Austenitic TRIP/TWIP Steels and Steel-Zirconia Composites

For a composite containing 5 vol% Mg-PSZ (consisting of S1 and Z1), a porosity < 10 vol% is required to form the maximum possible strain-induced α-martensite content [18]. The grain size of the steel matrix changes only for small Mg-PSZ contents with pressure.

Table 9.1 Chemical composition of the steel powders

Sintering of Functionally Graded Materials (FGM) by FAST

Thus, the temperature in the ceramic-rich zone converged to that of the steel-rich zones, as shown by temperature measurements within the matrix [46]. During the dwell time, the temperature of the ceramic-rich area even exceeded the temperature of the steel-rich area by 10 K [46].

Fig. 9.8 Current path (marked by white arrows) through the sintering tool and the FGM and necessary sintering temperatures

Conclusions

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Electron Beam Technologies

Introduction

To evaluate a potential application, the weldability of the material should be investigated. Due to evaporation and the lower grain size, the energy of the weld seam stacking defect is locally lower than in the base material.

Fig. 10.1 Weldability of TRIP-matrix composites dependent on the the material, the construction and the process (according to DIN ISO/TR 581:2007 [11])

Materials and Methodology

Electron Beam Facility and Temperature Measurements
Base Materials
Microstructural Characterization
Mechanical Characterization
Non-destructive Testing
Electron Beam Welding of Similar Joints Without Reinforcement

Influence of the Welding Parameters on the Seam Geometry In order to investigate the welding behavior of the base material without the influ-
Macroscopic Appearance of the Welding Joints
Influence of the Welding Parameters on Evaporation
Microstructure of the Welding Joints
Mechanical Behavior of the Welding Joints

Electron Beam Welding of Similar Joints with Reinforcement

The beam current required for complete penetration (IB,FP) is shown in Figure 10.4d as a function of welding speed. In the weld, the appearance of which was almost independent of the welding speed, dark etched areas were also visible (Figure 10.8d).

Fig. 10.2 Sample geometry and positioning of EB for a all joining methods with beam offset (if applied) and information about cross-sections (CS), longitudinal-sections (LS) and flat-sections (FS) with respect to the welding seam (WS); b electron beam braz

Electron Beam Welding of Dissimilar Joints with TWIP-Matrix Composites

Typical Microstructure of the Welded Zone
Influence of Beam Parameters on the Weld Quality
Verification of Welding Defects
Mechanical Characterization

Other ceramic particles were partially melted and deposited at the edge of the cavity (cf. Fig. 10.11d). The level of dilution (D) was used as a measure of the amount of Mg-PSZ introduced into the weld pool (Fig. 10.12a).

Table 10.6 Criteria for categorization of welding seams into evaluation groups (EG) according to DIN EN ISO 13919-1

Electron Beam Brazing of TWIP-Matrix Composites .1 Macroscopic Phenomena

Microscopic Characterization
Tensile Tests

Following the temperature gradient in the z direction, a transition zone (Fig. 10.23c, d) was detected, which decreased in width from the top (≈85 µm) to the bottom (≈18 µm) of the samples. Zr was primarily found in the Mg-PSZ particles of the base material and the transition zone in the filler (Fig. 10.24e–h).

Fig. 10.19 Influence of the energy distribution on the temperature distribution at the sample sur- sur-face during brazing

Summary

Strong ejection of the melt occurred and a cavity formed in the center of the weld zone. Thus, the current research focuses on the interaction of the Mg-PSZ with the electron beam.

Microstructure Aspects

Introduction

In fig. 11.1a, b, the temperature dependence of the mechanical behavior is illustrated respectively on the true stress-strain curves measured under tensile loading and on the strain hardening calculated from these stress-strain curves. Below 100 °C, the steel shows a drastic increase in strength and a further hardening in the strain range between approx. This difference in the temperature dependence of the mechanical properties can be explained by different deformation mechanisms in OFHC copper and in TRIP/TWIP steel, and by different activities of underlying microstructure defects and microstructure phenomena such as dislocations, stacking faults and deformation induced martensitic phase transformations.

Fig. 11.1 Mechanical properties of X3 CrMnNi16-6-6 TRIP/TWIP steel [4] and pure (OFHC) copper [5] obtained from tensile testing at different temperatures: a stress-strain behavior and b corresponding strain hardening depicted in form of the Kocks-Mecking p

Fundamental Microstructure Defects, Their Activity and Configurations in Austenitic Steels

Dislocations and Stacking Faults in fcc Materials
Dislocations and Stacking Faults in Austenitic Steels, Their Configurations and Interactions
Arrangement of the Stacking Faults in Austenite

Crystallographic and Thermodynamic Aspects of the Stacking Fault Arrangement in Fcc Materials
Detection of the Stacking Faults and Their Arrangements by Diffraction Methods
Coexistence of Different Stacking Fault Arrangements
Thermal Stability of the Deformation-Induced ε-Martensite
Phase Transformations in Austenitic Steels Under High Pressure

In crystals without defects, the stacking sequence of the close-packed planes is {111} along the respective perpendicular direction 111ABCABC (Fig. 11.2, left). An example of the Frank-Read source can be seen in area II in Fig.11.5a. The contribution from the strain fields of partials (σp) can be neglected, especially if the stacking faults are wide [21].

Fig. 11.2 Schematic arrangement of the atoms in the close-packed layers in an unfaulted fcc crystal (left), after the passage of the leading partial dislocation (middle) and beyond the stacking fault terminated by the trailing partial (right)

Formation of α -Martensite

The stages of the phase transformation are schematically depicted in the middle column and on the left side of Figure 11.16. In addition to local phase identification, FFT/HRTEM patterns from adjacent regions A, B and C in Figure 11.17 were used to analyze the orientation relationships (OR) between original austenite (phase 1), defective ε-martensite (phase 2) and α-martensite (phase 3). From the coincidence of different reciprocal grid points in the FFT/HRTEM patterns, the following parallelisms are obtained.

Fig. 11.16 Stages of the deformation-induced α -martensite transformation in an X2 CrMnNi16- CrMnNi16-7-6 steel as seen by ECCI (left-hand side), and depicted schematically for nucleation at deformation band intersections (centre) and inside deformation b

Quantification of Microstructure Features

Experimental Methods for Quantitative Microstructure Analysis

The stress-strain curve (solid line) shows the deformation state of PM X3CrMnNi16-7-6 steel. The effect of the limited lateral resolution of EBSD on the result of quantitative microstructure analysis is shown in Figure 11.19. In addition, the effect of twinning on line broadening and displacement is small compared to isolated folding faults [33].

Fig. 11.19 Difficulty of indexing twin bundles by EBSD due to the limited resolution and the superposition of Kikuchi patterns in the fine-scale microstructure

Methods for Determination of the Stacking Fault Energy (SFE) in fcc Crystals

In the computer routines used for the Rietveld refinement of the XRD patterns, the dependence of the dislocation-induced line broadening on the diffraction indices is calculated using the Popa model [88]. If the distance between the partials becomes 'infinite', the work due to the action of the shear stress component τzx on the partials must be in equilibrium with the stacking fault energy as described by (11.19). The last term in (11.24) describes the correction of the measured lattice parameters for instrumental aberrations, which cause a line shift θ.

In Situ Diffraction Studies on TRIP/TWIP Steels During Plastic Deformation

The reason for the observed discrepancies is the invalidity of the Warren model for short (Fig.11.21d) and crossing (Fig.11.21f) faults. Analogously, the apparently negative SF probability observed for εmechx < 2% is related to the early stage of SF formation (cf. Fig.11.21d). The 6% Ni steel (Fig. 11.30b) exhibits similar characteristics, which are finer-scaled due to the fully recrystallized initial microstructure.

Fig. 11.21 Left-hand side: Lattice parameters of austenite deformed at the strains in the outer fiber of 1% (a), 3.5% and 5.6% (b), and 8.3% (c)

Interplay of Deformation Mechanisms, Development of Deformation Microstructure

Interaction of Microstructure Defects in Deformation Bands
Orientation Dependence of the Stacking Fault and Deformation Band Formation
Dependence of the Deformation Mechanisms on Local Chemical Composition and Temperature

The mechanisms of the deformation band formation and the strengthening by the α-martensite nucleation are summarized in Fig.11.33. The corresponding dependence of the deformation mechanisms on the temperature is illustrated in Fig.11.36 for the steel composition X3CrMnNi16- 6-6. A high volume of α-martensite at low (deformation) temperatures is one of the reasons for a high ultimate tensile stress (UTS in Fig.11.38).

Fig. 11.31 TEM micrograph evidencing the obstacle effect of a deformation band for the dislocation motion on secondary slip systems.

Conclusions

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Investigations on the Influence of Strain Rate, Temperature and Reinforcement

Introduction

The folding fault energy mainly depends on the chemical composition of the steel and the deformation temperature [8]. On the other hand, the use of these metal matrix composites (MMCs) in cellular networks follows the example of wood or cork, which could have a favorable strength-to-weight ratio [11]. Detailed investigations of the developed microstructure were carried out in order to interpret the experimental results with special emphasis on the martensitic transformation.

High Strain Rate Deformation of Austenitic High-Alloy TRIP/TWIP Steel

Processing and Experimental Methods
Approaches to Rate-Dependent Constitutive Modeling
Microstructural Deformation Mechanisms at High Strain Rates

As expected, increasing strain rate led to crossing of the current curves (Fig. 12.3a) due to reduced strain hardening due to lower fraction of α-martensite formed (Fig. 12.3b). 23], the formation of α-martensite nuclei led to the obstruction of the dislocation motion within the deformation bands and the accumulation of dislocations at the phase boundaries. The shock loading at 200 °C caused only a small amount of planar defects to appear in the microstructure (see Figure 12.10b).

Fig. 12.1 Schematic setup of the plate impact experiments to determine the strength properties under planar shock loading at room temperature [20]

Honeycomb-Like Structures Made from TRIP-Steel and TRIP-Matrix-Composites

Deformation Behavior of Honeycomb-Like Structures

Out-of-Plane Direction
In-Plane Direction

Selection of Cell Wall Materials

Influence of Nickel Content
Effect of Particle Reinforcement

However, centered cell walls are limited to deform freely due to the constraining effect of the outer elements. The weighting of their influence is contained in the nickel equivalent equation, cf. Consequently, the enhancing effect of the martensitic phase transformation within the Mg-PSZ is relatively low in conventional sintered honeycomb-like structures [84].

Fig. 12.12 Stress-strain curves of square-celled structures with 64 cpsi and 196 cpsi made from steel batch 18-1-9, compressed at quasi-static strain rates and room temperature [62]

Conclusions

As a result, Mg-PSZ particles are arranged in unattached cluster chains, leading to a weakening of cell wall cross-section, cf. Nevertheless, the reinforcement of TRIP steel with Mg-PSZ particles can contribute to increased energy absorption capacity up to 50% technical strain, as the comparison of the non-reinforced and the 5 vol% Mg-PSZ-containing sample in Fig.12.18 shows. The strength level and strain hardening potential at this stage were controlled respectively by the relative density of the structure, the deformation mechanisms in the steel and the interactions between steel and Mg-PSZ particles.

Cyclic Deformation and Fatigue

Behavior of Metastable Austenitic Steels and Steel-Matrix-Composites

Introduction

First, the influence of the chemical composition and thus the SFE on the cyclic deformation behavior and especially on the fatigue life is presented in the section. Subsequently, in Sect.13.4, the influence of the manufacturing method for austenitic steels is discussed, including additively manufactured [14,15], cast [12] and ultrafine-grained material states [16]. In the case of the current high-alloy austenitic steels (cf. 13.3), particles of MgO partially stabilized zirconia (Mg-PSZ) are used for reinforcement [17].

Methodology

Materials
Manufacturing Methods
Fatigue Testing
Analytical Methods

A meandering scanning strategy in combination with a 90° rotation of the scanning direction after each layer was used. 21] requires subzero cooling in a second step (Fig. 13.2) to induce athermalα-martensite, the target fraction of which can be adjusted by variation of the cooling temperature Tc. Scanning electron microscopy (SEM) was used to characterize the microstructure of the different materials before and after cyclic deformation.

Table 13.2 Chemical composition of the delivered Mg-PSZ powder in wt.% and its designation in this book (5 and 10, respectively, indicate the volume fraction)

Influence of Chemical Composition on the Fatigue Behavior

Cyclic Deformation Behavior
Microstructure After Cyclic Deformation
Fatigue Life

Some individual fold faults are also present between the deformation bands, as shown in Figure 13.7c (white arrows). Near the tip of the crack, due to greater local plasticity, a higher density of the deformation band is observed (Fig. 13.8c). Fatigue life estimation for HP materials by Basquin-Manson-Coffin is shown in Figure 13.9a.

Fig. 13.3 Cyclic deformation curves for different strain amplitudes of the materials HP 16-6-4 (a), HP 16-6-6 (b) and HP 16-6-9 (c)

Influence of the Manufacturing Method on the Fatigue Behavior

Microstructure of the Undeformed State
Cyclic Deformation Behavior and α -Martensite Formation
Fatigue Life

The values given in Table 13.1 are valid for powders of EBM and HP states, respectively. As discussed above, the grain size of austenite and α-martensite is similar in the UFG condition (<1μm). In this context, the relatively high porosity of the material after EBM processing must be taken into account (cf. Section 13.4.1).

Fig. 13.10 EBSD orientation maps of the microstructures before cyclic deformation of the cast (a), EBM (b, build direction vertical), UFG (c) and HP (d) state, respectively

Influence of Particle Reinforcement

Cyclic Deformation Behavior of Particle Reinforced Steel-Matrix-Composites
Damage Evolution
Cyclically Deformed Microstructure
Fatigue Life

Thus, the deformation in the vicinity of the defects is quite pronounced at the beginning of the cyclic deformation. For variants based on HP 16-6-9 steel (Fig.13.18d–f) secondary hardening due to strain-induced twinning (cf. 13.3.1) is even completely suppressed in the case of composites. This section presents the microstructure of the HP 16-6-9+10Z composite with a focus on Mg-PSZ particles.

Fig. 13.18 Cyclic deformation curves of the composites reinforced with 5% Mg-PSZ (a, d) and 10% Mg-PSZ (b, e) based on the HP 16-6-6 (a, b) and HP 16-6-9 steels (d, e), respectively

Fatigue Properties of a Q&P Ultra-High Strength Steel

Microstructure After Q&P
Cyclic Deformation Behavior
Fatigue Life

While at medium and small strain amplitudes the overall α-martensite fraction after failure (Fig.13.24d) remains lower in the case of the. As discussed in Section.13.3.3, different austenite stabilities cause a crossover of the fatigue life curves in the case of high-alloy austenitic steels (Fig.13.9a). The BSE, ECCI and EBSD images of Fig.13.26 give an impression of the microstructure of the Q&P steel after cyclic deformation.

Fig. 13.23 BSE image (a) and corresponding EBSD phase map with superimposed band contrast (b, red—austenite, yellow— ε -martensite, blue— α -martensite) of the − 120 °C state after Q&P.

Conclusions

Consequently, the martensitic phase transformation occurs faster in the -20 °C condition due to its lower austenite stability. The grain size of the steel affects the cyclic deformation behavior and the martensitic phase transformation. Consequently, the fatigue life of the composites is reduced in the total-strain-controlled experiments, especially in the LCF regime.

Fig. 13.26 Microstructures of the − 120 °C (a–d, g, h) and − 20 °C (e, f, i) state cyclically deformed at ε t /2 = 0.3% (a, c–f, i), ε t /2 = 0.5% (b) and ε t /2 = 0.8% (g, h), respectively

Behaviour of Metastable and Stable Austenitic Stainless Steels Under

Introduction

As a result, for planar-biaxial crack growth experiments, the crack growth is described in several dependencies: (i) exclusive consideration of the fatigue crack growth ina-N plots, as e.g. 49], (iii) consideration of fatigue crack growth in da/dN plots, as done by Wang et al. In this paper, both a–Nplots and da/dN–Kplots are used for the study of fatigue crack growth.

Materials and Methods .1 Material

Quasi-static Loading
Low Cycle Fatigue
Fatigue Crack Growth

The slits in the "Triple-slit" and "Double-slit" samples were subsequently inserted by water jet cutting, cf. Double-slit' samples had the same length and were shorter than the slits on. In the highly stressed measurement area in the middle of the samples, there was a homogeneous stress distribution.