nach oben

Erschienen in:

Open Access 2024 | OriginalPaper | Buchkapitel

27. Effect of Non-metallic Inclusions on the Temperature and Strain-Rate-Dependent Strength, Deformation and Toughness Behavior of High-Strength Quenched and Tempered Steel

verfasst von : Kevin Koch, Sebastian Henschel, Lutz Krüger

Erschienen in: Multifunctional Ceramic Filter Systems for Metal Melt Filtration

Verlag: Springer International Publishing

Aktivieren Sie unsere intelligente Suche, um passende Fachinhalte oder Patente zu finden.

search-config

KI-gestützte Suche

Patentsuche

Aus

Abstract

This chapter presents results of investigations on the strength, deformation and toughness behavior of quenched and tempered 42CrMo4 steel. Intentional impurification and, afterwards, filtration by functionalized ceramic foam filters were applied in order to process cast steels with different amounts and distributions of non-metallic inclusions. As references, a hot-rolled steel batch and spark-plasma sintered materials were studied. The investigations focused on the loading rate and temperature effects. Both, tensile and fracture mechanics tests, were performed in order to investigate the damaging behavior due to non-metallic inclusions remaining after the melt processing of the steel. A further goal was to predict the fracture toughness of the material based on the combination of microstructural information on the inclusion distribution and the strain rate and temperature-dependent strength and deformation behavior. It was shown that the damaging effect of non-metallic inclusions, in particular agglomerated inclusions properties, is localized which leads to relatively low strain to fracture and fracture toughness, but also to crack path deflection. Furthermore, it could be observed that the small interparticle distances within agglomerated non-metallic inclusions determine the fracture toughness behavior of the materials. By analyzing the acoustic emissions, the onset of crack growth as well as the size of the plastic zone at the crack tip could be estimated.

27.1 Introduction

Non-metallic inclusions occur in various steps during steelmaking, including desulfurization and deoxidation [1]. However, the reaction products can only be transferred partially into the slag. The aim of the Collaborative Research Center (CRC) 920 “Multifunctional Filters for Molten Metal Filtration—A Contribution to Zero Defect Materials” consists of removing the remaining particles and dissolved components by providing active and reactive surfaces to trap impurities and thus to increase the purity of molten metals [2]. Since a complete removal of these impurities is not possible, the remaining impurities in the casting component occur as non-metallic inclusions [3].

Non-metallic inclusions represent stress concentration points in a mechanically loaded component. Under these conditions, the risk of cleavage fracture increases. This can be critical even in ductile materials. However, most ductile materials tend to show ductile fracture which consists of void nucleation, void growth and void coalescence [4]. With more stress concentration points, these stages of ductile fracture happen earlier, resulting in lower strength, deformability and toughness of the material. As a result, the functional properties, especially those of safety components are restricted. This includes safety against brittle fracture at low temperatures and crack resistance to stable crack propagation. In addition, failure occurs due to inhomogeneous distribution of non-metallic inclusions. This leads to a scatter of the mechanical properties deformability and toughness that complicates a mathematical analysis [5‐7].

During operation, components are often exposed to changing temperatures. High temperatures can increase the deformability and ductility, while low temperatures favor brittle material behavior [8, 9]. Additionally, sudden force initiation is particularly critical for safety components, since brittle fracture occurs particularly in hazardous situations (e.g. vehicle crash). Because such components are subject to previous operational stress, crack-like imperfections and their crack resistance behavior under sudden stress must also be known [10]. When considering the toughness behavior of high-strength steels, both temperature and loading rate effects have to be taken into account [11].

Various methods already exist for determining the resistance to crack initiation at high and very high stress rates, e.g. [12, 13]. These were only applied for materials with elastic–plastic material behavior. In addition to already established methods, a relatively new method for fracture mechanics testing under high stress rates is presented. This method uses the principle of the split Hopkinson pressure bar (Kolsky Bar) for detailed analysis of forces and displacements during sudden loading. Other non-contact methods (laser interferometry, high-speed photography) were applied to support this analysis.

Material damage can be examined by using various in situ methods, e. g. optical microscopy, thermography and acoustic emissions [14]. The latter was applied for investigations under high loading rates to determine the onset of crack initiation and to characterize deformation processes.

The experimental determination of the material behavior as a function of the inclusion characteristics leads to a deeper understanding of the influence of different non-metallic inclusions, especially manganese sulfides and aluminum oxides. The relationship between the inclusion characteristics and the temperature and loading-rate-dependent strength, deformation and toughness behavior are discussed.

27.2 Materials and Methods

Materials Quenched and tempered 42CrMo4 steel is used for various applications, e.g. in automotive engineering. Due to its favorable strength and toughness properties, it is suited for drive and transmission components such as crankshafts, gears and piston rods.

The investigated 42CrMo4 steel was processed by sand casting. An alumina-coated ceramic foam filter in the gating system was used to intentionally contaminate steel melt with non-metallic inclusions. A second filter was then used to purify the steel melt. Table 27.1 gives an overview of the filters used for contamination and purification. Two batches of filtered steel castings were investigated. In both cases, carbon-bonded alumina filters were used for purification. An uncoated filter was used for the processing of batch C1. For C2, a filter with mullite coating was used. More detailed information on the process is given by Emmel et al. [15].

Table 27.1

Characteristics of the multifunctional filters

Batch	Contamination	Purification
C1	Al₂O₃	Al₂O₃-C
C2	Al₂O₃	Al₂O₃-C + Al₂O₃·SiO₂

For comparison, two other batches of 42CrMo4 were used as a reference. Batch R represents a hot-rolled material. RS was processed by continuous casting and shows a higher sulfur content compared to the other batches. Details on the chemical composition of the investigated steels are given in Table 27.2. In order to remove shrinkage cavities, the material underwent hot isostatic pressing [16]. The machined samples were austenitized at 840 °C for 20 min in a vacuum atmosphere, quenched in a He stream and tempered at 560 °C for 60 min in a N₂ atmosphere.

Table 27.2

Chemical composition of investigated 42CrMo4 steels

Batch	C	Si	Al	S	P
R	0.44	0.04	0.04	0.006	0.012
RS	0.41	0.25	0.02	0.031	0.012
C1	0.41	0.53	0.09	0.007	0.011
C2	0.43	0.54	0.09	0.007	0.011

After heat treatment, the microstructure of the steels consists of martensite, see Fig. 27.1. Non-metallic inclusions of Al₂O₃ as well as MnS were observed. The characteristics of the non-metallic inclusions are shown in Table 27.3.

Table 27.3

Characteristics of non-metallic inclusions

Batch	Inclusion size [µm]	Inclusions/area [mm⁻²]	Inclusion content [10^–3 Vol.%]	Mean distance [µm]
R	4.24	5.1	8.0	125
RS	2.81	62.2	42	53
C1	2.00	13.1	5.0	88
C2	1.98	13.5	5.5	87

Another three batches of material were processed by field assisted sintering technology (FAST). For this purpose, 42CrMo4 steel powder was used. Different amounts of Al₂O₃ powder were added to simulate non-metallic inclusions. Further information on the process is given in Koch et al. [17]. Table 27.4 shows the chemical composition of the three different steel batches.

Table 27.4

Chemical composition of investigated sintered 42CrMo4 steel

Batch	C	Si	Al	S	P
S0	0.36	0.21	0.018	0.011	0.015
S1	0.37	0.23	0.460	0.011	0.014
S2	0.33	0.21	3.540	0.012	0.010

The microstructures of the steels processed by FAST are shown in Fig. 27.2. It can be observed that the added alumina particles are located along the boundaries of the steel particles. With an increasing amount of Al₂O₃, the distance between the non-metallic particles decrease. Within the areas which are occupied by the steel particles, no alumina particles were observed, see Fig. 27.2c.

After machining, the samples were austenitized at 840 °C for 20 min in a vacuum atmosphere, quenched at 15 bars in a He stream and tempered at 450 °C for 60 min in an N₂ atmosphere.

Tensile Testing The strength and deformation behavior was examined in tensile tests. For this purpose, quasi-static and dynamic strain rates as well as different temperatures were used. Table 27.5 gives an overview of the test program. Quasi-static tensile tests at strain rates of 10⁰ s⁻¹ and less were performed with a universal testing machine Zwick 1476.

Table 27.5

Investigated test temperatures and strain rates for tensile testing, (X): only selected materials

Strain rate [1/s]	Test temperature [°C]
Strain rate [1/s]	−60	−40	−20	20
0.4–1·10^–3	(X)	X	X	X
1·10⁰		X		X
1·10¹				X
1·10²				X
1·10³		(X)		(X)

For strain rates at approx. 10^–3 s⁻¹, the specimen geometry (ISO 6892 [18]) shown in Fig. 27.3a was used. In this case, the elongation of the specimen was determined by a clip-on extensometer. For tests at strain rates of 10⁰ s⁻¹ (ISO 26203–2 [19]) and more, specially-shaped specimens were used, see Fig. 27.3b. The force as well as the elongation in these tests were measured with strain gauges on the specimen. Additionally, tests at strain rates of 10² s⁻¹ and 10³ s⁻¹ were performed in a drop weight testing machine and a rotating wheel test machine, respectively.

Fracture Toughness Testing The aim of the fracture toughness tests was to determine the relationship between the crack resistance J and the stable crack propagation Δa. The material resistance to crack initiation is derived from the crack resistance curve. It can also be derived from the crack tip blunting (stretch zone width, SZW). The fracture mechanics tests were carried out over a wide range of loading rates and temperatures, see Table 27.6.

Table 27.6

Investigated test temperatures and loading rates for fracture toughness testing, (X): only selected materials

Loading rate [MPa $\sqrt{\text{m}}$/s]	Test temperature [°C]
Loading rate [MPa $\sqrt{\text{m}}$/s]	−60	−40	−20	20
10⁰	(X)	X	(X)	X
10⁵	(X)	X		X
10⁶		(X)		(X)

Quasi-static tests were carried out with a servo-hydraulic universal testing machine MTS 810 to determine the crack resistance curve. Single-edge notch bending (SENB) specimens (B = 10 mm) according to ASTM E 1820 [21] or ISO 12135 [22] were used, see Fig. 27.4. The length L of the specimen was 120 mm. Hence, two smaller Charpy-type bending specimens could be made from the broken samples.

The stress intensity factor K_I and the current crack length (a = a₀ + Δa) was determined from the elastic compliance according to ISO 12135. The J-integral is calculated by

$$J=\frac{{K}^{2}\left(1-{v}^{2}\right)}{E}+\frac{1.9{U}_{{\text{pl}}}}{{B}_{{\text{N}}}\left(W-{a}_{0}\right)}\cdot \left(1-\frac{\Delta a}{2\left(W-{a}_{0}\right)}\right)$$

(27.1)

Fracture toughness tests at loading rates of about ${10}^{5}\, {\text{MPa}}\sqrt{\text{m}}{{\text{s}}}^{-1}$ were performed in an instrumented pendulum impact test machine (PSd 300, WPM Leipzig). The test setup is shown in Fig. 27.5. The force was measured by an instrumented tup. The displacement of the specimen was measured in force direction by a laser doppler interferometer (Polytec OFV-525). Further information on this technique is given in Henschel et al. [23].

The dynamic fracture toughness tests have been performed with the low-blow method to create different amounts of stable crack growth. The impact energy was varied by setting the deflection angle of the pendulum prior to each test. Investigations were carried out at impact energies E₀ = 2.3 – 3.5 J which correspond to impact velocities v₀ = 0.48 – 0.59 m/s. The J-Integral was calculated from the force–displacement curve similar to quasi-static testing:

$${J}_{{\text{d}}}=\frac{{K}^{2}\left(1-{v}^{2}\right)}{E}+\frac{2{W}_{{\text{pl}}}}{{B}_{{\text{N}}}\left(W-{a}_{0}\right)}\cdot \left(1-\frac{\Delta a}{2\left(W-{a}_{0}\right)}\right)$$

(27.2)

With

$${W}_{{\text{pl}}}={\int }_{0}^{s}F{\text{d}}s\mathrm{^{\prime}}-\frac{{F}^{2}{C}_{0}}{2}$$

(27.3)

Here, W_pl is the plastic portion of deformation energy and C₀ is the initial compliance.

Loading rates of about ${10}^{6}\, {\text{MPa}}\sqrt{\text{m}}{{\text{s}}}^{-1}$ were achieved by applying the split Hopkinson pressure bar (Kolsky bar) technique. The setup shown in Fig. 27.6 consists of four different bars that are axially aligned.

The striker bar is propelled by compressed air and hits the incident bar and introduces a nearly rectangular pressure pulse of known amplitude ${\varepsilon }_{\text{I,max}}$:

$${\varepsilon }_{\text{I,max}}=\frac{{v}_{\text{St}}}{2{c}_{\text{B}}}$$

(27.4)

Here, ${v}_{\text{St}}$ and ${c}_{\text{B}}$ are the striker bar velocity and the bulk sound velocity of the bars, respectively. Equation (27.4) is only valid, if striker and incident bars consist of the same material and have the same diameter.

This pressure pulse propagates along the incident bar and is partially transmitted into the specimen (and the transmitted bar) and partially reflected as a tensile pulse. The pulse in the transmitted bar propagates further in the momentum trap that eventually separates from the transmitted bar, which leads to a single loading of the specimen. The elastic deformation of the bars that are made of high-strength aluminum (AA7075) is measured by strain gauges at the center of the bars. At the incident bar, a laser interferometer was used to measure the particle velocity. Since particle velocity and strain in the bar are proportional, an additional measurement site was established. This measurement site was applied if relatively long incident pulses were used. Details are found in Henschel et al. [24]. The signal acquisition (20 MSample/s) was triggered by two light barriers near the impact site of the striker bar. Furthermore, the light barriers were also used to calculate ${v}_{\text{St}}$.

From the incident, reflected and transmitted pulses (${\varepsilon }_{\text{I}}, {\varepsilon }_{\text{R}}$ and ${\varepsilon }_{\text{T}}$, respectively), the forces and displacements at the interfaces 1 and 2 were calculated. For the sake of brevity, the equations are not shown here, but can be found in Henschel [20].

The analysis of the force equilibrium is a crucial step. If the axial forces ${F}_{1}$ and ${F}_{2}$ at the interfaces 1 and 2 are approximately equal, inertial effects can be neglected. Consequently, there is a proportionality between the acting force $F={F}_{1}={F}_{2}$ and the stress intensity factor ${K}_{\text{I}}$ at the crack tip. Hence, equations for quasi-static calculation of $K$ and $J$ [24] can also be applied here. On the other hand, if there are significant inertia effects, the forces ${F}_{1}$ and ${F}_{2}$ are treated separately. In order to calculate the time-dependent stress-intensity factors, two-point loading from each side can be assumed. This method is described in detail in Henschel [20, 25]. The pressure pulse can be shaped by using small pieces of deformable material at the impact site, see Fig. 27.6. Consequently, this affects the rise time of the incident pulse.

Acoustic Emission Analysis The relation between different mechanisms of material damage and the acoustic emission (AE) signal was studied at different temperatures and loading rates. The main focus was to investigate ductile fracture at ambient temperature. However, tests at −40 °C were aimed to examine the transition from ductile to brittle behavior. To this end, quasi-static tensile tests as well as quasi-static and dynamic fracture toughness tests were carried out with one or more acoustic transducers mounted to the specimen to record acoustic emissions. Details on the test setup are given in Kietov et al. [26].

The AE signal was recorded continuously in order to identify possible events of material damage. In order to analyze the signal, first transient signals were identified by the threshold technique. Secondly, the continuous low-amplitude signal was analyzed with respect to signal energy and power. Usually, transient signals are associated with crack growth and fracture, while continuous signals correspond to plastic deformation or dislocation motion, see Fig. 27.7. Further information on the test method and evaluation is given in Kietov et al. [26, 27].

Mixed-Mode Testing In order to study the fracture toughness under mixed-mode loading, compact tension shear specimens (CTS) were used. Details can be found in Henschel et al. [28]. The specimen dimensions were based on recommendations of Richard [29] and scaled for a thickness of 8 mm. The heat treatment consisted of austenitization (840 °C, 20 min, vacuum), quenching in a He stream at 15 bars and tempering (450 °C, 1 h, N₂). The stress intensity factors K_I and K_II were determined by finite-element analysis. Equations used for the calculation of K_I and K_II are given in Henschel et al. [28]. The calculated stress intensity factors were compared to experimentally measured stress intensity factors. To this end, strain gauges were applied to analyze the stress field around the crack tip, according to Sarangi et al. [30].

The aim of the investigations is to determine the mixed-mode fracture toughness for a high-strength steel as well as the effect of the loading angle on the crack growth.

27.3 Temperature and Strain-Rate Dependent Material Behavior of High-Strength Steels

27.3.1 Strength and Deformability

The strength and deformation behavior was investigated in uniaxial tensile tests for various temperatures and strain rates. Figure 27.8 shows the effect of the temperature for the steels R and RS. The yield strength ${R}_{{\text{eL}}}$ increases at low temperatures. This can be explained by thermally activated dislocation motion. The plastic elongation at fracture A also increases with decreasing temperature. This is primarily due to an increasing uniform strain A_g, while necking is only slightly affected by temperature. Therefore, the ratio A_g/A increases with decreasing temperature. The reduction of area Z decreases slightly with decreasing temperature which indicates an embrittlement. The higher tensile strength ${R}_{{\text{m}}}$ of R compared to RS can be explained by the higher content of C. However, both steels show a similar strain hardening behavior.

Figure 27.9 shows the results of the tested steel castings. The strength values of C1 and C2 showed no significant differences. However, slight differences in tensile strength and yield strength can be explained by differences in the chemical composition.

The deformability shows significant scatter. However, the impurification with alumina particles only slightly affects the properties which were investigated in tensile tests. Compared to batches R and RS, the steel castings also show an increasing uniform strain at low temperatures. The elongation at fracture is unaffected by temperature, which can be seen for batch C1. On the other hand, C2 showed decreasing elongation at fracture at low temperatures.

The reduction of area decreased at very low temperatures. The embrittlement at low temperatures is affected by non-metallic inclusions. Figure 27.10 shows the fracture surfaces of the steels C1 and C2. The agglomerate in C2 is identified as the main reason for the relatively low deformability, see the fracture path lines in Fig. 27.10b. Due to low temperature and the notch effect of the agglomerate, a mixture of ductile fracture and cleavage fracture is observed.

Figure 27.11 shows the relationship between the size of the largest damage-relevant agglomerate and the elongation at fracture. It can be seen that an increasing agglomerate size reduces the elongation at fracture. Furthermore, the amount of agglomerates and their arrangement in the specimen volume have an effect. Figure 27.11b shows a fracture surface with several small agglomerates.

Due to small spatial distance, a relatively low strain during necking was observed (mark II). On the other hand, another sample shows a significantly greater elongation during necking, although the inclusion agglomerate is almost twice as large in diameter (mark I and Fig. 27.10b and c). However, there is no other agglomerate in the weakest cross-section. An effect of the temperature was not observed.

The effect of the strain rate on the strength of R and RS is shown in Fig. 27.12a. Analogously to a decreasing temperature, an increasing strain rate causes an increase of yield strength and tensile strength. Due to adiabatic heating, the increasing strain rate between 10² and 10³ s⁻¹ did not affect the tensile strength. This leads to softening of the material which is superposed by the strain hardening. Figure 27.12b shows that the deformation parameters of steel R are not depending on the strain rate. In contrast, the reduction of area of the steel RS shows a strong decrease up to strain rates of 10² s⁻¹. The increase of Z by a further increase of strain rate is associated with the increase in temperature and correlates with an increasing ductility.

The effect of the strain rate on the strength and deformability parameters are shown in Fig. 27.13. The steel castings show a similar behavior compared to the steels R and RS, as a positive strain rate effect occurs. The effect of the strain rate on yield and tensile strength is more significant at T = −40 °C than at 20 °C. However, the steel C2 is an exception to this. The deformability does not show a distinct dependence. At T = 20 °C, the strain at failure and the reduction of area decrease, in particular for strain rates greater than 10² s⁻¹. Additionally, the elongation at fracture is not temperature dependent at T = −40 °C. Furthermore, the behavior of the uniform elongation correlates with embrittlement at the highest strain rate.

In Henschel et al. [32], C1 and C2 were compared regarding the strength and deformation behavior. It was found that C2 has a higher ductility at quasi-static loading in the temperature range from −40 to 20 °C. Taking the deformability parameters of the other steel castings into account as well as the strain rate dependency, it can be assumed that structural differences within the cast plates are the most probable cause for apparent differences in mechanical behavior. In previous studies, it was additionally found, that the crucible material [33] and the usage of an additional immersion filter [34] may have a significant effect on the spatial inclusion distribution and, consequently, on the mechanical properties.

The results of the tensile tests of the sintered materials are shown in Fig. 27.14. The tensile strength and strain at failure decrease with an increasing content of alumina particles. It should be mentioned that the sintered material showed more brittle behavior in general compared to the previously mentioned 42CrMo4 batches. Batch S3 shows no plastic deformation, neither at quasi-static loading nor at dynamic loading. Hence, failure occurs before the yield strength is reached. Microscopic examinations underneath the fracture surfaces showed secondary cracks along the former steel particle boundaries, see Fig. 27.14. In this area, the ceramic particles are arranged with small distances to each other. Due to the increased notch effect, crack deflection was favored. The batch S0 already shows low strain at failure without the addition of aluminum oxide, compared to the 42CrMo4 batches mentioned before. This can be attributed to oxide layers on the steel particles. These layers are formed e.g. during powder processing. This leads to an internal notch effect in the sintered material, which results in a reduced deformability. This effect is intensified by the addition of the ceramic particles.

In order to examine the material damage by means of acoustic emissions (AE), the voltage signal of a piezo-electric transducer mounted to the specimen was recorded continuously during tensile tests. Only steel R was investigated for this purpose. For the analysis, damage relevant signals had first to be discriminated from the electrical and mechanical background noise that was also recorded. Transient signals were identified with a threshold-based technique. However, continuous signals show very low amplitudes and are more difficult to distinguish from background noise. Therefore, the signal was analyzed with respect to the frequency of the AE signals with a statistical approach. This initial test was performed under quasi-static loading, without any plastic deformation. Detailed information is given in Kietov et al. [27]. The results are shown in Fig. 27.15. It could be observed that transient signals occur mostly in a median frequency range between 250 and 600 kHz. However, continuous signals showed median frequencies between around 180 to 240 kHz. Both signals could be very well discriminated from the background noise which showed lower median frequencies.

The actual tensile tests with AE recording were carried out up to failure as mentioned earlier in this chapter. Figure 27.16 shows the stress–strain diagram and the corresponding AE signal for steel R. It can be observed from the rapid increase of the transient AE rate that the majority of transient signals occurred during elastic deformation.

At the beginning of the yield plateau (Lüders strain), the transient AE activity started to decrease. Shortly before fracture, another rapid increase of transient signals was observed. The continuous AE signals showed a different behavior. Until the beginning of yielding, the level of the continuous AE remains below the noise. During the yielding plateau a significant increase of the power of the continuous AE was observed. This was attributed to the high amount of dislocation motion which occurs at the first deformation of the material. It only happened at the yield plateau and no other further increase happened after the continuous AE dropped back to noise level. Effects that correspond to damage in context of non-metallic inclusions were not observed during these tensile tests.

27.3.2 Fracture Toughness

Figure 27.17 shows the J-Δa curves for the steels R, RS, C1 and C2. The difference between the steels R and RS for Δa < 0.2 mm is insignificant. However, at higher Δa, dJ/dΔa and J are significantly higher for steel R. In Fig. 27.17b one can see that the cast steels with the added non-metallic inclusions show both a lower toughness level as well as a lower slope of the J-Δa curve. Minor differences are observed between the industrially produced and contaminated steels C1 and C2.

The J_iBL and J_0.2BL values derived from the J-Δa curves are shown in Fig. 27.18. In Fig. 27.18a it can be seen that the material resistances to both physical and engineering crack initiation (J_iBL or J_0.2BL) do not differ significantly for the steels R and RS. Likewise, Fig. 27.18b shows no significant differences between the two cast steels. On the other hand, the material resistances J_iBL and J_0.2BL of the cast steels are significantly lower than the references R and RS. At T = 20 °C, an average of 30 N/mm is determined for J_iBL for the cast steels, while approx. 70–80 N/mm are measured for R and RS, respectively. The differences between the physical crack initiation (J_iBL) determined from the J-Δa curve and the crack initiation (J_iSZW) determined from the fracture surface are discussed later.

Figure 27.19 shows the effect of temperature on the J-Δa curves for the steels R and RS. The values of J_iBL, J_SZW and J_0.2BL have already been shown in Fig. 27.18. No significant relationship between crack resistance and temperature was observed. Multiple cases of pop-in behavior or completely unstable failure were observed at all temperatures. No J-Δa curve was obtained under these conditions. Nevertheless, the data points of the steel R for T = −60 °C are shown in Fig. 27.19. In addition, the characteristic values of crack initiation were estimated and shown in brackets in Fig. 27.18. In case of an unstable crack propagation after Δa > 0.2 mm, only the toughness J_u can be determined according to ISO 12135. However, this neither describes the crack initiation nor is it independent of the thickness.

Analogously to the steels R and RS, C1 and C2 show no significant temperature dependence of crack initiation, see Fig. 27.20. Large scatter was observed at low temperatures, which can be explained by the cleavage fracture surfaces and the associated higher scatter to be expected in the ductile–brittle transition region [35]. This scatter is not due to pop-ins. The fracture toughness is affected by the strength and deformability of the material. It was shown, that the strength of all investigated steels increases with decreasing temperature. An attempt to model the resistance to crack initiation at quasi-static loading rates was previously shown in Henschel and Krüger [36]. It was observed, that the distance of the particles within an agglomerate determines the toughness behavior.

The increase is about 30 MPa or 3% in the temperature range from 20 to −40 °C. On the other hand, there was no clear connection between temperature and deformability. Steels R and RS showed a slight increase in A and a slight decrease in Z with decreasing temperature (about +6% and −4%, respectively). No significant effect of temperature on the deformability parameters A and Z was observed for the cast steels. Due to small differences and scatter of the deformability, the fracture toughness is independent of the temperature.

The effect of the loading rate on fracture toughness is described and discussed below, including the temperature effect at high loading rates. The J_d−Δa curves for the impact tests in the pendulum impact test machine are shown in Figs. 27.21 and 27.22. The results for 20 °C are represented by full symbols and those for −40 °C by open symbols. The crack resistance curve was determined by the J−Δa values from the multiple specimen method (MSM) using the low-blow technique. Additionally, the normalization method (NM) according to ISO 26843 [37] enables the J−Δa values to be estimated up to the end of the respective test. This corresponds to the turning point in the low-blow test. A complete J−Δa curve is thus calculated from the force–displacement curve by using the NM.

A comparison of the measured correlations between J and Δa and those determined using the normalization method according to ISO 26843 shows that J is overestimated and Δa is underestimated, see Fig. 27.21. Lucon [38] also points to this circumstance using the example of two reactor pressure vessel steels (20MnMoNi5-5 and ASTM A533B Cl.1).

Analogously to quasi-static tests, no temperature dependency of the crack resistance curves was determined under dynamic loading. Due to the multiple specimen method and slightly different crack lengths, the J_d−Δa curves show relatively high scatter, see Fig. 27.21a. Different crack lengths can be caused by tolerances in the production of B and W.

The characteristic parameters of crack initiation (J_iBL, J_SZW and J_0.2BL) were determined based on the crack resistance curves, see Fig. 27.23. On the one hand, the crack resistance curve determined using several low-blow tests was evaluated for this purpose. On the other hand, the J−Δa curve determined using the normalization method analyzed. It was already established in Fig. 27.23 that the crack resistance curve determined using the normalization method shows relatively large values of J_d for small Δa. This circumstance also affects J_iBL and J_iSZW. The technical crack initiation parameter J_0.2BL, on the other hand, is hardly affected by the evaluation method. Furthermore, it can be seen that the difference between J_iBL and J_iSZW is smaller on average when using the normalization method than when using the multiple specimen method.

A slight temperature effect of the crack initiation parameters between −40 °C and 20 °C was observed for the steels RS, C1 and C2. With decreasing temperature, the fracture toughness increases. Occasionally, cleavage fracture was detected. The investigated temperature interval is therefore at the lower end of the upper-shelf toughness. The results of dynamic tests at very high loading rates (10⁶ MPa√ms⁻¹) are shown in Fig. 27.24.

For steel R a significant temperature effect can be observed. Regardless of the loading rate, a decreasing temperature leads to a decrease in fracture toughness of about 20 MPa√m. In the case of quasi-static loading, this correlates to a 23% reduction in toughness compared to room temperature. At very high loading rates, a significantly greater decrease in toughness of 44% is observed. The steels R and RS show a decreasing toughness with an increasing strain rate. However, the toughness loss for loading rates >10⁵ MPa√ms⁻¹ is less pronounced for steel RS.

The microscopic damage analysis shows ductile failure. Figure 27.25a shows that oxidic inclusions are detached from the metallic matrix even at low levels of stress. With a favorable spatial arrangement of the non-metallic inclusions, shear bands form between the inclusions or inclusion clusters. This phenomenon can be recognized by the steep flanks on the fracture surfaces see Fig. 27.26a.

The resistance to physical crack initiation is based on the blunting capacity of the crack tip. The blunting depends to a large extent on the distribution of the non-metallic inclusions in terms of size and position in front of the crack tip [39], see Fig. 27.26.

During quasi-static fracture toughness testing, the acoustic emissions have been analyzed for the steel R. The acoustic emissions were processed the same way as described for tensile testing. Figure 27.27 shows a load–displacement curve for a specimen tested at T = 20 °C. The number of transient AE increases almost linear over time. However, the continuous AE showed a different behavior, which is presented by a significant change of the power level.

Figure 27.28a shows the energy of the transient AE during fracture toughness testing. The emitted energy of the transient AE remains relatively low at the beginning, but becomes significantly stronger at a certain point. According to Roy [40, 41], this behavior could be attributed to large amounts of stable crack growth for this type of steels. By analyzing of the continuous AE, the size of the plastic zone could be estimated. This was done by calculating the volumetric AE energy density from the continuous AE power. The specific equations for this calculation are given in Kietov et al. [27]. Figure 27.28b shows the comparision of the plastic zone sizes estimated by the continuous AE and by the stress intensity factor. It shows good coincidence until the onset of crack growth which happened at 0.7 to 0.8 mm of displacement. After the onset of crack growth, both estimated plastic zone sizes begin to deviate, since the plastic zone size is understimated by K_I. Hence, the analysis of the AE gives a more realistic measure of the plastic zone. Furthermore, the beginning of the deviation of the plastic zone size can be used to identify the onset of crack growth. Consequently, the acoustic emission analysis shows its advantages during the study of both unstable fracture [42], and stable fracture.

The results of mixed-mode fracture toughness tests on steel RS will be discussed in the following. A total of four different loading angles was tested in order to investigate the effect on the fracture toughness and crack growth. The loading angle α = 0° resembles pure mode I loading. Hence, the ratio K_I/K_II decreases with increasing α. It could be observed that all tested loading angles showed nearly linear-elastic behavior. The maximum force until fracture increases with increasing α. Detailed information on the force measuring technique with strain gauges as well as the displacement measuring technique is given in Henschel et al. [28].

The results for K_IQ and K_IIQ at different loading angles are presented in Fig. 27.29a. The mode I fracture toughness decreased with increasing loading angle, while the mode II fracture toughness increased. This behavior can be described with the limiting curve shown in Fig. 27.29b. Within the area of the curve, the material will not fail. Flat as well as slant areas could be observed on the fracture surfaces, see Fig. 27.30. This indicates a mixture of plane strain and plane stress on the specimen. Crack tip blunting was observed at the tip of the fatigue pre-crack. The crack extension predominantly consisted of unstable ductile fracture, which was promoted by non-metallic inclusions.

27.4 Conclusions

This chapter contributes to the understanding of the effect of non-metallic inclusions on the mechanical material behavior of quenched and tempered 42CrMo4 steel. The strength, deformability and fracture toughness behavior were examined at different temperatures and loading rates under uniaxial loading. The most important findings are presented below.

All examined steels showed increasing strength with decreasing temperature and increasing strain rate. The relationship between deformability and strain rate is not clear. Due to the strain rate, adiabatic heating occurs, which is superimposed on the strain rate-related hardening. Non-metallic inclusions did not have a significant effect on the strength and the deformability.
Higher inclusion contents lead to a decreasing fracture toughness under quasi-static loading. However, the effect was not observed at very low temperatures or very high loading rates. The non-metallic inclusions lead to a deflection of the crack front. Inclusion agglomerates are particularly critical.
The material damage could be examined based on the AE analysis. Different AE events enabled different damage mechanisms to be identified. Hence, the onset of crack growth and the size of the plastic zone at the crack tip could be estimated.

Acknowledgements

The authors thank the German Research Foundation (DFG) for the financial support of the investigations in the Collaborative Research Center 920—Project-ID 169148856, subproject C05. The acoustic emission analysis of Wladimir Kietov and the experimental support of Birgit Witschel, Georg Maiberg, Sascha Graf, Lars Reichert and Florian Posselt are gratefully acknowledged.

Open Access This chapter is licensed under the terms of the Creative Commons Attribution 4.0 International License (http://creativecommons.org/licenses/by/4.0/), which permits use, sharing, adaptation, distribution and reproduction in any medium or format, as long as you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons license and indicate if changes were made.

The images or other third party material in this chapter are included in the chapter's Creative Commons license, unless indicated otherwise in a credit line to the material. If material is not included in the chapter's Creative Commons license and your intended use is not permitted by statutory regulation or exceeds the permitted use, you will need to obtain permission directly from the copyright holder.

Vorheriges Kapitel A Numerical Investigation of Heat Generation Due to Dissipation in Ultrasonic Fatigue Testing of 42CrMo4 Steel Employing Thermography Data

Nächstes Kapitel Influence of Filter Structure and Casting System on Filtration Efficiency in Aluminum Mold Casting

L. Zhang, B.G. Thomas, Metall. Mater. Trans. B 37, 733 (2006). https://doi.org/10.1007/s11663-006-0057-0CrossRef

C.G. Aneziris, U. Fischer, M. Emmel, J. Hubálková, H. Berek, Keram. Z. 64, 124 (2012)

H.V. Atkinson, G. Shi, Prog. Mater. Sci. 48(5), 457 (2003). https://doi.org/10.1016/S0079-6425(02)00014-2CrossRef

R.D. Thomson, J.W. Hancock, Int. J. Fract. 26, 99 (1984). https://doi.org/10.1007/BF01157547CrossRef

K. Wallin, P. Nevasmaa, A. Laukkanen, T. Planman, Eng. Fract. Mech. 71(16–17), 2329 (2004). https://doi.org/10.1016/j.engfracmech.2004.01.010CrossRef

W.C. Leslie, Trans. Iron Steel Soc. AIME 2, 1 (1983)

W.M. Garrison Jr., N.R. Moody, J. Phys. Chem. Solids 48(11), 1035 (1987). https://doi.org/10.1016/0022-3697(87)90118-1CrossRef

Y.-C. Lin, M.-S. Chen, J. Zhang, Mat. Sci. Eng. A 499(1–2), 88 (2009). https://doi.org/10.1016/j.msea.2007.11.119CrossRef

G. Quan, Y. Tong, G. Luo, J. Zhou, Comp. Mater. Sci. 50(1), 167 (2010). https://doi.org/10.1016/j.commatsci.2010.07.021CrossRef

10.

R. Chaouadi, J.L. Puzzolante, Int. J. Pres. Ves. Pip. 85(11), 752 (2008). https://doi.org/10.1016/j.ijpvp.2008.08.004CrossRef

11.

Y.C. Chi, S. Lee, K. Cho, J. Duffy, Mat. Sci. Eng. A 114, 105 (1989). https://doi.org/10.1016/0921-5093(89)90850-2CrossRef

12.

H.-W. Viehrig, J. Böhmert, J. Džugan, in From Charpy to Present Impact Testing. ed. by D. François, A. Pineau (Elsevier, Amsterdam, 2002), p.245CrossRef

13.

F. Jiang, K.S. Vecchio, Rev. Sci. Instrum. 78, 063903 (2007). https://doi.org/10.1063/1.2746630CrossRefPubMed

14.

A. Weidner, Deformation Processes in TRIP/TWIP Steels (2020). https://doi.org/10.1007/978-3-030-37149-4

15.

M. Emmel, C.G. Aneziris, G. Schmidt, D. Krewerth, H. Biermann, Adv. Eng. Mater. 15(12), 1188 (2013). https://doi.org/10.1002/adem.201300118CrossRef

16.

Bodycote plc, Hot Isostatic Pressing–Simple Equations for Better Materials (2014)

17.

K. Koch, V. Kietov, S. Henschel, L. Krüger, EPJ Web Conf. 250, 03001 (2021). https://doi.org/10.1051/epjconf/202125003001CrossRef

18.

International Organization for Standardization, Metallic Materials–Tensile Testing–Part 1: Method of test at room temperature (ISO 6892–1) (2019)

19.

International Organization for Standardization, Metallic Materials — Tensile Testing at High Strain Rates — Part 2: Servo-Hydraulic and Other Test Systems (ISO 26203–2) (2011)

20.

S. Henschel, Dissertation. TU Bergakademie Freiberg (2018). https://nbn-resolving.org/urn:nbn:de:bsz:105-qucosa2-331498

21.

American Society for Testing and Materials (ASTM), Standard Test Method for Measurement of Fracture Toughness. American Society for Testing and Materials, West Conshohocken, PA (ASTM E 1820) (2015)

22.

International Organization for Standardization, Metallic Materials–Unified Method of Test for the Determination of Quasistatic Fracture Toughness. International Organization for Standardization, Geneva, Switzerland (ISO 12135) (2016)

23.

S. Henschel, L. Krüger, Mater. Test. 57(10), 837 (2015). https://doi.org/10.3139/120.110785CrossRef

24.

S. Henschel, L. Krüger, Eng. Fract. Mech. 133(S1), 62 (2015). https://doi.org/10.1016/j.engfracmech.2015.05.020CrossRef

25.

S. Henschel, L. Krüger, Int. J. Fract. 201, 235 (2016). https://doi.org/10.1007/s10704-016-0127-5CrossRef

26.

V. Kietov, S. Henschel, L. Krüger, Eng. Fract. Mech. 188, 58 (2018). https://doi.org/10.1016/j.engfracmech.2017.07.009CrossRef

27.

V. Kietov, S. Henschel, L. Krüger, Eng. Fract. Mech. 210, 320 (2019). https://doi.org/10.1016/j.engfracmech.2018.06.035CrossRef

28.

S. Henschel, F. Posselt, S. Dudczig, T. Wetzig, C.G. Aneziris, L. Krüger, Proc. Struct. Int. 28, 1369 (2020). https://doi.org/10.1016/j.prostr.2020.10.108CrossRef

29.

H.A. Richard, Bruchvorhersagen bei überlagerter Normal- und Schubbeanspruchung von Rissen (VDI, Düsseldorf, 1985)

30.

H. Sarangi, K. Murthy, D. Chakraborty, Eng. Fract. Mech. 88, 63 (2012). https://doi.org/10.1016/j.engfracmech.2012.04.006CrossRef

31.

S. Henschel, L. Krüger, MATEC Web Conf. 12, 04007 (2014). https://doi.org/10.1051/matecconf/20141204007CrossRef

32.

S. Henschel, D. Krewerth, F. Ballani, A. Weidner, L. Krüger, H. Biermann, M. Emmel, C.G. Aneziris, Adv. Eng. Mater. 15(12), 1216 (2013). https://doi.org/10.1002/adem.201300125CrossRef

33.

S. Henschel, J. Gleinig, T. Lippmann, S. Dudczig, C.G. Aneziris, H. Biermann, L. Krüger, A. Weidner, Adv. Eng. Mater. 19(9), 1700199 (2017). https://doi.org/10.1002/adem.201700199CrossRef

34.

M. Seleznev, S. Henschel, E. Storti, C.G. Aneziris, L. Krüger, A. Weidner, H. Biermann, Adv. Eng. Mater. 22(2), 1900540 (2020). https://doi.org/10.1002/adem.201900540CrossRef

35.

A. Krabiell, Dissertation. RWTH Aachen (1982)

36.

S. Henschel, L. Krüger, Steel Res. Int. 87(1), 29 (2016). https://doi.org/10.1002/srin.201400567CrossRef

37.

International Organization for Standardization, Metallic Materials–Measurement of Fracture Toughness at Impact Loading Rates using Precracked Charpy-Type Test Pieces. International Organization for Standardization, Geneva, Switzerland (ISO 26843) (2015)

38.

E. Lucon, J. ASTM Int. 8(10), 103644 (2011). https://doi.org/10.1520/JAI103644CrossRef

39.

S. Henschel, L. Krüger, Fract. Struct. Integrity 9(34), 326 (2015). https://doi.org/10.3221/IGF-ESIS.34.35CrossRef

40.

H. Roy, N. Parida, S. Sivaprasad, S. Tarafder, K.K. Ray, Mat. Sci. Eng. A 486(1–2), 562 (2008). https://doi.org/10.1016/j.msea.2007.09.036CrossRef

41.

H. Roy, H.N. Bar, S. Sivaprasad, S. Tarafder, K.K. Ray, Int. J. Pres. Ves. Pip. 87(10), 543 (2010). https://doi.org/10.1016/j.ijpvp.2010.08.016CrossRef

42.

S. Henschel, V. Kietov, F. Deirmina, M. Pellizzari, L. Krüger, Mat. Sci. Eng. A 709, 152 (2018). https://doi.org/10.1016/j.msea.2017.10.053CrossRef

Titel: Effect of Non-metallic Inclusions on the Temperature and Strain-Rate-Dependent Strength, Deformation and Toughness Behavior of High-Strength Quenched and Tempered Steel
verfasst von: Kevin Koch
Sebastian Henschel
Lutz Krüger
Verlag: Springer International Publishing
Buch: Multifunctional Ceramic Filter Systems for Metal Melt Filtration
Print ISBN: 978-3-031-40929-5

Electronic ISBN: 978-3-031-40930-1

Copyright-Jahr: 2024
DOI: https://doi.org/10.1007/978-3-031-40930-1_27

Premium Partner

Marktübersichten

Die im Laufe eines Jahres in der „adhäsion“ veröffentlichten Marktübersichten helfen Anwendern verschiedenster Branchen, sich einen gezielten Überblick über Lieferantenangebote zu verschaffen.

Zur Marktübersicht