or aluminium titanate

Accepted Manuscript Compressive and tensile deformation behaviour of TRIP steel-matrix composite materials with reinforcing additions of zirconia and/or aluminium titanate C. Weigelt, G. Schmidt, C.G. Aneziris, R. Eckner, D. Ehinger, L. Krüger, C. Ullrich, D. Rafaja PII:

S0925-8388(16)33299-6

DOI:

10.1016/j.jallcom.2016.10.176

Reference:

JALCOM 39343

To appear in:

Journal of Alloys and Compounds

Received Date: 27 November 2015 Revised Date:

9 September 2016

Accepted Date: 18 October 2016

Please cite this article as: C. Weigelt, G. Schmidt, C.G. Aneziris, R. Eckner, D. Ehinger, L. Krüger, C. Ullrich, D. Rafaja, Compressive and tensile deformation behaviour of TRIP steel-matrix composite materials with reinforcing additions of zirconia and/or aluminium titanate, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.10.176. 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 proof before it is published in its final 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.

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ACCEPTED MANUSCRIPT

ACCEPTED MANUSCRIPT Compressive and tensile deformation behaviour of TRIP steel-matrix composite materials with reinforcing additions of zirconia and/or aluminium titanate Authors: Weigelt, C. Rafaja, D.

1,*)

1)

1)

2)

2)

2)

3)

; Schmidt, G. ; Aneziris, C. G. ; Eckner, R. ; Ehinger, D. ; Krüger, L. ; Ullrich, C. ;

3)

Institute of Ceramic, Glass and Construction Materials, Technische Universität Bergakademie Freiberg

2)

Institute of Materials Engineering, Technische Universität Bergakademie Freiberg

3)

Institute of Materials Science, Technische Universität Bergakademie Freiberg

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[*] Corresponding author Dr.-Ing. Christian Weigelt Institute of Ceramic, Glass and Construction Materials Technische Universität Bergakademie Freiberg

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Agricolastraße 17 09596 Freiberg, Germany E-mail: [email protected]

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Phone: +49 (0) 3731 394254 Fax: +49 (0) 3731 392419 Abstract

The mechanical properties and the related microstructure of metal-matrix composites (MMC) based on a high-alloyed CrMnNi steel with varying particle reinforcements (5 % or 10 vol.%) of magnesia partially stabilized zirconia and/or aluminium titanate were investigated. The powder metallurgical processing

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comprised the cold extrusion and conventional solid-state sintering at 1350 °C. The mechanical propert ies were examined by quasi-static compressive and tensile loading tests at ambient temperature. The microstructure characteristics contributing to significant changes in strength and ductility and affecting the failure mechanisms during deformation were characterized by scanning electron microscopy including

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energy-dispersive X-ray spectroscopy and electron backscatter diffraction, and by X-ray diffraction. For all compositions the stress-strain and the deformation behaviour were mainly controlled by dislocation hardening and α’-martensite formation in the steel matrix, but it was further improved by the particle

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strengthening of the ceramic reinforcement. Thus, the composite materials showed higher strength and work hardening than unreinforced steel specimens over a wide strain range. The pressureless sintering triggers several reactions at the metal/ceramic interface, which leads to pronounced destabilisation of the initial metastable zirconia particles with a lack of transformation toughening capability, and the formation of invalid olivine. These MMC suffer from early particle/matrix displacement and particle fracture under loading. A more positive effect of interfacial reactions was observed in the composites with addition of aluminium titanate. Here, the formation of a dense spinel structure and the reliable matrix/particle interface bonding during firing provides a significant particle strengthening effect. The combination of zirconia and aluminium titanate provides mechanical and microstructural benefits as well. Thus, aluminium titanate facilitates the consolidation of powder metallurgical processed steel-matrix MMCs via the conventional

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ACCEPTED MANUSCRIPT sintering for advanced load applications. The results of the study help to understand essential hardening mechanisms in composite materials and provide potential for future cellular MMC structures. Keywords Metal-matrix composite, Aluminium Titanate, Zirconia, TRIP steel

1 Introduction

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Metal-matrix composites (MMCs) gained in popularity, since they provide a large variety of tailored characteristics with a wide range of mechanical, thermal, corrosion-resistant, and wear-related properties. The volume proportion, the size, and the morphology of the reinforcement mainly affects the performance of the composite material. The addition of ceramic particles to a ductile metallic matrix increases the

like debonding, particle fracture or crack coalescence. [1-7]

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strength and the elastic modulus, but lowers the material’s deformability by inducing typical damage events

MMCs composed of a steel showing a transformation induced plasticity effect (TRIP) offer a matrix

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material with outstanding ductility, high strength, and reasonably high capacity for absorption of mechanical energy [8-10]. These properties originate from a strain-induced martensitic phase transformation from austenite (fcc) to bcc martensite under mechanical loading. The phase transformation can be overlapped by a twinning induced plasticity (TWIP) effect, which further improves the deformability and strain hardening of such steels. The deformation mechanisms are governed by the austenite stability and by the stacking fault energy γSF (SFE) that depend on the chemical composition of the alloy and on the temperature applied during deformation [8,11]. The SFE determining the dominant deformation mechanism

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can be estimated by various empirical equations. According to Weiß et al. [12], the presence of austenite and martensite can be predicted from the chemical composition by using the proportions and individual weighting factors for each element of the alloy. The nickel equivalent Nieq (Equation 1) is calculated from the austenite promoting elements (e.g. Ni, C, Mn), while the chrome equivalent Creq (Equation 2) considers the austenite supressing character of elements such as Cr, Ti, and Si, whereby an alloy composition with

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14-22 wt.% of Creq and with 8-16 wt.% of Nieq of is essential for TRIP steels [12]. Nieq = wt.%Ni + 30×wt.%C + 18×wt.%N + 0.5×wt.%Mn + 0.3×wt.%Co + 0.2×wt.%Cu – 0.2×wt.%Al

(1)

Creq = wt.%Cr + wt.%Mo + 4× wt.%Ti + 4×wt.%Al + 1.5×wt.%Si + 0.9×wt.%Nb + 0.5×wt.%W + 0.9×wt.%Ta + (2)

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1.5×wt.%V

Numerous MMCs have been investigated using a variety of manufacturing processes such as stir casting [13], squeeze casting [14], spray forming [15], preform infiltration [16,17] or powder manufacturing [2,3,18-22]. Despite some advantages, common MMC technologies suffer from the restriction to bulk structures or components. The fabrication of cellular structures from powder materials [23,24] has rarely been reported on in spite of prospects in automotive or engineering applications. The ceramics-derived extrusion technology at room temperature facilitates the shaping of both cellular and bulk structures of the same material. This route comprises the powder-blending, admixture of an aqueous organic binder system to form a plastic paste and pressing the paste through a rigid die. Thermal binder removal and sintering ensures the final consolidation. Conventional pressureless sintering is necessary in case of cellular structure to prevent extensive global densification.

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ACCEPTED MANUSCRIPT Different studies reported on the reinforcement of a steel matrix with zirconia particles [1,25]. The reinforcing mechanism of zirconia is due to the martensitic phase transformation from the tetragonal (tZrO2, space group P42/nmc) to the highly distorted monoclinic (m-ZrO2, space group P21/c) modification, which causes a volume expansion of 3–5 % and generates compressive stress in the surrounding matrix material [26]. The addition of stabilizers, e.g. MgO or Y2O3, retains the tetragonal and the cubic (c-ZrO2,

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space group Fm-3m) polymorph structures in a metastable state at ambient temperature. MMCs based on a TRIP steel with randomly dispersed magnesia partially stabilised zirconia (Mg-PSZ) particles showed improved performance up to a certain degree of deformation owing to the combined martensitic phase transformations in the metal and ceramic components [10,18,27-29]. The ceramic particles increased both the strength and the initial nucleation rate of α’-martensite in the steel matrix during quasi-static compressive deformation. After exceeding a critical strain, the material suffers from debonding at the

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steel/zirconia interface, from the rupture of the zirconia particles, and from the crack coalescence, which initiates an extended macroscopic failure [30]. In this context, it should be pointed out that the conventional

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pressureless sintering of TRIP steel-matrix composites reinforced with magnesia partially stabilised zirconia is accompanied by diffusion reactions between the material components, even if the specimens are sintered well below the normal sintering temperature of Mg-PSZ. The out-diffusion of Mg from the ZrO2 particles and the formation of precipitates (spinel of type (Mg,Mn)Al2O4 and silicate structures of type Mg0.38Mn0.62SiO3) at the steel/zirconia interface lead to an increase of non-transformable monoclinic zirconia during firing [28,31,32].

A new approach in the fabrication of MMCs for high load applications is the introduction of aluminium

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titanate (Al2TiO5) as a reinforcing phase. Aluminium titanate is a well-known refractory material offering excellent thermal properties such as low coefficient of thermal expansion, low heat conductivity, high thermal shock resistance, and appropriate refractoriness [33,34]. Commonly, alumina and titania are fired in a stoichiometric mixture forming aluminium titanate at elevated temperatures [35]. However, on cooling below 1280 °C the aluminium titanate is susceptible to astringent decomposition, forming α-Al2O3 and

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rutile-TiO2. Besides, the thermodynamic instability of aluminium titanate is enhanced during exposure to a low-oxygen atmosphere at elevated temperatures [36-38]. Analogously to zirconia, the presence of minor fractions of additions, e.g. MgO, SiO2, or Fe2O3, can crucially affect the material’s characteristics and its

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phase composition after cooling. In case of MgO the formation of complex phases containing Mg, Al, Ti, and O and the partial substitution of Al

3+

in the aluminium titanate phase effect the stabilizing efficiency

[33]. Due to its reactivity, the presence of titanium is beneficial in the formation of improved metal/ceramic interfaces during firing [16,32,39,40] and thus, aluminium titanate is a promising material for reinforcing TRIP-matrix materials.

The aim of the present study was to analyze and to understand the interplay between the matrix material and different reinforcing components in powder metallurgically processed MMCs based on a highalloyed CrMnNi steel-matrix with additions of Mg-PSZ and/or aluminium titanate. The reinforcing materials were used solely or coevally with total volume fractions of 5 and 10 %. The materials were microstructurally characterised and their deformation behaviour was tested under quasi static compressive and tensile loading at ambient temperature.

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ACCEPTED MANUSCRIPT 2 Material and Methods The matrix material used in the present work was a TRIP steel of type 16-7-3 (16 wt.%Cr, 7 wt.%Mn, and 3 wt.%Ni) prepared via gas atomization (TLS, Germany). The chemical composition of the spherically shaped steel powder and the characteristic values are given in Table 1. In reference [41], the stacking fault -2

energy of this steel was determined to be (8.1±0.9) mJm . This value agrees well with the SFE of -2

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(8.8±0.6) mJm , which was estimated by using the equation of Dai et al. [42] for this chemical composition. Table 1. Chemical composition and the characteristic values of the initial steel powder. Carbon concentration was determined by using combustion analysis, nitrogen concentration by using carrier hot gas extraction and the concentration of other elements by using inductively coupled plasma (ICP)

Cr

Mn

Ni

Si

N

C

in wt.% 7.1±0.1

Ti

Al

3.4±0.1 0.91±0.02 0.097±0.003 0.05±0.01 0.007±0.001 0.017±0.003

Nieq

Creq -

10.2

18.2

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16.7±0.1

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spectroscopy

Four MMC variants were prepared with reinforcing additions of magnesia partially stabilized zirconia and/or aluminium titanate. A commercially available Mg-PSZ (Saint Gobain, USA) with 8 mol% MgO was used. The most important impurities and their concentrations are listed in Table 2. The fused cast ceramic powder had an irregularly shaped and fissured morphology. According to the phase analysis performed by

and 33 vol.% c-ZrO2.

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using X-ray diffraction (XRD) [43], the Mg-PSZ contained approximately 35 vol.% m-ZrO2, 32 vol.% t-ZrO2,

Table 2. Chemical composition of the initial zirconia powder determined by ICP, in wt.% Mg

Hf

bal.

1.7

1.5

Si

Al

Ti

Ca

Fe

0.2

0.2

0.1

0.1

0.1

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Zr + others

The aluminium titanate was synthesized from a stoichiometric mixture of a high-purity 99.8 wt.%

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alumina powder (Martinswerk, Germany) and a titania powder with the purity of 98 wt.% (Sachtleben, Germany). The addition of 1.4 wt.% of a high-purity grade (98 %) MgO powder (neoLab, Germany) was used in order to improve the thermal stability and to prevent the aluminium titanate from decomposition during cooling after reactive sintering [33]. The solid mixture was fired at 1400 °C for 2 h in an oxi dising atmosphere. The phase composition of approximately 94 vol.% orthorhombic aluminium titanate with a low amount of residual trigonal alumina (~4 %) and tetragonal titania (~1 %) was estimated by XRD. The milled aluminium titanate exhibited a similar powder fineness as the used zirconia powder (see Table 3) in order to avoid differences in the distribution of the second phase within the matrix. Actually, the ratio of the reinforcement particle size to the matrix’ particle size is ≤0.3 for both ceramic variants, which is much smaller than the ideal ratio of one [5]. Thus, powder mixing and processing favours clustering of the

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Material

d10

d50

d90

TRIP steel

6.5

18.1

76.3

Mg-PSZ

0.2

3.0

20.4

Aluminium titanate

0.3

4.2

15.7

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Table 3. Particle size distribution of the metal and ceramic powders, in µm

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TRIP steel matrix composite variants were prepared with zirconia (Z) or aluminium titanate (AT) only, and with balanced mixtures of both (ZAT, see Table 4). Additionally, the 100 % steel composition (S100)

Table 4. Composition of mixtures, in vol.% Recipe TRIP steel (S) Mg-PSZ (Z) Aluminium titanate (AT)

S100

5Z

100 -

95 5

-

-

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without any ceramic additions provides a reference material.

5AT

2.5ZAT

5ZAT

95 -

95 2.5

90 5

5

2.5

5

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The compact specimens were fabricated by a ceramics-derived extrusion technology. The solid powders were mixed in their dry states for 90 min using a ball mill with a PET vessel. In order to prevent contamination of the materials, the compositions were mixed using high-purity yttria-stabilized zirconia balls (5Z/2.5ZAT/5ZAT) or alumina balls (5AT). Afterwards, an organic binder system was added to the mixed powders to form a plastic paste at room temperature. The binder system was composed of 1.3 wt.%

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cellulose derivative of type Zusoplast C28NEU (Zschimmer & Schwarz, Germany), 0.3 wt.% HPMC 874 (Ashland, USA), and 7.0 wt.% de-ionized water. Additionally, 0.25 wt.% of a surfactant (glycerine) and 0.15 wt.% of a lubricant (oleic acid) were necessary for sound extrusion. Continuous rods of 11 mm in

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diameter were extruded using a single-screw extruder of type LK III 2A (Linden, Germany) and cut to a length of 110 mm for further handling. The water was removed stepwise from the shaped specimens in an air-circulated dryer, applying a maximum temperature of 110 °C which was maintained for 24 hours. Then the specimens were preheated in an oxidising atmosphere in order to ensure sufficient elimination of the organics. The heating and cooling gradient was 1 Kmin-1 and the maximum temperature 450 °C. The inertgas atmosphere sintering of the binder-free specimens was conducted in a graphite furnace (Xerion, Germany) applying a heating gradient of 5 Kmin

-1

up to the maximum temperature of 1350 °C. This

temperature was maintained for 2 hours before cooling with a maximum cooling gradient of 5 Kmin

-1

to

room temperature. The bulk density was calculated from the geometrical dimensions and the weight of the fired rods. The pore size distribution of the as-fired material was determined by mercury intrusion. The total

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ACCEPTED MANUSCRIPT porosity was estimated by the ratio of the bulk density and the true density of the starting powders which was measured with a Helium pycnometer. The sintered products were processed to cylindrical specimens with 6 mm in diameter and height for compressive load testing. The test series were performed using a 200 kN servo-hydraulic testing machine -1

of type MTS 810 at a quasi-static nominal strain rate of 0.001 s at room temperature. Tensile tests were

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carried out in a universal testing device ZWICK 1476. Therefore, tensile specimens with 5 mm in diameter and a gauge length of 30 mm were machined. The experiments were carried out at room temperature at a -1

quasi-static strain rate of 0.001 s . For accurate strain measurement a clip gauge (MTS-Systems) was used. The α’-martensite content was measured non-destructively during testing using a Feritscope FMP-30 magnetic induction measurement device (Helmut Fischer GmbH). Consequently, these tensile tests were stopped in 2 % strain steps without unloading to measure the ferromagnetic content at five different

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positions evenly distributed over the gauge length. Fig. 1 shows a schematic illustration of the experimental setup. Due to the difference in permeability between ferrite and α’-martensite, the Feritscope measurement

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was converted to the martensitic volume fraction by using a multiplication factor of 1.7, as stated by Talonen et al. [44]. Furthermore, the measured values had to be corrected for the initial ferromagnetic content of the specimens as well as for the curvature of the specimen. In addition to these effects, the measured values are affected by the inverse magnetostriction, which is often attributed to the Villari effect [45,46]. The magnetic domains rotate under tensile straining which leads to a deviation in permeability [47]. The measured content of α’-martensite should therefore be different from those without stress. This

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influence is characteristic for such in situ tests and must be considered when comparing with other data.

Fig. 1. Schematic experimental setup for the tensile tests with in situ magnetic volume fraction measurement

The microstructural characterization of the material variants comprised analyses in the as-fired state and after tensile deformation tests. As known from previous studies, the effect of ceramic particles on the composite material’s deformation behaviour and the initiation of failure are more pronounced under tension than under compression. The deformed specimens were retrieved parallel to the load axis in the area of uniform elongation. The materials were cut and prepared by conventional grinding and polishing steps

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ACCEPTED MANUSCRIPT followed by vibrational polishing with colloidal silica. The microstructural characterization was conducted using scanning electron microscopy (SEM, Zeiss LEO1530 Gemini). Electron backscatter diffraction (EBSD) was utilised for the analysis of local crystal structures and orientation relationships. Energy dispersive X-ray spectroscopy (EDS) complemented the microstructural characterization by measuring local chemical compositions. The XRD measurements were performed for qualitative and quantitative

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phase analysis, which was carried out by using the software Maud [48]. The XRD analyses of the 100 % steel specimens were performed after an additional electrolytic polishing step. The homogeneity of reinforcement distribution was determined from SEM images using the Stream Image Analysis Software (Olympus) at 6-8 individual areas of 74·10³ µm² each. The evaluation of the nearest-neighbour-distance reflects the overall dispersion of the ceramic particles within the matrix.

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3 Results and Discussion 3.1 Characterization of the initial state

The as-fired specimens appeared with a glossy metallic surface due to the high-purity argon

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atmosphere inhibiting appreciable oxidation during sintering. All materials exhibited relative densities of >95.5 % irrespective the quality and quantity of ceramic addition. This value corresponds to conventionally (solid-state) sintered MMCs based on similar matrix steels [22,49]. However, the medium pore size of composite 5Z (41 µm) differs clearly from the pore system of the remaining materials (1.7-2.7 µm) which implies differences in the deformation behaviour. This deviation is more likely a reason of the extrusion process than the consolidation during thermal treatment, since all composite variants were prepared with the same starting powders and thermal regime.

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The as-delivered steel powder consisted of 53 % austenite. The remainder was α’-martensite and δferrite that cannot be reliably distinguished using XRD due to their bcc crystal structure. During firing, the phase composition altered; a highly austenitic matrix material formed that contained about 20 % α’martensite/δ-ferrite and roughly 8 % ε-martensite, as it can be expected for this chemical composition

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according to Schaefflers diagram [50]. The α’-martensite fraction in the 100 % steel specimens was (7±4) % as measured by Feritscope without loading. The alteration of the phase composition is a result of the re-transformation of martensite into austenite during sintering [51]. The phase composition of the fired

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steel was measured by XRD after electrolytic polishing (EP) the specimens’ surface since the used steel composition is prone to surface-martensite formation during conventional polishing (CP). Thus, the measurement of surfaces prepared by CP would be misleading (only about 20 % austenite) and not representative for the phase fractions in the specimens’ volume. However, the microstructure suffered during EP hampering the following microstructural analysis. The concentrations of selected elements are summarized in Table 5. The compositions of all materials exhibited several shifts as compared with the initial steel powder (cf. Table 1). These alterations, especially in the C, Cr, and Mn content, led to shifts of the characteristic values and of the stacking fault energy. Decreasing chromium and manganese contents and increasing carbon contents increase the austenite stability. It should be noted, that the wet-chemical digestion procedures (except for carbon and nitrogen) were designed for steels and, therefore, the resulting concentrations are mainly valid for the metal

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ACCEPTED MANUSCRIPT portion in all materials tested. Estimation of the 100 % steel specimens’ stacking fault energy (10.9±0.4) mJm-2 was done using the empirical relationship proposed by Dai et al. [42]. Such a low SFE implies that the austenite in this steel is highly metastable and that the predominant deformation mechanism will be the transformation induced plasticity (TRIP) and negligible twin formation [8,52]. The calculated SFE of the composite variant with solely zirconia addition ((10.4±0.4) mJm ) was lower as -2

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compared with the 100 % steel material. All composite variants with additions of primary aluminium titanate showed an increase in their matrix’ SFE as calculated from their chemical compositions: (11.7±0.4) mJm

-2

for the material 5AT, (12.6±0.4) mJm for the material 2.5ZAT, and (13.5±0.4) mJm for the material 5ZAT. -2

-2

Thus, the austenite is expected to undergo the martensitic transformation during the deformation tests at room temperature in all material compositions.

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Table 5. Chemical composition and the characteristic values of the steel matrix of as-fired specimens. Carbon concentration was determined by using combustion analysis, nitrogen by using carrier hot gas extraction, chromium by using potentiometry, nickel by gravimetry and the concentration of the other

Recipe S100

Cr

Mn

Ni

Si

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elements by using inductively coupled plasma spectroscopy N

C

Ti

Al

Nieq Creq

in wt.%

-

15.8±0.1 6.32±0.01 3.3±0.1 0.55±0.02 0.037±0.001 0.110±0.001 0.010±0.001 0.01±0.01 10.4 16.8

5Z

15.2±0.1 6.30±0.01 3.3±0.1 0.58±0.02 0.036±0.001 0.090±0.001 0.010±0.001 0.02±0.01

5AT

9.8

16.2

15.5±0.1 5.01±0.01 3.2±0.1 0.56±0.02 0.032±0.001 0.130±0.001 0.020±0.001 0.04±0.01 10.2 16.6

2.5ZAT 13.7±0.1 5.66±0.01 3.2±0.1 0.61±0.02 0.032±0.001 0.122±0.001 0.010±0.001 0.02±0.01 10.2 14.8 14.4±0.1 4.94±0.01 3.0±0.1 0.49±0.02 0.034±0.001 0.151±0.001 0.020±0.001 0.02±0.01 10.6 15.3

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5ZAT

Fig. 2 illustrates the characteristic coarse-grained microstructure of specimens prepared from 100 % steel after firing. The transformation of the highly metastable steel from austenite to martensite on the specimens’ surfaces is a result of mechanical polishing. The steel grains exhibited scattered segregations

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at their grain boundaries, with higher concentrations of chromium and lower concentrations of nickel as compared with the surrounding matrix, forming δ-ferrite. The local chemical composition of the common matrix determined by EDS, as given in Table 5, is in good agreement with the alloy composition derived

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from the overall chemical analyses (cf. Table 5). In addition to the metallic grains, the material contains randomly dispersed non-metallic sections that were most likely formed during firing. The influence of these particles on the mechanical properties is of minor significance, since all material variants were prepared from one steel powder batch and, therefore, they are also expected in all MMCs. The typical chemical composition of these randomly dispersed particles determined by EDS is given in Table 6. They are characterised by higher concentrations of light elements (particularly Si, O, and Mn as measured by EDS) as compared with the surrounding steel matrix, leading to significant changes in the greyscale intensities in backscattered electron (BSE) contrast mode. Therefore, the interpretation of the BSE micrographs is somehow misleading and the distinction of pores and precipitations requires the verification using EDS. However, the application of secondary electron (SE) contrast images is limited by the very small

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methods.

Fig. 2. Microstructure of a 100 % steel specimen after firing and conventional polishing: a) SEM (BSE)

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image illustrating the presence of martensite, δ-ferrite, and non-metallic particles in the austenitic matrix, and the cross section of EDS analysis (white dashed line); b-e) EDS mapping of the highlighted cross section in (a)

Table 6. Average chemical composition (EDS) of the S100 material, in wt.%

matrix particle

Fe + Cr Mn others 70.8-74.8 15.6-16.4 7.0-7.3 25.6-37.9 6.2-11.2 26.1-37.9

Ni

Si

Al

Ti

O

3.1-3.3 0.8-0.9

0.5-0.7 14.9-20.0

0.3-2.4

0.4-0.6

13.4-23.8

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Position

The formation of non-metallic phases during firing is more pronounced in the composite variants since the metal/ceramic interface is prone to diffusional interchange forming supplementary phases. Decker et al. [18] reported on the fabrication of similar materials by using SPS. In their study the higher

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heating gradient (90 K/min), the lower maximum temperature (1100 °C), and the lower dwell time (5 min) prevented the pronounced diffusion of elements in the composite materials and mutual influence on the

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chemical composition of both the matrix material and the reinforcing particles. Alterations in their chemical composition, especially in the Cr and Mn content, led to an increase of the matrix’ mean stacking fault energy in the composite variants as compared with the 100 % steel specimens. It should be noted that these effects are more pronounced in materials with additions of aluminium titanate as will be shown later. However, the matrix material in all composite variants is still expected to show the martensitic phase transformation (TRIP) under deformation [8,52]. The characteristic microstructure of composite specimens with zirconia is shown in Fig. 3. The ceramic particles are randomly dispersed and widely embedded in the surrounding steel matrix. Delamination and broken-out particles indicate the loose bonding at the metal/ceramic and ceramic/ceramic interfaces. This is most likely a consequence of the firing temperature of the composite material, which was considerably below the normal sintering temperature of such zirconia ceramics

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ACCEPTED MANUSCRIPT (≥1600 °C [26]). Extended regions of a lenticular sha ped microstructure within the zirconia indicate the presence of a large amount of monoclinic ZrO2 that was most probably formed during thermal processing of the specimens. The amount of Mg within the zirconia ranged between 1 % and 2 wt.% with a decreasing concentration from the core of the zirconia particles towards the metal/ceramic interface as measured by EDS. The diffusion of Mg is a well-known mechanism in such MMCs which leads to pronounced

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destabilisation during long-term firing at 1350 °C [31,53,54]. Hence, the possibility for the desirable tetragonal to monoclinic phase transformation under mechanical loading was restricted due the high amount of non-transformable monoclinic ZrO2 even just after sintering. Additionally, areas with high concentrations of Mg, Mn, and Si, without measureable quantities of Al were found by EDS analysis at the steel/zirconia interfaces. These were identified by EBSD as orthorhombic (Mg, Mn)Si2O4 with the olivine structure (space group Pbnm) that was certainly formed during sintering of the composite materials. The

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formation of these precipitations is likely to occur in such powder metallurgically processed materials during the long-term sintering by solid-state reactions at temperatures slightly below the melting temperature of the metal [10,32]. The (Mg, Mn)Si2O4 phase region acts as an intermediate layer which mainly promotes

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noticeably firm inclusion of the zirconia particles into the steel matrix.

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Fig. 3. Micrograph of material 5Z in the as-fired state: a) SEM (BSE) image showing largely embedded zirconia particles (light grey) in the steel matrix (dark grey); b) EDS mapping of the cross section highlighted in (a) showing the areas with olivine structure (arrows)

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Fig. 4 shows the microstructures of the initial aluminium titanate powder (after synthesis and milling) and of the composite material 5T after firing. Representative chemical compositions derived from EDS are given in Table 7. However, it must be considered that the oxygen quantification by EDS is inaccurate and involves large errors. In the composite, the ceramic particles are evenly distributed within the steel matrix, but are subject to a certain degree of agglomeration. Nevertheless, the embedded ceramic particles exhibited a smooth and homogeneously rounded shape as detected by SEM micrographs. Fractured or broken-out particles and porous interlayers as observed for the MMC variants with additions of zirconia particles could rarely be observed. By taking the alteration of the particle shape into consideration, several chemical reactions were assumed to have occurred between the metal and ceramic phase during thermal processing of the composite material. The randomly dispersed ceramic grains were identified as a (Mn, Mg)(Al, Ti, Cr)2O4 spinel (space group ‫݀ܨ‬3ത݉) by using EBSD/EDS analyses and verified by XRD, whereby

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ACCEPTED MANUSCRIPT no residual Al2TiO5 could be found. The decomposition of fine-grained aluminium titanate during firing was confirmed with the aid of XRD analysis. About 77 % of the initial aluminium titanate remained stable in the pure ceramic material (without any additions of steel), whereas no Al2TiO5 could be found in the composite material 5AT after sintering with the same heat treatment conditions. This phenomena is related to the thermal, chemical, and atmospheric conditions during sintering [55]. The spinel phase was further

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investigated by EBSD measurements. Their crystal orientations indicate that the spinel phase was arbitrary formed from fine-grained aluminium titanate particles without complete melting and recrystallization [55]. However, a liquid phase must be present during the thermal processing at 1350 °C, as it follows from t he phase diagram of the quasi-ternary Al2O3-MnO-TiO2 system [56]. Close to the composition of 47 at.% MnO, 28 at.% Al2O3, and 25 at.% TiO2, a particular melting is expected already at temperatures below 1300 °C. A second ceramic phase consisting of Al and O was found rarely in the centre of some large agglomerates,

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and was identified as corundum (α-Al2O3, space group R-3c) with the aid of EBSD. The corundum stems from the initial aluminium titanate powder, which contained ~4 vol.% unreacted α-Al2O3. Again, it must be

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kept in mind that the conventional sintering of the TRIP steel causes several shifts of the chemical composition compared with the initial steel powder (cf. Table 1) [53] and that these effects are more pronounced in the composite variants compared with the 100 % steel material. Manganese is particularly susceptible to evaporation and diffusion, which leads to a decrease of the Mn fraction within the metallic phase whose intensity is enhanced with increasing volume fraction of ceramic particles and is more

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pronounced in composite variants with aluminium titanate than in specimens containing solely Mg-PSZ.

Fig. 4. Microstructure of the aluminium titanate: a) initial powder after synthesis and milling; b) sintered composite material with 5 vol.% of initial aluminium titanate showing the metal matrix, the (Mn, Mg)(Al, Ti, Cr)2O4 spinel, and individual alumina grains; the markings denote the positions measured by EDS (see Table 7)

Table 7. EDS analysis of the material 5AT (cf. Fig. 4), in wt.% Position

Al

Ti

1

-

0.3±0.1

2

Cr

Mn

Fe

Ni

Si

15.1±0.4 6.5±0.2 74.2±2.0 3.2±0.1 0.8±0.1

19.1±0.9 14.7±0.4 10.2±0.4 31.8±0.9

-

11

-

-

Mg

O -

0.8±0.1 23±10.3

ACCEPTED MANUSCRIPT The microstructures of composite variants prepared with coeval additions of Mg-PSZ and initial aluminium titanate are shown in Fig. 5. Both types of ceramic are randomly dispersed in the steel matrix, and the coarser grained zirconia agglomerates are partially fringed by (Mn, Mg)(Al, Ti, Cr)2O4 spinel structures with a similar composition compared with the spinel found in 5AT. These spinel intermediate layers apparently improve the interfacial bonding between the TRIP matrix and the zirconia particles

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leading to less voids and broken-out particles after polishing, just as the olivine structures in 5Z specimens (cf. Fig. 3). Additionally, a second phase with high concentrations of Zr, Mn, Ti and O was detected in the composite specimens of type 5ZAT as an interlayer between the steel matrix and the zirconia particles by EDS (see Figure 6). The interfacial phase contains 5-10 wt.% titanium, while Ti can rarely be found in the zirconia particles (≤1 wt.%). Beyond 60-65 wt.% Zr, 4-7 % Mn, and 18-20 % O were measured. The high diffusivity and mutual solubility of Ti and Zr on crystal lattice site promote the formation of several solid

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solutions in the system Zr-Ti-O [57,58]. This local seam is similar to the Zr–Ti–O solid solution found in TRIP steel/zirconia MMCs doped with minor volume fractions of Ti [32] and which hampered the distinctive diffusion of Mg out of the zirconia. Similar Zr-Ti-O-solid solutions have been reported to show no

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deformation-triggered phase transformation [32]. The orthorhombic ZrTiO4 could not be identified by EBSD in the frame of the present experimental setup although the TiO2 fraction exceeds the limited solubility in the binary ZrO2-TiO2 system [57]. As a consequence of the random distribution of both ceramic components (aluminium titanate and zirconia) in the metal matrix during preparation, the thickness of the Zr–Ti–O interlayer is strongly fluctuating between 0 and 5 µm. Despite the significant loss of Mn and Cr in the steel and the elevated stacking fault energy in the material variants with additions of aluminium titanate, the metal matrix is still expected to undergo the martensitic phase transformation under deformation. It

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among all materials tested.

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should be noted, that cracked particles from the processing of the composite materials could not be found

Fig. 5. Microstructure of MMC specimens with combinations of zirconia and aluminium titanate in their asfired state showing the particularly inclusion of zirconia particles by aluminium titanate/spinel (light grey arrows) and a solid solution of Zr-Ti-O also containing Mn (dark grey arrows): a) 2.5ZAT; b) 5ZAT

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elements

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Fig. 6. Microstructure of specimen 5ZAT: a) SE-SEM image; b-i) appendant EDS maps for selected

The homogeneity of reinforcement dispersion was determined from SEM images (BSE mode), which provides sufficient material contrast between the matrix, the reinforcement, and auxiliary regions (e.g.

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pores and spinel/silicate precipitates). The procedure is exemplary shown in Fig. 7 for material 5ZAT showing the indexed particles. The nearest-neighbor-distance (see Table 8) of the MMC variants with 5 vol.% total reinforcement fraction reveals a uniform particle arrangement within the matrix among the

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materials tested. Naturally, the MMC with 10 vol.% total particle fraction (5ZAT) exhibited a lower particle distance. Nevertheless, the average radius of the detected particles corresponds with the initial particle size of both reinforcement types (cf. Table 3) which indicates the correct identification of single particles, even within the agglomerated regions. Suitable reinforcement dispersion involves a narrow and adapted particle size distribution of the ceramic powders to avoid clustering of the fine ceramic particles in the interstices of the coarser grained steel particles during powder mixing and processing. However, a coarser grained MgPSZ was not commercially available yet. Contrarily, high-energy milling of the powder mixtures would initiate the desired phase transformation in the metal and in the ceramic component already during sample preparation.

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identified

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Fig. 7 SEM (BSE) image of material 5ZAT after firing showing the matrix, pores, and the particles: red - (Mn, Mg)(Al, Ti, Cr)2O4 spinel (initial aluminium titanate), green - zirconia

Recipe APR NND

5Z 2.1 ± 1.9 14.7 ± 8.9

5AT 2.7 ± 1.8 14.1 ± 5.6

3.2 Mechanical properties

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Table 8 Average particle radius (APR) and nearest-neighbor-distance (NND), in µm 2.5ZAT 2.4 ± 2.1 11.4 ± 7.0

5ZAT 2.2 ± 1.8 9.8 ± 5.1

The flow stress and work hardening behaviors of the matrix and the various composite variants under quasi-static compressive deformation are shown in Fig. 8. The stress-strain level and the

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deformation behaviour of all specimens tested were mainly controlled by the characteristics of the matrix material. The matrix strain hardening was driven by the dislocation movement, by the accumulation of dislocations, the formation of stacking faults and deformation bands, and by the strain-induced α’martensite formation (TRIP effect) [27]. The work hardening rates, as shown in Fig. 8b, illustrate the

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strengthening effect of the ceramic particles in a certain range of strain followed by a decreasing work hardening rate at further deformation. The stress level in the composites is significantly higher than in the unreinforced matrix material at moderate deformation. The compressive yield strength at 0.2% plastic

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compressive strain was ≈70 MPa higher for the MMCs with 5 vol.% Mg-PSZ, whereas the other composite variants showed an increase of ≈55-60 MPa as compared with the 100 % steel specimens. This is in agreement with the results reported by Decker et al. [18] for similar materials prepared via SPS. The lower strengthening effect of the present specimens originates from the higher porosity in the matrices, the larger austenite grain size, and the distinctive reactions at the steel/zirconia interfaces. As a consequence of damage initiation and propagation, the compression stress of composite materials dropped below the values measured for the unreinforced specimen at higher strain levels, which depend on the type of ceramic and its volume fraction. However, the materials 5AT and 5ZAT revealed a pronounced strengthening effect under quasi-static compressive loading up to a deformation degree of approximately 45 %. In contrast, the strengthening effect of material 5Z was exhausted when exceeding 18 % compressive strain. Various phenomena, either individually or superimposed, control the failure mechanism

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ACCEPTED MANUSCRIPT of metal-matrix composites. Material 5Z exhibited appx. 20 times larger initial pores than the other materials tested. Pore growth under loading and the coalescence of these voids results in crack initiation and consequently the stress level decreases at a medium stage of deformation. Conversely, the highest total particle volume fraction within this study (10 % instead of 5 %) modifies the deformation behaviour of material 5ZAT. Dislocation movement is impeded due to geometric constraints caused by the reinforcing

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particles, which results in higher stresses and the apparent work hardening during early deformation stages [4,27]. Additionally, the stress concentration initiates the austenite to martensite transformation in the surrounding of the reinforcement at a sacrifice of ductility and decline of the stress level on further deformation. Comparing materials 5AT and 2.5ZAT at similar particle volume fractions and initial pore sizes reveals a distinct advance in the presence of aluminium titanate as compared with pure zirconia addition

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firstly indicating the meaning of the metal/ceramic interface bonding [30,59].

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Fig. 8. Mechanical behaviour of the compressed specimens indicating the reinforcing effect of Mg-PSZ and/or aluminium titanate on the compressive yield strength and the work hardening: a) flow stress curves;

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b) work hardening rates

A similar material behaviour was observed under quasi-static tensile loading (see Fig. 9). As expected, 100 % steel specimens exhibited the highest ultimate tensile strength and ductility (23 to 28 % fracture strain) among the materials tested, whereas all types of MMC showed a higher stress level until rupture as compared to the matrix material [4]. However, the effect of ceramic reinforcement on the MMCs’ tensile yield strengths is negligible small. Similar to compressive deformation, the highest ultimate tensile strength were obtained for the materials 5ZAT and 5AT at the expense of their ductility. The highest fracture strain of the composite variants was measured for the material with additions of 2.5 vol.% both zirconia and initial aluminium titanate (18 to 20 % fracture strain). The particle strengthening effect of aluminium titanate exceeds the less pronounced transformation toughening mechanism of zirconia in the present study. Further improvements of zirconia reinforcement require elevated proportions of metastable

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ACCEPTED MANUSCRIPT tetragonal ZrO2 even after firing. The long thermal impact during conventional sintering, which is required for future cellular structures, restricts the processibility of the zirconia reinforced composites as it

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destabilizes the transformable tetragonal phase.

Fig. 9. Tensile deformation behaviour of TRIP steel and the composite variants showing the strengthening and strain-depressing effect of certain ceramic fractions: a) flow stress curves; b) work hardening rates As already mentioned before, the conventional pressure-less sintering process facilitates chemical reactions, segregations, and the formation of δ-ferrite and α’-martensite both in the unreinforced TRIP steel

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specimens and in the composite variants during firing [55]. Thus, the ferromagnetic volume fraction in the as-fired state strongly depends on the material’s composition and particularly on the presence of primary aluminium titanate in the material variants (see Fig. 10). The in situ tests revealed the kinetics of straininduced α’-martensite during deformation. The S100 specimens exhibited the highest overall percentage of

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strain-induced α’-martensite at rupture. The lower formation of strain-induced α’-martensite in the composite variants 5Z/2.5ZAT was in accordance with the less pronounced sigmoidal shape of their strain hardening response. The reduction of the austenite stabilizing alloying element Mn was caused by its

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evaporation, diffusion, and the formation of phases containing Mn at the metal/ceramic interfaces like (Mg, Mn)Si2O4 and/or (Mn, Mg)(Al, Ti, Cr)2O4, and led to a distinct shift of the initial metallic phase composition (cf. Table 5). This mechanism is more pronounced in MMC with aluminium titanate as compared with steel/zirconia mixtures. Since the degree of the expected phase transformation of t-ZrO2 to m-ZrO2 is low [28,53], the reinforcing mechanism of the ceramic particles obviously depends on the bonding strength of the metal/ceramic interface [59], which is apparently more reliable in material variants with aluminium titanate. Thus, the evolution of strain-induced α’-martensite and the work hardening rate at low strain values is higher in these MMCs. Furthermore, the maxima of the work hardening rates in the MMCs are shifted to lower deformation degrees as compared with the non-reinforced steel. This is due to the more pronounced strain hardening of the steel matrices in the MMCs imposed by the presence of ceramic particles developing an early saturation level of the dual-phase austenite/martensite microstructure [4].

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beginning

of

plastic

deformation.

Then,

the

constrained

matrix

of

the

work

hardened

austenitic/martensitic steel suffers from the lack of strain relaxation, which promotes crack initiation and

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propagation.

Fig. 10. Evolution of the ferromagnetic phase fraction under tensile deformation at room temperature, normalized by the metal fraction in the material 3.3 Microstructure after deformation

The microstructure of the specimens after tensile deformation tests was analysed in order to explain

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the mechanical behaviour as well as the characteristic failure mode of the material (see Fig. 11). The small strengthening effect of zirconia in the MMCs is a result of the loose embedding of the Mg-PSZ particles, of the irregular formation of silicates as an intermediate layer between steel/zirconia and zirconia/zirconia particles, and of the lack of pronounced metastable ZrO2 phase regions. The zirconia agglomerates are broadly crushed and fissured after tensile deformation; considerable delamination occurred at the

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steel/ceramic interface, which indicates loose interface boundaries [4,30]. Thus, sufficient load transfer from the TRIP steel matrix can be assumed. Still, the sintering temperature of the composite material was

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well below the normal sintering temperature of Mg-PSZ, thus the clustered particles are expected to fail at lower stresses as compared with the commercially available bulk material. Since the majority of the zirconia particles exhibited the non-transformable monoclinic crystal structure already after sintering, the probability of the desirable tetragonal to monoclinic phase transformation is low. The significance of minute zirconia phase transformations on the mechanical properties of similar MMCs at low deformation (εtrue≤ 0.25) has already been reported in [27,28,60]. Taking the higher stress level in the material 5Z as compared with the 100 % steel and the contrary effect, which is the retarded α’-martensite formation in the composite’s matrix into account, the reinforcing mechanism of the zirconia particles is most likely assisted by the expansive tetragonal to monoclinic ZrO2 phase transformation. However, a large amount of voids in the ceramic sections, caused by broken-out particles, prevents the material from a detailed and representative phase analysis. It is obvious that the zirconia particles with a silicate intermediate layer and these silicate

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ACCEPTED MANUSCRIPT structures itself remained reliably bonded in the matrix material, whereas a large amount of zirconia particles was broken out after specimen preparation. The tensile deformed microstructure of MMCs doped with 5 vol.% initial aluminium titanate powder is characterised by the spinel particles that maintained their pronounced interface bonding to the steel matrix despite the extended multiple-crack formation in the interior of the ceramic particles. In contrast, debonding

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was the dominant damage initiation mechanism in previous TRIP steel/Mg-PSZ composites [3,30,61]. The intact particle/matrix interfaces indicate a higher shear strength at the interface than the particle fracture strength which is a further indication for the improved interface-related performance of the MMCs with initial Al2TiO5 [4,59]. Still, some of the small grained spinel particles were nearly intact and exhibited neither cracking nor debonding. Although these composites underwent internal damage by crack initiation and the ceramic particles show no phase transformation or decomposition as compared with composite materials

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comprising metastable zirconia particles [27,28], the addition of aluminium titanate results in a higher strain-induced α’-martensite formation and these MMCs sustain a higher stress level than the pure matrix

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material. Conversely, the addition of alumina particles in similar MMCs has shown no reinforcing effect on the mechanical properties of the composites due to the absence of a pronounced interface bonding [27]. Thus, the interfacial (de-)bonding between the TRIP steel and the ceramic particles is more essential for the mechanical properties of the MMCs than the phase transformation in zirconia. The microstructure of the composite materials with concurrent additions of Mg-PSZ and aluminium titanate largely complied with the observations of 5Z and 5AT and the specific features for each type of ceramic particles. Here, the zirconia particles were subject to intense cracking but less to debonding due to

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the formation of the Ti-enriched spinel interlayer between the ZrO2 particles and the steel matrix. The mutual interference of the zirconia particles and the initial aluminium titanate particles was apparently more pronounced in the MMCs with the higher ceramic volume fraction due to the coincidentally distribution of

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these particles within the steel matrix.

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Fig. 11. SEM image of the composite materials at fracture strain showing broadly crushed ceramic particles

4 Conclusions

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and isolated cracks in the surrounding steel matrix: a) 5Z; b) 5AT; c) 2.5ZAT; d) 5ZAT

The effect of additions of Mg-PSZ and/or aluminium titanate on the microstructure, on the transformation characteristics and on the mechanical behavior of TRIP-matrix composites was investigated. The non-deformed matrix material showed a highly austenitic microstructure with minor

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amounts of α’-martensite and δ-ferrite. Generally, the ceramic particles were homogeneously dispersed within the metal matrix with a certain degree of agglomeration. To improve the particle dispersion, the size

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distribution of the metal and the ceramic component should be narrow and compliant to avoid clustering of fine reinforcing particles. The diffusional interchange of the alloying elements of the steel and the stabilising agent of the zirconia led to the formation of a (Mg, Mn)Si2O4 olivine structure during firing. Intense reaction diffusion between the steel matrix and the aluminium titanate particles led to the decomposition of the aluminium titanate and the formation of a (Mn, Mg)(Al, Ti, Cr)2O4 spinel structure during firing. This phenomenon provides a more dense and less fissured particle/matrix interface than in the case of zirconia addition. In early deformation stages, the deformation behaviour of the materials was mainly controlled by initial dislocation hardening of the austenitic matrix. The presence of zirconia and/or phases derived from the initial aluminium titanate improved the stress-level under tensile and under compressive loading. At higher deformations, α’-martensite formed in the steel matrix, which further increased the strength of MMCs. Finally, all composite materials failed at a lower strain than the unreinforced matrix material by internal damage after cracks have initiated in the steel matrix surrounding the ceramic particles. The

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ACCEPTED MANUSCRIPT improved performance of MMCs is crucially associated with the particle strengthening mechanism of the ceramic particles, and which depends on the interface formation between the metal and ceramic component that proceeded during sintering. Because of the predominantly monoclinic crystal structure of ZrO2 and the weak bonding at the steel/zirconia and zirconia/zirconia interfaces the addition of zirconia particles resulted in a moderate particle strengthening of the TRIP-matrix composites at low deformation

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(<20 %), both under compressive and tensile loading. The composite material with aluminium titanate offered a significant strengthening effect and a distinctive strain hardening under quasi-static compressive (up to 45 % strain) and tensile (until rupture) loading compared with the matrix material, also involving an elevated rate of α’-martensite formation in the early stage of plastic deformation. The void-free interface between the steel and the spinel structures of the initial microstructure offered optimal conditions for the load transfer; the intergranular bonding strength and the toughness of the metal/ceramic interfaces are

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assumed to be higher than in composites with Mg-PSZ particles. The combination of Mg-PSZ and Al2TiO5 further improved the stress level of the MMCs during both tensile and compressive deformation.

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Consequently, the fabrication of TRIP-matrix composites with addition of aluminium titanate via the conventional sintering process provides advanced materials for high mechanical load applications at ambient temperature, e.g. cellular structures for crash-absorber that cannot be fabricated by other techniques. Acknowledgements

The authors gratefully acknowledge the financial support of the German Research Foundation (DFG)

References

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for funding the research project under the Collaborative Research Center 799: TRIP-Matrix Composites.

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ACCEPTED MANUSCRIPT Compressive and tensile deformation behaviour of TRIP steel-matrix composite materials with reinforcing additions of zirconia and/or aluminium titanate Authors: Weigelt, C. 3)

1,*)

1)

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2)

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; Schmidt, G. ; Aneziris, C. G. ; Eckner, R. ; Ehinger, D. ; Krüger, L. ; Ullrich,

3)

C. ; Rafaja, D.

• Composite materials based on a TRIP-steel matrix with ceramic reinforcing particles • Novel MMCs with combinations of zirconia and aluminium titanate in MMCs

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• High interfacial bonding strength in composites variants with aluminium titanate

• Improved mechanical characteristics under tensile and compressive deformation

• Characteristics of MMCs suitable for future high load applications/ crash absorber

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Kind regards,

AC C

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M AN U

Christian Weigelt