cu precipitatesin the HAZ, the strength contributions arising from Cuprecipitation are difficult to model [12]. Many factors, suchas misfit strengthening, chemical strengthening, modulusdifference strengthening and dislocation core–precipitateinteraction strengthening contribute to the overall Cu pre-cipitate strengthening in steels [30]. Since modulusstrengthening based upon the model of Russell and Brown[14] plays the most important role among all the factors,the current focus is on evaluating the component of Cuprecipitation strengthening provided by modulus strength-ening. The Russell–Brown model assumes that the modulusstrengthening effect is due to the relative difference in dislo-cation energy between the matrix and Cu precipitate as aresult of the modulus difference, as a dislocation passesfrom the matrix through the Cu precipitate and back intothe matrix. The critical shear stress increase caused bythe Cu precipitates can be expressed as:

s ¼ 0:8GbL

1� E2P

E2M

� �1=2

; sin�1 EP

EM6 500

s ¼ GbL

1� E2P

E2M

� �3=4

; sin�1 EP

EMP 500

EP

EM¼ E

1 log rr0

P

E1M log Rr0

þlog R

r

log Rr0

ð5Þ

where G is the shear modulus in the matrix, assumed tobe 77 GPa; L is the interprecipitate spacing; b is the Bur-gers vector, which is 0.25 nm; EP is the dislocation en-ergy in the precipitate and EM is the dislocation energyin the matrix; E1P is the dislocation energy per unitlength in the precipitate; E1M is the dislocation energyper unit length in matrix; r is the average radius of pre-cipitates; r0 is the inner cut-off radius, which is 1.2 nm;and R is the outer cut-off radius, which is 1000 r0. It isassumed that the shear modulus for bcc Cu is equivalentto the shear modulus for fcc Cu. The estimated disloca-tion energy ratio per unit length, E1P =E1M , is approxi-mately 0.6 [14]. Substituting the precipitate propertiesobtained from the atom probe analysis into the modulusdifference equation, an estimate of the Cu precipitatestrengthening can be provided.

It is realized that the Russell–Brown strengtheningmodel has some limitations. The shear modulus of bccCu is different from shear modulus of fcc Cu and couldhave a negative value [31]. The inner and outer cut-off radiiare based on estimation. However, since the variation ofCu size and distribution is large, the Russell–Brown modelwas used in the present research for simplicity.



Fig. 8. Predicted strength increment due to martensite sub-structure andCu precipitate compared with experimentally measured hardness changesin the corresponding regions.


Fig. 8 shows the predicted strength difference change,using Eq. (2), for the various HAZ regions using the basemetal microhardness as a reference, Dr ¼ rHAZ � rBM . Anear linear relationship between the predicted strengthand measured microhardness implies that the currentstrengthening models can be used to evaluate the strength-ening of the HAZ regions. It is important to note that themodel could be improved if an accurate bcc Cu shear mod-ulus and effective interprecipitate distance could be esti-mated. This will require measurement of the Cuprecipitate modulus using in situ scattering techniques [32].

5. Suggested Cu precipitate evolution in HAZ

In the HAZ of BA-160, the solid-state transformationsinvolved are complex due to on heating transformationof martensite to austenite, precipitate growth and coarsen-ing, precipitate dissolution, re-formation of precipitates

Fig. 9. (a) Predicted solubility of Cu in austenite, ferrite is shown as a functiostages in microstructure evolution in CGHAZ region.


during cooling, and on-cooling transformation of austeniteto martensite. Fig. 9 is a schematic of the thermal history ofthe CGHAZ divided into five stages. The Cu solid solubil-ity in austenite and ferrite is also plotted as a function oftemperature, evaluated with ThermoCalc� software usingthe TCFE5 thermodynamic database [33]. The Cu solidsolubility in a-ferrite was calculated by considering theequilibrium calculation between a-ferrite, bcc Cu andM3C phases for the BA-160 composition. The Cu solid sol-ubility in c-austenite was calculated by considering equilib-rium between c-austenite, fcc Cu and M3C phases for theBA-160 composition. The proposed stages of Cu precipi-tate evolution are as follows.

Stage 1: At a temperature below the Ac1, on heating, nomartensite is transformed to austenite. Cu precipitatesare expected coarsen.Stage 2: At temperatures between the Ac1 and Ac3, onheating, some martensite is transformed to austenite.The bcc Cu precipitates located in the austenite regionsmay transform to the fcc structure in order to lowertheir free energy. Due to the higher solubility of Cu inaustenite, fcc Cu precipitates will start dissolving withan increase in temperature. The bcc Cu in the untrans-formed martensite may continue to coarsen. It is alsopossible for some partial dissolution of these precipitatesin a-ferrite since the solubility of Cu in a-ferrite alsoincreases (see Fig. 9).Stage 3: At a temperature above the Ac3, all martensitetransforms to austenite. All bcc Cu precipitates willtransform to fcc Cu precipitates and may start dissolv-ing into the austenite phase.Stage 4: Thermodynamic equilibrium calculations pre-dict reducing solubility of Cu in the c-austenite matrixbelow 800 �C. As a result, on-cooling below 800 �Cand above the Ms temperature, Cu precipitates will re-

n of temperature and (b) schematic illustration of hypothesized different




precipitate out of the c-austenite. The Cu precipitatesthat form within the c-austenite are expected to be ofthe fcc crystal structure.Stage 5: At a temperature below the Ms, Cu is notexpected to re-precipitate out of martensite since the dif-fusivity of Cu at that temperature is very low(<10�18 cm2 s–1) [34].

Although Cu precipitation phenomena in Fe–Cu alloyshave been modeled, the above precipitate and matrix evo-lution in BA-160 cannot be predicted by existing computa-tional thermodynamic and kinetic models. In order todevelop comprehensive material and strengthening models,in situ tracking of the crystal structure and modulus of Cuprecipitates is necessary on the time scale of the HAZ ther-mal cycle.

The implications of the above research on the weldabil-ity of Cu-bearing low-carbon high-strength steels are sum-marized below. The research suggests that welding thermalcycles alter both the Cu precipitates and martensite matrixin the HAZ regions of BA-160, affecting both the strengthand the ductility. The segregation of Cu to martensite lathboundaries during the first welding thermal cycle maycause liquation cracking during the second thermal cycleduring multi-pass welding.

6. Conclusions

1. Observed hardening and softening behavior in the differ-ent HAZ regions of BA-160 steel was explained by therelative strength contributions from lath martensitemorphology and Cu precipitates.

2. The martensite packet size and block width play animportant role in the strengthening of martensitic steels.Since prior austenite grain size plays an important roleon block and packet size, prior austenite grain size con-trol in the HAZ is critical.

3. The evolution of Cu precipitates in the HAZ was char-acterized by APT. The Cu precipitates were shown tocoarsen substantially in the SCHAZ and dissolve par-tially in the ICHAZ. A low number density and smallsize of Cu precipitates in the FGHAZ indicates thatCu precipitates are almost completely dissolved duringheating to a peak temperature of 900 �C. A high numberdensity of very fine Cu precipitates in the CGHAZimplies possible reprecipitation of Cu during cooling,which is to be evaluated in future work.

4. A suggested Cu precipitate evolution sequence is pro-posed for describing the dynamics of Cu precipitatesduring continuous heating and cooling.


Acknowledgements

The authors would like to acknowledge financial sup-port from the Office of Naval Research; J. Christodoulou,grant officer. APT measurements were performed at theNorthwestern University Center for Atom-Probe Tomog-raphy (NUCAPT). The LEAP tomograph was purchasedand upgraded with funding from NSF-MRI (DMR-0420532) and ONR-DURIP (N00014-0400798, N00014-0610539, NOOO14-0910781) grants. Michael D. Mulhol-land is thanked for his help with APT data evaluation withthe IVAS software.

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