one-step quenching and partitioning heat treatment of medium carbon low alloy steel

One-Step Quenching and Partitioning Heat Treatmentof Medium Carbon Low Alloy Steel

Fawad Tariq and Rasheed Ahmed Baloch

(Submitted January 4, 2013; in revised form December 2, 2013)

This paper presents the results of novel one-step quenching and partitioning (Q&P) heat treatment con-ducted on medium carbon low alloy steel sheet. Samples were austenitised at 1193 K followed by inter-rupted quenching at 473 K for different partitioning times and finally they were quenched in water.Dilatometry was employed for selection of treatment temperatures. Optical and scanning electronmicroscopy was carried out to examine the microstructural changes. Volume fraction of retained austenitewas measured by x-ray diffraction technique. Resulting microstructures were correlated with themechanical properties such hardness, tensile strength, elongation, impact absorbed energy, etc. The notchtensile and fracture toughness properties of Q&P steels are still lacking therefore notch tensile strength andplain strain fracture toughness tests were conducted and results are reported here. Results of Q&Ptreatments were also compared with the properties obtained by conventional Quenching and Tempering(Q&T) and normalizing treatments. Optimum strength-ductility balance of about 2000 MPa tensilestrength with 11% elongation was achieved in samples quenched at 473 K and isothermally partitioned for100 s. Higher ductility of Q&P steel was attributed to the presence of 6.8% film-type interlath retainedaustenite. Fine-grained martensitic structure with high density of interphase boundaries imparted ultrahighstrength. It was further noted that the impact toughness, notch tensile strength and fracture toughness of1000 s partitioned samples was higher than 100 s partitioned samples. Possible reasons for high toughnessare synergetic effect of recovery of dislocations, partial loss of martensite tetragonality and precipitation offine transition carbides.

Keywords fracture toughness, lath martensite, microstructure,quenching and partitioning, retained austenite

1. Introduction

Increase in the demand for high strength combined withhigh ductility and formability by the automotive and otherindustries has led to the development of Third Generation ofAdvanced High Strength Steels (AHSS) (Ref 1). At present,Third Generation AHSS are in conceptual or pre-developmentstage and extensive research projects are underway to exploredifferent steel compositions as well as the processing tech-niques to achieve the desired strength-ductility combinations.Among various competitive approaches (Ref 2), recent andnovel technique of quenching and partitioning (Q&P) heattreatment has gained wide popularity among the researchers.This innovative non-equilibrium heat treatment route hasattracted attention of many scientists and industrialists becauseof its ability to produce composite microstructure withoutemploying heavy alloying elements and complex treatmentcycles. Further, it exhibits superior properties, particularly

strength, over conventional transformation induced plasticity(TRIP) steels (Ref 3).

The Q&P process was originally proposed by Speer et al.(Ref 4, 5) in 2003 to produce martensitic structure with certaincontrolled amount of retained austenite (RA). The intention wasto achieve the high strength of martensite along with enhancedductility and toughness because of TRIP effect of RA (Ref 6-10). Although, Quenched and Tempered (Q&T) steels are alsomartensitic in structure and often contain little amount of RAbut they differ from Q&P steels in the sense that in Q&P steelconsiderable amount of RA is deliberately induced withinmartensite laths or packets to augment ductility and workhardening rate without significantly decreasing the strength. Inconventional Q&T steels, carbon of supersaturated martensite ismainly utilized in carbide formation during tempering; as aresult less amount of carbon is left for austenite stabilization(Ref 11). Moreover, it is also known that RA present in Q&Tsteel decompose on tempering (Ref 12) contrary to RA in Q&Pwhich is highly enriched with carbon (1-2% by weight) andtherefore thermally stabilized (Ref 13, 14). In addition to this,Si or Al is also added in Q&P to suppress cementiteprecipitation which adversely affects the toughness of theQ&P steels (Ref 15).

Quenching and Partitioning treatment involves rapidlyquenching the partially (ferrite + austenite) or fully austenitisedsteel to below theMs temperature and above the Mf temperatureto produce the known amount of martensite then isothermalsoaking (partitioning step) at either quench temperature (TQ) orabove it to allow migration of excess carbon from supersatu-rated martensite to adjacent untransformed austenite. Hence

Fawad Tariq and Rasheed Ahmed Baloch, Materials Research andTesting Laboratory, Pakistan Space and Upper Atmosphere ResearchCommission (SUPARCO), Karachi 75270, Pakistan. Contact e-mail:[email protected].

JMEPEG �ASM InternationalDOI: 10.1007/s11665-014-0902-2 1059-9495/$19.00

Journal of Materials Engineering and Performance

enriching it and retaining it at room temperature upon finalquench. In this way, the supersaturated austenite is thermallystabilized; transformation to martensite through quenchingbecomes difficult. If the partitioning step is done at TQ then theprocess is referred to one-step Q&P and in case of partitioningabove the TQ it is called two-step Q&P (Ref 16, 17). In bothone and two-step Q&P process the resulting microstructureconsists of multiple constituents: carbon depleted lath martens-ite (formed on initial quench at TQ), carbon-enriched RA(formed as a result of carbon partitioning), high carbon plate ortwin martensite (also called fresh martensite and formed due totransformation of insufficiently stabilized high carbon austeniteon subsequent final cooling) and ferrite (in case of partialaustenitising). Depending upon the carbon content and hard-ening elements present in the alloy, precipitation of finetransition carbides is possible during partitioning step (Ref 18).Sometimes, the precipitation of finely dispersed alloy carbidesis deliberately promoted by tempering after Q&P. This resultsin an augmented strength due to precipitation strengthening.The modified Q&P method is therefore referred to as‘‘Quenching-Partitioning-Tempering (Q-P-T)’’ (Ref 19-22).The formation of brittle cementite is avoided in all conditionsby alloying with significant percentage of Si or Al (Ref 23).

Quite recently, the focus in on designing and producinglarge steel structures based on nano-structural constituents. Themodern microstructure design philosophy relies on 3 M(Metastability-Multiphase-Multiscale) concept (Ref 19). Theexceptional feature of Q&P process is that this heat treatmentprocess satisfies 3 M microstructure criteria. Combination ofsubmicron lath decarburized martensite, plate-type carbon-enriched martensite, nano-size thin-film retained austenite andnanoscale transition carbides offer outstanding set of properties.The required set of mechanical properties can be attained bytailoring the alloy composition, volume fraction, morphologyand size of martensite/RA phases by optimizing the processparameters.

The innovative Q&P approach is designed on thermody-namic model called ‘‘Constrained-Carbon Paraequilibrium(CCE)’’ which states that the full partitioning of excess carbonsolute atoms occurs between martensite and austenite underrestricted equilibrium conditions, i.e., without interface migra-tion, substitutional elements diffusion and carbide precipitation(Ref 4, 17). The maximum amount of metastable RA after finalquench to room temperature is then predicted by OptimumQuench Temperature selection methodology devised by Speerand coworkers (Ref 24) and Clarke and coworkers (Ref 25).The methodology assumes that all of the carbon escaped frommartensite is employed in adjoining untransformed austenite.But practically, the final microstructure seldom containstheoretically predicted amount of RA due to various competingreactions occurring during the partitioning process discussed inliterature. Readers are suggested to review articles (Ref 2, 17,26) for details.

Several research articles have appeared on Q&P and Q-P-Tprocess in last few years but most of the work was related totwo-step Q&P process. Mostly low carbon steels having

chemical composition close to TRIP steels were investigated(Ref 17-19, 23, 27-34). Relatively less attention was given toQ&P heat treatment of medium carbon steels. Few compre-hensive studies have been found in literature presentingexperimental results of one-step Q&P on medium carbon lowalloy steels (Ref 19, 35). The majority of the authors focused ontensile strength and ductility of one-step Q&P steel. Notchtensile strength, fatigue crack growth and fracture toughness ofone-step Q&P steel has seldom reported. Ivo et al. hasdiscussed the fatigue properties of medium carbon low alloysteels treated by Q&P process (Ref 36). Nevertheless, mediumcarbon low alloy steels are widely used in engineeringapplications requiring high or ultra-high strength. It wastherefore considered valuable to investigate and explore theeffects of one-step Q&P process on microstructures andmechanical properties (particularly notch tensile strength andfracture toughness) of medium carbon low alloy steel. In thepresent study, one-step Q&P heat treatments are conducted onmedium carbon low alloy steel sheet having 1.3% Si, 1% Mnand 1% Cr as major alloying elements. The chosen alloy isactually modified C-Si-Mn-Cr type sheet steel suitable forproducing carbide-free bainite/martensite (CFB/M) duplexmicrostructure upon normalizing. Silicon is added to suppresscementite precipitation (solubility of Si in cementite is verylow) thereby stabilizing the interlath RA in the martensitematrix (Ref 23). Manganese is added to increase the hardena-bility and lower the Ms temperature which results in an increasein volume fraction of RA at room temperature. Moreover, Mn

retards the pearlitic reactions. Chromium is a hardening elementwhich increases the strength of steel and shifts the bainiteformation to longer times. Different partitioning times (PT) at473 K (200 �C) quenching temperature are employed tomonitor the affect of varying PT on amount of RA withinmartensitic structure. Microstructures evolved during Q&Ptreatments are characterized and correlated with the observedmechanical properties. Tensile strength, elongation, strainhardening exponent, Charpy v-notch impact toughness, notchtensile strength and plain strain fracture toughness are deter-mined. The primary aim of this investigation is to get ultra-highstrength steel with ‡ 10% elongation and ‡ 50 MPa

ffiffiffiffi

mp

fracture toughness which can find numerous applications incivil as well as the defense industry.

2. Experimental Procedure

The investigation was conducted on C-Si-Mn-Cr steel sheetof 6 mm thickness. The chemical composition was determinedthrough spark emission spectrometer and given in Table 1. Thesheet was received in annealed condition having ferrite + finepearlite microstructure and hardness of 250 HV. The grain sizeof the sheet was measured to be 23 lm (grain size no. 8) as perASTM standard E 112-96.

The TTT/CCT diagram for the alloy was constructed usingMUCG-83 program (Ref 37) as shown in Fig. 1. From

Table 1 Elemental composition (in wt.%) of steel used in this study

Alloy C Si Mn Cr Mo P S

35CrMnSiA 0.36-0.38 1.23-1.25 0.84-0.87 1.17-1.19 0.02 0.01 0.01


calculated TTT/CCT diagram, the Ms temperature was found tobe 608 K.

In order to determine the critical transformation tempera-tures for selecting the heat treatment cycle, dilatometry (Orton1600D) was done on 20 mm long and 3 mm diameter solidcylindrical normalized sample. The composition of steel is suchthat, CFB/M mixed microstructure formed on air cooling (i.e.,normalizing). Sample was heated from room temperature to1273 K at a rate of 10 K/s and cooled in still air. Thedilatometric curve is shown in Fig. 2. Th represents the onset ofeither cementite formation or austenite decomposition. Sincedilatometer did not have the quenching capability thereforemartensite start temperature (Ms) was taken from TTT/CCTdiagram. Ms was also calculated from the empirical relationship(Ref 38) and found to be approximately 601 K.

Prior to any heat treatment, sheet was normalized at 1193 Kfor 1200 s. Specimens of about 259 259 6 mm were cut fromthe sheet for Q&P heat treatments. The initial quenchtemperature (TQ) of 473 K was selected on the basis ofoptimum quench temperature methodology devised by A.J.Clarke (Ref 31). Moreover, lower TQ is favorable for enhancedstabilization of austenite. The Volume fraction of martensite (atTQ) was calculated to be 0.886 by Modified Koistinen-Marburger relationship proposed by Van Bohemen et al. (Ref39).

The specimens were fully austenitised at 1193 K for 1200 s,followed by rapid quenching in salt bath (NaOH + KOHsolution) maintained at 473 K. Cooling rate of about 60 K/swas employed to avoid bainitic reaction. Partitioning was alsocarried out at 473 K for 10, 100, and 1000 s to allow migrationof C from supersaturated martensite to neighboring austenite.Specimens were finally water quenched to room temperature.K-type thermocouple was spot welded on specimens for precisetemperature monitoring and control. In the partitioning process,the average time for carbon escape from aM fi c and itshomogenization in the adjacent c can be estimated by theequation:

t ¼ x2= 6D; ðEq 1Þ

where t is time required for carbon escape, x is average diffu-sion distance, and D is given by Ref 40:

D ¼ Do exp Q/RTð Þ: ðEq 2Þ

Time for partition of excess C from aM to c at 473 K wascalculated to be less than 5 s and about 79,121 s were requiredfor complete homogenization of untransformed c.

In parallel to Q&P treatments, specimens were alsoquenched and tempered for comparison. Samples were auste-nitised at 1193 K for 1200 s followed by quenching in oil.Tempering was done at 473 K for 1 h followed by air cooling.All heat treatments were carried out in electric resistanceheating muffle furnace. Temperature variation was within±5 K.

Vickers macrohardness measurements were taken on suit-ably prepared samples after each heat treatment cycle. The testswere conducted using a 100 kg load, 10 s dwell time andaverage of at least 5 indents per sample is reported here.Detailed microstructural investigation was carried out usingoptical microscope (Optika, Italy) and scanning electronmicroscope (JEOL JSM6380A) equipped with EDS. Metallo-graphic samples were prepared by adopting standard mounting,mechanical grinding, and polishing techniques. Observationswere made after etching with 2% nital solution.

The volume fraction of RA in Q&P and Q&T specimenwere measured by x-ray diffraction (Bruker AXS D8-Advancediffractometer) using Direct Comparison method (Ref 41). Allthe x-ray measurements were carried out on polished samplesin step scan mode with a step size of 0.02, time per step of 2 s,and 2h range from 35� to 95�. Copper Ka radiation filtered withNi was used at 40 kV and 40 mA. The integrated intensity of(200)a, (211)a peaks were compared with (200)a, (220)c and(311)c peaks for obtaining average value of RA. Carboncontent of RA was calculated from Eq. 3 (Ref 6). The RAlattice parameter was estimated from an average based on the(220)c and (311)c peak position.

ac ¼ 3:555þ 0:044 %wt: Cð Þ ðEq 3Þ

Tension tests were performed on standard sub-size sheetspecimens of rectangular cross-section sliced from directionparallel to the rolling direction. Gage length of 25 mm, width6.3 and 6 mm thick specimens were machined in accordancewith ASTM E 8-99. All the tension tests were performed on150 kN universal tensile testing machine at a crosshead speedof 5 mm/min. Yield strength was determined by 0.2% offsetmethod and percentage elongation was precisely measured byoptical digital image correlation (DIC) method. The detail of

Fig. 1 TTT/CCT diagram generated by MUCG-83 program

Fig. 2 Dilatometric curve of alloy used in this study


non-contact 2D-DIC technique is given in (Ref 42). Tensilestrain hardening exponent (n) was also evaluated as per ASTME 646-98. Average of at least three tensile tests for eachcondition is reported here.

Charpy V-notched (CVN) impact test was conducted onsub-size specimens (5 mm9 10 mm9 55 mm) as per ASTM E23-98. Notch tensile strength was also evaluated using sharp-edge notched sub-size specimen (80 mm length, 4 mm thick-ness, 30 mm width, 5.4 mm notch depth and 60� notch tipangle) according to the guidelines of ASTM E 338-03. Sampleswere tested at a crosshead speed of 2 mm/min and ratio ofNTS/UTS was calculated. Plain strain fracture toughness wasdetermined by using 3-point single edge notch bend (SENB)specimen (99 mm length, 88 mm span length, 22 mm width,6 mm thickness) in accordance with ASTM E 1820-06. Themachined notch length was 9.9 mm, width 1.38 mm, 60�v-notch and tip radius of about 0.21 mm. Fatigue pre-crackingwas carried out at a frequency of 3 Hz and load ratio of 0.1.Once the fatigue crack length of 1.5 mm was achievedspecimens were monotonically loaded till fracture. Candidatefracture toughness (KQ) was calculated from load (P) versusdisplacement (v) curve. After checking all the validity criteriadevised by ASTM E 1820, the KQ was declared valid KIC.

Fracture surfaces of CVN impact specimen and fracturetoughness specimens were examined under stereomicroscope(Meiji Techno, Japan) and SEM from the central location ofeach specimen.

3. Results and Discussion

3.1 Microscopy and XRD

Figure 3 shows the representative optical micrographs ofQ&P steel quenched at 200 �C and partitioned for 10, 100, and1000 s, respectively. Micrograph of Q&T specimen (473 K for1 h) is also presented in Fig. 3. The Fig. 3 shows typicalmartensitic structure containing both lath (ML) and plate-type(MP) martensite. Retained austenite is not seen in 10 and 1000 sQ&P specimens (Fig. 3a, c). However, some islands of RA areclearly visible in 100 s Q&P specimen (Fig. 3b). These RAislands are mostly formed at prior austenite grain boundaries(PAGs). Depending upon the carbon concentration, ML etcheddifferently from MP, i.e., varies in contrast. Larger size plate aMwith high band contrast is likely formed at TQ. The light etchedmatrix is the initial aM formed in the quenching step. Dark-etched plates of martensite (lesser width) in Fig. 3 are eitherformed at TQ or apparently fresh martensite (MF) having highcarbon content. Fresh aM is formed during final quenching inwhich carbon depletion through partitioning did not take placeto a significant extent. However the volume percentage of MF

must be very low because 88.6% aM is estimated to beproduced at TQ. This 88.6% aM contains both ML and MP. It isvery difficult to distinguish between MP and MF because ofidentical morphology. Another type of MF was formed fromless carbon-enriched metastable RA. The martensite formed at

Fig. 3 Optical micrographs showing structure obtained after (a) Q&P at 473 K for 10 s, (b) Q&P at 473 K for 100 s, (c) Q&P at 473 K for1000 s, and (d) Q&T at 473 K for 1 h (RA is in white color)


TQ and MF was transformed from c at different stages havingdifferent carbon contents. In case of Q&T specimen, structuremainly consists of conventional tempered aM. Both morphol-ogies of martensite (ML and MP) are visible in opticalmicrograph in Fig. 3(d). Austenite is not detected in Q&Tspecimens.

Figure 4 shows the selected SEM micrographs of Q&P andQ&T microstructures. Typical acicular aM morphology isobserved in all microstructures. It is noted that multiple phasesare present in the Q&P specimens (Fig. 4a-d). The specimenpartitioned for 100 s contains ML and MP as well as thin-film-like RA, as shown in Fig. 4(b) and (c). The aM laths aregrouped into packets (Fig. 4b). Retained c is mainly located

in-homogeneously between these laths as thin films (Fig. 4c)and occasionally found on packet boundaries. It is witnessed inall microstructures of Q&P steel that the size of blocks andpackets is pretty fine with inherent high dislocation density.Finer sub-grain constituents resulted in higher number ofinterfaces. Significant portion of structure is occupied by thelow angle grain boundaries and interphase boundaries. This isone of the main strengthening sources in martensitic steels.

During partitioning step, excess carbon migrates fromsupersaturated BCT martensite into neighboring untransformedFCC austenite (acting as a potential sink for carbon) assumingCCE condition. This leads to increase in carbon concentrationas well as lattice parameter of c. The c high in carbon is

Fig. 4 SE micrographs showing microstructural constituents in specimens (a) Q&P at 473 K for 10 s, (b) Q&P at 473 K for 100 s, (c) highmagnification pic of Q&P specimen at 473 K for 100 s, (d) Q&P at 473 K for 1000 s, and (e) Q&T at 473 K for 1 h (ML lath martensite, MP

plate martensite, and RA Retained austenite)


chemically stabilized and has lower Ms temperature. Therefore,it is retained upon final quenching. In contrast, less enriched cis unstable and transformed into MF. Another observation is thepresence of very fine metastable carbides inside thick plates ofmartensite (Fig. 4c). The possibility of carbide precipitation inlow carbon ML is very low. The aM formed on initial quench istempered (i.e., autotempering) during partitioning step as inconventional tempering process. Some of the carbon escapingfrom supersaturated ferrite (i.e., martensite) was utilized informing metastable transition carbides. Carbides are assumed tobe transition (e) carbides and their volume fraction is very low;i.e., they are occasionally found in some martensitic plates(Fig. 4c). Transition carbides form in medium carbon steelseven in the presence of Si, as reported by Matlock et al. (Ref43). In fact, higher Si content promotes the e carbide formationat lower temperatures. It is now well recognized that the Si iseffective mainly in suppression of cementite. It is documentedthat various competing reactions occurs during partitioning stepin Q&P treatment (Ref 2, 4, 44). In the initial stages ofpartitioning at lower temperatures, i.e., between 373 and573 K, carbon is more expected to partition into c or trapped ondislocations. Moreover, diffusion of carbon in ferrite (a) ismuch higher than in c, so carbon leaves aM very quickly at aM/c interface (calculated to be 5 s). Segregation of carbon atomsoccurs initially on aM/c interface. Grain boundaries (GBs) actas a reservoir for carbon. This fact is also confirmed recentlythrough AES measurement of carbon concentration on GBs ofnanograined hypoeutectoid Fe-C alloys (Ref 45). In nanometersize grains, strong interface segregation of carbon atoms hasbeen observed. Upon isothermal holding, the excess interfacecarbon either partitions into adjoining c or formed transitioncarbides. The mobility of aM/c interface and carbon saturationat interface decides the relative amount of carbon-enriched RAand transition carbides in the final structure. Under conditionsof CCE, it is assumed that almost all available carbon goes intothe c. The diffusion of carbon in a is faster than in c. Thereforecarbon left aM speedily and build up at the interface. However,it takes time to completely homogenize the c (79,121 s). Thetime for homogenization throughout the c depends on its size,place in the structure and partitioning temperature. However, atintermediate temperatures metastable transition carbide forma-tion becomes more favorable until higher temperatures where itis unstable and is replaced by cementite (Ref 46). In fact, aM/cinterface (where carbon concentration is higher than in bulkgrains) or dislocations forests are favorable site for carbidenucleation. Clarke (Ref 31) and Edmonds et al. (Ref 25) havealso showed from atom probe microscopy that at lowerpartitioning temperatures the c stabilization is less effectiveand e carbide precipitation occurs. He et al. (Ref 47) haveobserved dispersion of transition carbides under transmissionelectron microscope in AISI 9260 steel bar during Q&Pprocess. X-ray diffraction studies by Streicher et al. (Ref 3)have also confirmed the presence of carbides in microstructureproduced by Q&P treatment. Silicon is known to inhibitformation of cementite at high tempering temperatures but haslittle or no effect in suppressing transition carbide precipitationat low tempering temperatures (Ref 17). In fact, Nayaka et al.(Ref 48) have showed that the high Si content can promote ecarbide precipitation at lower temperature. Hence, the presentstudy has further reinforced the thought that the J.G Speer�sideal CCE theory do not holds in real scenarios. However,various contending reactions (for instance, formation of lowerbainite in two-step Q&P process) are operative in Q&P process.

Some researchers have exploit this concept in getting precip-itation strengthening affect by intentionally promoting carbideprecipitation and named this heat treatment as Quenching-Partitioning-Tempering (QPT); another variant of Q&P process(Ref 19-22, 49). The carbon content in QPT steels has beenincreased to compensate the carbon consumption caused bycarbide formation. G. Krauss has also reported that thedispersion of fine-scale transition carbides beneficially contrib-utes toward hardness (Ref 50). In addition, fine transitioncarbides are usually not considered detrimental, whereascementite can be of more concern (Ref 51). Since carbideformation consumes escaping carbon, therefore very lessamount of carbon is available for c stabilization. That is whymore carbide and less amount of RA were detected inspecimens partitioned for 1000 s (Fig. 4d). Comparison ofFig. 4(d) with Fig. 4(a) shows that there has not been anysubstantial tempering of the martensitic microstructure duringpartitioning for less than 100 s. Micro-examination suggeststhat the PT should not be very long because it would lead toextensive carbide precipitations with less percentage of RA.

X-ray diffraction was conducted on Q&P as well as Q&Tspecimens to determine the RA volume fraction. Figure 5presents the selected XRD pattern of Q&P specimen (quenchedat 473 K and partitioned at 473 K for 100 s) and Q&Tspecimen. The maximum RA of 6.8% was calculated in 100 spartitioned specimen. The diffraction pattern in Fig. 5(a) showsoverlapping (111)c and (110)a peaks due to rolling texture.Retained c peaks are absent in Q&T specimen and this depictsthat the % of RA must be less than 3% (Fig. 5b). Least amountof RA (3%) was in 10 s partitioned specimen. It was observedthat %RA increases with increasing PT up to 100 s after whichit decreases on further increasing PT. Moreover, lower thanpredicted RA fraction is measured. Similar trend is alsoreported by De. Moor et al. and Li et al. in a study on CMnSi(Ref 52) and 40SiMnNiCr steels respectively (Ref 35). Thecalculation assumes full partitioning of carbon from aM into cin the absence of carbide precipitation but SEM microscopyverified that the precipitation of transition carbides hasoccurred. This resulted in lower than expected % of RA. Otherpossible reasons for lower % RA are c decomposition byinterface migration or isothermal aM formation (Ref 2). Kimet al. and Santofimia et al. have discussed the possibility ofinterface migration and isothermal formation of aM duringpartitioning step (Ref 53, 54). It should be noted that in eithercase (aM/c interface motion or isothermal aM formation) thefinal phase fractions would be altered. This phenomenon is alsoobserved by some authors (Ref 53-55).

Figure 6 shows the experimentally determined carbonconcentration (in wt.%) of RA after Q&P at 473 K for 10,100, and 1000 s. It is noted that the highest carbon concen-tration of 1.05% (above the To but below paraequilibrium Ac3¢is calculated in 100 s partitioned specimen. It seems that uponinitial quenching to TQ from austenitising temperature andholding at TQ the carbon starts escaping from aM to c. It isbelieved that at low partitioning temperature, a small proportionof carbon atoms are bound with dislocations, at interface aM/cand lath boundaries whereas the large proportion of carbonatoms remain in the normal interstitial positions. With increas-ing holding time above 10 s at partitioning temperature thecarbon concentration (as well as %RA) further increases to1.05% at 100 s PT. Up to this stage the carbide precipitation isslow rather carbon enrichment of c is more active but onprolong partitioning, the kinetics of transformation shift in such


a way that carbon preferentially accumulate more in carbideinstead of RA. For this reason 0.7% carbon is estimated in RAat 1000 s partitioned specimen. Analogous behavior is docu-mented in (Ref 35).

3.2 Mechanical Properties

The true stress-strain curve of normalized, Q&T and Q&Pspecimens are shown in Fig. 7. Normalized specimen has thehighest true strain of all the heat treated conditions but thestrength is comparatively low. It is noteworthy that amongQ&P specimens, the highest level of true strain (0.1150) isobtained in specimen partitioned for 100 s. It is furtherobserved that increasing the PT up to 1000 s resulted in lossof ductility (%elongation). Quenched and tempered specimenexhibits highest true stress with true strain of about 0.085(Fig. 7).

Changes in mechanical properties of Q&P specimens as afunction of PT at 473 K are depicted in Fig. 8. Figure 8(a)illustrates that the ultimate tensile strength (UTS) as well as

Fig. 5 XRD pattern of specimen (a) Q&P at 473 K for 100 s and (b) Q&T at 473 K for 1 h

Fig. 6 Calculated To¢/To phase boundary diagram showing carbonconcentration of RA determined by XRD in Q&P specimens


yield strength (YS) decreases with increasing PT, whereasductility increases up to 100 s after which it declines. Themaximum ductility (i.e., 11.5% tensile elongation) was found in100 s partitioned specimen at UTS of 1975 MPa and YS of1425 MPa. The highest elongation (EL) in 100 s partitionedspecimen is obvious because of presence of about 6.8% RA.Moreover, uniform elongation (UEL) is also highest for 100 spartitioned specimen. Decline in EL and UEL with increasingPT is attributed to less % of RA and high volume fraction oftransition carbides. Additionally, the MF formed as a result ofpartitioning step is untempered and brittle which means thatthis will add less toward ductility. Bhadeshia (Ref 56) has alsopointed out, while working on TRIP assisted steels, that 4-5%of RA contributes very little to ductility. Figure 8(b) demon-strates the change in hardness and product of strength-ductility(PSE) with increasing PT. In early stage of partitioning (i.e.,10 s) the hardness was 615 HV which increases to 630 HV at100 s PT and then decrease upon prolong partitioning at1000 s. Similar trend is noticed in PSE. Highest PSE of22,910 MPa% exists for 100 s partitioned specimen. Highhardness in 100 s specimen is attributed to the precipitation oftransition carbides. Wilson and Russell (Ref 57) have alsoshowed that the hardness and strength of steels can markedlyincrease during the early stages of transition carbide precipi-tation. Reduction in hardness at extended PT is owed to therecovery of dislocation structure, loss of interstitial carbon fromBCT aM and the reduction in aM tetragonality. As conventionaltempering reactions also occur during partitioning, thereforereduction in hardness with PT is rational. This behavior is inagreement with other studies (Ref 35). One question arise hereis that with increasing PT the precipitation of transition carbidesincreases which are supposed to raise the hardness (andstrength also), but here decline in hardness is observed after100 s (Fig. 8b). This reason for such observation is because thecarbon is enriching RA and at the same time carbon is alsoconsumed in carbide formation. Hence the martensite (i.e.,source of carbon) is heavily decarburized resulting in lower

hardness. Hardness and strength are cumulative effects ofvarious microstructural features because different reactions aretaking place concurrently. In general, product of strength andelongation is better in Q&P treated specimens in comparisonwith normalized and Q&T specimens. Ultrahigh strength isowed to very fine lath martensitic/RA multiphase microstruc-ture with high density of interphase boundaries and superiorductility is credited to the presence of blocky as well as thin-film-type RA.

Modification in YS/UTS ratio and strain hardening value (nvalue) with varying %RA and carbon concentration of RA areplotted in Fig. 8(c). It is quite surprising that %RA has morepronounced effect on YS as compared to UTS. Ratio of YS/UTSis 0.87 for 10 s partitioned specimen which decrease to 0.69 for1000 s specimen (Fig. 8c). With increasing carbon concentra-tion of RA, UEL and n value also increases. Choi et al. (Ref 58)have also reported similar observations that the n value of steelwith a higher volume fraction of RA is comparatively higherthan the steel with a lower volume fraction of RA. The increasedstrain hardening with PT is attributed to TRIP effect of carbonstabilized RA. As the mechanical stability of RA increases with%C, the transformation of metastable c to aM occurs at higherstrain levels.

Figure 8(d) shows monotonic increase in notch tensilestrength (NTS) and NTS/UTS ratio with increasing PT. Thespecimen partitioned for 10 s has the lowest NTS/UTS ratio(i.e., 0.33) because of the lowest %RA. Most of the carbon stillaccumulates within the ML and MP and material is highly notchsensitive. Nevertheless, increase in PT resulted in moreretention of c and depletion of martensitic structure, conse-quently brittleness as well as notch sensitivity reduced(Fig. 8d). Charpy v-notch impact toughness (CVN) andfracture toughness (FT) also continuously increases withincreasing PT, as shown in Fig. 8(e). Although 100 s parti-tioned specimen has highest %RA but it is interesting toobserve that the highest CVN (30 J/cm2) and FT (62 MPaffiffiffiffi

mp

) is achieved in 1000 s partitioned specimen. The CVN and

Fig. 7 True stress-strain curves of medium carbon steel in different heat treated conditions


FT of 100 s partitioned specimen should be higher than 1000 sdue to high %RA but they were not found in the experiments.Toughness could be influenced by various parameters. Thepresence of relatively high volume fraction of transitioncarbides in 1000 s partitioned specimen yields better toughness.This idea is also supported by the work of Ooi et al. (Ref 59)who reported the increase in ductility due to autotempering.Furthermore, the ML (having low %C and dislocation sub-structure) is tougher than MP (high in %C). Microstructuralexamination and XRD results explain that with increasing PTthe depletion of carbon from martensitic structure and recoveryof aM occurs with the corresponding increase in RA andtransition carbides. This observation is in conformance with the

previous studies (Ref 57, 60). Further, it is well established thatthe ductility (which is highest for 100 s partitioned specimen) isa property related to the capability of deformation which itselfprimarily influenced by resistance to dislocations and slipsystem. FCC (c) structure has more slip systems than BCC (a).On the other hand, fracture toughness property is related tostructure�s resistance to crack propagation in the presence ofinitial crack. This is highly influenced by the morphology ofphases (like lath or plate morphology in case of martensite,lamellar or spheroidize cementite in case of pearlite, etc.).Dispersion of second phase particles also play vital role indeciding crack path propagation. It is therefore assumed that thehigh CVN and FT of 1000 s partitioned specimen are a result of

Fig. 8 Illustrates the changes in mechanical properties of Q&P specimens as a function of partitioning time


synergetic effect of inter-lath film-like RA, carbon depletion ofmartensite matrix, presence of primarily sub-micro size lathmartensite, carbon depleted plate martensite, recovery ofdislocations and dispersion of fine-scale transition carbides.The mechanism of fracture operative in Q&P steel andinvolvement of microstructure is further discussed under theheading of Fractography. In order to conduct fracture tough-ness test, fatigue cracking is a pre-requisite. The affect of RAon fatigue resistance is essential to talk about. Figure 8(f)presents the relationship between number of fatigue cycles and%RA as a function of PT. One of the key findings of thisinvestigation is that the 100 s partitioned specimen (havinghighest %RA) consume higher fatigue cycles (13,000) toinitiate the crack as compared to other Q&P specimens. It iswell establish that strain-induced transformation of RA intoaM ahead of the crack tends to close the crack because of tworeasons. Firstly, RA absorbs relatively more energy because itis a ductile phase, and at certain local strain level, ittransformed into brittle, high strength aM. This untemperedaM will acts as a strong barrier for the crack. Secondly, thetransformation of c fi aM causes approx. 4.1% expansionwhich compresses the surrounding material (i.e., any discon-tinuity) and inhibits crack opening. In addition to this, ctransformation reduces stress concentration. The beneficialeffect of RA on fatigue crack closure is confirmed from theliterature (Ref 61, 62).

All mechanical properties of Q&P are compared withproperties of same material in normalized and Q&T condition,as listed in Table 2. Normalized structure (containing CFB/M)has high ductility, impact toughness, NTS and FT in compar-ison to Q&P steel (partitioned at 473 K for 100 s) but the UTSand YS are fairly low. Furthermore, strain hardening ofnormalized specimen is noticeably lower (i.e., 0.19) thanQ&P specimen (>0.22) due to less %RA. The Q&T specimenoffers high strength level but with relatively less ductility,strains hardening and impact toughness. Moreover, FT of Q&Tis also significantly lower than the Q&P specimen. It is clearfrom Fig. 8 that excellent combinations of strength, ductilityand toughness have been achieved through Q&P process in thepresent work, when compared to other heat treatmentapproaches. Quenching and Partitioning (Q&P) steels offergreater flexibility in mechanical properties which can be easilytailored by altering process parameters to satisfy differentstrength requirements. Additional thorough investigation isconsidered necessary to fully understand the relationshipbetween microstructure and fracture toughness behavior ofQ&P steels.

Figure 9 exhibits the change in instantaneous strain hard-ening value as a function of true strain for normalized, Q&Tand Q&P treated conditions. Continuously decreasing instan-taneous n value curves were observed for all heat treatments(Fig. 9). However, careful examination reveals that in 100 sQ&P specimen instantaneous n value rises in the range of0.022-0.027 true strain before completely declining. Thisplateau is associated with the strain-induced transformation ofmetastable RA into aM. This TRIP effect is responsible forhigher elongation in 100 s specimen (Ref 63, 64). For all otherQ&P specimens, continuous decline in instantaneous n valuereflects weak TRIP effect due to less %RA (Fig. 9). Normalizedand Q&T specimens also do not exhibit TRIP phenomenon.

3.3 Fractography

Fractured surfaces of CVN impact specimen were observedunder SEM to reveal the nature of fracture. SEM fractographsof selected specimens are shown in Fig. 10. Figure 10(a) showsmainly quasi-cleavage shinny fracture surface of 10 s parti-tioned specimen. Facets of random size with very few dimplesare observed in Fig. 10(a). The specimen partitioned for 100 sshowed greater impact toughness (22.5 J/cm2), as seen inFig. 10(b). Various voids and equiaxed dimples (produced bymicrovoid nucleation, growth and coalescence) are clearlyvisible confirming ductile nature of fracture. Another observa-tion is the presence of elongated, film-like band between the

Table 2 Mechanical properties of steel in normalized and Q&T condition

Property Normalized Q&T Q&P (100 s PT)

Hardness (HV) 540 650 6300.2% offset yield strength (MPa) 1100 1845 1425Ultimate tensile strength (MPa) 1659 2318 1975Total elongation (%) 11 8.2 11.6Uniform elongation (%) 6 4.7 6Strain hardening exponent (n) 0.190 0.178 0.24Product of strength x elongation (MPa%) 19,908 19,008 22,910CVN impact toughness (J/cm2) 27.5 20 22.5Notched tensile strength (MPa) 971 1300 790NTS/UTS 0.46 0.56 0.40Plain strain fracture toughness (MPa

ffiffiffiffi

mp

) 75 50 57

Fig. 9 Instantaneous strain hardening (n value) plotted as a func-tion of true strain for different heat treated conditions


facets (Fig. 10b). The morphology of these bands suggests thatthese were mostly islands of RA which upon deformationtransformed into aM. The fractographs of Q&T specimen do notdisplay any sign of RA. Moreover, the fracture surface showsrelatively less ductility, shallow dimples and primarily cleavagefacets (Fig. 10c).

Figure 11 represents the stereomicroscope photographs ofFT specimens revealing fatigue pre-cracking zone and finalfracture zone. It can be seen from Fig. 11(a) that the fatiguepre-cracking zone has uniform curvature followed by a ductile(dull) final fracture zone in Q&P specimen partitioned for1000 s. Contrary to this, fatigue pre-cracking zone of Q&Tspecimen is relatively non-uniform at crack front and finalfracture is quite brittle (bright surface) (Fig. 11b).

The difference between the microstructure and fracturemorphology in fatigue pre-crack zone of Q&P fracturetoughness specimens partitioned for 100 and 1000 s, respec-tively, is illustrated in Fig. 12. In case of 100 s partitionedspecimen (having FT of 57 MPa

ffiffiffiffi

mp

) it is evident that theregion near crack consists of both ML and MP and the crackfollows the path between laths and plates (Fig. 12a). In fact,crack selectively propagated between the laths and plates, i.e.,

most probably regions of RA. In Fig. 12(a), the crackmorphology reflects the interlath and island-like RA thatwas transformed into fresh brittle aM (during cyclic loading).These brittle regions of MF provided preferential sites forcrack propagation. On the other hand, the crack morphologyis clearly distinguished in Fig. 12(b). Mostly MP is visible inthe vicinity of crack and crack follow transcrystalline path.Broken halves of various MP are shown with arrows inFig. 12(b). Since the MP are brittle (less tough) and of highstrength due to high carbon content, they are relatively lesseffective in resisting crack propagation. However, it is notconfirmed whether this MP is formed on initial quench orformed on final quench to room temperature, i.e., MF. Further,the tempering and recovery of aM formed on initial quenchhas occurred which itself influence the crack growth. Themultiphase microstructure of Q&P steel provides a goodtoughness and effectively prohibits crack propagation. Find-ings of this study are expected to provide basic data for thedesign of suitable heat treatment procedure to combat againstfatigue and fracture. However, additional research work isrequired to clearly understand the fracture mechanism of Q&Psteel.

Fig. 10 SEM fractographs of specimen (a) Q&P at 473 K for 10 s, (b) Q&P at 473 K for 100 s, and (c) Q&T at 473 K for 1 h


4. Conclusions

The one-step Quenching and Partitioning heat treatment wasapplied on medium carbon low alloy steel and mechanicalproperties were evaluated and correlated with microstructures.Results were also compared with the properties in Q&T andnormalized condition.

1. Optical and SEM examination revealed that the Q&Psteel consists of multiphase microstructures containingcarbon depleted lath martensite, fresh plate-type martens-ite, islands of RA, film-type RA and nanometer sized

transition carbides. However, the RA was not visible inspecimens partitioned at 10 and 1000 s. The % of RAwas determined through XRD and maximum amount of6.8% with a carbon content of 1% was achieved in 100 spartitioned specimen. Partitioning for higher time periodsresulted in decrease in RA fraction and carbon concentra-tion due to competing reactions.

2. The tensile test results indicated that the best combinationof strength and elongation (2000 MPa and 11.5%) wasattained in 100 s partitioned specimen. Excellent ductilitywas attributed to the presence of interlath RA within mar-tensite matrix.

Fig. 11 Photographs of broken halve of fracture toughness specimens (a) Q&P at 473 K for 1000 s and (b) Q&T at 473 K for 1 h

Fig. 12 SE Micrographs of Q&P fracture toughness specimens taken from the fatigue pre-crack region adjacent to the specimen edges and par-allel to the direction of crack propagation: (a) partitioned for 100 s and (b) partitioned for 1000 s


3. Although the % of RA was low in 1000 s partitionedspecimen even then it was noted that the impact tough-ness, notch tensile strength and fracture toughness washigher as compared to 100 s specimen.

4. Another interesting observation was that the 100 s parti-tioned specimen exhibits higher number of fatigue cyclesto initiate crack than other specimens. The better fatigueresistance of 100 s was attributed to presence of film-likemetastable RA. The RA present ahead of the crack trans-formed into martensite during cyclic loading therebycausing crack closure effect.

5. The high impact and fracture toughness of 1000 s parti-tioned specimen were a cumulative effect of TRIP as-sisted transformation of metastable RA, annihilation ofdislocations, carbon depleted lath & plate martensite andpresence of fine transition carbides. However, this behav-ior needs further detailed investigation.

6. Comparison of properties with Q&T and normalized steelsuggests that Q&P treatment yields superior ductility andtoughness at high strength levels. Moreover, Q&P treat-ment offers greater control over evolution of microstruc-ture. By altering the microstructures, attractive propertycombinations could be tailored.

Acknowledgments

This research project is financially supported by the PakistanSpace and Upper Atmosphere Research Commission (SUPARCO).Authors would like to thank Mr. Ahmed Bilal (Chairman,SUPARCO) for approval and provision of facilities. The authorsalso gratefully acknowledge Dr. Sajid Mirza for his valuablesuggestions throughout this work. Further acknowledgement is dueto Mr. Waqas Hussain and Mr. Quaisar Khan for helping in thespecimen preparation, testing and microstructural examinations.

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