On Factors Affecting the Phase Transformation and Mechanical ...

Tag : low-alloy steel

On the other hand, the variation in the ferrite grain size thatoccurs when the hot-rolling schedule is varied is caused by thedifferent austenite grain morphology that forms the ferrite. Forthe P schedule, the grains are produced from work-hardened(pancaked) austenite through an austenite-to-ferritetransformation. The sizes of these ferrite grains are smaller thanthose produced from the recovered and recrystallized austenitegrains of the R schedule. It is well known that ferrite tends tonucleate at the austenite grain boundaries.[28'29] Thus, theferrite grain size developed after the transformation stronglydepends upon the austenite grain structure that appears just beforethe start of the transformation. During the straining of theaustenite in the nonrecrystallization region, deformation bands andtwinning boundaries form and the dislocation density insideaustenite grains is greatly increased; the increased dislocationdensity provided favorable nucleation sites and an enhancednucleation rate. The grain refinement effect of deformation in thenonrecrystallization region is greater than that in therecrystallization region.[30,31] The higher austenite decompositionkinetics causes the faster nucleation rate of ferrite. Thenonrecrystallized austenite transforms to ferrite at a faster ratethan does recovered or recrystallized austenite, due to twoeffects:[31]
(a) the higher internal energy of the deformed and, thus, lessstable austenite, and
(b) the larger number of nucleation sites provided by defects.
B. Intercritical Annealing Conditions
During the current investigation, the transformation kinetics werestudied using the dilatometer. All the dilatometric measurementswere performed on samples produced from the cold- rolled stripswith the long specimen side oriented parallel to the rollingdirection. Unless mentioned, the dilatometric tests were performedon samples from the RP materials.
1. Continuous heating transformation
In order to define the intercritical region, dilatometricmeasurements were applied by heating specimens up to 1373 K. TheAc1 and Ac3 temperatures are plotted in the continuous heatingtransformation (CHT) diagram in Figure 6.
Figure 6 shows that the increased heating rate resulted inincreased Ac1 and Ac3. The current results indicate that thehealing rate exerts a stronger influence on the finish temperaturewhen all the ferritic phase is exhausted than on the starttemperature at the beginning of carbide dissolution. On the otherhand, the dependence of the transformation temperature on theheating rate decreases as the heating rate decreases. At a veryslow heating rate, the Ac1 and Ac3 are independent of the heatingrate. This feature is associated with the transformationtemperatures at equilibrium (slow heating and cooling). Theequilibrium Ae1 and Ae3 transformation temperatures depend only onthe chemical composition and are unaffected by the heating orcooling rates.[32] This infers that the Ac1 and Ac3 are close tothe Ae1 and Ae3, respectively, when the transformation temperaturesbecome independent of the heating rate. In this context, fine grainsize resulted from the hot- and cold- rolling processes; thisfavors the transformation process in such a way that a very slowheating rate (less than 0.05 K/s), is not required to trace theequilibrium points. This is because the greater density ofnucleation sites, resulting from the fine- grained structure,enhances the transformation kinetics.
2. Phases in equilibrium
From the variation of the relative change in length as a functionof temperature, the lever rule was employed to calculate the formedaustenite fractions (f^sub gamma^). Figure 7 shows the results forboth steels. Depending on the previous discussion, for bettertracing of the equilibrium points, the 0.05 K/s dilatation curvewas used for these calculations.
Figure 8 compares the measured formed austenite fraction with theone predicted using the THERMOCALC TCW3 software. This figure alsoshows the predicted dependence of the ferrite and cementite on thetemperature. During the continuous heating of steels, thetransformation reportedly takes place by the initial growth ofaustenite into the carbide-rich areas, and subsequently by thegrowth of the austenite into ferrite and the redistribution ofcarbon between the former and the latter phases.[33-34] Figure 8 ismarked by T^sub Cp^ for the predicted and by T^sub Cm^ for themeasured kinetics, and shows a clear demarcation of these twosteps. The change in slope of the line represents how f^sub gamma^depends on the annealing temperature, which corresponds to thechange in the austenite formation mechanisms. On the other hand, atthe beginning of transformation (near the Ae1), the measuredkinetics are slower than the predicted rates for both alloys.However, this difference decreases gradually with an increasingtransformation temperature. The predicted and the measuredtransformation kinetics are almost identical at the end of thetransformation, close to the Ae3. The reason for this differencecould be the sluggish kinetics of the austenite formation at thelow temperature, close to the Ae1.[34- 351
C. Heat Treatment and Microstructure Formation
The intercritical annealing temperature was chosen on the basis ofthe dilatometric measurement shown in Figure 8. A phase content of70, 50, and 30 pct PF was required at the end of intercriticalannealing. In equilibrium conditions, this phase distribution isobtained in the intercritical annealing temperatures (T^sub A^)shown in Table III. It should be mentioned that all theintercritical annealing temperatures used are above the measuredcementite dissolution temperature T^sub Cm^. Note that cementite isa cleavage and void-initiating phase that is best eliminated fromstrong steels.
For heat-treatment processes, the salt baths mentioned in SectionII-B were used. The intercritical annealing temperatures listed inTable III were implemented in combination with austemperingtemperatures, T^sub B^, of 365 [degrees]C, 400 [degrees]C, and 435[degrees]C. The isothermal holding time for each of the two stepswas 8 minutes. After austempering, the samples were quenched inwater.
1. Effect of heat-treatment conditions
Figure 9 shows representative microstructural results using thesame magnification for the two steels, after applying the two-stepannealing conditions. In all specimens, the minor microstructure isthe martensiteaustenite (MA) structure (white areas). Figures 9(a)through (c) show the influence of the intercritical annealingtemperature on the microstructure of steel 2, for specimensaustempered at 673 K. An increase in the intercritical annealingtemperature resulted in a decrease in the PF amount toapproximately reach the expected amount (70, 50, and 30 pct). Aconcurrent increase in the amount of the MA phase was observed.
Geib et al.[36] reported that intercritical annealing produceswell-developed precipitate distributions in ferrite retained duringintercritical annealing, whereas the austenite that forms onintercritical annealing dissolves any carbonitride precipitatesinitially present. Austenite formation at ferrite-ferriteboundaries during partial austenitization would lead to thedissolution of these stable carbides/carbonitrides residing atferrite boundaries, making the austenite with niobium in solutionsluggish to transform on quenching.[37] It is reported thatmicroalloying steel with Nb increases the amount of retainedaustenite.[38]
On the other hand, the intercritical annealing temperature has nosignificant effect on the grain size of ferrite and retainedaustenite. A similar observation has been reported by Shi etal.[39] The grain sizes of the ferrite and retained austenitevaried between 3 and 7 [mu]m and 0.7 and 3 [mu]m, respectively.This study detected no significant influence of the chemicalcomposition and austempering temperature on the location andmorphology of the retained austenite.
Additionally. Figure 9 shows that the structure is fullyrecrystallized. Consistent with that result, Petrov et al. observedno interaction between recrystallization and transformationphenomena, because the SRX was already completely finished beforethe start of the alpha [arrow right] gamma phasetransformation.[40]
Figure 10 shows the dependence of the retained austenite content(V^sub gamma^) on the heat-treatment parameters for both alloys. Ingeneral, the retained austenite content measured at RT is dependenton two phenomena:[41]
(a) the isothermal transformation of austenite to bainite at agiven temperature, and
(b) the a thermal transformation of austenite to martensite oncooling to RT after isothermal holding.
Figure 10 shows that decreasing the PF amount (increasing theintercritical annealing temperature) leads to a higher amount ofretained austenite (V^sub gamma^) in the final microstructure forthe two alloys. This could be attributed to the higherintercritical austenite content.
For the alloys annealed to the expected PF levels of 50 and 70 pct,increasing the T^sub B^ temperature from 638 to 708 K resulted inincreasing the V^sub gamma^. For these intercritical annealingconditions, the authors also noted that the richer aluminum alloy(steel 2) had a lower V^sub gamma^ compared with steel 1. Thelatter observation is consistent with the light optical microscopicone.
These two observations were not recorded for the materialsintercritically annealed for an expected PF level of 70 pct. Suchresults for V^sub gamma^ can be justified by tracking the austeniteand its composition throughout the heal-treat ment process asfollows: during the intercritical annealing stage, the austenite isenriched with carbon, due to the partitioning of carbon between theferrite grains and the intercritical austenite (gamma^sub i^)during the formation of gamma^sub i^. Figure 11 uses THERMO-CALCand gives the predicted dependence of the elemental concentrationsin gamma^sub i^ on the PF content (i.e., T^sub A^). During thecourse of the second isothermal holding (austempering), theaustenite is further enriched with the carbon rejected from thebainitic ferrite. This reaction can occur until the point at whichthe free energy of ferrite is equal to the free energy of austenite(G^sub alpha^ = G^sub gamma^); thus, at this point, no furthertransformation of austenite to ferrite can occur. This residualaustenite, because of its high carbon content, has a martensitestart temperature (M^sub S^) below RT; thus, a certain amount ofaustenite from this process can be retained at RT. The amount ofcarbon that can enrich the austenite during this process has beenfound to depend on the isothermal austempering temperature,according to the T^sub 0^ concept.[27] Figure 12 shows the T^sub 0^curve drawn for both steels I and 2 using THERMO-CALC software(TCCQ). However, the amount of carbon that enriched the V^subgamma^ does not necessarily match the carbon content predicted bythe T^sub 0^ curve. For example, if the holding time isinsufficient to reach the T^sub 0^ curve, especially for lowaustempering temperatures that require a long holding time, theexpected carbon concentration value cannot be reached. On the otherhand, for a high holding temperature, the carbides can form easilyand the carbon in the austenite decreases.[42] Reportedly, carbidestarts to precipitate between 723 and 748 K, and this temperaturerange is hardly affected by chemical composition.[43]
In the current case, considering the same amount of gamma^sub i^ atthe beginning of the bainitic transformation, the bainitic reactionwill last to a lower concentration of carbon at a highertemperature. Thus, the final austenite will be of a lower carbonconcentration (Figure 12), but of a higher content (Figure 10). Theuse of the relatively narrow austempering temperature range around673 K is the reason why the results could be justified in the lightof the T^sub 0^ curve. Hashimoto reported that the best combinationof bainite transformation rate and carbon content at T^sub 0^ curveoccurs at this temperature.[42] However, the trend may not continuefor lower or higher austempering temperatures, at which the carbonin the retained austenite may not fit with the T^sub 0^ curve, dueto the insufficient holding time or the carbide precipitation,respectively, as explained previously.
On the other hand, the V^sub gamma^ of the 70 pct PF material doesnot follow the trend expected from the T^sub 0^ curve. In thiscase, the gamma^sub i^ is enriched with the carbide former elements(C and Mn), whereas the carbide suppression element (Al) is at alower level (Figure 11). This motivates the carbide formation inbainite. Carbide formation withdraws carbon from the austenite, soits stability drops noticeably. This may proceed to the extent thatsome of the austenite formed during the isothermal holding cannotbe stabilized down to RT and, thus, transforms to martensite.Furthermore, Kim et al. have reported that the improvement of theaustenite hardenability because of the high Mn content (2.52 wtpct, in their case) results in martensite transformation duringcooling after the isothermal bainite holding.[44] For the materialannealed to 70 pct PF, a similar excessive Mn content is predictedin gamma^sub i^ (Figure 11).
The occurrence of the martensite phase in the 70 pct PF materialwas confirmed with investigation using a SEM. The samples for theSEM were prepared following the procedure proposed by Girault etal.; in that manner, distinguishing martensite from retainedaustenite in the micrograph is possible.[16] Figure 13 showsrepresentative SEM micrographs.
For the material annealed to 70 pct PF, the martensite phase wasobserved. Figure 13(c) shows a representative example. Themartensite grain is characterized by its well-delineatedsubstructure, while the retained austenite looks rather smooth.Using transmission electron microscopy, Sakuma et al[41] observed ahigh dislocation density in the ferrite surrounding the martensiteregion, whereas the ferrite surrounding the retained austenite andbainite regions showed a very low dislocation density. On the otherhand, increasing the aluminum content leads to a lower V^sub gamma^(Figure 10). The ability of steel 2 to proceed to the higher carbonenrichment of austenite during bainitic holding (Figure 12)consequently results in lowering V^sub gamma^, as discussedearlier. Therefore, increasing the aluminum content not onlyreduces the cementite stability but also motivates the enrichmentof austenite with carbon and leads to a lower retained austenitevolume fraction (with a higher carbon content and greaterstability).
2. Bainite transformation kinetics
The dilatometric measurements were used to compare the bainitetransformation kinetics of the two steels. During the dilatometrictests, a cooling rate of 50 K/s was used to quench the samples fromthe T^sub A^ temperature (Table III) to 673 K.
Figure 14 compares the dilatation curves for steels 1 and 2 afterbainitic transformation at 673 K. The bainite reaction proceeds toa higher final amount of bainite for the alloy with a higher Alcontent. Figure 14 also demonstrates that bainite formation isaccelerated by increasing the aluminum content (assuming the samePF content). THERMO-CALC calculations of the driving force for thetransformation of austenite into ferrite (DeltaG^sup gammaalpha^)for steels 1 and 2 (Figure 15) confirm the thermal dilatometricresults. Thus, the micro structure evolution throughout theisothermal bainite transformation demonstrated that both thebainite formation rate and the total bainite amount increase withan increase in the aluminum content.
The dilatometric experiments have shown that a cooling rate of 50K/s is sufficient for avoiding the formation of the allotriomorphicferrite during cooling to the T^sub B^. This was detected from thelinearity of the temperature-dilatation curve observed duringcooling.
3. Effect of hot-roiling conditions
As explained in Section A-2, the different hot-rolling schedules(Figure 2) resulted in pronounced differences in the hot-rolledstructure size (Figure 4). Due to the latter variation, differentfinal TRIP-aided steel structure sizes are observed. Figure 16shows representative microstructures that result due to applyingdifferent hotrolling schedules using the same magnification. TableIV compares the final structure grain sizes of ferrite and retainedaustenite. Making use of the beneficial effect of the rolling belowT^sub nRX^, a pronounced finer cold-rolled TRIP-aided steelstructure was produced. Based on the current result (as well as ona previous report[39]), the employed heat-treatment parameters haveno significant effect on the ferrite and retained austenite grainsize. The current study infers that the structure fineness dependsonly on the hot-rolling conditions. Thus, controlling thehot-rolling schedule, prior to the cold rolling and heat treatmentof TRIP- aided steel, has a decisive effect on the structurerefinement and results in a more convenient structure size. On theother hand, the observed effect of the prior hot-rolling scheduleon the V^sub gamma^ was very limited.
Additionally, dilatometric investigations showed that thedifference in the hot-rolling schedules influences the bainitetransformation kinetics in a pronounced way. Figure 17 presents theevolution of DeltaL/L^sub 0^ as a function of the transformationtime for the different hot-rolling schedules applied on steel 2.Figure 17 shows that the transformation in smaller grains startsfaster but proceeds at a slower rate. Indeed, the grain-sizereduction causes an increase in the grain-boundary area, at whichthe first bainitic ferrite subunits nucleate. Thus, thetransformation starts more quickly, due to an enhanced nucleationrate. The transformation then proceeds by means of the nucleationand the growth of new subunits from the tip of the previous onestoward the interior of the austenite grain; thus, when theaustenite grain size is reduced, the transformation proceeds at aslower rate.[45,46] The influence of the austenite grain size onthe transformation rate was already shown and modeled by Rees andBhadeshia.[47]
In addition to the effect of the chemical composition on thebainite transformation kinetics (Figure 14), another featurehighlights the transformation occurring in the cold-rolled TRIP-aided steels: the prior hot-rolling schedule.
On the other hand, instead of differing transformation rates, thebainite reaction proceeds to the same final amount of bainite(Figure 17). Accordingly, by the mass-balance relationship amongthe material phases, the V^sub gamma^ in the final microstructureshould remain unaffected by the prior structure fineness. Thisdilatometric observation is consistent with the fact that themeasured V^sub gamma^ is only insignificantly affected by thehot-rolling schedule.
D. Mechanical Properties
Figures 18 and 19 show the effect of the processing route on theultimate tensile strength (UTS), yield strength (YS), and the totalelongation percent (pct El) for steels 1 and 2, respectively. TheYS values plotted in the histograms are the lower yield strengthsor the 0.2 pct offset YS (0.2 pct YS) values, in the case of theabsence of a yield point. In Figure 20, the pct El values areplotted vs the UTS values for all combinations of the rollingconditions and heat-treatment parameters (3 structure sizes x 3annealing temperatures x 3 austempering temperatures).
It can be concluded from Figures 18 to 20 that it is the PF percent(the intercritical annealing condition) that has the mostpronounced effect on the mechanical properties of both steels. TheUTS values increase with a decrease in the PF percent. This can beexplained by the higher austenite content that formed during theintercritical annealing. A higher amount of gamma^sub i^, resultsin a greater amount of high-strength bainite, after the isothermalbainitic transformation step.
On the other hand, increases in the PF percent from 30 to 50 pctresults in El values with increasing percents. In addition to thefact that an increase in the soft PF content affects the pct El,the expected higher amount of retained austenite surrounded by thebainite phase in the 30 pct PF microstructure can also contributeto the observed lower pct El. That is because the retainedaustenite grains with the film shape located at the bainite are nottransformed to martensite, even when a considerable amount ofdeformation is applied; thus, no contribution is made to theductility improvement.[44]
An additional increase in the PF percent to 70 pct has no furtherenhancement effect on the percent of El. This observation can bejustified by the retained austenite characteristics associated withthe 70 pct PF material. Principally, the total elongation of TRIP-aided steel is controlled by the volume fraction and stability ofthe retained austenite and the difference in strength between thematrix and second phase.[48] In the 70 pct PF material, the V^subgamma^ is obviously lower than that of the 50 pct PF material(Figure 10) and has lower stability, as well, as discussed earlier.
Furthermore, Timokhina et al. have reported that the PF in the C-Mn-Si-Nb-steel showed a ferrite deformation lower than that in thenon-Nb steels; the PF also did not flow around the bainitic regionduring straining. This has been justified by the fact that the Nbaddition has the effect of promoting precipitation hardening of theferrite, either directly or through grain refinement, which leadsto the strengthening of the ferrite.[49]
The materials annealed to 70 pct PF have recorded relatively highYS/UTS values. For this case, the retained austenite stability isvery low, as explained earlier. Consequently, transformation insuch low austenite stability would take place even in the elasticregion.[38] Thus, the TRIP effect is most likely to take place inthe elastic region and would eventually increase the YS values, butsubsequent transformation in the plastic region would be limited.Consequently, the contribution of the TRIP effect in enhancing theUTS and the percent of El is also limited.
Figures 18 and 19 show that the mechanical properties do not varysimilarly with respect to the isothermal bainitic holdingtemperature, T^sub B^. The complex nature of the microstructuresand the interaction among their phases result in an unpredictableeffect of T^sub B^, especially in the relatively narrow rangeinvestigated (638 to 708 K).
Since TRIP-aided steels combine high strength and high ductility,their mechanical properties are often characterized by the productof the tensile strength and the total elongation (UTS x pct El),which is known as the formability index.[42,50] At each combinationof PF content and hot-rolling condition, the average form-abilityindex was calculated from the values recorded at the threecorresponding T^sub B^ temperatures. Figure 21 shows a histogramthat compares these averages. This figure shows that the highestformability index was obtained for the steels annealed to 50 pctPF. This result is associated with the strength-ductility balancerecorded at this PF content (Figure 20).
Based on Figure 21, the hot-rolling schedule has a pronouncedeffect on the formability index for both steels- It is interestingthat the hot-rolling schedule has a very limited effect on theV^sub gamma^ values. Thus, the improvement in properties seemsentirely due to microstructural refinement (Figure 16). In additionto the well-known effect of grain refinement on improvingmechanical properties, the greater stability of the smallerretained austenite grains has an additional improving effect.Brandt et al.[51] have stated that smaller retained austeniteparticles contain fewer potential nucleation sites for thetransformation to martensite; consequently, these require a largertotal driving force for the nucleation of martensite. Thus, bycontrolling the deformation temperature and the degree ofdeformation below T^sub nRX^ during the hot-rolling process, it waspossible to improve the strength- ductility balance of thecold-rolled TRIP-aided steel. The deformation below the T^sub nRX^results in the distinct refinement of the final TRIP-aided steelmicrostructure. A poorer formability index for the lower-Al-contentalloy (steel 1) in spite of its higher austenite content is relatedto lower carbon enrichment during isothermal bainite transformation(Figure 12). This results in less stable austenite, which is thentransformed to martensite in an earlier stage of plasticdeformation. In cases in which the austenite stability is very low,the transformation occurs even in the elastic region, so that thedual-phase behavior can be observed. Nevertheless, a minimum amountof retained austenite should be present, in order to guarantee TRIPbehavior and enhanced strength and formability.[9]
For additional analysis, the average mechanical properties ofsteels 1 and 2 were calculated from all the values recorded underthe different employed processing routes. Table V lists thecalculated values. The values are highly reliable, because theywere obtained from many samples processed under differentconditions, but these conditions are quite similar for the twoalloys. It can be concluded that increasing the aluminum contenthas improved the formability of steel 2 by improving its ductility.The higher ductility of steel 2 is correlated with its higherbainite content (Figure 14) and higher retained austenite stability(Figure 12). Sakuma et al. stated that more stable austenitetransforms at high strains and increases strain hardening atnecking, thereby improving ductility.[41,52]
IV. CONCLUSIONS
This work provided an investigation of the microstructure formationand mechanical properties obtained after different hot- rolling andtwo-step annealing treatment conditions of two Mn-Si-Al cold-rolledTRIP-aided steels alloyed with Mo and Nb. From the presentinvestigation, the following conclusions can be drawn.
1. By making use of the beneficial effect of the hot-rolling belowT^sub nRX^, a pronounced finer cold-rolled TRIP-aided steelstructure was produced. This resulted in an improvement in thestrength-ductility balance of the steel.
2. A comparison of the transformation kinetics using thethermodynamic calculations and the dilatometric method is onlyrelevant when using a heating rate at which the transformationtemperatures are independent of the heating rate; using a higherheating rate is then misleading.
3. As long as the bainite transformation in TRIP steel can complywith the T^sub 0^ curve, the retained austenite content increasesby increasing the bainite transformation temperature.
4. It is the composition of intercritical austenite that controlsits behavior during the bainite transformation, and not the bulkcomposition of the alloy. During the current study, it is proposedthat the retained austenite has low stability, due to thepartitioning of elements during intercritical annealing in such away that, at a low intercritical annealing temperature, gamma^subi^ is enriched with the carbide former elements (C,Mn); the carbidesuppression element (Al) concentration, however, is at a lowerlevel.
5. Increasing the aluminum content not only reduces the cementitestability but also achieves accomplishes the following.
(a) Motivates the formation of bainitic-ferrite and leads to ahigher fraction of bainite.
(b) Lowers the fraction of retained austenite.
(c) Increases the carbon content in retained austenite and improvesits stability.
(d) Enhances the formability though increasing the ductility. (Thisis a result of the latter effect.)
6. For the given alloying elements, the most promisingmicrostructures with respect to the strength-ductility balance arethose containing 50 pct PF.
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MOHAMED SOLIMAN, Doctor, and HEINZ PALKOWSKI, Professor of MetalForming, are with the Institute of Metallurgy, Clausthal Universityof Technology, 38678 Clausthal-Zellerfeld, Germany. Contact e-mail:mohamed.soliman@tu-clausthal.de Manuscript submitted February 24,2008.
Article published online July 15, 2008
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