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Effect of carbon on the damping capacity and mechanical properties of thermally trained Fe-Mn based high damping alloys
Author’s Accepted Manuscript Effect of Carbon on the Damping Capacity and Mechanical Properties of Thermally Trained Fe-Mn Based High Damping Alloys Won Seok Choi, Bruno C. De Cooman www.elsevier.com/locate/msea
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S0921-5093(17)30784-0 http://dx.doi.org/10.1016/j.msea.2017.06.020 MSA35158
To appear in: Materials Science & Engineering A Received date: 30 January 2017 Revised date: 31 March 2017 Accepted date: 5 June 2017 Cite this article as: Won Seok Choi and Bruno C. De Cooman, Effect of Carbon on the Damping Capacity and Mechanical Properties of Thermally Trained FeMn Based High Damping Alloys, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2017.06.020 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Effect of Carbon on the Damping Capacity and Mechanical Properties of Thermally Trained Fe-Mn Based High Damping Alloys Won Seok Choi1 and Bruno C. De Cooman2
1
Max-Planck-Institut für Eisenforschung GmbH Düsseldorf, Germany
2
Graduate Institute of Ferrous Technology, Pohang University of Science and Technology, Pohang, Republic of Korea
e-mail: [email protected]
1
ABSTRACT The effect of C on the damping and mechanical properties of thermally trained Fe-17 wt.%Mn-X wt.-%C (0
Keywords: Fe-Mn alloys; phase transformation; damping; internal friction; thermal cycling
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1. Introduction High-damping metallic materials (HDMM) have been developed for applications which require the reduction of noise and vibration, the prevention of fatigue problems, and the improvement of tool quality [1, 2]. A high damping capacity can be obtained in alloys with a heterogeneous microstructure, which have a thermo-elastic martensitic transformation, contain mobile magnetic domain walls, and/or mobile dislocations [3]. Point defect relaxations are not suitable damping mechanisms for HDMM since the corresponding damping peaks are usually located in a narrow frequency and temperature range [4]. The Fe17 wt.-%Mn alloy is a promising HDMM candidate due to its high damping capacity and mechanical properties which make it suitable for constructional and automotive applications. The alloy is also more cost efficient as compared to the Mn-Cu [5, 6], Cu-Zn-Al [7], Cu-AlNi [8, 9], and Ni-Ti [10] HDMMs.
The damping mechanism leading to the high damping capacity of Fe-Mn based alloys is known to be due to the hysteretic movement of the partial dislocations associated with stacking faults and -austenite/-martensite phase boundaries [2, 11-16]. The damping capacity should therefore depend on the density of stacking faults and the volume fraction of ɛ -martensite [13]. Huang et al. [15] however reported that the internal friction (IF) of Fe-Mn alloys was not related to the phase fraction of ɛ -martensite. They argued that the damping capacity is only determined by the motion of partial dislocations associated with stacking faults. The effect of interstitial C on the IF of Fe-Mn based HDMM has been reviewed in two reports [17, 18]. Both indicate that C has a negative influence on the damping capacity for two reasons: (a) the interaction between the damping sources and the solute C atoms, and (b) the reduction of ɛ -martensite volume fraction. 3
The ɛ -martensite phase fraction in a Fe-17 wt.-%Mn HDMM which has not been trained is lower as compared to the thermally trained material. Thermal cycling or training refers to the repeated forward transformation (upon heating) and backward transformation (upon cooling) to obtain a microstructure which reversibly transforms between a stable high temperature microstructure and a stable low temperature microstructure. Lee et al. [17] reported that cyclic thermal treatments increased the dislocation density and the ɛ -martensite volume fraction, and resulted in a reduction of the damping capacity. They also reported that an increase of the ɛ -martensite volume fraction obtained by quenching had the opposite effect. Shin et al. [19] however reported a higher damping capacity for a Fe-17 wt.-%Mn HDMM after thermal cycling.
The austenite start temperature (As) of the Fe-17 wt.-%Mn-X wt.-%C alloys used in the present work is approximately 200°C. Their martensite start temperature (Ms) is slightly above room temperature. This temperature range is similar to the one in the in-service conditions of industrial machinery, such as e.g. internal combustion engines. Fe-Mn based HDMMs are most likely to be utilized in these applications due to their combination of high damping capacity and mechanical properties. Small additions of C are expected to affect the phase transformation temperatures, the strength and the damping capacity of the alloy. The present study therefore focuses on the effect of the C content on the damping capacity and the mechanical properties of Fe-17 wt.-%Mn-X wt.-%C alloys after cyclic thermal treatments.
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2. Experimental procedure Fe-17 wt.-%Mn-X wt.-%C alloys were prepared by vacuum induction melting. The hot-rolled materials were annealed at 990°C for 30 minutes in a protective N2 atmosphere and subsequently water quenched to room temperature. The composition of the investigated alloys was Fe-17 wt.-%Mn-X wt.-%C, with X equal to 0.0, 0.01, 0.02, 0.03, 0.05, and 0.06.
The phase transformation temperatures and microstructural evolution during the training treatment was carried out in a Bähr 805 pushrod dilatometer operating in vacuum. The dilatometer specimens were prepared with their length direction parallel to the rolling direction. The specimens, 10 × 5 × 1 mm3 in size, were heated using a heating rate of +5°C/s to 300°C and cooled down to room temperature at a cooling rate of −5°C/s.
Phase fractions and grain size were determined by electron back-scatter diffraction (EBSD) in a FEI Quanta 3D scanning electron microscope equipped with a field emission electron source (FE-SEM). The TEM observations were carried out in a JEOL JEM-2100F operated at 200 kV. TEM foils were prepared by mechanically polishing samples to a 150 μm thickness, and subsequently electro-polishing the thin foils at room temperature using a solution of 10 % perchloric acid (HClO4) and 90 % acetic acid (CH3COOH).
Mechanical properties were obtained by means of uniaxial tensile tests using ASTM E8 standard sub-size tensile specimens with a gauge length of 25 mm. The tests were carried out on a Zwick/Roell Z100 universal tensile testing machine using a strain rate of 10-3 s-1.
The measurement of the internal friction (IF) spectra was done at resonant frequency using an 5
IMCE RFDA LTVP800 automated impulse excitation measurement system in which the test samples vibrate in the free flexural vibration mode. The resonant frequency f was approximately 800 Hz. The samples were 80 mm in length and 20 mm in width. The samples for IF measurements were cooled with liquid nitrogen to -70°C at the start of the IF measurements and heated under vacuum in an infra-red furnace using a heating rate of +5°C/min. It should be noted that the measuring frequency has an influence on the damping capacity i.e. the damping is lower a higher frequencies. As the measurement frequency in the present study was about 800 Hz, i.e. much higher than the 1 Hz frequency used in the conventional torsion pendulum equipment, direct comparisons of the present IF data with IF results obtained at other frequencies should be done carefully.
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3. Results Repeated thermal cycles were conducted by means of dilatometry to determine the number of training cycles required to obtain a microstructure which reversibly transforms between a stable high temperature microstructure and a stable low temperature microstructure. Fig. 1a shows the dilatometry data for the second thermal cycle for a sample of Fe-17 wt.-%Mn-0 wt.-%C.
The length reduction (h) observed upon cooling form high temperature is
indicative of an increase of ɛ -martensite phase fraction and the shape memory effect due to the thermo-elastic martensite transformation [20]. Note that a thermal cycle to 300°C is sufficient to fully austenitize the alloy. The dilatometry curves gradually form a stable closed loop with increasing number of thermal cycle. A negligible room temperature length change is obtained after the fifth thermal cycle as shown in Fig. 1b. All the microstructural observations, IF measurements, and mechanical properties determinations reported hereafter were conducted with alloys that had been trained five times.
The repeated dilatometer experiments also allow for the determination of the phase transformation temperatures As and Ms. Fig. 2 shows that the As temperature decreases from 189°C to 171°C with increasing C content, and that the Ms temperature decreases from 147°C to 100°C with increasing C content. This decrease is consistent with the fact that C is a strong austenite stabilizer. The C additions decrease the As and Ms temperatures by -298°C/wt.-%C and -758°C/wt.-%C, respectively. The reported C content dependence of the Ms temperature based on thermodynamics calculation is -710°C/wt.-%C [21]. This is in good agreement with the present experimental result [21].
More than 300 grains were analyzed to calculate -austenite and-martensite phase fractions 7
and their average grain size by EBSD (Fig. 3). The results are summarized in Table 1.The austenite and-martensite phase fractions and their grain size were found to be independent of the C content of the Fe-17 wt.-%Mn-X wt.-%C alloys. The phase fractions and the grain sizes remained almost constant. The phase fraction and the grain size of the ɛ -martensite were 0.91 and 1.31 µm, respectively. The phase fraction and the grain size of the -austenite was 0.09 and 0.62 µm, respectively. These results are in agreement with the X-ray diffraction (XRD) test results reported for a thermally trained Fe-17 wt.-%Mn alloy [17]. The effect of the C on the damping capacity of Fe-17 wt.-%Mn-X wt.-%C alloys could therefore be analyzed by assuming that the phase fraction and the grain size of -austenite andmartensite were not influenced by an increase in the C content up to 0.06 wt.-%.
The microstructure of the thermally trained alloy as observed in TEM consisted only of austenite and-martensite. Fig. 4 shows TEM micrographs and the corresponding selected area diffraction patterns (SADP) taken from ɛ -martensite grains (Fig. 4g), alternating layers of -austenite and-martensite (Fig. 4h), and alternating layers of -austenite, twinned austenite and-martensite layers (Fig. 4i) in the thermally trained Fe-17 wt.-%Mn-0 wt.-%C alloy. The TEM bright field (BF) micrograph shown in Fig. 4a and the corresponding dark field (DF) micrographs shown in Fig. 4b, c, and d reveal the presence of ɛ -martensite, γaustenite, and twinned γ-austenite. The diffraction spots used for the dark field micrographs are indicated by colored arrows in Fig. 4i.
The grains always contained stacking faults, as shown in Fig. 4e. The characteristic microstructure of defect-free ɛ -martensite is readily visible in the high resolution lattice 8
image shown in Fig. 4f. The measured lattice parameters of γ-austenite, aγ, and ɛ -martensite, aɛ and cɛ , are found to be approximately to 0.362 nm, 0.254 nm, and 0.417 nm, respectively. The c/a ratio of the ɛ -martensite is 1.642. This is very close to the ideal c/a packing ratio of 1.633.
Fig. 5 shows the engineering stress-strain curves of the Fe-17 wt.-%Mn-X wt.-%C alloys. The alloy C content dependence of the yield strength (YS) and the ultimate tensile strength (UTS) is shown in Fig. 5b. While the YS increases by 24 MPa for an addition of 0.06 wt.-%C, corresponding to a solid solution strengthening effect of +533 MPa/wt.-%C, the increase in UTS resulting from the addition of C is much more pronounced, i.e. +4967 MPa/wt.-%C.
There was no clear C content dependence of the elongation, but the uniform elongation was always larger than 20 %, regardless of C content. The alloys used in the present work are characterized by a high content of ɛ -martensite, which has a hexagonal close-packed (hcp) crystal structure. Metals or alloys with a hcp crystal structure are known to have a poor plasticity due to the fact that fewer slip systems are activated during deformation as compared to metals and alloys with a cubic crystal structure. In a recent study by Kim and De Cooman, accommodation of c-axis deformation was identified in a Fe-17 wt.-%Mn-0.03 wt.-%C alloy [22]. This mechanism is believed to enable the large homogeneous deformation of ɛ martensite.
It is also noteworthy that serrations were present on the tensile stress-strain curves for alloys with a C content ≥0.03 wt.-%C. Koyama et al. [23] reported that the serrated flow curves, observed for the austenitic Fe-17 wt%Mn-0.6 wt.-%C and Fe-17 wt%Mn-0.8 wt.-%C alloys, 9
were due to dynamic strain aging (DSA) rather than deformation twinning or strain-induced martensite transformation. The combined analysis of the temperature dependence of the dilatation, Young’s modulus, and IF of the fully trained Fe-17 wt.-%Mn-X wt.-%C alloys was carried out to identify the effect of an increasing content of solute C. Fig.6 compares the results for Fe-17 wt.-%Mn-0 wt.-%C and Fe-17 wt.-%Mn-0.06 wt.-%C. The temperature dependence of Young’s modulus can be divided in three stages: (i) a linear temperature dependence at temperatures below 0°C, (ii) a step-like increase at the Néel temperature (TN ~ 4°C) [24], and (iii) a negative temperature dependence with a small reduction at the As temperature for the γ→ɛ phase transformation. In stage (ii) the change of Young’s modulus is associated with the antiferromagnetic transition in the γ phase. The size of the modulus anomaly is affected by the volume fraction of γ-austenite [19]. In the present study, it was clearly visible despite the fact that there was only a small phase fraction of γ-austenite in the alloy microstructure (~ 0.09). The modulus anomaly is caused by a magneto-volume effect [25]. Magnetic transitions may result in relaxation peaks in IF spectra due to the effect of stress on the magnetic ordering [3]. A small IF peak, very likely related to the anti-ferromagnetic transition, was observed for the Fe-17 wt.-%Mn-0.06 wt.-%C alloy in the present study.
The internal friction spectra of the Fe-17 wt.-%Mn-X wt.-%C alloys are compared in Fig. 7. The phase fractions were calculated by means of dilatometer data. In the heating segment of the IF spectrum for the Fe-17 wt.-%Mn-0 wt.-%C alloy (Fig. 7a), there are three different potential contributions to the damping: (a) the intrinsic internal friction of ɛ -martensite phase (IFint. ɛ ), (b) the anti-ferromagnetic transition, and (c) the ɛ →γ phase transformation. It is not possible to quantify the contribution of each relaxation effect because the three processes 10
overlap and this results in a broad IF peak in the 0ºC-250ºC temperature range.
The effect of the C addition is already very apparent in the IF spectrum of the Fe-17 wt.%Mn-0.01 wt.-%C alloy. There is a clear reduction of the intrinsic damping capacity of the ɛ -martensite phase during both the heating and cooling stages. The intrinsic damping capacity is gradually reduced with increasing C content and the effect saturates for a C content ≥0.03 wt.-%.
The contribution of the magnetic transition and the ɛ →γ phase transformation to the damping is clearly visible as a small internal friction peak. The overall damping capacity is reduced by C additions. The internal friction peak induced by ɛ →γ phase transformation was not identified when the C content exceeded 0.03 wt.-%C.
At temperatures above the austenite finish temperature (Af) the microstructure consists entirely of γ-austenite. The segment of the internal friction spectrum above Af therefore corresponds to the intrinsic internal friction of γ-austenite (IFint. γ). The IF measurement results clearly show that IFint. γ is not influenced by C. The IF spectrum measured in the cooling stage (Fig. 7b) also shows that IFint. γ is independent of the C content, as there is only an increase of damping capacity at the γ→ɛ martensite transformation start temperature.
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4. Discussion The damping capacity of metallic materials is usually to be reversely proportional to the tensile strength, i.e. high strength metallic materials have a low damping capacity [26]. HDMM such as Fe-Mn based HDMMs are unusual in that they combine high strength properties and a high damping capacity [1, 2]. As C is known to be a very effective strengthening addition in most ferrous alloys, the addition of C to Fe-Mn based HDMMs is an obvious approach to further increase the strength of Fe-Mn based HDMMs [27]. The experimental results reviewed in the previous section however show that the addition of C has a detrimental effect on the damping properties. The solute C atoms must therefore prevent the activation of the relaxation processes which are responsible for damping in Fe-Mn based HDMMs.
In the present work, Fe-17 wt.-%Mn-X wt.-%C alloys with six different C contents in the range of 0.00-0.06 wt.-% were subjected to identical thermos-mechanical processes and cyclic thermal treatments, which resulted in similar C-content independent microstructures. The average grain size and the phase fractions were both similar in the six alloys, as shown in Fig. 3 and Table 1. The C content was therefore the only parameter affecting the mechanical and damping properties of the material.
The following is known about solute C atoms in γ-austenite Fe-Mn alloys: 1. C increases the stability of the austenite relative to -martensite. 2. C promotes paramagnetism, i.e. C decreases the Néel temperature of γ-austenite. 3. C increases the stacking fault energy substantially. Abbasi et al. [28] reported a significant increase of the stacking fault energy with increasing C content. They also reported a 12
strong dependence of the stacking fault energy on the position of the C atoms relative to the stacking fault. Experimental evidence for the nano-diffusion of C away from the stacking fault has been observed by Hickel et al. [29]. 4. C has a strong solid solution hardening effect in γ-austenite alloys. The solid solution strengthening effect of C, i.e. approximately +279 MPa/wt.-%C, in γ-austenite Fe-Mn-C TWIP steel [30] is similar to the solid solution strengthening effect of solute C in stainless steel. Allain reported a similar solid solution hardening effect of +250 MPa/wt.-%C [31]. 5. C occupies octahedral interstitial sites in -Fe and alloys. The small lattice distortion caused by C is isotropic. The nearest neighbor (NN) and next nearest neighbor (NNN) interactions between C-C pairs in austenite have been found to be repulsive in -Fe-C alloys [32, 33]. This result implies that the C distribution is expected to be homogeneous. This reduces the probability that an interstitial octahedral site is occupied in the vicinity of an octahedral site already occupied by a C atom, and leads to an increased C diffusivity with increasing C content [34, 35]. First principle analysis of austenitic alloys has shown that there is a pronounced repulsion between solute C atoms [28, 36-38]. 6. The strong attractive NN C-Mn (+26kJ/mol) and NNN C-Mn (+13kJ/mol) interaction between Mn and C atoms in γ-austenite Fe-Mn-C alloys result in the formation of Mn-C pairs [36, 37]. 7. Shun et al. [39], who studied the temperature dependence of the properties of austenitic Fe-30 wt.-%Mn-1.0 wt.-%C, reported that the apparent activation energy for the onset of serration in the flow curve was much lower than the activation energy for C bulk diffusion. This suggests that the serrations are related to the short range diffusion of C in the dislocation core. In the present work serrations were observed for the Fe-17 wt.%Mn-X wt.-%C alloys with a C content ≥0.03 wt.-%. The C-Mn complexes are expected 13
to be formed in a well-annealed, undeformed high Mn alloys due to the C-Mn attractive force. Lee et al. [40] proposed that DSA in high Mn γ-austenite Fe-Mn-Al-C TWIP steel results from the interaction of C-Mn complexes with staking faults by a single diffusive jump, rather than C diffusion to dislocation core.
Much less is known about the effects of solute C atoms in -martensitic Fe-Mn alloys, but if the high -martensite phase fraction is considered in the present Fe-Mn alloys, it is clear that the room temperature effects observed when the C content is increased are due to the effect of solute C on the mechanical properties and the damping of the -martensite. Lee et al. [17] have reported that C additions significantly deteriorated the damping capacity of Fe-Mn alloys. They mentioned that a 0.06 wt.-%C addition decreased the damping capacity of asquenched Fe-Mn alloys by approximately 1 %. In the present study, the observed decrease in damping capacity at room temperature was 64 % for a 0.01 wt.-%C addition and 80 % for a 0.06 wt.-%C addition. The peak damping at 150°C decreases by 76 % for a 0.01 wt.-%C addition, and by 93 % for a 0.06 wt.-%C addition.
Similarly to the case of γ-austenite, solute C or C-substitutional atom point defect complexes can affect the motion of mobile lattice defects in -martensite, i.e. dislocations and interface boundaries, responsible for the damping. The damping will be reduced when these relaxation processes are inhibited. It is very likely that the cyclic thermal treatment allows for both (a) C atoms to diffuse to mobile interface boundary segregation sites, and (b) C atoms to form stable C-Mn complexes in the matrix. In both cases, the mobility of the interfaces could be impeded and the damping capacity reduced. 14
It is well known that IFint. γ is low as compared to IFint. ɛ [3]. This has also been confirmed in the present experimental result for alloys with a low C content. With increasing C content IFint. ɛ decreased and its value became similar or even lower than IFint. γ. IFint. γ was not influenced by the C content. The cooling segment of the IF spectrum of the Fe-17 wt.-%Mn0.06 wt.-%C alloy (Fig. 7b) shows a steady decrease in damping capacity with decreasing temperature even after the start of the martensitic transformation.
Grain boundary related damping has been reported to be the major source of IFint. γ in ultrafine grained Fe-17 wt.-%Mn [19], nano-crystalline copper and stainless steel [41]. Grain boundary sliding and the disordering and reordering of atom groups in grain boundaries by thermal excitation at high temperature have been proposed as the damping mechanisms [3]. In the present study the grain size contribution to damping was not affected by the C content as care was taken to study alloys with a very similar grain size.
An alternative origin for the damping in Fe-17 wt.-%Mn-X wt.-%C alloys could be the ↔ phase transformation. The damping Q -1 associated with a solid-state phase transformation can be described by the De Jonghe-Delorme equation [42], 3 m T 1 4 C m Q k 0 1 0 T t f 3 -1
here k ,
(1)
m T m , , 0 , , and C are a material constant, the quantity of T t
transformed material per unit temperature change, the heating or cooling rate, the stress15
amplitude during the internal friction experiment, the quantity of transformed material per unit stress, and the critical stress required for the stress-induced transformation or the defect reorientation, respectively. The explanation for the difference in the amplitude of the internal friction peaks during heating and cooling is given by the De Jonghe-Delorme equation. The heating rate used for internal friction measurement was +5 °C/min, whereas the specimen was slowly cooled naturally without the use of a cooling gas. The transformation rate per a unit temperature change is therefore lower during the cooling (Fig. 7). As a consequence, the damping is reduced due to the smaller value of the factors
m T and during cooling. T t
The De Jonghe-Delorme equation also explains two observations made for alloys with a higher C content: (a) the absence of an IF peak related to phase transformation at high C contents, and (b) the fact that IFint. γ and IFint. ɛ become similar at high C contents.
As discussed above, C atoms can diffuse to preferential segregation sites e.g. mobile partial dislocations associated with stacking faults, mobile γ/ɛ phase boundaries and other mobile interfaces, immobilizing them. The
m term in the De Jonghe-Delorme equation is related
to the mobility of boundaries and interfaces [42]. The applied stress range, at which the damping was measured, was constant in the present experiments. If solute C atoms reduce the mobility of both interfaces in ɛ -martensite interfaces and the γ/ɛ phase boundaries interfaces, IFint. ɛ should decrease and no phase transformation-related peak should be observed. This appears to be the case in the alloys with C ≥0.03 wt.-%C. When the C content is such that only the C segregation sites at interfaces in ɛ -martensite are occupied, IFint. ɛ should be 16
reduced, but the ↔phase transformation-related internal friction peak should still be present. This is the case for the alloys with a C content of 0.01 wt.-% and 0.02 wt.-%.
If the γ/ɛ -phase boundaries were preferred segregation sites for solute C atoms, no peak would be obtained in the temperature interval of the phase transformation, and there should be a clear difference between IFint. γ and IFint. ɛ , which should not be affected by the C content. This was however not observed in the present case. The experimental results therefore suggest that ɛ /ɛ -martensite interfaces were the preferred segregation sites for the solute C atoms.
In the absence of C additions, i.e. for the Fe-17 wt.-%Mn alloy, the motion of interfaces in ɛ martensite and ↔phase boundaries were not impeded. This resulted in a higher value of IFint. ɛ as compared to IFint. γ and a very pronounced ↔phase transformation peak.
Comparison of the results of the present study with a previous study by Shin et al. [19], offers the opportunity to analyze the effect of cold-rolling on the damping capacity of Fe-17 wt.%Mn. Shin et al. [19] analyzed the effect of grain size on the damping spectrum of an Fe-17 wt.-%Mn alloy. Their Fe-17 wt.-%Mn was cold-rolled, recrystallization annealed at 1000°C for 30 minutes, and trained four times using a thermal cycle involving reheating to 400°C for 10 minutes and cooling to room temperature. The reported phase fractions were similar to those in the present work. The average grain size was larger, i.e. 9.6 µm for ɛ -martensite and 2.1 µm for γ-austenite.
The internal friction spectra published by Shin et al. [19] showed a clear damping peak to the 17
anti-ferromagnetic transition and the ↔ phase transformation. The 150°C damping peak size for the hot rolled and trained Fe-17 wt.-%Mn alloy used in the present study was three times larger than the damping peak observed at the same temperature for their cold-rolled Fe17 wt.-%Mn alloy.
In the present work the maximum damping peak was located at a temperature between the modulus anomaly temperature and the phase transformation temperature, clearly indicating that the motion of interfaces in ɛ -martensite was enhanced in the hot-rolled microstructure as compared to the cold-rolled and recrystallized microstructure. This results in a broad IF peak by the combined effect of the motion of interfaces in ɛ -martensite, ↔phase boundaries and the anti-ferromagnetic transition in the hot-rolled material. The difference between the IF spectra of the same alloy in the hot-rolled and the cold-rolled state is very likely due to an unfavorable crystallographic texture developed after cold rolling and recrystallization annealing. This effect is currently being analyzed in more detail for reporting at a later date.
Fig. 8 summarizes the effect of the C content on damping capacity and mechanical properties of the Fe-17 wt.-%Mn-X wt.-%C (0
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5. Conclusions The effect of C on damping capacity in thermally trained Fe-17 wt.-%Mn-X wt.-%C (0
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Acknowledgements The authors sincerely acknowledge the support of Dr Tae Jin Song of the POSCO Technical Research Laboratories in Gwangyang, Republic of Korea.
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354. 22. J.-K. Kim, B.C. De Cooman, Scripta Mater. 128 (2017) 78-82. 23. M. Koyama, T. Sawaguchi, T. Lee, C.S. Lee, K. Tsuzaki, Mater. Sci. Eng. A 528 (2011) 7310-7316. 24. A. Dumay, J.P. Chateau, S. Allain, S. Migot, O. Bouaziz, Mater. Sci. Eng. A 483–484 (2008) 184–187. 25. D. Music, T. Takahashi, L. Vitos, C. Asker, I.A. Abrikosov, J.M. Schneider, Appl. Phys. Lett. 91 (2007) 191904. 26. H. Mizubayashi, S. Murayama, H. Tanimoto, J. Alloys Comp. 330-332 (2002) 389-392. 27. O. Bouaziz, S. Allain, C.P. Scott, P. Cugy, D. Barbier, Curr. Opin. Solid State Mater. Sci. 15 (2011) 141-168. 28. A. Abbasi, A. Dick, T. Hickel, J. Neugebauer, Acta Mater. 59 (2011) 3041-3048. 29. T. Hickel, S. Sandlöbes, R.K.W. Marceau, A. Dick, I. Bleskov, J. Neugebauer, D. Raabe, Acta Mater. 75 (2014) 147-155. 30. W.S. Choi, B.C. De Cooman, S. Sandlöbes, D. Raabe, Acta Mater. 98 (2015) 391-404. 31. S. Allain, These d’Habilitation, Université de Lorraine, INPL, Grenoble, (2012). 32. K. Oda, H. Fujimura, H. Ino, J. Phys.: Condens. Matter 6 (1994) 679-692. 33. H.K.D.H. Bhadeshia, J. Mater. Sci. 39 (2004) 3949–3955. 34. R.B. McLellan, C. Ko, Acta Metall. 36 (1988) 531-537. 35. S.S. Babu, H.K.D.H. Bhadeshia, J. Mater. Sci. Lett. 14 (1995) 314-316. 36. N.I. Medvedeva, M.S. Park, D.C. Van Aken, J.E. Medvedeva, J. Alloys Comp. 582 (2014) 475–482. 37. K.R. Limmer, J.E. Medvedeva, D.C. Van Aken, N.I. Medvedeva, Comput. Mater. Sci. 99 (2015) 253-255. 38. D.J. Hepburn, D. Ferguson, S. Gardner, G.J. Ackland, Phys. Rev. B 88 (2013) 024115. 39. T. Shun, C.M. Wan, J.G. Byrne, Acta Metall. et Mater. 40 (1992) 3407-3412. 40. S-J. Lee, J. Kim, S.N. Kane, B.C. De Cooman, Acta Mater. 59 (2011) 6809-6819. 41. R.R. Mulyukov, A.I. Pshenichnyuk, Metal Sci. Heat Treat. 54 (2012) 244–248. 22
42. W. Dejonghe, R. De Batist, L. Delaey, Scr. Metall. 10 (1976) 1125-1128.
23
Fig. 1 – Dilatometer measurements used to determine the number of thermal cycles required to obtain a microstructure in Fe-17 wt.-%Mn-X wt.-%C high-damping alloys which reversibly transforms between a stable high temperature microstructure and a stable low temperature microstructure. The stable dilatation loop was obtained after the 5th training 24
cycle.
25
Fig. 2 – C content dependence of the austenite start temperature (As) and the martensite start temperature (Ms).
26
Fig. 3 – Combined EBSD phase map and image quality (IQ) map of the thermally trained Fe17 wt.-%Mn-X wt.-%C high-damping alloys. The average grain size of ɛ -martensite and γaustenite is 1.31 µm and 0.62 µm, respectively. The average ɛ -martensite phase fraction is 0.91.
27
28
Fig. 4 – TEM micrographs of the trained Fe-17 wt.-%Mn alloy. (a) Bright field image and the corresponding dark field images of (b) ɛ -martensite, (c) γ-austenite, and (d) twinned γaustenite (γ'). (e) Bright field image of a stacking fault network. (f) High resolution micrograph of a ɛ -martensite plate. Selected area diffraction pattern of (g) ɛ -martensite, (h) overlapping layers of ɛ -martensite and γ-austenite, and (i) overlapping layers of ɛ martensite, γ-austenite and twinned γ-austenite (γ'). The diffraction spots used for dark field images (b), (c), and (d) are indicated by colored arrows in (i).
29
Fig. 5 – (a) Engineering stress-strain curves for the Fe-17 wt.-%Mn-X wt.-%C alloys. Serrations occur in the engineering stress-strain curves of the alloys with a C content ≥0.03 wt.-%C. (b) Solute C strengthening effect for the yield stress and the ultimate tensile strength of Fe-17 wt.-%Mn-X wt.-%C alloys. 30
Fig. 6 – Temperature dependence of the dilatation, Young’s modulus, and damping capacity of (a) Fe-17wt%Mn-0.0wt%C and (b) Fe-17wt%Mn-0.06wt%C alloys.
31
32
Fig. 7 – Comparison of internal friction spectra measured during (a) heating and (b) cooling. The modulus change and phase transformation start temperature range are indicated. The phase fraction evolution during the phase transformation of the Fe-17wt%Mn-0.0wt%C alloy is also indicated.
33
Fig. 8 – The effect of the C content on the mechanical properties of Fe-17 wt.-%Mn-X wt.%C alloys and the intrinsic internal friction of ɛ -martensite (IFint. ɛ ) and γ-austenite (IFint. ).
34
Table 1 –C content dependence of the phase fraction and grain size of ɛ -martensite and γ after thermal training.
ɛ
γ
Sample
Phase fraction
Grain size (µm)
Phase fraction
Grain size (µm)
0.0C
0.95
1.45
0.05
0.41
0.01C
0.89
1.17
0.11
0.58
0.02C
0.91
1.44
0.09
0.73
0.03C
0.88
1.33
0.12
0.70
0.05C
0.89
1.25
0.11
0.66
0.06C
0.91
1.2
0.09
0.64
35
PII: DOI: Reference:
S0921-5093(17)30784-0 http://dx.doi.org/10.1016/j.msea.2017.06.020 MSA35158
To appear in: Materials Science & Engineering A Received date: 30 January 2017 Revised date: 31 March 2017 Accepted date: 5 June 2017 Cite this article as: Won Seok Choi and Bruno C. De Cooman, Effect of Carbon on the Damping Capacity and Mechanical Properties of Thermally Trained FeMn Based High Damping Alloys, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2017.06.020 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Effect of Carbon on the Damping Capacity and Mechanical Properties of Thermally Trained Fe-Mn Based High Damping Alloys Won Seok Choi1 and Bruno C. De Cooman2
1
Max-Planck-Institut für Eisenforschung GmbH Düsseldorf, Germany
2
Graduate Institute of Ferrous Technology, Pohang University of Science and Technology, Pohang, Republic of Korea
e-mail: [email protected]
1
ABSTRACT The effect of C on the damping and mechanical properties of thermally trained Fe-17 wt.%Mn-X wt.-%C (0
Keywords: Fe-Mn alloys; phase transformation; damping; internal friction; thermal cycling
2
1. Introduction High-damping metallic materials (HDMM) have been developed for applications which require the reduction of noise and vibration, the prevention of fatigue problems, and the improvement of tool quality [1, 2]. A high damping capacity can be obtained in alloys with a heterogeneous microstructure, which have a thermo-elastic martensitic transformation, contain mobile magnetic domain walls, and/or mobile dislocations [3]. Point defect relaxations are not suitable damping mechanisms for HDMM since the corresponding damping peaks are usually located in a narrow frequency and temperature range [4]. The Fe17 wt.-%Mn alloy is a promising HDMM candidate due to its high damping capacity and mechanical properties which make it suitable for constructional and automotive applications. The alloy is also more cost efficient as compared to the Mn-Cu [5, 6], Cu-Zn-Al [7], Cu-AlNi [8, 9], and Ni-Ti [10] HDMMs.
The damping mechanism leading to the high damping capacity of Fe-Mn based alloys is known to be due to the hysteretic movement of the partial dislocations associated with stacking faults and -austenite/-martensite phase boundaries [2, 11-16]. The damping capacity should therefore depend on the density of stacking faults and the volume fraction of ɛ -martensite [13]. Huang et al. [15] however reported that the internal friction (IF) of Fe-Mn alloys was not related to the phase fraction of ɛ -martensite. They argued that the damping capacity is only determined by the motion of partial dislocations associated with stacking faults. The effect of interstitial C on the IF of Fe-Mn based HDMM has been reviewed in two reports [17, 18]. Both indicate that C has a negative influence on the damping capacity for two reasons: (a) the interaction between the damping sources and the solute C atoms, and (b) the reduction of ɛ -martensite volume fraction. 3
The ɛ -martensite phase fraction in a Fe-17 wt.-%Mn HDMM which has not been trained is lower as compared to the thermally trained material. Thermal cycling or training refers to the repeated forward transformation (upon heating) and backward transformation (upon cooling) to obtain a microstructure which reversibly transforms between a stable high temperature microstructure and a stable low temperature microstructure. Lee et al. [17] reported that cyclic thermal treatments increased the dislocation density and the ɛ -martensite volume fraction, and resulted in a reduction of the damping capacity. They also reported that an increase of the ɛ -martensite volume fraction obtained by quenching had the opposite effect. Shin et al. [19] however reported a higher damping capacity for a Fe-17 wt.-%Mn HDMM after thermal cycling.
The austenite start temperature (As) of the Fe-17 wt.-%Mn-X wt.-%C alloys used in the present work is approximately 200°C. Their martensite start temperature (Ms) is slightly above room temperature. This temperature range is similar to the one in the in-service conditions of industrial machinery, such as e.g. internal combustion engines. Fe-Mn based HDMMs are most likely to be utilized in these applications due to their combination of high damping capacity and mechanical properties. Small additions of C are expected to affect the phase transformation temperatures, the strength and the damping capacity of the alloy. The present study therefore focuses on the effect of the C content on the damping capacity and the mechanical properties of Fe-17 wt.-%Mn-X wt.-%C alloys after cyclic thermal treatments.
4
2. Experimental procedure Fe-17 wt.-%Mn-X wt.-%C alloys were prepared by vacuum induction melting. The hot-rolled materials were annealed at 990°C for 30 minutes in a protective N2 atmosphere and subsequently water quenched to room temperature. The composition of the investigated alloys was Fe-17 wt.-%Mn-X wt.-%C, with X equal to 0.0, 0.01, 0.02, 0.03, 0.05, and 0.06.
The phase transformation temperatures and microstructural evolution during the training treatment was carried out in a Bähr 805 pushrod dilatometer operating in vacuum. The dilatometer specimens were prepared with their length direction parallel to the rolling direction. The specimens, 10 × 5 × 1 mm3 in size, were heated using a heating rate of +5°C/s to 300°C and cooled down to room temperature at a cooling rate of −5°C/s.
Phase fractions and grain size were determined by electron back-scatter diffraction (EBSD) in a FEI Quanta 3D scanning electron microscope equipped with a field emission electron source (FE-SEM). The TEM observations were carried out in a JEOL JEM-2100F operated at 200 kV. TEM foils were prepared by mechanically polishing samples to a 150 μm thickness, and subsequently electro-polishing the thin foils at room temperature using a solution of 10 % perchloric acid (HClO4) and 90 % acetic acid (CH3COOH).
Mechanical properties were obtained by means of uniaxial tensile tests using ASTM E8 standard sub-size tensile specimens with a gauge length of 25 mm. The tests were carried out on a Zwick/Roell Z100 universal tensile testing machine using a strain rate of 10-3 s-1.
The measurement of the internal friction (IF) spectra was done at resonant frequency using an 5
IMCE RFDA LTVP800 automated impulse excitation measurement system in which the test samples vibrate in the free flexural vibration mode. The resonant frequency f was approximately 800 Hz. The samples were 80 mm in length and 20 mm in width. The samples for IF measurements were cooled with liquid nitrogen to -70°C at the start of the IF measurements and heated under vacuum in an infra-red furnace using a heating rate of +5°C/min. It should be noted that the measuring frequency has an influence on the damping capacity i.e. the damping is lower a higher frequencies. As the measurement frequency in the present study was about 800 Hz, i.e. much higher than the 1 Hz frequency used in the conventional torsion pendulum equipment, direct comparisons of the present IF data with IF results obtained at other frequencies should be done carefully.
6
3. Results Repeated thermal cycles were conducted by means of dilatometry to determine the number of training cycles required to obtain a microstructure which reversibly transforms between a stable high temperature microstructure and a stable low temperature microstructure. Fig. 1a shows the dilatometry data for the second thermal cycle for a sample of Fe-17 wt.-%Mn-0 wt.-%C.
The length reduction (h) observed upon cooling form high temperature is
indicative of an increase of ɛ -martensite phase fraction and the shape memory effect due to the thermo-elastic martensite transformation [20]. Note that a thermal cycle to 300°C is sufficient to fully austenitize the alloy. The dilatometry curves gradually form a stable closed loop with increasing number of thermal cycle. A negligible room temperature length change is obtained after the fifth thermal cycle as shown in Fig. 1b. All the microstructural observations, IF measurements, and mechanical properties determinations reported hereafter were conducted with alloys that had been trained five times.
The repeated dilatometer experiments also allow for the determination of the phase transformation temperatures As and Ms. Fig. 2 shows that the As temperature decreases from 189°C to 171°C with increasing C content, and that the Ms temperature decreases from 147°C to 100°C with increasing C content. This decrease is consistent with the fact that C is a strong austenite stabilizer. The C additions decrease the As and Ms temperatures by -298°C/wt.-%C and -758°C/wt.-%C, respectively. The reported C content dependence of the Ms temperature based on thermodynamics calculation is -710°C/wt.-%C [21]. This is in good agreement with the present experimental result [21].
More than 300 grains were analyzed to calculate -austenite and-martensite phase fractions 7
and their average grain size by EBSD (Fig. 3). The results are summarized in Table 1.The austenite and-martensite phase fractions and their grain size were found to be independent of the C content of the Fe-17 wt.-%Mn-X wt.-%C alloys. The phase fractions and the grain sizes remained almost constant. The phase fraction and the grain size of the ɛ -martensite were 0.91 and 1.31 µm, respectively. The phase fraction and the grain size of the -austenite was 0.09 and 0.62 µm, respectively. These results are in agreement with the X-ray diffraction (XRD) test results reported for a thermally trained Fe-17 wt.-%Mn alloy [17]. The effect of the C on the damping capacity of Fe-17 wt.-%Mn-X wt.-%C alloys could therefore be analyzed by assuming that the phase fraction and the grain size of -austenite andmartensite were not influenced by an increase in the C content up to 0.06 wt.-%.
The microstructure of the thermally trained alloy as observed in TEM consisted only of austenite and-martensite. Fig. 4 shows TEM micrographs and the corresponding selected area diffraction patterns (SADP) taken from ɛ -martensite grains (Fig. 4g), alternating layers of -austenite and-martensite (Fig. 4h), and alternating layers of -austenite, twinned austenite and-martensite layers (Fig. 4i) in the thermally trained Fe-17 wt.-%Mn-0 wt.-%C alloy. The TEM bright field (BF) micrograph shown in Fig. 4a and the corresponding dark field (DF) micrographs shown in Fig. 4b, c, and d reveal the presence of ɛ -martensite, γaustenite, and twinned γ-austenite. The diffraction spots used for the dark field micrographs are indicated by colored arrows in Fig. 4i.
The grains always contained stacking faults, as shown in Fig. 4e. The characteristic microstructure of defect-free ɛ -martensite is readily visible in the high resolution lattice 8
image shown in Fig. 4f. The measured lattice parameters of γ-austenite, aγ, and ɛ -martensite, aɛ and cɛ , are found to be approximately to 0.362 nm, 0.254 nm, and 0.417 nm, respectively. The c/a ratio of the ɛ -martensite is 1.642. This is very close to the ideal c/a packing ratio of 1.633.
Fig. 5 shows the engineering stress-strain curves of the Fe-17 wt.-%Mn-X wt.-%C alloys. The alloy C content dependence of the yield strength (YS) and the ultimate tensile strength (UTS) is shown in Fig. 5b. While the YS increases by 24 MPa for an addition of 0.06 wt.-%C, corresponding to a solid solution strengthening effect of +533 MPa/wt.-%C, the increase in UTS resulting from the addition of C is much more pronounced, i.e. +4967 MPa/wt.-%C.
There was no clear C content dependence of the elongation, but the uniform elongation was always larger than 20 %, regardless of C content. The alloys used in the present work are characterized by a high content of ɛ -martensite, which has a hexagonal close-packed (hcp) crystal structure. Metals or alloys with a hcp crystal structure are known to have a poor plasticity due to the fact that fewer slip systems are activated during deformation as compared to metals and alloys with a cubic crystal structure. In a recent study by Kim and De Cooman, accommodation of c-axis deformation was identified in a Fe-17 wt.-%Mn-0.03 wt.-%C alloy [22]. This mechanism is believed to enable the large homogeneous deformation of ɛ martensite.
It is also noteworthy that serrations were present on the tensile stress-strain curves for alloys with a C content ≥0.03 wt.-%C. Koyama et al. [23] reported that the serrated flow curves, observed for the austenitic Fe-17 wt%Mn-0.6 wt.-%C and Fe-17 wt%Mn-0.8 wt.-%C alloys, 9
were due to dynamic strain aging (DSA) rather than deformation twinning or strain-induced martensite transformation. The combined analysis of the temperature dependence of the dilatation, Young’s modulus, and IF of the fully trained Fe-17 wt.-%Mn-X wt.-%C alloys was carried out to identify the effect of an increasing content of solute C. Fig.6 compares the results for Fe-17 wt.-%Mn-0 wt.-%C and Fe-17 wt.-%Mn-0.06 wt.-%C. The temperature dependence of Young’s modulus can be divided in three stages: (i) a linear temperature dependence at temperatures below 0°C, (ii) a step-like increase at the Néel temperature (TN ~ 4°C) [24], and (iii) a negative temperature dependence with a small reduction at the As temperature for the γ→ɛ phase transformation. In stage (ii) the change of Young’s modulus is associated with the antiferromagnetic transition in the γ phase. The size of the modulus anomaly is affected by the volume fraction of γ-austenite [19]. In the present study, it was clearly visible despite the fact that there was only a small phase fraction of γ-austenite in the alloy microstructure (~ 0.09). The modulus anomaly is caused by a magneto-volume effect [25]. Magnetic transitions may result in relaxation peaks in IF spectra due to the effect of stress on the magnetic ordering [3]. A small IF peak, very likely related to the anti-ferromagnetic transition, was observed for the Fe-17 wt.-%Mn-0.06 wt.-%C alloy in the present study.
The internal friction spectra of the Fe-17 wt.-%Mn-X wt.-%C alloys are compared in Fig. 7. The phase fractions were calculated by means of dilatometer data. In the heating segment of the IF spectrum for the Fe-17 wt.-%Mn-0 wt.-%C alloy (Fig. 7a), there are three different potential contributions to the damping: (a) the intrinsic internal friction of ɛ -martensite phase (IFint. ɛ ), (b) the anti-ferromagnetic transition, and (c) the ɛ →γ phase transformation. It is not possible to quantify the contribution of each relaxation effect because the three processes 10
overlap and this results in a broad IF peak in the 0ºC-250ºC temperature range.
The effect of the C addition is already very apparent in the IF spectrum of the Fe-17 wt.%Mn-0.01 wt.-%C alloy. There is a clear reduction of the intrinsic damping capacity of the ɛ -martensite phase during both the heating and cooling stages. The intrinsic damping capacity is gradually reduced with increasing C content and the effect saturates for a C content ≥0.03 wt.-%.
The contribution of the magnetic transition and the ɛ →γ phase transformation to the damping is clearly visible as a small internal friction peak. The overall damping capacity is reduced by C additions. The internal friction peak induced by ɛ →γ phase transformation was not identified when the C content exceeded 0.03 wt.-%C.
At temperatures above the austenite finish temperature (Af) the microstructure consists entirely of γ-austenite. The segment of the internal friction spectrum above Af therefore corresponds to the intrinsic internal friction of γ-austenite (IFint. γ). The IF measurement results clearly show that IFint. γ is not influenced by C. The IF spectrum measured in the cooling stage (Fig. 7b) also shows that IFint. γ is independent of the C content, as there is only an increase of damping capacity at the γ→ɛ martensite transformation start temperature.
11
4. Discussion The damping capacity of metallic materials is usually to be reversely proportional to the tensile strength, i.e. high strength metallic materials have a low damping capacity [26]. HDMM such as Fe-Mn based HDMMs are unusual in that they combine high strength properties and a high damping capacity [1, 2]. As C is known to be a very effective strengthening addition in most ferrous alloys, the addition of C to Fe-Mn based HDMMs is an obvious approach to further increase the strength of Fe-Mn based HDMMs [27]. The experimental results reviewed in the previous section however show that the addition of C has a detrimental effect on the damping properties. The solute C atoms must therefore prevent the activation of the relaxation processes which are responsible for damping in Fe-Mn based HDMMs.
In the present work, Fe-17 wt.-%Mn-X wt.-%C alloys with six different C contents in the range of 0.00-0.06 wt.-% were subjected to identical thermos-mechanical processes and cyclic thermal treatments, which resulted in similar C-content independent microstructures. The average grain size and the phase fractions were both similar in the six alloys, as shown in Fig. 3 and Table 1. The C content was therefore the only parameter affecting the mechanical and damping properties of the material.
The following is known about solute C atoms in γ-austenite Fe-Mn alloys: 1. C increases the stability of the austenite relative to -martensite. 2. C promotes paramagnetism, i.e. C decreases the Néel temperature of γ-austenite. 3. C increases the stacking fault energy substantially. Abbasi et al. [28] reported a significant increase of the stacking fault energy with increasing C content. They also reported a 12
strong dependence of the stacking fault energy on the position of the C atoms relative to the stacking fault. Experimental evidence for the nano-diffusion of C away from the stacking fault has been observed by Hickel et al. [29]. 4. C has a strong solid solution hardening effect in γ-austenite alloys. The solid solution strengthening effect of C, i.e. approximately +279 MPa/wt.-%C, in γ-austenite Fe-Mn-C TWIP steel [30] is similar to the solid solution strengthening effect of solute C in stainless steel. Allain reported a similar solid solution hardening effect of +250 MPa/wt.-%C [31]. 5. C occupies octahedral interstitial sites in -Fe and alloys. The small lattice distortion caused by C is isotropic. The nearest neighbor (NN) and next nearest neighbor (NNN) interactions between C-C pairs in austenite have been found to be repulsive in -Fe-C alloys [32, 33]. This result implies that the C distribution is expected to be homogeneous. This reduces the probability that an interstitial octahedral site is occupied in the vicinity of an octahedral site already occupied by a C atom, and leads to an increased C diffusivity with increasing C content [34, 35]. First principle analysis of austenitic alloys has shown that there is a pronounced repulsion between solute C atoms [28, 36-38]. 6. The strong attractive NN C-Mn (+26kJ/mol) and NNN C-Mn (+13kJ/mol) interaction between Mn and C atoms in γ-austenite Fe-Mn-C alloys result in the formation of Mn-C pairs [36, 37]. 7. Shun et al. [39], who studied the temperature dependence of the properties of austenitic Fe-30 wt.-%Mn-1.0 wt.-%C, reported that the apparent activation energy for the onset of serration in the flow curve was much lower than the activation energy for C bulk diffusion. This suggests that the serrations are related to the short range diffusion of C in the dislocation core. In the present work serrations were observed for the Fe-17 wt.%Mn-X wt.-%C alloys with a C content ≥0.03 wt.-%. The C-Mn complexes are expected 13
to be formed in a well-annealed, undeformed high Mn alloys due to the C-Mn attractive force. Lee et al. [40] proposed that DSA in high Mn γ-austenite Fe-Mn-Al-C TWIP steel results from the interaction of C-Mn complexes with staking faults by a single diffusive jump, rather than C diffusion to dislocation core.
Much less is known about the effects of solute C atoms in -martensitic Fe-Mn alloys, but if the high -martensite phase fraction is considered in the present Fe-Mn alloys, it is clear that the room temperature effects observed when the C content is increased are due to the effect of solute C on the mechanical properties and the damping of the -martensite. Lee et al. [17] have reported that C additions significantly deteriorated the damping capacity of Fe-Mn alloys. They mentioned that a 0.06 wt.-%C addition decreased the damping capacity of asquenched Fe-Mn alloys by approximately 1 %. In the present study, the observed decrease in damping capacity at room temperature was 64 % for a 0.01 wt.-%C addition and 80 % for a 0.06 wt.-%C addition. The peak damping at 150°C decreases by 76 % for a 0.01 wt.-%C addition, and by 93 % for a 0.06 wt.-%C addition.
Similarly to the case of γ-austenite, solute C or C-substitutional atom point defect complexes can affect the motion of mobile lattice defects in -martensite, i.e. dislocations and interface boundaries, responsible for the damping. The damping will be reduced when these relaxation processes are inhibited. It is very likely that the cyclic thermal treatment allows for both (a) C atoms to diffuse to mobile interface boundary segregation sites, and (b) C atoms to form stable C-Mn complexes in the matrix. In both cases, the mobility of the interfaces could be impeded and the damping capacity reduced. 14
It is well known that IFint. γ is low as compared to IFint. ɛ [3]. This has also been confirmed in the present experimental result for alloys with a low C content. With increasing C content IFint. ɛ decreased and its value became similar or even lower than IFint. γ. IFint. γ was not influenced by the C content. The cooling segment of the IF spectrum of the Fe-17 wt.-%Mn0.06 wt.-%C alloy (Fig. 7b) shows a steady decrease in damping capacity with decreasing temperature even after the start of the martensitic transformation.
Grain boundary related damping has been reported to be the major source of IFint. γ in ultrafine grained Fe-17 wt.-%Mn [19], nano-crystalline copper and stainless steel [41]. Grain boundary sliding and the disordering and reordering of atom groups in grain boundaries by thermal excitation at high temperature have been proposed as the damping mechanisms [3]. In the present study the grain size contribution to damping was not affected by the C content as care was taken to study alloys with a very similar grain size.
An alternative origin for the damping in Fe-17 wt.-%Mn-X wt.-%C alloys could be the ↔ phase transformation. The damping Q -1 associated with a solid-state phase transformation can be described by the De Jonghe-Delorme equation [42], 3 m T 1 4 C m Q k 0 1 0 T t f 3 -1
here k ,
(1)
m T m , , 0 , , and C are a material constant, the quantity of T t
transformed material per unit temperature change, the heating or cooling rate, the stress15
amplitude during the internal friction experiment, the quantity of transformed material per unit stress, and the critical stress required for the stress-induced transformation or the defect reorientation, respectively. The explanation for the difference in the amplitude of the internal friction peaks during heating and cooling is given by the De Jonghe-Delorme equation. The heating rate used for internal friction measurement was +5 °C/min, whereas the specimen was slowly cooled naturally without the use of a cooling gas. The transformation rate per a unit temperature change is therefore lower during the cooling (Fig. 7). As a consequence, the damping is reduced due to the smaller value of the factors
m T and during cooling. T t
The De Jonghe-Delorme equation also explains two observations made for alloys with a higher C content: (a) the absence of an IF peak related to phase transformation at high C contents, and (b) the fact that IFint. γ and IFint. ɛ become similar at high C contents.
As discussed above, C atoms can diffuse to preferential segregation sites e.g. mobile partial dislocations associated with stacking faults, mobile γ/ɛ phase boundaries and other mobile interfaces, immobilizing them. The
m term in the De Jonghe-Delorme equation is related
to the mobility of boundaries and interfaces [42]. The applied stress range, at which the damping was measured, was constant in the present experiments. If solute C atoms reduce the mobility of both interfaces in ɛ -martensite interfaces and the γ/ɛ phase boundaries interfaces, IFint. ɛ should decrease and no phase transformation-related peak should be observed. This appears to be the case in the alloys with C ≥0.03 wt.-%C. When the C content is such that only the C segregation sites at interfaces in ɛ -martensite are occupied, IFint. ɛ should be 16
reduced, but the ↔phase transformation-related internal friction peak should still be present. This is the case for the alloys with a C content of 0.01 wt.-% and 0.02 wt.-%.
If the γ/ɛ -phase boundaries were preferred segregation sites for solute C atoms, no peak would be obtained in the temperature interval of the phase transformation, and there should be a clear difference between IFint. γ and IFint. ɛ , which should not be affected by the C content. This was however not observed in the present case. The experimental results therefore suggest that ɛ /ɛ -martensite interfaces were the preferred segregation sites for the solute C atoms.
In the absence of C additions, i.e. for the Fe-17 wt.-%Mn alloy, the motion of interfaces in ɛ martensite and ↔phase boundaries were not impeded. This resulted in a higher value of IFint. ɛ as compared to IFint. γ and a very pronounced ↔phase transformation peak.
Comparison of the results of the present study with a previous study by Shin et al. [19], offers the opportunity to analyze the effect of cold-rolling on the damping capacity of Fe-17 wt.%Mn. Shin et al. [19] analyzed the effect of grain size on the damping spectrum of an Fe-17 wt.-%Mn alloy. Their Fe-17 wt.-%Mn was cold-rolled, recrystallization annealed at 1000°C for 30 minutes, and trained four times using a thermal cycle involving reheating to 400°C for 10 minutes and cooling to room temperature. The reported phase fractions were similar to those in the present work. The average grain size was larger, i.e. 9.6 µm for ɛ -martensite and 2.1 µm for γ-austenite.
The internal friction spectra published by Shin et al. [19] showed a clear damping peak to the 17
anti-ferromagnetic transition and the ↔ phase transformation. The 150°C damping peak size for the hot rolled and trained Fe-17 wt.-%Mn alloy used in the present study was three times larger than the damping peak observed at the same temperature for their cold-rolled Fe17 wt.-%Mn alloy.
In the present work the maximum damping peak was located at a temperature between the modulus anomaly temperature and the phase transformation temperature, clearly indicating that the motion of interfaces in ɛ -martensite was enhanced in the hot-rolled microstructure as compared to the cold-rolled and recrystallized microstructure. This results in a broad IF peak by the combined effect of the motion of interfaces in ɛ -martensite, ↔phase boundaries and the anti-ferromagnetic transition in the hot-rolled material. The difference between the IF spectra of the same alloy in the hot-rolled and the cold-rolled state is very likely due to an unfavorable crystallographic texture developed after cold rolling and recrystallization annealing. This effect is currently being analyzed in more detail for reporting at a later date.
Fig. 8 summarizes the effect of the C content on damping capacity and mechanical properties of the Fe-17 wt.-%Mn-X wt.-%C (0
18
5. Conclusions The effect of C on damping capacity in thermally trained Fe-17 wt.-%Mn-X wt.-%C (0
19
Acknowledgements The authors sincerely acknowledge the support of Dr Tae Jin Song of the POSCO Technical Research Laboratories in Gwangyang, Republic of Korea.
20
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Fig. 1 – Dilatometer measurements used to determine the number of thermal cycles required to obtain a microstructure in Fe-17 wt.-%Mn-X wt.-%C high-damping alloys which reversibly transforms between a stable high temperature microstructure and a stable low temperature microstructure. The stable dilatation loop was obtained after the 5th training 24
cycle.
25
Fig. 2 – C content dependence of the austenite start temperature (As) and the martensite start temperature (Ms).
26
Fig. 3 – Combined EBSD phase map and image quality (IQ) map of the thermally trained Fe17 wt.-%Mn-X wt.-%C high-damping alloys. The average grain size of ɛ -martensite and γaustenite is 1.31 µm and 0.62 µm, respectively. The average ɛ -martensite phase fraction is 0.91.
27
28
Fig. 4 – TEM micrographs of the trained Fe-17 wt.-%Mn alloy. (a) Bright field image and the corresponding dark field images of (b) ɛ -martensite, (c) γ-austenite, and (d) twinned γaustenite (γ'). (e) Bright field image of a stacking fault network. (f) High resolution micrograph of a ɛ -martensite plate. Selected area diffraction pattern of (g) ɛ -martensite, (h) overlapping layers of ɛ -martensite and γ-austenite, and (i) overlapping layers of ɛ martensite, γ-austenite and twinned γ-austenite (γ'). The diffraction spots used for dark field images (b), (c), and (d) are indicated by colored arrows in (i).
29
Fig. 5 – (a) Engineering stress-strain curves for the Fe-17 wt.-%Mn-X wt.-%C alloys. Serrations occur in the engineering stress-strain curves of the alloys with a C content ≥0.03 wt.-%C. (b) Solute C strengthening effect for the yield stress and the ultimate tensile strength of Fe-17 wt.-%Mn-X wt.-%C alloys. 30
Fig. 6 – Temperature dependence of the dilatation, Young’s modulus, and damping capacity of (a) Fe-17wt%Mn-0.0wt%C and (b) Fe-17wt%Mn-0.06wt%C alloys.
31
32
Fig. 7 – Comparison of internal friction spectra measured during (a) heating and (b) cooling. The modulus change and phase transformation start temperature range are indicated. The phase fraction evolution during the phase transformation of the Fe-17wt%Mn-0.0wt%C alloy is also indicated.
33
Fig. 8 – The effect of the C content on the mechanical properties of Fe-17 wt.-%Mn-X wt.%C alloys and the intrinsic internal friction of ɛ -martensite (IFint. ɛ ) and γ-austenite (IFint. ).
34
Table 1 –C content dependence of the phase fraction and grain size of ɛ -martensite and γ after thermal training.
ɛ
γ
Sample
Phase fraction
Grain size (µm)
Phase fraction
Grain size (µm)
0.0C
0.95
1.45
0.05
0.41
0.01C
0.89
1.17
0.11
0.58
0.02C
0.91
1.44
0.09
0.73
0.03C
0.88
1.33
0.12
0.70
0.05C
0.89
1.25
0.11
0.66
0.06C
0.91
1.2
0.09
0.64
35