Today, most coil springs for automotive applications areade of quenched and tempered, medium carbon high-strength
teels [1–4]. In order to increase the hardenability, elementsuch as chromium, manganese and silicon are added to theseteels [2]. Furthermore, silicon retards the conversion of the �-arbide to cementite during tempering. It refines the carbidesnd improves the sag resistance significantly [4,5]. For springteels, the emphasis in materials research has been focused onncreasing the strength while maintaining good ductility, tough-ess and fatigue properties. The conventional heat treatmentoes not exploit the maximum potential of existing steel grades.or a fixed composition the increase in strength requires furthereduction of tempering temperature. But with decreasing tem-ering temperature the ductility decreases. The loss in ductilitys additionally dependent on impurity element concentration.
It has been shown that embrittlement of high-strengthuenched and tempered martensitic steels occurs when thin car-ide films form at the prior austenite grain boundaries (PAGB),hat are enriched with impurities such as phosphorous [6–9]. The
tempering temperature, where the embrittlement is observed,overlaps with the lower range of the third stage of tempering, inwhich cementite precipitation takes place. During tempering inthis temperature range the PAGB are preferential nucleation sitesfor cementite [10]. The thin plate-like carbides can act as slip bar-riers and can initiate intergranular cracks at the already impurityweakened PAGB [11]. At tempering temperatures below 673 Kthe diffusivity of impurity elements is relatively low. Banerjiet al. [12] showed that enrichment of the PAGB occurs duringthe austenitization before quenching and not in the temperaturerange where tempered martensite embrittlement was observed.In order to avoid or minimize the embrittlement, plate-like car-bides should be fragmented or avoided at the PAGB, especiallywhen impurity elements are present.
According to the early proposal of Smith [13], grain boundaryprecipitates generating by diffusion have a specific orientationrelationship with respect to one of the adjacent matrix grains.Furuhara et al. reported that among the possible crystallographicvariants only one single variant tends to precipitate on a flatgrain boundary [14]. A carbide film along the grain boundaryis then formed by further growth and coalescence of small indi-
vidual precipitates. The selection of the variant depends on theactivation energy for the formation of the critical nucleus. Itis assumed that similar assumptions must apply to the precip-itation of cementite during tempering of martensite at PAGB.For quenched and tempered Fe–2Mn–0.36C (mass%) Yusa et
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al. [15] showed that refinement of grain boundary cementite ispossible by modified ausforming. They showed that deforma-tion of austenite prior to quenching produced serrated austen-ite grain boundaries. After tempering the PAGB are decoratedwith cementite precipitates of different morphologies and vari-ants. Furthermore, modified ausforming can also produce prioraustenite grain boundaries free from cementite.
In this study we examine the effect of austenite deforma-tion on the carbide precipitation at the PAGB of a quenchedand tempered Fe–0.56C–1.50Si–0.67Cr–0.59Mn (mass%) steel,which is widely used for automotive coil springs. The austenitedeformation was carried out in such a way to produce both, arecrystallized and a non-recrystallized austenite structure. Thesteel under the investigation had a copper and tin content abovethe usual content observed in steels produced by the electric arcprocess, which is about 0.3 mass%. In the electrical arc processan increased amount of scrap is used. Copper and tin are themost widely quoted examples of impurity pickup from scrap insteel production.
2. Experimental
The chemical composition of the alloy studied is given inTable 1. Copper and tin were added to investigate the influence ofthese elements on the mechanical properties of steels produceda
An ingot with a cross section of 140 mm × 140 mm and aweight of 70 kg was prepared by vacuum induction melting. Theingot was cut into pieces with a cross section of 70 mm × 70 mm.After homogenization at 1373 K for one hour the samples wererolled to 30 mm thickness and air cooled. After reheating to1373 K and holding for one hour again the samples were thenrolled to 21 or 15 mm thickness, respectively and air cooled toroom temperature. From these bars samples for conventionalheat treatment and thermomechanical treatment were machinedas shown in Fig. 1.
The conventional heat treatment (CHT) and the thermome-chanical treatment (TMT) were conducted using a large scale2.5MN hot press at the Max-Planck-Institut fur Eisenforschung[16,17]. The treatments are presented schematically in Fig. 2.The temperature of the specimen was determined with a ther-mocouple inserted in a bore (Ø 2 mm and 12 mm depth) athalf thickness directly. Specimens were heated to austenitiza-tion temperature by induction heating. The thermomechanicallytreated samples were air cooled to the deformation tempera-ture, deformed in plane compression to a thickness of 13 mm(see Fig. 1) and quenched after 15 s of quench delay. Thequench delay tQD is the time between the end of austenitiza-tion (for conventional heat treatment) or the end of deformation(for thermomechanical treatment) and the start of quenching.The strain rate was 5 s−1 and the nominal logarithmic strainwas 0.4. The temperature of the oil bath for quenching was
TC
S
M 0.021A 0.035
Fd
t a given phosphorous level.
able 1hemical analysis of the investigated spring steel
atment (b), and tensile testing. RD, TD, and ND rolling, transverse and normal
196 A. Ardehali Barani et al. / Materials Science and Engineering A 426 (2006) 194–201
Fig. 2. Schematic diagram for conventional heat treatment (a) and for (b) thermomechanical treatment.
333 K. Tempering was performed in a furnace for 1 h at 623 or673 K.
Round tensile specimen were machined from the tem-pered CHT or TMT samples in the rolling direction (RD) asshown in Fig. 1. Microstructural observations, optical and SEMmicroscopy were performed on the plane normal to the trans-verse direction after standard sample preparation and etching innital (1%) or in alkaline sodium picrate.
3. Results
Fig. 3 presents micrographs of TMT samples tempered at673 K. The sample deformed at 1123 K shows equiaxial prioraustenite grains. This observation indicates that the austenitewas recrystallized before the quenching. Therefore, this treat-ment will be addressed as “recrystallized TMT variant”. Thesample deformed at 1023 K exhibits elongated, pan-cake struc-tured prior austenite grains. It will be referred to as “ausformedTMT variant”. The deformation at 1123 K refines the austenitegrain size from approx. 25 �m (without deformation) to approx.10 �m.
Compared to CHT, deformation of the austenite prior toquenching improves the ductility (Fig. 4). Conventionally heattreated samples that are tempered at 623 K fracture before neck-ing starts (Table 2, Fig. 5). The CHT samples tempered at 673 KftTw
Fig. 4. Reduction of area and ultimate tensile strength of conventionally heattreated and thermomechanically treated samples. Each value is an average ofthe results of three tensile tests.
is equivalent or higher. For this variant the improvement dueto austenite deformation in strength is higher at the lower tem-pering temperature 623 K. The same applies to the 0.2%-offsetyield strength (Table 2).
For a CHT sample tempered at 673 K the fracture surfaceindicates the occurrence of intergranular fracture with a low
rmed at 1123 K (a) and 1023 K (b) and tempered at 673 K. Etched in 1% nital.
racture at strains slightly larger than the uniform strain. For bothempering temperatures the tensile strength of the recrystallizedMT variant (TD = 1123 K) is slightly below that of the CHT,hereas that of the non-recrystallized TMT route (TD = 1023 K)
Fig. 3. Optical micrographs of thermomechanically treated samples defo
A. Ardehali Barani et al. / Materials Science and Engineering A 426 (2006) 194–201 197
Table 2Tensile properties, standard deviation in parenthesis
a Deformation temperature.b Tempering temperature.
amount of ductile fracture (Fig. 6a). This corresponds to a lowductility (reduction in area RA = 15%). Additionally, grain sep-aration along the loading direction takes place between prioraustenite grains. In contrast to conventionally heat treated sam-ples the thermomechanically processed samples never failed byintergranular fracture. The TMT samples always exhibited thehigher ductility (RA > 26%). Even after tempering at lower tem-perature where the conventionally heat treated samples revealeda minimum in ductility, the fracture surface of the thermome-chanically processed sample shows a ductile fracture withoutintergranular fracture (Fig. 6b).
The cementite morphology of the CHT sample tempered at623 K is presented in Fig. 7. Carbides form preferentially atPAGBs and at interfaces between neighboring martensite laths.Carbide precipitation within individual lath is observed as well.After tempering at 623 K almost all PAGBs are occupied withcontinuous carbide films.
After tempering at 673 K the carbide film consists of a chainof small spherical carbide particles that are still at some areasconnected to each other (Fig. 8). For both tempering temper-atures disc-like and spherical carbides were observed withinindividual laths.
F(a
Deformation at 1123 K results in refined carbides at thePAGBs. After tempering at 623 K the films seem to be continu-ous at small magnifications. At higher magnification it becomesclear that most of these PAGB are decorated with small spherical
Fig. 6. Fractograph of conventionally heat treated sample (a) tempered at 673 Kand exhibiting separation along prior austenite grain boundaries and fracto-graph of thermomechanically treated sample (b) deformed at TD = 1023 K andtempered at TT = 623 K.
ig. 5. Engineering stress–strain curves (tensile) for conventional heat treatmentno deformation), and for thermomechanically treated samples deformed at 1023nd 1123 K.
198 A. Ardehali Barani et al. / Materials Science and Engineering A 426 (2006) 194–201
Fig. 7. SEM micrograph showing carbide morphology and distribution of conventionally heat treated specimen tempered at 623 K. Etched in alcaline sodium picrate.
Fig. 8. SEM micrograph showing carbide structures of conventionally heat treated specimen tempered at 673 K. Etched in 1% nital.
carbides. After tempering at 673 K the carbides at PAGBs areall spherical or ellipsoidal (see Fig. 9). The typical feature of thePAGB after TMT was the occurrence of pronounced curvaturesin the micrographs which was in contrast to the flat appearanceof the PAGBs after CHT.
After ausforming (TD = 1023 K) and tempering at 673 K thedecoration of the PAGBs is similar to that of samples deformedat 1123 K. Hence after tempering at the 673 K the grain bound-
ary carbide morphology is similar for both TMT variants andaustenite structures prior to quenching, i.e. for the recrystallizedand for the non-recrystallized austenite. Whereas, most PAGBsof ausformed samples tempered at 623 K were free of any pre-cipitations. The carbide morphology at the interfaces betweenlaths and packets was changed too (Fig. 10). Compared to theother treatments no pronounced carbide precipitation could beobserved after etching.
Fig. 9. SEM micrographs of TMT sample deformed at 1123 K and tempered at 673 K showing chain of round or ellipsoidal carbide particles at prior austenite grainboundary. Etched in 1% nital.
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Fig. 10. SEM micrographs of ausformed sample tempered at 623 K showing no carbides at PAGB.
4. Discussion
For Fe–2%Mn–0.36%C Yusa et al. [15] showed that grainboundary carbides can be refined by ausforming. They per-formed the deformation of the austenite with 80% reductionand immediate quenching after deformation. They related thecarbon refinement to the serrated structure of the prior austen-ite grain boundaries. In Fig. 11 the carbide distribution, whichwas observed in our study after tempering, is shown schemat-ically for conventionally heat treated and thermomechanicallyprocessed samples. In our study we show that grain boundarycarbides can be refined by high temperature thermomechanicaltreatment HTMT (recrystallized austenite prior to quenching)and by a low temperature thermomechanical treatment LTMTor modified ausforming (forming in metastable austenite regionwithout recrystallization). The employed strain rate in this work
was lower (5 s−1) than those used by Yusa et al. [15] (10 s−1).Thus due to the lower strain rate no pronounced serration wasobserved. However, the ausforming resulted in former austenitegrain boundaries free of carbides and in a higher ductility. Forthe same etching conditions (temperature, time, etchant used),less topology effects were developed during etching for the aus-formed samples. Less carbides were observed within the matrixat magnifications up to 100 000, because etching did not producea structure with clear, sharp edges that can be distinguished eas-ily as a separate phase. Hence, we conclude that the resultantdislocation substructure in the austenite after the ausforming isstabilized by the diffusion of carbon atoms prior to quenchingto the dislocation cores. During transformation this dislocationsubstructure together with the segregated carbon atoms at thedislocation core is inherited from the parent austenite phase tothe final martensite product phase. This leads to either smaller
Fp
ig. 11. Schematic illustration showing the different types of carbide morphologies arior austenite grain boundaries observed for conventional heat treatment and thermo
nd their distribution within individual lath, at lath and packet interfaces and atmechanical treatment after tempering at 623 K.
200 A. Ardehali Barani et al. / Materials Science and Engineering A 426 (2006) 194–201
carbides after tempering at low temperatures or delays the pre-cipitation reaction, because diffusion of carbon atoms out ofthe dislocation core is not favorable and increases the internaldistortion energy.
Deformation at 1123 K leads to equiaxed austenite grainsprior to quenching. Due to this fact and the fact that the austen-ite grains observed were smaller than those after CHT, it wasdeduced that the austenite had recrystallized prior to quenching.Observation with SEM revealed that the prior austenite grainboundaries were curved. Thus the observed variety in carbidemorphology and size after TMT and deformation at TD = 1123 Kcan be related to the curvature of the PAGBs. In addition tothe grain boundary energy the activation energy for the criticalnucleus formation of the precipitates depends on the orientationof the grain boundary plane [18,19]. If the boundaries are flat, aswas observed for conventional heat treatment, only the variantwith the smallest critical nucleus energy will form. As a resultseveral carbides will form with the same orientation relation-ship to the neighboring grains and these particles will grow atthe former austenite grain boundary until they touch each otherand coagulate. At a curved boundary either the interface prop-erties change or the growing carbide variant changes direction.In both cases the carbide particles do not coagulate with eachother to form a continuous film at the PAGB. A curved boundarycould therefore be the reason for the carbide refinement at thePAGB.
m1stopasGgtgbftdptpmrtcbs
5
F
with and without deformation of austenite before quenching tomartensite.
After conventional heat treatment consisting of austeniti-zation, quenching and tempering, the subsequent observationswere made:
(1) After tempering at 623 K most of the prior austenite grainboundaries (PAGB) were occupied with a thin carbide film.
(2) At tempering temperature of 673 K the carbide films startto spheroidize but still form a continuous film at the PAGB.
The study of the microstructure and the mechanical propertiesafter thermomechanical treatment consisting of austenitization,deformation, quenching and tempering leads to the followingconclusions:
(3) Deformation of austenite prior to quenching improves theductility significantly.
(4) Deformation of austenite at 1123 K results in a recrystallizedaustenite prior to quenching, while deformation at 1023 Kto elongated, pan-cake austenite grains.
(5) Deformation at 1123 K leads to fine spherical carbides atthe PAGBs for the investigated tempering conditions.
(6) After deformation at 1123 K and quenching the PAGBs arecurved. This curvature can be a reason for the observed
(
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1otGTG
R
The increase in ductility cannot only be ascribed to the refine-ent of carbides. The steel in this study contains in addition to
20 ppm phosphorous a high content of copper and tin. Our testshow that for conventional heat treatment and for constant P con-ent the ductility of samples tempered at 673 K strongly dependsn the copper and tin concentration [20]. Because in the tem-ering range of 623–673 K the diffusivity of tin is very small,nd the holding time is very short, equilibrium grain boundaryegregation during tempering can be neglected. According torabke’s investigation [21] there is a rather low tendency forrain boundary segregation of tin in alpha iron. Furthermore, ashere is a limited number of vacant sites at the prior austeniterain boundary to which the atoms can segregate, and as the car-on diffusivity is much faster, the tin is replaced by carbon atomsrom grain boundaries as the carbon content increases. Never-heless, the solid phase reactions like recovery, recrystallization,iffusion and segregation taking place after the deformation andrior to quenching should not be neglected. It has been shownhat recrystallization of austenite can lead to an enrichment ofhosphorous at the grain boundaries [22]. From the observationade here, it can only be deduced that austenite deformation
efines the grain boundary carbides. But it should be noted, thathe improved ductility is not only a result of grain boundaryarbide refinement, but might as well be enhanced by the redistri-ution of the impurity elements and refinement of the martensitetructure.
. Conclusion
The microstructure and tensile properties of the spring steele–0.56C–1.50Si–0.67Cr–0.59Mn (mass%) were determined
fragmentation of the carbides and the variety of their sizeand morphology.
7) If the austenite is deformed at 1023 K and tempered at 623 Kthe carbides disappear at the PAGBs. At 673 K carbides atthe PAGB are fine and spherical.
cknowledgements
Results presented here are from the research project AiF3409N. Funding was provided by the German Federal Ministryf Economics and Labour (BMWA) via the German Federa-ion of Industrial Cooperative Research Associations “Otto vonuericke” (AiF) and the Steel Forming Research Society (FSV).he long version of the final report can be ordered from the FSV,oldene Pforte 1, D-58093 Hagen.
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