In-situ measurement of phase transformation kinetics in austempered ductile iron

M A TE RI A L S CH A RACT ER IZ A TI O N 85 (2 0 1 3 ) 1 2 4– 1 3 3

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In-situ measurement of phase transformation kinetics in austempered ductile iron Leopold Meiera , Michael Hofmannb,⁎, Patrick Saala , Wolfram Volka , Hartmut Hoffmanna a

Technische Universität München, Lehrstuhl für Umformtechnik und Gießereiwesen, Walther-Meißner-Straße 4, 85748 Garching, Germany Technische Universität München, Forschungsneutronenquelle Heinz Maier-Leibnitz (FRM II), Lichtenbergstraße 1, 85748 Garching, Germany

b

AR TIC LE D ATA

ABSTR ACT

Article history:

Austempered ductile iron (ADI) alloyed with 0.42% Mn and 0.72% Cu was heat treated in a

Received 22 June 2013

mirror furnace and the phase transitions were studied in-situ by neutron diffraction. The

Received in revised form 6

heat treatment consisted of austenitisation at 920 °C and isothermal austempering at

September 2013

400 °C, 350 °C and 300 °C, respectively. Due to the growth of ferrite platelets, the austenite

Accepted 10 September 2013

content decreases rapidly at all temperatures within the first 15–20 min and reaches a stable plateau after 35 min (400 °C) to 80 min (300 °C).

Keywords:

The carbon content of the residual austenite, which was monitored and characterised by

In-situ neutron diffraction

the change of the lattice parameter, increases up to 1.6 wt.% caused by redistribution from

Austempered ductile iron

the newly formed ferrite. While at higher austempering temperatures this takes place

Phase transformation

almost parallel to the phase transformation, at 300 °C the redistribution of carbon to

Retained austenite

austenite lags behind considerably. Furthermore the neutron data revealed an austenite peak asymmetry during austempering which is attributed to successive phase transformation. It results temporarily in two fractions of austenite, an initial low-carbon and an enriched high-carbon modification. © 2013 Elsevier Inc. All rights reserved.

1.

Introduction

Austempered ductile iron (ADI) is a heat treated nodular cast iron which combines excellent mechanical properties with the relatively low manufacturing costs and weight saving potential of a cast material. It offers engineers an alternative to steel and aluminium alloys and has been used in many applications in the automotive, agricultural and heavy engineering industries [1–9]. The heat treatment of ADI is started by an austenitisation in the range of 850–950 °C (Tγ), where the as-cast ferritic and/or pearlitic matrix transforms into austenite and gets enriched with carbon from the graphite nodules (label (1) in Fig. 1). After quenching to the austempering temperature TAUS in the range of 250–450 °C (label (2) in Fig. 1), a two-stage phase transformation takes place in ADI (label (3) in Fig. 1) [10–12]. In stage I, ferrite

platelets grow into the austenite grains. Since the solubility for carbon in ferrite is very low and the high silicon content typically present in cast irons delays the precipitation of carbides [13–15], the excess carbon is redistributed into the surrounding residual austenite. The austenite gets stabilised and the ADI material can be cooled to room temperature without formation of martensite. After prolonged times of austempering, the high carbon austenite eventually decomposes into ferrite and iron carbides (stage II reaction). Both martensite and decomposition carbides are detrimental to the mechanical properties of ADI [11,16–18]. The best heat treatment results are therefore achieved after the end of stage I (maximisation of austenite carbon enrichment) and before the beginning of stage II (onset of carbide precipitation). The time interval between these two stages is called the heat treatment window (HTW) of the ADI

⁎ Corresponding author. Tel.: + 49 89 289 14744; fax: + 49 89 289 14666. E-mail addresses: [email protected] (L. Meier), [email protected] (M. Hofmann), [email protected] (P. Saal), [email protected] (W. Volk), [email protected] (H. Hoffmann). 1044-5803/$ – see front matter © 2013 Elsevier Inc. All rights reserved. http://dx.doi.org/10.1016/j.matchar.2013.09.005

M A TE RI A L S C HA RACT ER I ZA TI O N 85 ( 20 1 3 ) 1 2 4–1 3 3

Fig. 1 – Schematic TTT-diagram of the ADI heat treatment.

reaction [10–12]. The final ADI microstructure consists of graphite nodules embedded in a matrix of ferrite and retained high carbon austenite. It is called “ausferrite” to distinguish it from the “bainitic” microstructure of steels heat treated similarly, which – besides the missing graphite – solely consists of ferrite and carbides [19,20]. The proportion and morphology of these phases, which determine the balance between ductility and strength in ADI, depend on the carbon redistribution which is strongly affected by chemical composition and the heat treatment parameters (Tγ, TAUS). Thus, the key to understand the ADI reaction kinetics is to monitor the phase transformation and its linkage to the redistribution of carbon during the reaction. Over the last 30 years research has focused on how microstructure and mechanical properties depend on chemical composition and heat treatment parameters [11,17,21–28]. Most of the studies are based on “interrupted austempering” samples, which were austempered for different times, quenched to room temperature and analysed by metallography and X-ray diffractometry. Metallography yields useful data about the general phase transformation and the microstructure morphology. However, the calculation of quantitative values for phase fractions of austenite, ferrite and martensite involves high uncertainties, especially in fine grained ADIs. Since metallography does not give any information about carbon redistribution, traditionally the carbon concentration xC,γ in retained austenite is calculated from lattice parameter measurements by X-ray diffraction (XRD) at room temperature. However, there is some uncertainty and a general tendency to overestimate carbon contents by analysing interrupted austempering samples [29,30]. Furthermore, for an adequate measurement of the phase transformation kinetics a large number of samples is necessary. Techniques such as in-situ neutron or synchrotron X-ray diffraction can help to overcome these experimental limitations. They offer quasi-continuous, phase-specific data including the lattice parameters which correlate to the carbon

125

redistribution. Furthermore, by analysing the bulk material at the austempering temperature, the measurements are neither affected by the quenching process nor the condition of the sample surface. However, only few studies have been published dealing with low temperature, two-stage decomposition of austenite measured in-situ by synchrotron X-ray diffraction [31–33] or Bragg edge neutron transmission [34–36]. Meggers [34] was the only one to study the austempering of cast iron. However, the main intention of this investigation was to show the capability of the relatively new Bragg edge neutron transmission technique with only the change in phase fractions being analysed. All other in-situ diffraction studies dealt with high-silicon steels, which show similar two-stage phase transformations. However, since these steels contain much higher levels of alloying metals (Mn, Mo, Cr) but less carbon and silicon than cast irons, the austempering reaction kinetics differ significantly. The aim of the present investigation is to follow the austempering of ductile cast iron in-situ by neutron diffraction and to demonstrate the potential of this technique to get detailed, phase-specific data helping to deepen the knowledge of phase transformation kinetics in ADI.

2.

Materials and Methods

2.1.

Base Material and Mirror Furnace

The material studied in this investigation is low alloyed ductile iron with a chemical composition of 3.65 wt.% C, 2.73 wt.% Si, 0.42 wt.% Mn, 0.72 wt.% Cu and Fe balance. After the nodulising treatment with Mg the melt was poured into bentonite sand moulds producing slabs in the dimensions of 225 × 150 × 25 mm3. The as-cast microstructure consisted of nodular graphite embedded in a predominantly pearlitic metallic matrix. Cylindrical heat treatment specimens 6 mm in diameter and 15 mm in length were machined from the cast slabs. To implement the ADI heat treatment steps of austenitising, quenching and austempering a mirror furnace as shown in Fig. 2 was used. In this furnace a cast iron sample is mounted in the centre of the nearly spherical chamber where the focal points of four parabolic mirror halogen spot lights coincide, each with a maximum output power of 150 W. With this set-up, heating rates up to 15 K/s can be achieved, resulting in heating times of approx. 100 s to an austenitising temperature of 920 °C. Quenching rates up to 35 K/s (between 800 °C and 500 °C) are realised by an argon gas stream running through a small tube pointed to the sample. In order to minimise oxidation and decarburisation, the gas-tight furnace was thoroughly flushed with argon gas before each heat treatment and kept sealed with a small overpressure of 300–500 mbar. The temperature during the experiments was measured and controlled using a tolerance class 1 type K sheath thermocouple 1.5 mm in diameter placed inside the sample. The temperature stability provided by the furnace controller was within ± 1 K during heat treatments. Characterisation measurements of the mirror furnace set-up showed that the temperature gradient inside the neutron measurement

126

M A TE RI A L S CH A RACT ER IZ A TI O N 85 (2 0 1 3 ) 1 2 4–1 3 3

Fig. 2 – Schematic design of the mirror furnace used in this investigation (left) and experimental set-up at the neutron diffractometer STRESS-SPEC (right).

volume of the heat treatment sample was within ± 9 K at 920 °C and ± 4 K at 350 °C, respectively.

2.2.

Neutron Diffraction

For the in-situ isothermal transformation studies the mirror furnace, whose aluminium walls are easily penetrable for neutrons, was placed at the STRESS-SPEC instrument at the FRM II neutron source in Garching, Germany [37]. A 2-dimensional 3He-PSD detector covering an angular range Δ2Θ of about ~10° was used to detect the scattered neutrons. The wavelength was set to ~ 2.1 Å (1 Å = 0.1 nm) by means of a Ge monochromator and was determined by measuring the standard silicon powder sample NBS-640c with an uncertainty of ± 0.0005 Å. Neutron diffraction data were collected during complete heat treatment cycles with varying measurement times between 20 s and 5 min adapted to the respective heat treatment stage. To monitor the phase transformation, the detector was set to a diffraction angle 2Θ ~ 60° where the high-intensity austenite (111) and ferrite (110) Bragg reflections could be measured simultaneously. Peak positions, integrated peak intensities and peak shapes were calculated using the curve-fitting technique with a Gaussian function by means of the software tools StressTextureCalculator [38] and PeakFit (Systat Software, Inc.). The interplanar spacings were calculated by Bragg's law from the fitted peak positions corresponding to the diffraction angles of the respective lattice planes. The volume fractions of austenite (Xγ) and ferrite (Xα) were calculated by the direct comparison method [39,40] using the integrated intensities of the austenite (111) and ferrite (110) peaks. Neutron diffraction texture measurements using the set-up at STRESS-SPEC [41] showed that no preferred orientation can be found in the ADI microstructures analysed in this investigation (multiples of random density, MRD < 1.10). With this almost random grain orientation the use of just two reflection peaks is sufficient for the analysis.

2.3.

Heat Treatment

The complete in-situ ADI heat treatments were performed in the mirror furnace by austenitising as-cast ductile iron

samples for 30 min at 920 °C, followed by quenching within 20–25 s to the austempering temperatures 400 °C, 350 °C and 300 °C, respectively. The samples were isothermally austempered for 128 min and subsequently cooled to room temperature.

2.4.

Metallography and Optical Microscopy

The samples for optical microscopy were cut perpendicular to the cylinder axis producing a circular cross section coinciding with the centre of the heat treated volume measured by neutron diffraction. The specimens were ground with SiC papers up to P2400 grade and polished with diamond and 0.05 μm alumina suspensions. After that they were etched with Beraha–Martensite etchant following a procedure given in [42]. Optical microscopy pictures were taken in a Zeiss Axioplan 2 microscope using an AxioCam MRc 5 CCD camera.

3.

Results and Discussion

3.1.

Austenitisation

Fig. 3 shows the change in lattice parameter of austenite, which was calculated from the measured interplanar spacing, during heating and austenitisation at 920 °C as an average of two treatments. The decomposition of pearlite during heating took place in less time than needed to reach the final austenitisation temperature. Thus, the change in austenite lattice parameter due to carbon enrichment is superimposed with temperature expansion in the first ~100 s. Furthermore, the cementite layers in the pearlite matrix contain a significant amount of carbon giving the austenite a high initial carbon content directly after the transformation. However, Fig. 3 shows that after a sharp increase within the first 5 min the lattice parameter reaches an almost constant value after approx. 10 min of austenitisation. This can be attributed to the enrichment of carbon up to the equilibrium content x0C,γ (in wt.%). At 920 °C this is 0.78 wt.% following the empirical Eq. (1) proposed by Voigt [43] which describes the solubility curve Am for carbon in austenite as a function of

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M A TE RI A L S C HA RACT ER I ZA TI O N 85 ( 20 1 3 ) 1 2 4–1 3 3

3.675

3.670

3.665

3.660 0

5

10

15

20

25

30

35

Heating time [min] Fig. 3 – Increase of austenite lattice parameter during austenitisation at 920 °C.

austenitisation temperature Tγ (in °C) and silicon content xSi (in wt.%) for cast iron.
x0 C;γ ¼ T γ =420  C −0:17ðxSi Þ−0:95

ð1Þ

In an austenitisation kinetics study with ductile irons of comparable chemical composition Batra et al. [44] determined enrichment times of 10 min and 2 min for materials with 80% and 95% pearlitic microstructure, respectively. However, in most studies on ADI, austenitisation times of 60–120 min were applied [11,21,23,25]. This might be appropriate for ferritic or higher alloyed grades, but involves the risk of austenite grain growth and leads to high heat treatment costs in industrial processes.

3.2.

Austempering

3.2.1.

Phase Fractions

A typical result from in-situ neutron diffraction of the austempering after quenching from the austenitising temperature (label (3-I) in Fig. 1) is shown in Fig. 4. It can be seen qualitatively that during austempering the intensity of the

austenite (111) reflections decreases due to the formation of ferrite, indicated by the increase of the ferrite (110) reflections. Furthermore, the peak shapes differ significantly, as the increasing ferrite peak is much broader than the initial austenite peak. Fig. 5 shows the volume fractions of austenite and ferrite calculated from the diffraction patterns during isothermal transformations at 400 °C, 350 °C and 300 °C, respectively. It should be noted that the volume fractions relate only to the metallic matrix since the change in nodular graphite fraction during austempering is negligible. At all austempering temperatures, the amount of austenite decreases sharply in the first 15 to 20 min with an almost identical transformation rate. Subsequently, the phase transformation begins to cease and the volume fractions reach a more or less stable plateau after about 35 min (400 °C), 50 min (350 °C) and 80 min (300 °C), respectively. Regarding the residual austenite, the corresponding volume fractions are approx. 50%, 40% and 25%, respectively. However, at TAUS = 400 °C the transformation of austenite to ferrite never ceases completely and seems to increase again after an austempering time of about 100 min. This can be attributed to a smooth transition between the end of the stage I reaction, which is characterised by the plateau, and the onset of stage II where the residual austenite decomposes to ferrite and carbides. In industrial austempering treatments, the main aim is to guarantee a high fraction of stabilised austenite and the prevention of martensite formation [21]. Interrupted austempering experiments with the same material reported earlier [45] showed that at TAUS = 400 °C no martensite can be found after austempering times of 20 min indicating that the heat treatment could be stopped even before reaching the plateau. The phase fraction results show good agreement with data from X-ray diffraction measurements found in the literature, 1.0

Austenite phase fraction

Austenite lattice parameter [Å]

3.680

TAUS = 400°C

0.8

TAUS = 350°C TAUS = 300°C

0.6

0.4

Ferrite phase fraction

0.8

0.6

0.4 TAUS = 400°C TAUS = 350°C

0.2

TAUS = 300°C

0.0 0

Fig. 4 – Typical neutron diffraction patterns during austempering at TAUS = 300 °C showing the evolution of austenite (111) and ferrite (110) peak intensities during austempering.

30

60

90

120

Transformation time [min] Fig. 5 – Evolution of phase fractions during austempering at 400 °C, 350 °C and 300 °C.

128

M A TE RI A L S CH A RACT ER IZ A TI O N 85 (2 0 1 3 ) 1 2 4–1 3 3

3.2.2.

Lattice Parameters and Carbon Redistribution

The lattice parameters of ferrite and austenite after quenching are presented in Fig. 6 and Fig. 7. At the austempering temperatures 400 °C and 350 °C the lattice parameters of ferrite show constant values within the error limits during the whole heat treatment (Fig. 6). In contrast, at TAUS = 300 °C the ferrite starts with a larger lattice parameter despite the lower temperature. It decreases within 25 min to a constant value and remains unchanged afterwards. This effect can be attributed to supersaturation of ferrite with carbon which was also observed in other studies at low temperature austempering of silicon steels [32,33] and ADI [23]. It supports the theory that bainitic ferrite initially forms in a displacive reaction and that carbon subsequently redistributes [15]. However, ferrite lattice expansion could not be observed at the higher austempering temperatures 400 °C and 350 °C. This can be attributed to higher carbon diffusion rates and the relatively high uncertainty of fitting the low intensity ferrite peaks at the beginning of austempering. The austenite lattice parameters increase considerably during austempering caused by the uptake of excess carbon 2.890 TAUS = 400°C

Ferrite lattice parameter [Å]

2.890 TAUS = 350°C

2.880

2.890 TAUS = 300°C

2.880

2.870 30

60

90

Transformation time [min] Fig. 6 – Evolution of ferrite lattice parameter during austempering at 400 °C, 350 °C and 300 °C.

3.650

3.640 TAUS = 400°C TAUS = 350°C

3.630

TAUS = 300°C

3.620 0

30

120

60

90

120

Transformation time [min] Fig. 7 – Evolution of austenite lattice parameter during austempering at 400 °C, 350 °C and 300 °C.

from the newly formed ferrite (Fig. 7). This carbon enrichment happens slightly later than the phase transformation at the higher austempering temperatures 400 °C and 350 °C due to the simple fact that the carbon redistributes to a relatively large fraction of residual austenite at the beginning and a decreasing fraction later on. The maximum lattice parameters are reached after about 40 min (400 °C) and 70 min (350 °C), respectively. It should be mentioned that there is no further carbon enrichment at 400 °C despite the continuing decrease in austenite fraction. This supports the assumption that the onset of stage II reaction lies at around 40 min of austempering where the stage I reaction has not finished yet. At TAUS = 300 °C there is an incubation time of about 15 min before the austenite lattice parameter starts to increase. Furthermore, the carbon enrichment lags behind the phase transformation considerably and the plateau stage is not reached even after an austempering time of 128 min. There have been several investigations in the last 60 years trying to find a relation between the carbon content of austenite xC,γ and the corresponding lattice parameter aγ,RT at room temperature. As a result, many equations are available in the literature, mostly in the form of aγ;RT ¼ a0;RT þ kC  xC;γ )=wt:%

2.880

0

3.660

Austenite lattice parameter [Å]

even though the amount of retained austenite found in our austempering experiments is in the upper range of differently alloyed ADIs [11,21–28]. This can be attributed to the elevated amounts of Cu and Mn, both elements promoting austenite stability in the austempering reaction [17]. The austempering times needed to finish the stage I reaction are comparable to other low alloyed ADIs [21,23,25–27], while the kinetics is generally slowed down in grades with higher amounts of the alloying elements Mn, Ni and Mo [11,21,24,28]. It should be mentioned that results within each of the studies cited above show relatively large scatter for the austempering treatment, presumably due to the use of interrupted austempering samples. Furthermore, in contrast to the in-situ results of this investigation, there is generally little data available on the early austempering stage up to 15 min, where the main part of phase transformation takes place in low alloyed ADIs.




ð2Þ

where a0,RT is the theoretical austenite lattice parameter in Å at room temperature in the absence of carbon and kC is an empirical constant. In some studies, kC was determined by X-ray diffraction of retained austenite in quenched steels of various carbon contents [46,47]. However, the lattice parameters of these experiments might be influenced by strains arising from the high fraction of formed martensite [48,49]. For in-situ studies other relations seem more appropriate which were derived from unstrained material resulting in a kC factor of 0.033 [48,50]. Assuming that during quenching from the austenitisation temperature no change in carbon content occurs, the reference lattice parameter a0,RT in Eq. (2) can be replaced by the initial lattice parameter a0,T(AUS) at the austempering temperature corresponding to the initial carbon content x0C,γ. Using these values and neglecting the very small influence of thermal expansion to the kC factor (<1%),

M A TE RI A L S C HA RACT ER I ZA TI O N 85 ( 20 1 3 ) 1 2 4–1 3 3

0.03 1.60

1.40

0.02

1.20

TAUS = 400°C

0.01

TAUS = 350°C

1.00

TAUS = 300°C

Austenite carbon content [wt.%]

Increase in lattice parameter [Å]

the carbon content can be calculated from the changes in austenite lattice parameter of the ADI samples during austempering as shown in Fig. 8. It increases to 1.45 wt.%, 1.6 wt.% and 1.5 wt.% at the austempering temperatures 400 °C, 350 °C and 300 °C, respectively. In contrast to the phase fraction results, these values are significantly lower than those reported in the literature, where maximum austenite carbon contents mostly range from 1.8 to 2.0 wt.% [11,21–28]. However, in almost all studies Eq. (2) with a0,RT = 3.548 Å and kC = 0.044 was used. This is a misinterpretation of values proposed by Roberts [46], who originally expressed the equation in the somewhat antiquated “kX units”. To obtain values in Å, the kX units have to be multiplied by the factor 1.00206 [51]. Thus, the carbon contents calculated with the misinterpreted values are about ~ 0.16 wt.% too high. Still, the combination of Mn- and Cualloying of the ADI studied in this investigation apparently causes quite low carbon contents which correspond to the comparably high residual austenite fractions shown in Fig. 5. At tAUS = 128 min the austenite accounts with 0.69 wt.% (400 °C), 0.62 wt.% (350 °C) and 0.33 wt.% (300 °C), for the total carbon content xC,tot. of the metal matrix, calculated from the product (Xγ * xC,γ). While at TAUS = 400 °C this value is just slightly lower than the initial carbon content from austenitisation (x0C,γ = 0.78 wt.%), there's a considerable gap at the lower austempering temperature of 300 °C. This indicates that at the latter temperature a considerable amount of carbon has precipitated as carbides during stage I of austempering or is still trapped inside the ferrite, either in solution or accumulated at grain boundaries, dislocations and vacancies [30]. The precipitation of transition carbides inside the growing ferrite sheaves has been demonstrated in many studies both in bainitic steels and in ADI [52,53]. It is believed to occur at austempering temperatures < 350 °C with increasing intensity at decreasing temperatures [15]. The amount of carbon trapped in ferrite is quite difficult to measure, so that there is a wide spread of published data ranging from <0.03 wt.% to 0.3 wt.% as summarised in Ref. [30]. Given the latter value, this would be more than enough to compensate the gap in xC,tot. for the austempering temperatures 350 °C and 400 °C even without

0.80

0.00 0

30

60

90

120

129

any precipitation of carbides during stage I. However, at TAUS = 400 °C the maximum austenite carbon content might rather be limited due to the onset of stage II. In contrast, at TAUS = 300 °C the precipitation of carbides is obvious in view of the large gap between the xC,tot. and the solute carbon in austenite after 128 min of austempering. Furthermore, the carbon redistribution evidently is restricted given the ferrite supersaturation and the considerable delay of increase in austenite lattice parameter at the beginning of austempering. However, the remarkable extent of the latter effect should be mentioned, since after 15 min approx. 50% of austenite has decomposed without a significant increase in lattice parameter.

3.2.3.

Microstructure

The corresponding microstructures formed in the heat treatments are depicted in Fig. 9. The microsections were made from samples that were cooled down after 128 min of austempering. Besides the nodular graphite, the metallic phases can be clearly distinguished from each other due to the Berahamartensite etching process. Ferrite appears in dark grey (originally brown) while the retained austenite is white. Ferrite both in bainitic steels and ADI consists of aggregates of ferrite platelets, formed by successive sympathetic nucleation [54]. Following terms employed by Aaronson and Wells [54], the aggregates are called “sheaves” and the platelets forming it are the “sub-units”. It can be seen from Fig. 9 that increasing undercooling has a significant effect on size and morphology of the ADI microstructure. At TAUS = 400 °C (Fig. 9a) the ferrite forms 5–20 μm long and 2–4 μm wide sheaves consisting of several platelets at least one magnitude smaller. The austenite is divided by the intersecting ferrite sheaves in “blocky” regions several μm in diameter and much smaller slivers which partly separate the subunits from each other. However, due to the high fraction of retained austenite, it presumably forms a matrix well interconnected three-dimensionally. Lowering the austempering temperature to 350 °C (Fig. 9b) increases both the ferrite sheaf density and the number of subunits at each sheaf giving them a feathery appearance. While the sheaves' lengths seem not to be affected, the width is reduced to ≤1 μm. The austenite matrix fragmentation increases correspondingly as size and number of blocky areas decrease and the separating austenite films between the ferrite subunits are most certainly reduced to sub-μm scale. At TAUS = 300 °C (Fig. 9c) the heat treatment results in a very fine ferrite structure of acicular appearance. With the resolution of light microscopy there are no longer sub-units in the ferrite observable. Thus, the microstructure rather resembles a nanoscale mesh consisting of bundles of parallel ferrite needles. Accordingly, the largest austenite grains have a diameter of approximately 1 μm while the films inside the ferrite bundles are hardly discernible.

Transformation time [min]

3.2.4. Fig. 8 – Increase in austenite lattice parameter during austempering at 400 °C, 350 °C and 300 °C with corresponding austenite carbon content starting with x0C,γ = 0.78 wt.% from austenitisation.

Evolution of Peak Widths

Additional information on the microstructure and the transformation kinetics can be extracted from the reflection peak shapes. Fig. 10 shows the evolution of the peak width (full width at half maximum, FWHM) of the austenite (111) and

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M A TE RI A L S CH A RACT ER IZ A TI O N 85 (2 0 1 3 ) 1 2 4–1 3 3

Fig. 9 – Microstructures of ADI samples after an austempering time of 128 min at temperatures (a) 400 °C, (b) 350 °C and (c) 300 °C.

ferrite (110) Bragg reflections. The peak widths of ferrite decrease slightly during the phase transformation and reach constant values of 0.55°, 0.7° and 0.9° at the austempering temperatures 400 °C, 350 °C and 300 °C, respectively. In contrast, the peak widths of austenite all start at ~0.4°, which corresponds to the instrumental resolution, and increase to a maximum of 0.6°, 0.85° and 1.0° at the austempering temperatures 400 °C, 350 °C and 300 °C, respectively. Subsequently, the peak widths decrease and reach constant values similar to those of the ferrite peaks.

Austenite FWHM [°]

0.9

0.7

0.5 TAUS = 400°C TAUS = 350°C

Ferrite FWHM [°]

TAUS = 300°C

0.9

0.7

0.5 0

30

60

90

βsize ¼ λ=ðd  cos θÞ

ð3Þ

Following Eq. (3), the average ferrite particle size d would decrease after subtraction of the instrumental resolution function to just ~ 40 nm at TAUS = 400 °C and to ~ 15 nm at TAUS = 300 °C. It should be emphasised that the term particle size in this context means “diffracting entity” which can be much smaller than seemingly continuous grains at microsections. Bhadeshia [56] reports sub-unit widths for bainitic steels of 100–300 nm and ~ 50 nm at austempering temperatures 400 °C and 300 °C, respectively. Line broadening due to particle size comes into effect from particle sizes approx. <100 nm [39]. Thus, given the microstructures shown in Fig. 9 it appears unlikely that the increase in peak width is caused solely by crystallite fragmentation, at least at higher austempering temperatures. Non-uniform lattice strains εnu are a further important source for line broadening (integral breadth βstrain in radians of 2θ). This effect can be expressed by Eq. (4) where β(εnu) is the integral breadth of the non-uniform strain distribution [39,55].

1.1

1.1

The different levels of FWHM as a function of austempering temperature, which are in good agreement with literature data [22], can be attributed to line broadening caused by decreasing crystallite sizes and non-uniform lattice strains [39]. The former can be described by the “Scherrer-Equation” (3) where βsize is the broadening of the diffraction line (integral breadth in radians of 2θ) and d is the diameter of the crystal particles [55].

120

βstrain ¼ 2  βðε nu Þ  tan θ

ð4Þ

Transformation time [min] Fig. 10 – Evolution of peak widths (FWHM) during austempering fitted by a single Gaussian model.

Following solely Eq. (4) as a reason for line broadening and subtracting the instrumental resolution function, the austempering at 400 °C and 300 °C would result in an

131

M A TE RI A L S C HA RACT ER I ZA TI O N 85 ( 20 1 3 ) 1 2 4–1 3 3

6 min 16 min 80 min

Intensity [Counts]

Austenite

6000

Ferrite

(a)

4000

(b)

2000

(c)

0 59

60

61

γLC

(a)

Intensity [Counts]

average non-uniform strain distribution of 0.6% and 1.3% in ferrite, respectively. The formation of ferrite is believed to take place following an invariant-plane strain displacive mechanism. It is associated both with a dilatational strain resulting in a density mismatch and a shear component leading to a geometric mismatch between the phases [57]. While the difference in density between austenite and ferrite calculated from the actual crystal lattice parameters during austempering is only ~ 1.7% at TAUS = 400 °C and ~ 2.0% at TAUS = 300 °C, the shear component of the invariant-plane strain reaction is around one magnitude higher [57]. Especially the latter leads to interphase stress, partially plastic accommodation and large local non-uniform strains [15,57] which results presumably in a significant contribution to the overall line broadening. However, especially at lower austempering temperatures a combination of both effects (Eqs. (3) and (4)) is most conceivable, regarding the increasing fineness of the developing microstructures. The small change in ferrite peak width despite the general broadening effect reflects the theory that due to the successive nucleation the single ferrite subunits just grow to a certain size depending on the austempering temperature [15]. Correspondingly, the level of non-uniform strain per sub-unit also just increases to a certain level. The overall increase in austenite FWHM is caused by the same reasons. The relatively large grains from the austenitisation treatment are cut and fragmented by the growing ferrite sheaves which leads to mosaicity and results in decreasing crystallite sizes. Particularly, this affects the austenite at TAUS = 300 °C where the former austenite grains have been completely replaced by fine, presumably heavily strained films between the ferrite sheaves. However, line broadening effects cannot explain the interim maximum in austenite peak width. A closer inspection of the neutron data reveals an asymmetry in all peak shapes of austenite at medium austempering times from 10 to 20 min. As an example, Fig. 11 shows the peak profiles after 6 min, 16 min and 80 min of austempering at TAUS = 300 °C. The reflections at the beginning (6 min, Fig. 11a) and at the end (80 min, Fig. 11c) of austempering can be fitted exactly by a single Gaussian model function and just differ significantly in intensity and peak width. Furthermore, the shift to lower scattering angles corresponds to the change from low-carbon austenite (γLC) to the enriched high-carbon austenite (γHC). At tAUS = 16 min an improved fit of the asymmetric reflection is obtained with the combination of two Gaussian peaks (Fig. 11b). The maximum in austenite peak width is therefore a consequence of the single peak fitting. The result, however, indicates that during austempering two modifications of austenite exist: a low-carbon fraction (γLC), which is the “untransformed” initial austenite, and the high-carbon austenite (γHC) between the sheaves and subunits of ferrite. This inhomogeneity in austenite carbon content leads to the observed peak asymmetry. It disappears after completion of the phase transformation, where just high-carbon austenite (γHC) remains and the carbon is distributed homogeneously by diffusion. However, the inhomogeneity is enhanced and prolongated by segregation of the alloying elements during

6000

62

γLC + γHC γLC γHC

(b)

63

64

γHC

(c)

4000

2000

0 60

61

60

61

60

61

Scattering angle 2θ [°] Fig. 11 – Peak shapes during austempering at 300 °C after 6 min, 16 min and 80 min, respectively. At the beginning of austempering (detail (a), 6 min) and at long austempering times (detail (c), 80 min) the reflections are symmetric and can be fitted by a single Gaussian model. At medium times (detail (b), 6 min) the peak shows an asymmetric shape caused by low-carbon (γLC) and high-carbon (γHC) austenite fractions.

solidification. Cu and Si segregate at the vicinity of the graphite nodules reducing the austenite carbon content which increases transformation kinetics. In contrast, Mn segregates at the last solidifying areas between the graphite, the so called “cell boundaries”, and delays the phase transformation [17]. In our investigation, this results just in a temporary coexistence of differently enriched austenite fractions, demonstrated by the austenite peak asymmetry. However, in higher alloyed cast irons the segregation of Mn and Mo can reach levels, where the phase transformation and thus the stabilisation of austenite at the cell boundaries is delayed so long, that a certain “untransformed austenite volume (UAV)” transforming to martensite upon cooling is inevitable [58]. The peak asymmetry effect was also observed during in-situ measurements of the austempering of high-silicon steels [32,33] and at the analysis of an interrupted austempering sample of ADI [23].

4.

Conclusion

The results show that in-situ neutron diffraction measurements are an excellent extension to interrupted austempering studies on ADI and give detailed insight to phase transformation kinetics and carbon redistribution during austempering. The isothermal holding temperature has great influence on microstructure, resulting phase fractions, kinetics and carbon redistribution:

132

M A TE RI A L S CH A RACT ER IZ A TI O N 85 (2 0 1 3 ) 1 2 4–1 3 3

• At higher austempering temperatures the reaction sets in almost immediately after quenching and the redistribution of carbon practically coincides with the phase transformation. A smooth transition to the stage II austenite decomposition takes place at TAUS = 400 °C. • At TAUS = 300 °C the phase transformation proceeds with almost the same rate than at higher austempering temperatures. However, the redistribution of carbon to austenite lags behind considerably, setting in only 15 min after quenching. One likely cause for this behaviour is that the ferrite is supersaturated with carbon on forming as seen from the initial lattice parameter expansion. The extent of this effect is remarkable and has not been observed in this detail so far. Further experiments are needed to clarify the underlying mechanism. • At all austempering temperatures the carbon enrichment is comparable to literature data, although at TAUS = 400 °C the maximum content might be limited by the onset of stage II decomposition. Furthermore at TAUS = 300 °C the total carbon content in the metallic matrix is significantly lower than the initial carbon content from austenitisation, presumably due to stage I transition carbide precipitation. • Reflection peak asymmetries revealed by the in-situ neutron diffraction experiments during austempering can be attributed to successive phase transformation resulting in two fractions of austenite modifications with different carbon contents.

Acknowledgements This work was supported by a grant (HO 2165/40-1 — PE 580/11-1) from the German Research Foundation/Deutsche Forschungsgemeinschaft (DFG). The provision of beam time to carry out the neutron diffraction experiments at the neutron source FRM II (TU München, Germany) is gratefully acknowledged. Finally, the authors would like to thank Ingo Schmidt (ACO Guss GmbH, Kaiserslautern, Germany) for the casting of the ductile iron.

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