Abstract
The influence of a thermal-treatment simulating a galvannealing-process (GA) on the microstructure of a 0.2C–4Mn–1.5Si and 0.2C–4Mn–1.5Al (in wt-%) medium-Mn-Q&P-steel was investigated in detail using dilatometry. After Q&P-process, the retained austenite (RA) fraction is higher and mechanical properties are superior for the Si-concept. Owing to galvannealing, strength decreases substantial for both concepts. For the Si-concept however, RA starts to decompose into pearlite upon galvannealing temperatures of 520°C, leading to a significant decrease in elongation. For the Al-concept, the RA-fraction remains constant up to 560°C, which even results in increased elongations values compared to the galvannealed Si concept. This contribution shows that Al as an alloy-element in Q&P-steel is more suitable for RA stabilisation at higher temperature than Si.
Keywords
Introduction
The need to improve passenger safety and fuel-efficiency has led to growing demand and increasing enhancement for advanced-high-strength steels (AHSS) for the automotive industry [1,2]. Among these, the quenching and partitioning (Q&P) steels, a third generation AHSS, have received much attention, due to their high tensile strength combined with excellent ductility. In the quenching step of the Q&P-process, which was originally proposed by Speer et al. [3], certain amount of athermal martensite (
) is formed, while cooling to temperatures (Tq) between martensite start temperature (Ms) and martensite finish temperature (Mf) after full austenitization. During a subsequent partitioning treatment, commonly between 300 and 450°C [4-6], the carbon from the supersaturated martensite enriches and stabilises the neighbouring austenite [4,7-10]. During this isothermal holding at partitioning temperature (Tp), an additional formation of bainitic ferrite (αB) can take place [4,5,11,12]. The stability and amount of retained austenite (RA) after final quenching to room temperature is especially affected by its C and Mn content. Depending upon that stability of the enriched austenite after the partitioning step, the final cooling to room temperature potentially decreases its fraction on behalf of fresh martensite (
) [5,13]. To suppress and retard carbide precipitation in tempered 
during the partitioning step and thereby enable sufficient carbon partitioning into the remaining austenite, additions of Si and/or Al are employed, since these elements are considered insoluble in cementite [14-16]. However, these elements cannot prevent the precipitation of ϵ- or η-transition-carbides in tempered 
(
), before the partitioning step [17-20]. As reported in [16,21], Al compared to Si has a minor effect on retarding these transition-carbides and, according to [20], even promotes the formation of these carbides at lower temperatures by increased carbon flux to carbides in 
.
In order to provide an excellent compromise between corrosion protection and good weldability/paintability [22,23] within the annealing cycle, sheet steels are produced by continuous inline hot dip galvannealing facilities. However, AHSS steels containing especially high Si contents are prone to major problems during the galvanising (GI) and galvannealing (GA) process by reasons of a delayed GA-reaction, bad wettability due to selective oxidation and liquid metal embrittlement (LME) during spot welding. In contrast, Al addition is expected to have little influence on coatability and therefore shows, according to [23,24], a better galvanizability. Furthermore, increasing GA-temperatures bear the risk of RA decomposition owing to carbide precipitation, or transformation into pearlite, respectively [25,26]. Thereby, the replacement of Si by Al should provide a positive effect, as Al retards austenite decomposition into pearlite [20,21] and cementite coarsening in tempered martensite at higher temperatures (<450°C) more effectively than Si [16,21,25].
In the present contribution, the microstructural evolution of a 0.2C–4Mn–1.5Si and a 0.2C–4Mn–1.5Al steel during a galvannealing heat treatment-process was compared with focus on the RA transformation behaviour.
Materials and experiments
Chemical composition (in wt-%) and transformation temperatures of the investigated cold rolled steel concepts.
As shown in Figure 1, the specimens were fully austenitized at 850°C (0.2C–4Mn–1.5Si) and 900°C (0.2C–4Mn–1.5Al) for 120 s with a heating rate of 10 K s−1, on a Bähr 805 A/D dilatometer. Austenitizing temperatures were chosen based on the A3-temperature of the material, which was also determined at a heating rate of 10 K s−1. Subsequently, samples were quenched to a Tq below Ms-temperature (Table 1), between 150 and 290°C (0.2C–4Mn–1.5Si) and 190–320°C (0.2C–4Mn–1.5Al), at a cooling rate of −50 K s−1. After a short isothermal hold at 10 s to provide equal heat distribution within the specimen, samples were reheated at 20 K s−1 to a Tp of 350°C and isothermal hold at tp for 40, 120 and 200 s before quenching to room temperature with a cooling rate of −50 K s−1 (Q1).
Schematic representation of the heat-treating cycle.
To investigate the distinctive influence of a subsequent GA-treatment, additional samples were processed. After partitioning, the samples were reheated at 70 K s−1 to 460°C and kept for 3 s (simulating the time-temperature schedule of steel strip during in-line galvanising). Afterwards the samples were additionally reheated from 460°C at 20 K s−1 to GA-temperatures (TGA) of 480, 520 or 560°C for 20 s and subsequently quenched to room temperature at −50 K s−1 (Q3). Scanning electron microscopy (SEM) performed on a Zeiss SUPRA 35, was used to perform a qualitative study of the respective microstructural constituents. Therefore the samples were electrochemically polished, whereby a topography correlating to the corresponding phase was generated. Furthermore, the hardness was measured according to Vickers (HV1), using an Emco Test DuraScan device. The RA content was determined by means of saturation magnetisation measurement (SMM). The phase fractions of 
, αB and 
were determined on the basis of the lever rule from the dilatometry curves as described in [27]. In addition, cold rolled material was heat treated with selected parameters in a multipurpose annealing simulator equipped with gas jet cooling. Afterwards, tensile test was conducted according to the international standard ISO 6892-1 using a specimen oriented longitudinal to the cold rolling direction with a gauge length of 25 mm on a ZWICK-Roell BTC-FR020TN tensile testing machine.
Results and discussion
Microstructural evolution during Q&P treatment
The microstructural development of Q&P steels is determined by Tq and the respective martensite fraction. This in turn has a strong influence on the chemical composition of the remaining austenite. As a result, the dilatation during isothermal holding at Tq, shown in Figure 2, which is mainly related to formation of αB the literature [4,5,12,24,28] is strongly dependent on this parameter. The dilatation is also slightly affected by carbon partitioning from Dilatation as a function of tp at Tp = 350°C depending on Tq: (a) 0.2C–4Mn–1.5Si; (b) 0.2C–4Mn–1.5Al.
to the remaining austenite. Substituting Si with Al shows a significantly accelerated αB-formation, which has also been reported in [12,25]. Furthermore, both concepts show an intensified αB-transformation with increasing Tq i.e. decreasing 
content, caused by the above mentioned alteration of the austenite composition (lower C-content). Owing to a higher difference in Gibb's energy between the remaining austenite and αB, according to [29], this results in a higher driving force for the αB-formation.
Solely at Tq just below MS (290 and 320°C, respectively), αB-transformation slightly decreases. This is consistent with the results of A. Navarro-Lopez et al. [30], who investigated the effect of prior athermal martensite on subsequent transformation kinetics for low C-, high Si-steel. Apparently αB-kinetics retard due to the lower amount of 
acting as nucleation points for the transformation. Therefore, it is shown for these Tq just below Ms, that dilatation curves do not reach a saturation level, respectively αB-formation was not completed during the investigated period of time.
Figure 3 depicts the phase fractions of the Si- and Al-concept depending on Tq, evaluated from the dilatometer curves and by SMM respectively. For both alloy concepts tempered martensite fractions ( Phase fraction and hardness after Q&P as a function of Tq and tp for Tp = 350°C: (a) 0.2C–4Mn–1.5Si; (b) 0.2–4Mn–1.5Al.
, orange) decline with increasing Tq, whereas the amount of bainitic ferrite (αB, green) formed during the partitioning step increases. These findings are in good correlation with the trend of αB-kinetics in Figure 2. Furthermore higher fractions of αB are formed with longer tp, especially when the previously formed 
-content is low.
The RA-fraction (blue) has, according to previous studies on conventional Q&P-steels [5,12,13,17,28], a maximum at a specific Tq (Figure 3), which is confirmed in this work. For the Al-concept this RA-maximum is about 40°C higher (at 270°C) due to its higher Ms-temperature. Below that maximum, 
formed during the initial quench, limits the amount of austenite still available for retention. Above this maximum, the remaining austenite cannot be fully stabilised by the present carbon content and fresh 
(grey) is formed during the final quench [9,10]. Furthermore, a higher fraction of RA can be stabilised after the Q&P-process by Si additions, compared to the Al-concept. This is in good agreement with the results in [6,12]. As mentioned in the introduction, Al compared to Si has a minor effect on preventing transition carbides during tempering of martensite. This was also confirmed in this work and will be demonstrated afterwards. Thus the increased precipitation of transition carbides in 
acts as carbon sinks, which according to [17,18] consequently leads to a decrease of carbon available for partitioning into austenite. According to [10,17,18], the increased αB-formation of the Al-concept, additionally leads to a decomposition of the remaining austenite. However, especially at higher Tq, RA- and αB-fraction increases significantly with longer tp. As reported in [19] from in-situ X-ray diffraction measurements, C-partitioning from supersaturated 
into the remaining austenite takes place immediately after the first quench and completes within seconds. The subsequent αB-formation, formed from unstable austenite and the coherent C-redistribution leads to further stabilisation of RA at the expense of 
(grey) formed during the final quench, especially for the Al-concept. This is in good agreement with the findings of [10,17]. With increasing Tq, more 
is formed during final quenching, based on the lower chemical enrichment of the remaining austenite, resulting also in an increased hardness (Figure 3) in a range of 450–550 HV1 with the Si-concept and 400–500 HV1 with the Al-concept. In general, the lower hardness for the Al-concept can be explained by the lower potential for solid solution hardening of the alloying concept. Owing to longer tp, i.e. more αB, the hardness likewise decreases.
Influence of an additional GA-treatment
Examining the influence of an additional GA-treat-ment, two Tq with the highest RA-values after the Q&P-process, Tq = 190 and 230°C for the 0.2C–4Mn–1.5Si and Tq = 230 and 270°C, respectively for the 0.2C–4Mn–1.5Al after tp = 40 s, were chosen. Figure 4(a) shows the RA- (solid line) and RA-fraction (solid lines) and fresh martensite-fraction (dashed lines) after Q&P depending on GA-treatment for:(a) 0.2C–4Mn–1.5Si Tq = 190°C (triangle) & 230°C (square); (b) 0.2C–4Mn–1.5Al Tq = 230 (triangle) & 270°C (square) with tp = 40 s. Dilatation as a function of tGA during GA treatment for: (a) 0.2C–4Mn–1.5Si Tq = 230°C; (b) 0.2C–4Mn–1.5Al Tq = 270°C.
-fraction (dashed line) of the Si-concept for the selected Tq, after Q&P and varying GA-temperatures. With the Si-concept, GA-treatment up to 480°C tends to increase the amount of RA at the expense of 
, especially for higher Tq (230°C), where more austenite is available after partitioning (different chemical composition of remaining austenite). This could arise from a further C-redistributing ferritic bainite formation during the reheating. However, hardly any expansion in dilatation (Figure 5(a)) was shown at TGA of 480°C. As a consequence, it is reasonable to suggest that the increase of RA-fraction potentially stems from a further C partitioning from 
to the remaining austenite, promoted by the increase in temperature.
The SEM investigation of the 480°C GA sample (Figure 6(a)), exhibits a microstructure consisting predominantly of SEM analyses after Q&P-process (Tp = 350°C, tp = 40 s) following a GA-treatment; top: 0.2C-4Mn-1.5Si Tq = 230°C, bottom: 0.2C–4Mn–1.5Al Tq = 270°C; (a,d) TGA = 480°C; (b,e) TGA = 520°C; (c,f) TGA = 560°C.
(∼60%) and αB. According to [17-20], Si and Al cannot completely prevent the precipitation of (transition-) carbides in tempered 
, as this happens under paraequilibrium condition. For the Si-concept, hardly any of this carbides can be dissolved in 
via SEM. Furthermore, blocky shaped RA-islands with a size of up to 1 µm can be found. In contrast, an increase in GA-temperature of 520 and 560°C yields in a decrease in RA-fraction for both Tq. At 520°C where the decrease of RA-fraction (Figure 4(a)) is still moderate, large RA-island can be found in the microstructure but first signs of decomposition in the centre (Figure 6(b)) start to appear. At these temperatures, dilatation shows a significant contraction (Figure 5(a)) at the beginning of GA. After tGA > 5 s this contraction is followed by a transition to a saturation for 520°C and an expansion at 560°C. In [26,31] this contraction was attributed to the precipitation and coarsening of cementite (θ) in tempered martensite, which resonates with the findings of the SEM investigations (Figure 6(b,c)). Here, spherical θ can be found predominantly as accumulations at the prior austenite grain boundaries (PAGBs) and heterogeneously distributed in the tempered martensite (
), which is getting more pronounced at 560°C. It shows that the targeted effect of Si on preventing cementite precipitation and coarsening does not suffice at temperatures ≥520°C, which corroborates the findings of [16,20,21]. The expansion in the dilatation curve in Figure 5(b) at 560°C can, when taking into account the SEM image (Figure 6(c)), be attributed to a decomposition of the remaining austenite into pearlite (P) or upper bainite. This transformation behaviour is also shown in [25] during isothermal holding at bainite forming temperatures higher than 450°C and in [31], after long time partitioning at 500–600°C during a Q&P process. In addition [26] reported that the shrinkage of prior martensite owing to tempering can further lead to a destabilisation of austenite due to the reduction of the hydrostatic pressure between 
and RA. The reduction of the aforementioned mechanical stability and the decrease in chemical stability caused by the observed cementite formation, leads at a Tq of 230°C to a slight increase of 
during quenching from 560°C to room temperature (Figure 4(a)).
The Al-concept in Figure 4(b) shows a similar increase in RA fraction after a GA-process up to 480°C as the Si-concept, predominantly at the higher quench temperature (230°C). In contrast to the Si-concept dilatation during GA at 480°C in Figure 5(b) exhibits a significant expansion for the Al-concept. When focusing on the SEM-investigation in Figure 6(d), a matrix of bainite (αB) and tempered martensite (
) can be seen. Compared to the Si-concept, the matrix microstructure is coarser but the RA is significantly finer and has a more film-like shape. In addition, numerous needle-type carbides can be spotted inside the tempered martensite (
) confirming the weaker effect of Al on preventing carbide formation in tempered martensite at lower temperatures. As no pearlite structure can be found, the expansion in relative length change during GA at 480°C as well as 520°C can be attributed to the formation of bainitic ferrite. This may also stabilise the remaining austenite further by C-partitioning. In [25] it was found, that Al compared to Si not only accelerates bainite formation, but also increases the bainite start temperature. Considering this, the different behaviour concerning the dilatation at GA can be explained by altering αB-transformation temperatures of the two alloying concepts.
Opposed to the SEM-images of the Si-concept (Figure 6(b)) at GA 520°C, no spherical carbides can be found in the 
and the formation of the carbides at the PAGB is less developed for the Al-concept (Figure 6(e)). This behavior is coherent with the result of the dilatation in Figure 5(b), which depicts an expansion and confirms findings in [16,20,21] that Al compared to Si has an enhanced effect on retarding carbide precipitation at high temperatures. The most distinct difference between the Al- and the Si-concept can be shown when the GA-temperature is increased to 560°C (Figure 4). While the decrease of RA is significant for the Si-concept, the Al-concept shows only a slight reduction with no increase in 
for both Tq.
The dilatation at 560°C for the Al-concept in Figure 5(b) faces a contraction, showing that at this temperature the positive effect on preventing carbide precipitation and coarsening in tempered martensite and at PAGBs does not suffice even with the Al-concept. The contraction in dilatation at 560°C is reaching a saturation level after 10 s. The subsequent expansion, which is associated with the pearlite formation for the Si-concept, is significantly delayed. The SEM-image of the Al-concept after GA at 560°C (Figure 6(f)) illustrates similar to the Si-concept at 520°C (Figure 6(b)), that preferably some of the larger islands of RA transform into a pearlitic structure, whereas the film-like retained austenite is still present. The slight expansion of the dilatation curve can therefore be related to the pearlite formation of bigger RA islands, which is consequently the reason for the slight decrease of retained austenite displayed in Figure 4(b) at 560°C. It is shown that Al has, compared to Si, beside the benefit on preventing cementite formation in tempered martensite, a positive effect on retarding pearlite formation when a Q&P process is followed by a GA treatment higher than 520°C. When focusing on the RA morphology were pearlite transformation starts, the more lath like shaped RA of the Al-concept, compared to globular characteristics of the Si-concept, could have an additional benefit against RA-decomposition. More specifically, a higher amount of targeted austenite can be retained to room temperature after a galvannealing heat treatment.
The displayed hardness in Figure 7 (error bars smaller than the symbols), shows for both concept a decrease with increasing TGA independent from the trend of the RA content in Figure 4. Therefore, it is reasonably to argue that the hardness decrease is predominantly caused by an intensified tempering of the Hardness after Q&P process depending on GA-treatment for 0.2C–4Mn–1.5Si (solid line) and 0.2C-–4Mn–1.5Al (dashed lines). UTS, YS, TE, and UE after Q&P depending on GA-treatment for: (a) 0.2C–4Mn–1.5Si Tq = 230°C; (b) 0.2C–4Mn–1.5Al Tq = 270°C tp = 40 s, TP = 350°C.
matrix. The trend of the hardness values displays also good correlation with the ultimate tensile strength (UTS) depicted in Figure 8(a,b). UTS decreases gradually for both concept due to increased tempering of the microstructure. For the Si-concept the UTS decreased from 1575 to 1300 MPa. As mentioned above the Al-concept exhibits the same overall trend, but UTS is in general 200 MPa lower, due to the missing solid solution strengthening, which was also found in [12]. The yield strength (YS) rather decreases with increasing galvannealing temperature for both concepts, due to additional tempering effects [8]. Considering the elongation after Q&P-process, it can be shown that the Si-concept displays considerably higher valves. Owing to gavannealing especially higher 480°C, total elongation (TE) as well as uniform elongation (UE) decreases significantly for the Si-concept. This can be explained due to the decomposing RA-fraction (Figure 4(a)), which reduces according to [1] the potential for the transformation induced plasticity – (TRIP) effect. Compared to the Si concept, Al shows an increase of UE and TE. As the RA does not decompose during galvannealing, the potential TRIP effect and its enhancement on the combination of strength and elongation can be preserved. Furthermore tempering of the martensite increases the elongation. Finally, it can be concluded that by conventional Q&P-process, in fact higher elongation values can be obtained via the Si-concept. But by applying higher temperatures between 480–560°C as needed by the galvannealing process, the Al-concept results in improved mechanical properties.
Conclusion
In this contribution, the microstructural evolution of the steels 0.2C–4Mn-1.5Si and 0.2C–4Mn–1.5Al in a Q&P-treatment and the effect of a subsequent GA-processing was determined by precise dilatometer study.
With the Si concept, higher amounts of RA subsequent to the Q&P heat treatment can be accomplished, arising from the stronger effect on preventing carbide precipitation and growth in tempered martensite at lower temperatures, compared to the Al-concept. Blocky shaped RA islands were dominantly found with the Si concept, whereas RA appeared more film-like with the Al concept, which has a significant effect on its thermal stability. Compared to Si, Al showed a more efficient effect on retarding the carbide formation in prior martensite and at the PAGBs for GA-temperature of ≥520°C. With the Si-concept, a significant decomposition of RA at temperatures >480°C into pearlite and fresh martensite was detected, whereas Al possesses a retarding effect on pearlite formation and consequently a more sufficient effect on stabilising R during a GA-treatment. A positive effect of Al can also be linked to the mechanical properties in tensile test, especially for the elongation. While the UE and TE decrease at higher temperature for the Si-concept due to the decomposition of RA, it even increases for the Al-concept because of the still existing high TRIP-potential and stronger tempering of the martensite.
Footnotes
Acknowledgements
The authors sincerely acknowledge the support of the Austrian Research Promotion Agency (FFG) related to the frontrunner project No. 860188 ‘Upscaling of medium Mn-TRIP steels’.
Disclosure statement
No potential conflict of interest was reported by the author(s).