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Materials Science & Engineering A 761 (2019) 137974
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Influence of high-energy ball milling on Mg-PSZ-reinforced TRIP steelmatrix composites synthesized by FAST / SPS
T
S. Decker∗, L. Krüger Institute of Materials Engineering, TU Bergakademie Freiberg, Gustav-Zeuner-Straße 5, 09599, Freiberg, Germany
A R T I C LE I N FO
A B S T R A C T
Keywords: TRIP-steel Mg-PSZ Spark plasma sintering Phase transformation
A composite with a high-alloy steel matrix was reinforced with varying amounts of Mg-PSZ and synthesized using Spark Plasma Sintering (SPS) technology. To prepare the powders for sintering, they were mixed in a planetary ball mill at two different rates of 100 rpm and 250 rpm. The influence of the rotation speed on the microstructures and compressive strengths of the sintered composites was investigated. Due to the high degree of deformation of the steel powder particles during milling at 250 rpm, the steel phase underwent grain refinement during SPS. Furthermore, the Mg-PSZ was more homogeneously distributed within the steel matrix and mechanically interlocked with the steel matrix. Thus, the compressive strength of the composite increased with increasing Mg-PSZ content and increasing rotation speed during milling of the powder due to grain refinement and the improved strength of the steel/ceramic interface. Furthermore, a larger amount of Mg-PSZ underwent stress-induced transformation under compressive loading.
1. Introduction The demand for high-performance materials is growing, and light materials with high strength and sufficient ductility are increasingly specified for industrial applications. There are several options in attempting to satisfy such requirements. On the one hand, the chemistry of the material is important, and a promising option is to combine several materials into a composite. This provides the opportunity to integrate the best properties of each component into one material. Composites consisting of steel matrices and partially stabilized zirconia have been investigated for many years [1–5]. In particular, composites with a TRIP steel matrix (TRansformation Induced Plasticity) reinforced with Mg-PSZ (MgO partially stabilized zirconia) are the focus of considerable research efforts [1,6–8]. TRIP steels exhibit a metastable austenitic phase that has the capacity to undergo deformation-induced transformation to α′-martensite [9–14]. Due to this phase transformation, such steels exhibit both high strength and high ductility. Mg-PSZ exhibits tetragonal, cubic and, to some extent, monoclinic phases at room temperature. When placed under stress, the tetragonal phase can transform into the monoclinic phase. This phase transformation is associated with a certain volume expansion, and is exploited to improve the toughness of ceramics [15]. In steel-matrix composite materials, this effect is used to increase the strength of the material [16].
∗
On the other hand, the properties of a material can be influenced by the technology used to produce the material. To obtain a material with high strength, a fine-grained microstructure is favored. If a powder metallurgical route is used, Spark Plasma Sintering (SPS) offers the facility to densify powders within only a few minutes [17–19]. Such technologies, therefore, leave insufficient time for diffusion-driven processes such as grain growth to unfold fully. The SPS technology is similar to hot pressing: The powder is filled into a graphite die before subsequent uniaxial compression during sintering. With this method, heat is generated within and/or close to the powder in the punches and the die by means of a direct current that passes through the punches, the powder and/or the die [17]. Composites consisting of partially stabilized zirconia as a reinforcement and steel as their matrix have been successfully sintered using SPS [16,20]. Due to the high heating rates and the short dwell time, these materials exhibit fine-grained steel matrices and high compressive strengths [16,20]. However, the fine Mg-PSZ particles tend to form clusters in the gaps between the much coarser steel particles. These clusters are the initiation points for failure under loading, and must be avoided. The tendency to form clusters increases with (a) an increasing amount of reinforcing phase; (b) a decreasing reinforcingphase particle size in comparison to the particle size of the matrix; and (c) with increasingly irregular particle geometries in the reinforcing phase. Furthermore, the tendency to form clusters depends on the
Corresponding author. E-mail addresses: [email protected] (S. Decker), [email protected] (L. Krüger).
https://doi.org/10.1016/j.msea.2019.05.104 Received 8 February 2019; Received in revised form 28 May 2019; Accepted 29 May 2019 Available online 30 May 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 761 (2019) 137974
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particle size distribution and on the equipment used to mix the composite powder [21,22]. Failure also occurs at steel/ceramic interfaces [16]. One possibility to achieve more homogeneous particle size distribution is high-energy ball milling. As recent studies show, the formation of clusters can be reduced by high-energy ball milling of the powder before SPS. Additionally, this method can be used to increase the degree of interlocking between steel and ceramic [23]. Furthermore, grain refinement and hardness can also be improved by highenergy ball milling [23–26]. Deirmina et al. demonstrate the homogeneous distribution of the reinforcing phase, an increase in hardness, and the distinct refinement of grains after high-energy ball milling and subsequent SPS for a composite consisting of 10 vol% and 20 vol% MgPSZ as reinforcing phases in combination with a tool steel [26]. However, the deformability, the strength and the possibility of stress-induced phase transformation of and in these composites were not investigated. Sufficient ductility and stress-induced phase transformation was reported for a TRIP matrix composite reinforced with 5 vol% MgPSZ. Yet only a slight increase of 60 MPa in the 1% offset compressive yield strength was achieved by the improved distribution of the reinforcing phase associated with milling at 250 rpm [23]. A distinct effect on compressive strength and stress-induced phase transformation is to be expected for larger quantities of reinforcing phase. In composites with low fractions of reinforcing phase, the probability of cluster formation is lower than in composites with high fractions of reinforcing phase. In the present study, Mg-PSZ-reinforced TRIP matrix composites with high fractions of reinforcing phase were prepared using high-energy ball milling and subsequent SPS. To clarify the effects of milling, the parameters of the milling process carried out prior to SPS were varied. Furthermore, two composites were prepared to investigate the influence of the ceramic content on the material strength – one with 5 vol% Mg-PSZ and one with 40 vol% Mg-PSZ, which probably exceeded the percolation threshold. Hence, one very brittle and one quite ductile composite were investigated. To ensure that a composite between these two extremes was also tested, a composite powder with 10 vol% Mg-PSZ was also mixed at high rotation speed. All of the composites produced were investigated with respect to their ductility and strength.
Table 2 Chemical composition of the steel powder. [wt.%]
Mn
Ni
Si
N
C
Al
Fe
16.4
6.31
6.30
0.96
0.06
0.03
0.1
Bal.
Table 3 Particle size of the initial powders, as determined by laser diffraction analysis. Powder
d10 [μm]
d50 [μm]
d90 [μm]
Steel Mg-PSZ
7.4 0.1
20.7 1.3
41.3 10.8
powders and to prevent deposition of the PCAs inside the SPS apparatus. The term ‘milling’ implies the refinement of powder particles. However, particle refinement was not emphasized in either case (100/ 250 rpm). The term ‘milling’ is used in the following in correlation to the equipment utilized in this study. To simplify the following description of results, the composite names were shortened according to their Mg-PSZ content and the rotation speed used during milling (xCy: x = vol% Mg-PSZ, y = milling velocity in rpm) with 5C100, 5C250, 10C250, 40C100 and 40C250. For comparison, an unreinforced steel sample was prepared and abbreviated with ‘S’. The composite powders were sintered under vacuum conditions in an SPS device (HP D 25 from FCT Systeme). All of the powders were heated at 100 K/min to 1100 °C and held at this temperature for 5 min. During the heating cycle, the pressure to uniaxially compress the powder was increased linearly to 51 MPa and held at 51 MPa during the dwell time. The density of the sintered samples was measured using Archimedes method. Based on these density values, the densification during sintering was calculated by applying equations (1) and (2) [27]. In this approach, the increase in density was correlated with the punch movement during sintering. However, such punch movement is a function of shrinkage, thermal expansion and the elastic deformation of powder and punches. To identify the fraction of punch movement due to thermal expansion and elastic deformation, the punch movement was recorded during heating of a dense sample while a: applying a constant pressure, and b: while increasing pressure. Accordingly, the change in sample height due to shrinkage Δli was calculated. The shrinkage εzi was determined by dividing the change in sample height by the initial sample height l0. The density present during sintering and the initial density were described by ρi and ρ0, respectively.
2. Materials and methods A steel powder which was gas atomized by TLS Technik GmbH & Co. Spezialpulver KG was used as the matrix material. A zirconia powder partially stabilized with 3.5 wt% MgO (Mg-PSZ) from Saint Gobain was utilized as a reinforcement phase. The chemical compositions of the initial powders and the particle-size distributions are shown in Tables 1–3. The steel powder was fully austenitic, with 10 vol% of the Mg-PSZ exhibiting the monoclinic phase, 53 vol% the cubic and 37 vol% the tetragonal phase. Based on these initial materials, composite powders were prepared that contained 5 vol%, 10 vol% and 40 vol% Mg-PSZ. To this end, the powders were milled in the appropriate ratios for 4 h using a planetary ball mill (Pulverisette 6, Fritsch GmbH) and a ball-mass to powder-mass ratio of 5:1. The milling balls were made of steel and had a diameter of 10 mm. Some of the powders were milled in isopropanol at 100 rpm to produce a thorough mixture. Further composite powders were prepared by dry milling at 250 rpm to bring about intense mechanical interlocking between the powder particles. Process control agents (PCAs) were not used to avoid any change in the chemical composition of the
ρi = ρ0 ⋅e εzi
(1)
εzi = Δl i / l 0
(2)
Samples of 6 mm in height and 6 mm in diameter were machined for quasi-static compression tests. The samples were compressed at a strain rate of 10−3 s−1 to a strain value of 60% if fracture did not occur at lower degrees of deformation. The microstructure of both the as-sintered material and the compressed samples was investigated by X-ray diffraction, light optical microscopy, optical inclusion analysis using AnalySIS Particle Inspector software, and scanning electron microscopy (SEM) in combination with electron backscatter diffraction (EBSD). From EBSD analysis, the grain size of the steel matrix was determined by means of the Line Intercept Method. 3. Results and discussion
Table 1 Chemical analysis of the Mg-PSZ powder. [wt.%]
Cr
3.1. Initial material
SiO2
MgO
Al2O3
CaO
TiO2
Na2O
Fe2O3
ZrO2
4.23
3.37
0.63
0.21
0.14
0.09
0.01
Bal.
Fig. 1 shows the initial powders as well as the powders after milling. High-energy ball milling at 250 rpm led to intensive deformation of the steel powder and mechanical interlocking between the steel and 2
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Fig. 1. Initial powders: a) TRIP steel; b) Mg-PSZ; composite powders after ball milling: c) 5C100; d) 5C250; e) 5C250, ground particle; f) 40C250.
the results obtained for a composite with a similar steel [23], the powder mixture prepared at 250 rpm exhibited very coarse powder particles (Fig. 1) and, therefore, a low apparent density at the onset of densification (Fig. 3). However, the densification rate of the powder mixed at 250 rpm was enhanced, which was illustrated by the higher slope in Fig. 3. Hence, the densification process was almost finished at 1000 °C for 5C250, while 5C100 was almost dense at 1100 °C. The increased densification rate resulted from the higher number of vacancies and dislocations due to high-energy ball milling [23]. Thus, the driving force for the densification was enhanced. These results show that it is possible to decrease the sintering temperature necessary for such composites by increasing the rotation speed during milling. Due to deformation during milling, recrystallization took place during sintering and resulted in grain refinement (Fig. 4). The average grain size of the steel matrix is shown in Table 4. It was obvious that the grain growth during SPS was constrained by the ceramic particles. Hence, the unreinforced steel sample exhibited much larger grains than the composites. For composites mixed at 100 rpm, however, the grain size did not vary with increasing ceramic content. Due to the high degree of deformation combined with the recrystallization process, the grain size of the composites mixed at 250 rpm was slightly smaller and decreased with decreasing ceramic content. Ceramic particles are rigid and absorb impact energy from milling balls during milling. Therefore, the steel particles experienced less deformation with increasing ceramic content. Hence, the grains were coarser with increasing ceramic content after milling at 250 rpm and subsequent SPS. Optical inclusion analysis indicated a median particle diameter for Mg-PSZ of 2.5 μm in the 5C250 material. Thus, the median reinforcement diameter was 1 μm larger than d50 of the Mg-PSZ. Therefore, intensive refinement of the Mg-PSZ due to milling can be ruled out. After SPS, the steel phase was completely austenitic, while the monoclinic content of the Mg-PSZ phase increased with respect to the milled powder in composites 5C100 and 40C250 (Table 4). The slight decrease in monoclinic content in 40C100 was within the acceptable measurement error range. According to EBSD analysis (Fig. 4), some 40%–60% of the Mg-PSZ exhibited the monoclinic phase. In particular, the smaller Mg-PSZ particles and the edges of the coarse Mg-PSZ
ceramic powder particles. Furthermore, the initial steel particles were kneaded together, and formed larger particles due to cold welding (Fig. 1e and f). However, the milling process at 100 rpm resulted merely in a homogeneous distribution of the powders without considerable deformation of the powder particles (Fig. 1c). Due to the deformation during milling, the austenitic steel was transformed to α′-martensite to a certain degree. However, the α′-martensite was fully converted to austenite during SPS. XRD did not indicate any phase transformation within the Mg-PSZ due to milling for the composite powders 5C100, 5C250 and 10C250. Approximately 8% of the Mg-PSZ transformed to the monoclinic phase during milling for 40C100, and 15% of the MgPSZ in 40C250. The ball milling parameters strongly affected the microstructure of the composite after SPS (Fig. 2). Due to the low rotation speed of 100 rpm, the originally spherical steel particles were clearly visible in the micrographs of the sintered samples (Fig. 2a and c), where these steel particles were now framed by Mg-PSZ. However, the Mg-PSZ formed clusters in the gaps of the initial steel particles. If the composite contained 40 vol% Mg-PSZ, the Mg-PSZ particles formed a continuous network. Thus, it was not expected that a direct current would run through the sample during SPS. Voids were visible within the clusters to some extent (Fig. 2e). The porosity measured within the composites was very low, with pores making up less than 2% of the samples. Thus, these voids within the clusters were locations where ceramic particles had fallen out. The sintering temperature applied was too low to allow sintering of the MgPSZ. Thus, the interfacial bonding between the Mg-PSZ particles was weak. If the composite powder was milled at 250 rpm, however, the number of clusters was reduced considerably. Furthermore, the Mg-PSZ particles were very well embedded in the strongly deformed steel matrix (Fig. 2f). However, a longer milling time would have been necessary to have achieved perfect homogeneity. Large areas of formerly coarse steel particles without the appearance of Mg-PSZ were visible in Fig. 2b, d and f. Nevertheless, the number of Mg-PSZ/Mg-PSZ interfaces was reduced significantly. Fig. 3 illustrates the densification behavior for the composite with 5 vol% of Mg-PSZ with increasing temperature during SPS. Similar to 3
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Fig. 2. Microstructure of the composites after SPS (light phase: steel, dark phase: Mg-PSZ); a) 5C100, b) 5C250, c) and e) 40C100 (white arrows indicate voids), d) and f) 40C250.
destabilized during SPS. In this study, however, XRD analysis did not find any monoclinic content in 5C250 and 10C250, which was contrary to the results shown in Fig. 4. Even with EBSD, it was not possible to indicate all ceramic particles (dark grey areas in Fig. 4), which may have been due to the strong deterioration of the lattice structures of the ceramic particles during the intensive milling process. 3.2. Deformation behavior If the composites with 40 vol% Mg-PSZ were not taken into account, all of the other composites exhibited high degrees of ductility and were deformable to nominal engineering compressive strains of 60% without failure (Fig. 5). Because the percolation threshold was exceeded in the composites with 40 vol% Mg-PSZ, these composites were very brittle under compressive deformation, while their degree of deformability would have been reduced even further under tension. Cracks occurred on the radial surfaces of the samples mixed at 250 rpm. The composites had a higher 1% offset compressive yield strength (σ1%) than the unreinforced steel. Due to the small grain size of the steel resulting from the SPS process, the steel exhibited a high σ1%, which was twice as high as the value for the cast steel type [31]. This effect is well known from the Hall-Petch relation [32]. However, σ1% deviated by approximately 70 MPa from the published results (Fig. 6), where the Hall-Petch relation was applied for the steel alloy used [30]. The σ1% increased further through the addition of the ceramic particles. Taking into account the deviation of the σ1% of the steel by 70 MPa and assuming that the slope of the Hall-Petch relation was in accordance with the slope given by Ref. [30], it was possible to calculate the strength σS of the steel matrix of the composite as a function of the grain size (Fig. 6, Table 5). It was obvious that the refinement of the steel grains caused an increase in the σ1%. The remaining amount of strength was related to the ceramic reinforcement. Assuming that the increase in strength followed the rule of mixture (equation (3)), the increase in strength σMg-PSZ caused by the
Fig. 3. Densification with increasing temperature during SPS for 5C100 and 5C250.
particles were of the monoclinic phase. This was due to destabilization of the Mg-PSZ during SPS [28,29]. In this case, the Mg ions diffused from the ceramic particles into the steel matrix. However, it is known that the results of XRD analysis can deviate strongly from EBSD analysis [16]. According to XRD measurements, the destabilized Mg-PSZ content after SPS seemed to decrease with increasing Mg-PSZ content and lower rotation speed during milling (Table 4). If a high milling speed was used, the Mg-PSZ particles were better distributed within the steel matrix, and fewer Mg-PSZ clusters were formed. Thus, more steel/ ceramic interfaces existed if a high rotation speed was selected than if a low rotation speed was used. Hence, the interfacial area where Mg-ions could leave the Mg-PSZ during SPS increased. In Ref. [23], it is reported that due to the higher proportion of steel/ceramic interfaces with increasing rotation speed during milling, a larger amount of Mg-PSZ is 4
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Fig. 4. Grains within a) the steel; followed by the steel matrix of the composites and indexing of phases by EBSD after SPS: b) 5C250, c) 5C100, d) 10C250, e) 40C100, f) 40C250 (grey: austenite; green: cubic zirconia, yellow: tetragonal zirconia, red: monoclinic zirconia). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.) Table 4 Microstructural characteristics; monoclinic content of the Mg-PSZ determined by XRD. Material
S 5C100 5C250 10C250 40C100 40C250
Relative Density [%]
97.6 98.8 98.0 98.6 98.1 98.8
Average Grain Size [μm]
5.6 2.4 1.4 1.8 2.5 2.0
Monoclinic Content of Mg-PSZ [%] Before SPS
After SPS
– 0 0 0 18 25
– 28 0 0 15 41
reinforcement could be calculated (Table 5). If the ceramic content increased, the amount of strength σMg-PSZ that was due to the ceramic reinforcement increased as well. Furthermore, the rotation speed during milling and the combined homogeneous distribution of the ceramic particles influenced the σ1% value. This effect was more pronounced for high ceramic contents. While an increase in the rotation speed during milling led to a rise in strength due to the ceramic of Δσ = 534 MPa in composites with 40 vol% ceramic material, the strength increased by approx. 20 MPa in composites with 5 vol% of ceramic powder. Due to the higher rotation speed, the ceramic particles were distributed more homogeneously within the steel matrix and mechanically interlinked with the steel. Furthermore, the amount of weak ceramic/ceramic interfaces was reduced dramatically. In
Fig. 5. Compressive deformation behavior.
addition, the application of the high rotation speed may have resulted in refinement of the Mg-PSZ particles. Thus, the number of reinforcing particles increased and the mean distance between them decreased, which served to improve the degree of dispersion strengthening. Hence, these effects were more distinct in composites with a high ceramic content and influenced the σ1% value.
σ1% = VS⋅σS + (1 − VS)⋅σMg − PSZ
(3)
As indicated by the slopes of the true compressive stress – 5
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40C250 exhibited a higher strength than 40C100. However, the cracks along the steel/ceramic interfaces in this composite ran along the network of reinforcement, were very long and were not stopped by the steel phase. Hence, the 40C250 material failed at a lower compressive deformation value. Due to the low deformability of 40C100 and 40C250, hardly any α′-martensite was formed. The 10C250 composite did not exhibit a dense network of reinforcement, such that cracks at interfaces were very short and were stopped by the steel matrix. Due to the high amount of steel, the composite exhibited a high degree of deformability. However, the amount of ceramic and the interfacial strength were high enough to facilitate load transfer to the ceramic particles and, therefore, to ensure the high strength of the composite. Hence, stress-induced phase transformation and deformation-induced phase transformation took place. 4. Conclusions Fig. 6. Calculated strength σS of the steel matrix and Hall-Petch fit, as reported by Ref. [30].
The results presented demonstrate that high-energy ball milling has the potential to improve the mechanical properties of Mg-PSZ-reinforced TRIP-matrix composites. Due to the high energy impact, the ceramic particles were more homogeneously distributed and were kneaded into the steel matrix. Therefore, fewer ceramic clusters appeared in the composite. In future research, the milling time should also be increased to improve the distribution of the Mg-PSZ particles further. Due to the absence of PCAs, coarse particles remained after high-energy ball milling. However, the sintering process was not negatively influenced by the coarse particles. Due to the interlocking between the Mg-PSZ and the steel phase, the transmission of the mechanical load from the steel matrix to the ceramic reinforcement improved. This effect led to an increase in compressive strength and was more pronounced with increasing ceramic contents up to 40 vol%. Furthermore, the steel was strongly deformed during high-energy ball milling and recrystallized during SPS. The resulting fine-grained steel matrix improved the compressive strength further. However, the grain refinement was diminished by large amounts of the reinforcing phase. Due to the fine-grained steel matrix and the reduced number of ceramic clusters, therefore, the compressive strength was improved without a distinct reduction in deformability – except for an Mg-PSZ content of 40 vol%. When a content of 40 vol% Mg-PSZ was used as the reinforcement, the percolation threshold was exceeded. Thus, the composite became very brittle and failure occurred – primarily in the Mg-PSZ phase. In contrast, an Mg-PSZ content of 5 vol% seemed to be low, but increased σ1% by approx. 100 MPa. Due to the fine-grained steel matrix, Mg-PSZ made a significant contribution to compressive strength, and by not exceeding the percolation threshold, the composite 10C250 exhibited a high σ1%, a high degree of deformability and a large degree of stress-induced and deformation-induced phase transformation. It was assumed that the ceramic content could have been increased further without exceeding the percolation threshold.
Table 5 1% offset compressive yield strength σ1% and calculated contribution of reinforcement phase as a function of the volume fraction of steel Vs. Material [rpm]
σ1% [MPa]
Strength σs of Fine-Grained Unreinforced Steel Matrix [MPa]
Increase in Strength (1-VS)∙σMg-PSZ due to Mg-PSZ [MPa]
S 5C100 5C250 10C250 40C100 40C250
400 540 636 681 776 1328
400 488 568 528 483 513
0 76 96 206 486 1020
logarithmic strain curve, the degrees of work hardening of 5C100 and 5C250 were similar, and they exhibited the same type of damage. As shown in Fig. 7, it was primarily the steel matrix that deformed. In Fig. 7d, deformation bands are clearly visible (marked by the white ellipse), and these deformation bands were the sites where α′-martensite was formed. The α′-martensite nuclei constrained the dislocation movement in a manner similar to grain boundaries and, therefore, intensified the work-hardening effect [30]. Furthermore, delamination occurred at the steel/ceramic interfaces (marked by white arrows in Fig. 7). Due to the high deformability of these composites, a large amount of α′-martensite was formed (Table 6). Furthermore, the ceramic material underwent stress-induced transformation (Table 6). In contrast, the deformation behavior of the composites with 40 vol % Mg-PSZ varied. Due to the exceedance of the percolation threshold in these composites, the damage was concentrated within the ceramic material (Fig. 7e), and it was mainly large ceramic particles and clusters that fractured. Due to the sintering temperature, which was quite low for zirconia, the Mg-PSZ clusters exhibited low strength and failed at a low degree of deformation of the composite. In some regions, deformation of the steel phase was visible. The cracks in the ceramic particles and clusters were stopped by the ductile steel matrix. Despite the failure of the Mg-PSZ particles, stress-induced phase transformation took place (Table 6). Furthermore, the strength of the interface between the steel phase and the ceramic phase was weak due to stress caused by their differences in thermal expansion during cooling. Thus, delamination occurred at these interfaces. In 40C250, some deformation bands formed during deformation within the steel phase. Due to a smaller number of clusters and large particles, their fracture was not the main type of damage, as it was mainly the steel/ceramic interfaces that failed. Furthermore, fewer Mg-PSZ particles experienced stress-induced phase transformation (Table 6). However, the strength of the interfaces seemed to be higher than the strength of the ceramic clusters. Hence,
Data availability The raw/processed data required to reproduce these findings cannot be shared at this time, as the data also forms part of an ongoing study. Acknowledgements The authors would like to thank the German Research Foundation (DFG) for supporting the investigations, which were part of the Collaborative Research Center TRIP-Matrix Composites (SFB 799, subproject A06). Furthermore, the authors would like to thank the Institute of Ceramic, Glass and Construction Materials and the Institute of Iron and Steel technology at TU Bergakademie Freiberg for the chemical characterization of the Mg-PSZ and steel powders. 6
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Fig. 7. Deterioration in the microstructure of the composites after 60% compressive deformation, a) 5C100, b) and d) 5C250, c) 10C250; and after failure, e) 40C100, f) 40C250 (light grey: ceramic particles, dark grey: steel, black: cracks and holes). Table 6 Phase contents due to deformation up to 60% or until fracture; amount of monoclinic content refers only to the ceramic material, and not the entire composite. Material
α′-Martensite [%]
Stress-Induced Monoclinic Content [%]
S 5C100 5C250 10C250 40C100 40C250
50 51 47 42 5 0
– 22 57 77 33 10
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[24] J. Yang, Z. Liu, R. Yang, Q. Zhi, M. Yang, Y. Zhang, E. Liang, Influence of highenergy milling and heat treatment on the thermal and mechanical properties of a ZrMgMo3O12p/2024Al composite, Ceram. Int. 45 (10) (2019) 13464–13471. [25] H. Kotan, Microstructural evolution of 316L stainless steels with yttrium addition after mechanical milling and heat treatment, Mater. Sci. Eng. A 647 (2015) 136–143. [26] F. Deirmina, M. Pellizzari, Production and characterization of a tool steel-PSZ composite by mechanical alloying and spark plasma sintering, J. Alloy. Comp. 709 (2017) 742–751. [27] M.N. Rahaman, Ceramic Processing and Sintering, Seconde Éd, CRC Press (Taylor & Francisr), Boca Raton, 2003. [28] C. Weigelt, S. Giersberg, C. Wenzel, C.G. Aneziris, Screening of the interactions between Mg-PSZ and TRIP-steel and its alloys during sintering, Adv. Eng. Mater. 12 (2010) 486–492. [29] S. Decker, K. Lange, T. Dubberstein, L. Krüger, Influence of the sulfur content on the mechanical properties of a TRIP/TWIP-Matrix composite reinforced with 10 vol% Mg-PSZ, Adv. Eng. Mater. 17 (2015) 1351–1356. [30] S. Martin, S. Wolf, S. Decker, L. Krüger, U. Martin, Deformation bands in high-alloy austenitic 16Cr6Mn6Ni TRIP steel: phase transformation and its consequences on strain hardening at room temperature, Steel Res. Int. 86 (2015) 1187–1196. [31] L. Krüger, S. Wolf, S. Martin, U. Martin, A. Jahn, A. Weiß, P.R. Scheller, Strain rate dependent flow stress and energy absorption behaviour of cast CrMnNi TRIP/TWIP steels, Steel Res. Int. 82 (2011) 1087–1093. [32] E.O. Hall, The deformation and ageing of mild steel: III Discussion of results, Proc. Phys. Soc. B 64 (1951) 747–753.
toughness relationship of zirconia toughened ceramics, Acta Metall. 33 (1985) 2083–2091. S. Martin, S. Richter, S. Decker, U. Martin, L. Krüger, D. Rafaja, Reinforcing mechanism of Mg-PSZ particles in highly-alloyed TRIP steel, Steel Res. Int. 82 (2011) 1133–1140. Z.A. Munir, U. Anselmi-Tamburini, M. Ohyanagi, The effect of electric field and pressure on the synthesis and consolidation of materials: a review of the spark plasma sintering method, J. Mater. Sci. 41 (2006) 763–777. O. Guillon, J. Gonzalez-Julian, B. Dargatz, T. Kessel, G. Schierning, J. Räthel, M. Herrmann, Field-Assisted sintering technology/spark plasma sintering: mechanisms, materials, and technology developments, Adv. Eng. Mater. 16 (2014) 830–849. R. Orru, R. Licheri, A.M. Locci, A. Cincotti, G. Cao, Consolidation/synthesis of materials by electric current activated/assisted sintering, Mater. Sci. Eng. R Rep. 63 (2009) 127–287. L. Krüger, S. Grützner, S. Decker, I. Schneider, Spark plasma sintering and strength behavior under compressive loading of Mg-PSZ/Al2O3-TRIP-steel composites, Mater. Sci. Forum 825–826 (2015) 182–188. W. Schatt, Pulvermetallurgie: Technologie und Werkstoffe, second., bearb. und erw. Aufl, Springer, Berlin, 2007. K.U. Kainer, Metal Matrix Composites: Custom-Made Materials for Automotive and Aerospace Engineering, Wiley-VCH, Weinheim, Chichester, 2010. S. Decker, L. Krüger, Improved mechanical properties by high-energy milling and Spark Plasma Sintering of a TRIP-matrix composite, J. Compos. Mater. 50 (2015) 1829–1836.
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Influence of high-energy ball milling on Mg-PSZ-reinforced TRIP steelmatrix composites synthesized by FAST / SPS
T
S. Decker∗, L. Krüger Institute of Materials Engineering, TU Bergakademie Freiberg, Gustav-Zeuner-Straße 5, 09599, Freiberg, Germany
A R T I C LE I N FO
A B S T R A C T
Keywords: TRIP-steel Mg-PSZ Spark plasma sintering Phase transformation
A composite with a high-alloy steel matrix was reinforced with varying amounts of Mg-PSZ and synthesized using Spark Plasma Sintering (SPS) technology. To prepare the powders for sintering, they were mixed in a planetary ball mill at two different rates of 100 rpm and 250 rpm. The influence of the rotation speed on the microstructures and compressive strengths of the sintered composites was investigated. Due to the high degree of deformation of the steel powder particles during milling at 250 rpm, the steel phase underwent grain refinement during SPS. Furthermore, the Mg-PSZ was more homogeneously distributed within the steel matrix and mechanically interlocked with the steel matrix. Thus, the compressive strength of the composite increased with increasing Mg-PSZ content and increasing rotation speed during milling of the powder due to grain refinement and the improved strength of the steel/ceramic interface. Furthermore, a larger amount of Mg-PSZ underwent stress-induced transformation under compressive loading.
1. Introduction The demand for high-performance materials is growing, and light materials with high strength and sufficient ductility are increasingly specified for industrial applications. There are several options in attempting to satisfy such requirements. On the one hand, the chemistry of the material is important, and a promising option is to combine several materials into a composite. This provides the opportunity to integrate the best properties of each component into one material. Composites consisting of steel matrices and partially stabilized zirconia have been investigated for many years [1–5]. In particular, composites with a TRIP steel matrix (TRansformation Induced Plasticity) reinforced with Mg-PSZ (MgO partially stabilized zirconia) are the focus of considerable research efforts [1,6–8]. TRIP steels exhibit a metastable austenitic phase that has the capacity to undergo deformation-induced transformation to α′-martensite [9–14]. Due to this phase transformation, such steels exhibit both high strength and high ductility. Mg-PSZ exhibits tetragonal, cubic and, to some extent, monoclinic phases at room temperature. When placed under stress, the tetragonal phase can transform into the monoclinic phase. This phase transformation is associated with a certain volume expansion, and is exploited to improve the toughness of ceramics [15]. In steel-matrix composite materials, this effect is used to increase the strength of the material [16].
∗
On the other hand, the properties of a material can be influenced by the technology used to produce the material. To obtain a material with high strength, a fine-grained microstructure is favored. If a powder metallurgical route is used, Spark Plasma Sintering (SPS) offers the facility to densify powders within only a few minutes [17–19]. Such technologies, therefore, leave insufficient time for diffusion-driven processes such as grain growth to unfold fully. The SPS technology is similar to hot pressing: The powder is filled into a graphite die before subsequent uniaxial compression during sintering. With this method, heat is generated within and/or close to the powder in the punches and the die by means of a direct current that passes through the punches, the powder and/or the die [17]. Composites consisting of partially stabilized zirconia as a reinforcement and steel as their matrix have been successfully sintered using SPS [16,20]. Due to the high heating rates and the short dwell time, these materials exhibit fine-grained steel matrices and high compressive strengths [16,20]. However, the fine Mg-PSZ particles tend to form clusters in the gaps between the much coarser steel particles. These clusters are the initiation points for failure under loading, and must be avoided. The tendency to form clusters increases with (a) an increasing amount of reinforcing phase; (b) a decreasing reinforcingphase particle size in comparison to the particle size of the matrix; and (c) with increasingly irregular particle geometries in the reinforcing phase. Furthermore, the tendency to form clusters depends on the
Corresponding author. E-mail addresses: [email protected] (S. Decker), [email protected] (L. Krüger).
https://doi.org/10.1016/j.msea.2019.05.104 Received 8 February 2019; Received in revised form 28 May 2019; Accepted 29 May 2019 Available online 30 May 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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particle size distribution and on the equipment used to mix the composite powder [21,22]. Failure also occurs at steel/ceramic interfaces [16]. One possibility to achieve more homogeneous particle size distribution is high-energy ball milling. As recent studies show, the formation of clusters can be reduced by high-energy ball milling of the powder before SPS. Additionally, this method can be used to increase the degree of interlocking between steel and ceramic [23]. Furthermore, grain refinement and hardness can also be improved by highenergy ball milling [23–26]. Deirmina et al. demonstrate the homogeneous distribution of the reinforcing phase, an increase in hardness, and the distinct refinement of grains after high-energy ball milling and subsequent SPS for a composite consisting of 10 vol% and 20 vol% MgPSZ as reinforcing phases in combination with a tool steel [26]. However, the deformability, the strength and the possibility of stress-induced phase transformation of and in these composites were not investigated. Sufficient ductility and stress-induced phase transformation was reported for a TRIP matrix composite reinforced with 5 vol% MgPSZ. Yet only a slight increase of 60 MPa in the 1% offset compressive yield strength was achieved by the improved distribution of the reinforcing phase associated with milling at 250 rpm [23]. A distinct effect on compressive strength and stress-induced phase transformation is to be expected for larger quantities of reinforcing phase. In composites with low fractions of reinforcing phase, the probability of cluster formation is lower than in composites with high fractions of reinforcing phase. In the present study, Mg-PSZ-reinforced TRIP matrix composites with high fractions of reinforcing phase were prepared using high-energy ball milling and subsequent SPS. To clarify the effects of milling, the parameters of the milling process carried out prior to SPS were varied. Furthermore, two composites were prepared to investigate the influence of the ceramic content on the material strength – one with 5 vol% Mg-PSZ and one with 40 vol% Mg-PSZ, which probably exceeded the percolation threshold. Hence, one very brittle and one quite ductile composite were investigated. To ensure that a composite between these two extremes was also tested, a composite powder with 10 vol% Mg-PSZ was also mixed at high rotation speed. All of the composites produced were investigated with respect to their ductility and strength.
Table 2 Chemical composition of the steel powder. [wt.%]
Mn
Ni
Si
N
C
Al
Fe
16.4
6.31
6.30
0.96
0.06
0.03
0.1
Bal.
Table 3 Particle size of the initial powders, as determined by laser diffraction analysis. Powder
d10 [μm]
d50 [μm]
d90 [μm]
Steel Mg-PSZ
7.4 0.1
20.7 1.3
41.3 10.8
powders and to prevent deposition of the PCAs inside the SPS apparatus. The term ‘milling’ implies the refinement of powder particles. However, particle refinement was not emphasized in either case (100/ 250 rpm). The term ‘milling’ is used in the following in correlation to the equipment utilized in this study. To simplify the following description of results, the composite names were shortened according to their Mg-PSZ content and the rotation speed used during milling (xCy: x = vol% Mg-PSZ, y = milling velocity in rpm) with 5C100, 5C250, 10C250, 40C100 and 40C250. For comparison, an unreinforced steel sample was prepared and abbreviated with ‘S’. The composite powders were sintered under vacuum conditions in an SPS device (HP D 25 from FCT Systeme). All of the powders were heated at 100 K/min to 1100 °C and held at this temperature for 5 min. During the heating cycle, the pressure to uniaxially compress the powder was increased linearly to 51 MPa and held at 51 MPa during the dwell time. The density of the sintered samples was measured using Archimedes method. Based on these density values, the densification during sintering was calculated by applying equations (1) and (2) [27]. In this approach, the increase in density was correlated with the punch movement during sintering. However, such punch movement is a function of shrinkage, thermal expansion and the elastic deformation of powder and punches. To identify the fraction of punch movement due to thermal expansion and elastic deformation, the punch movement was recorded during heating of a dense sample while a: applying a constant pressure, and b: while increasing pressure. Accordingly, the change in sample height due to shrinkage Δli was calculated. The shrinkage εzi was determined by dividing the change in sample height by the initial sample height l0. The density present during sintering and the initial density were described by ρi and ρ0, respectively.
2. Materials and methods A steel powder which was gas atomized by TLS Technik GmbH & Co. Spezialpulver KG was used as the matrix material. A zirconia powder partially stabilized with 3.5 wt% MgO (Mg-PSZ) from Saint Gobain was utilized as a reinforcement phase. The chemical compositions of the initial powders and the particle-size distributions are shown in Tables 1–3. The steel powder was fully austenitic, with 10 vol% of the Mg-PSZ exhibiting the monoclinic phase, 53 vol% the cubic and 37 vol% the tetragonal phase. Based on these initial materials, composite powders were prepared that contained 5 vol%, 10 vol% and 40 vol% Mg-PSZ. To this end, the powders were milled in the appropriate ratios for 4 h using a planetary ball mill (Pulverisette 6, Fritsch GmbH) and a ball-mass to powder-mass ratio of 5:1. The milling balls were made of steel and had a diameter of 10 mm. Some of the powders were milled in isopropanol at 100 rpm to produce a thorough mixture. Further composite powders were prepared by dry milling at 250 rpm to bring about intense mechanical interlocking between the powder particles. Process control agents (PCAs) were not used to avoid any change in the chemical composition of the
ρi = ρ0 ⋅e εzi
(1)
εzi = Δl i / l 0
(2)
Samples of 6 mm in height and 6 mm in diameter were machined for quasi-static compression tests. The samples were compressed at a strain rate of 10−3 s−1 to a strain value of 60% if fracture did not occur at lower degrees of deformation. The microstructure of both the as-sintered material and the compressed samples was investigated by X-ray diffraction, light optical microscopy, optical inclusion analysis using AnalySIS Particle Inspector software, and scanning electron microscopy (SEM) in combination with electron backscatter diffraction (EBSD). From EBSD analysis, the grain size of the steel matrix was determined by means of the Line Intercept Method. 3. Results and discussion
Table 1 Chemical analysis of the Mg-PSZ powder. [wt.%]
Cr
3.1. Initial material
SiO2
MgO
Al2O3
CaO
TiO2
Na2O
Fe2O3
ZrO2
4.23
3.37
0.63
0.21
0.14
0.09
0.01
Bal.
Fig. 1 shows the initial powders as well as the powders after milling. High-energy ball milling at 250 rpm led to intensive deformation of the steel powder and mechanical interlocking between the steel and 2
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Fig. 1. Initial powders: a) TRIP steel; b) Mg-PSZ; composite powders after ball milling: c) 5C100; d) 5C250; e) 5C250, ground particle; f) 40C250.
the results obtained for a composite with a similar steel [23], the powder mixture prepared at 250 rpm exhibited very coarse powder particles (Fig. 1) and, therefore, a low apparent density at the onset of densification (Fig. 3). However, the densification rate of the powder mixed at 250 rpm was enhanced, which was illustrated by the higher slope in Fig. 3. Hence, the densification process was almost finished at 1000 °C for 5C250, while 5C100 was almost dense at 1100 °C. The increased densification rate resulted from the higher number of vacancies and dislocations due to high-energy ball milling [23]. Thus, the driving force for the densification was enhanced. These results show that it is possible to decrease the sintering temperature necessary for such composites by increasing the rotation speed during milling. Due to deformation during milling, recrystallization took place during sintering and resulted in grain refinement (Fig. 4). The average grain size of the steel matrix is shown in Table 4. It was obvious that the grain growth during SPS was constrained by the ceramic particles. Hence, the unreinforced steel sample exhibited much larger grains than the composites. For composites mixed at 100 rpm, however, the grain size did not vary with increasing ceramic content. Due to the high degree of deformation combined with the recrystallization process, the grain size of the composites mixed at 250 rpm was slightly smaller and decreased with decreasing ceramic content. Ceramic particles are rigid and absorb impact energy from milling balls during milling. Therefore, the steel particles experienced less deformation with increasing ceramic content. Hence, the grains were coarser with increasing ceramic content after milling at 250 rpm and subsequent SPS. Optical inclusion analysis indicated a median particle diameter for Mg-PSZ of 2.5 μm in the 5C250 material. Thus, the median reinforcement diameter was 1 μm larger than d50 of the Mg-PSZ. Therefore, intensive refinement of the Mg-PSZ due to milling can be ruled out. After SPS, the steel phase was completely austenitic, while the monoclinic content of the Mg-PSZ phase increased with respect to the milled powder in composites 5C100 and 40C250 (Table 4). The slight decrease in monoclinic content in 40C100 was within the acceptable measurement error range. According to EBSD analysis (Fig. 4), some 40%–60% of the Mg-PSZ exhibited the monoclinic phase. In particular, the smaller Mg-PSZ particles and the edges of the coarse Mg-PSZ
ceramic powder particles. Furthermore, the initial steel particles were kneaded together, and formed larger particles due to cold welding (Fig. 1e and f). However, the milling process at 100 rpm resulted merely in a homogeneous distribution of the powders without considerable deformation of the powder particles (Fig. 1c). Due to the deformation during milling, the austenitic steel was transformed to α′-martensite to a certain degree. However, the α′-martensite was fully converted to austenite during SPS. XRD did not indicate any phase transformation within the Mg-PSZ due to milling for the composite powders 5C100, 5C250 and 10C250. Approximately 8% of the Mg-PSZ transformed to the monoclinic phase during milling for 40C100, and 15% of the MgPSZ in 40C250. The ball milling parameters strongly affected the microstructure of the composite after SPS (Fig. 2). Due to the low rotation speed of 100 rpm, the originally spherical steel particles were clearly visible in the micrographs of the sintered samples (Fig. 2a and c), where these steel particles were now framed by Mg-PSZ. However, the Mg-PSZ formed clusters in the gaps of the initial steel particles. If the composite contained 40 vol% Mg-PSZ, the Mg-PSZ particles formed a continuous network. Thus, it was not expected that a direct current would run through the sample during SPS. Voids were visible within the clusters to some extent (Fig. 2e). The porosity measured within the composites was very low, with pores making up less than 2% of the samples. Thus, these voids within the clusters were locations where ceramic particles had fallen out. The sintering temperature applied was too low to allow sintering of the MgPSZ. Thus, the interfacial bonding between the Mg-PSZ particles was weak. If the composite powder was milled at 250 rpm, however, the number of clusters was reduced considerably. Furthermore, the Mg-PSZ particles were very well embedded in the strongly deformed steel matrix (Fig. 2f). However, a longer milling time would have been necessary to have achieved perfect homogeneity. Large areas of formerly coarse steel particles without the appearance of Mg-PSZ were visible in Fig. 2b, d and f. Nevertheless, the number of Mg-PSZ/Mg-PSZ interfaces was reduced significantly. Fig. 3 illustrates the densification behavior for the composite with 5 vol% of Mg-PSZ with increasing temperature during SPS. Similar to 3
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Fig. 2. Microstructure of the composites after SPS (light phase: steel, dark phase: Mg-PSZ); a) 5C100, b) 5C250, c) and e) 40C100 (white arrows indicate voids), d) and f) 40C250.
destabilized during SPS. In this study, however, XRD analysis did not find any monoclinic content in 5C250 and 10C250, which was contrary to the results shown in Fig. 4. Even with EBSD, it was not possible to indicate all ceramic particles (dark grey areas in Fig. 4), which may have been due to the strong deterioration of the lattice structures of the ceramic particles during the intensive milling process. 3.2. Deformation behavior If the composites with 40 vol% Mg-PSZ were not taken into account, all of the other composites exhibited high degrees of ductility and were deformable to nominal engineering compressive strains of 60% without failure (Fig. 5). Because the percolation threshold was exceeded in the composites with 40 vol% Mg-PSZ, these composites were very brittle under compressive deformation, while their degree of deformability would have been reduced even further under tension. Cracks occurred on the radial surfaces of the samples mixed at 250 rpm. The composites had a higher 1% offset compressive yield strength (σ1%) than the unreinforced steel. Due to the small grain size of the steel resulting from the SPS process, the steel exhibited a high σ1%, which was twice as high as the value for the cast steel type [31]. This effect is well known from the Hall-Petch relation [32]. However, σ1% deviated by approximately 70 MPa from the published results (Fig. 6), where the Hall-Petch relation was applied for the steel alloy used [30]. The σ1% increased further through the addition of the ceramic particles. Taking into account the deviation of the σ1% of the steel by 70 MPa and assuming that the slope of the Hall-Petch relation was in accordance with the slope given by Ref. [30], it was possible to calculate the strength σS of the steel matrix of the composite as a function of the grain size (Fig. 6, Table 5). It was obvious that the refinement of the steel grains caused an increase in the σ1%. The remaining amount of strength was related to the ceramic reinforcement. Assuming that the increase in strength followed the rule of mixture (equation (3)), the increase in strength σMg-PSZ caused by the
Fig. 3. Densification with increasing temperature during SPS for 5C100 and 5C250.
particles were of the monoclinic phase. This was due to destabilization of the Mg-PSZ during SPS [28,29]. In this case, the Mg ions diffused from the ceramic particles into the steel matrix. However, it is known that the results of XRD analysis can deviate strongly from EBSD analysis [16]. According to XRD measurements, the destabilized Mg-PSZ content after SPS seemed to decrease with increasing Mg-PSZ content and lower rotation speed during milling (Table 4). If a high milling speed was used, the Mg-PSZ particles were better distributed within the steel matrix, and fewer Mg-PSZ clusters were formed. Thus, more steel/ ceramic interfaces existed if a high rotation speed was selected than if a low rotation speed was used. Hence, the interfacial area where Mg-ions could leave the Mg-PSZ during SPS increased. In Ref. [23], it is reported that due to the higher proportion of steel/ceramic interfaces with increasing rotation speed during milling, a larger amount of Mg-PSZ is 4
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Fig. 4. Grains within a) the steel; followed by the steel matrix of the composites and indexing of phases by EBSD after SPS: b) 5C250, c) 5C100, d) 10C250, e) 40C100, f) 40C250 (grey: austenite; green: cubic zirconia, yellow: tetragonal zirconia, red: monoclinic zirconia). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.) Table 4 Microstructural characteristics; monoclinic content of the Mg-PSZ determined by XRD. Material
S 5C100 5C250 10C250 40C100 40C250
Relative Density [%]
97.6 98.8 98.0 98.6 98.1 98.8
Average Grain Size [μm]
5.6 2.4 1.4 1.8 2.5 2.0
Monoclinic Content of Mg-PSZ [%] Before SPS
After SPS
– 0 0 0 18 25
– 28 0 0 15 41
reinforcement could be calculated (Table 5). If the ceramic content increased, the amount of strength σMg-PSZ that was due to the ceramic reinforcement increased as well. Furthermore, the rotation speed during milling and the combined homogeneous distribution of the ceramic particles influenced the σ1% value. This effect was more pronounced for high ceramic contents. While an increase in the rotation speed during milling led to a rise in strength due to the ceramic of Δσ = 534 MPa in composites with 40 vol% ceramic material, the strength increased by approx. 20 MPa in composites with 5 vol% of ceramic powder. Due to the higher rotation speed, the ceramic particles were distributed more homogeneously within the steel matrix and mechanically interlinked with the steel. Furthermore, the amount of weak ceramic/ceramic interfaces was reduced dramatically. In
Fig. 5. Compressive deformation behavior.
addition, the application of the high rotation speed may have resulted in refinement of the Mg-PSZ particles. Thus, the number of reinforcing particles increased and the mean distance between them decreased, which served to improve the degree of dispersion strengthening. Hence, these effects were more distinct in composites with a high ceramic content and influenced the σ1% value.
σ1% = VS⋅σS + (1 − VS)⋅σMg − PSZ
(3)
As indicated by the slopes of the true compressive stress – 5
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40C250 exhibited a higher strength than 40C100. However, the cracks along the steel/ceramic interfaces in this composite ran along the network of reinforcement, were very long and were not stopped by the steel phase. Hence, the 40C250 material failed at a lower compressive deformation value. Due to the low deformability of 40C100 and 40C250, hardly any α′-martensite was formed. The 10C250 composite did not exhibit a dense network of reinforcement, such that cracks at interfaces were very short and were stopped by the steel matrix. Due to the high amount of steel, the composite exhibited a high degree of deformability. However, the amount of ceramic and the interfacial strength were high enough to facilitate load transfer to the ceramic particles and, therefore, to ensure the high strength of the composite. Hence, stress-induced phase transformation and deformation-induced phase transformation took place. 4. Conclusions Fig. 6. Calculated strength σS of the steel matrix and Hall-Petch fit, as reported by Ref. [30].
The results presented demonstrate that high-energy ball milling has the potential to improve the mechanical properties of Mg-PSZ-reinforced TRIP-matrix composites. Due to the high energy impact, the ceramic particles were more homogeneously distributed and were kneaded into the steel matrix. Therefore, fewer ceramic clusters appeared in the composite. In future research, the milling time should also be increased to improve the distribution of the Mg-PSZ particles further. Due to the absence of PCAs, coarse particles remained after high-energy ball milling. However, the sintering process was not negatively influenced by the coarse particles. Due to the interlocking between the Mg-PSZ and the steel phase, the transmission of the mechanical load from the steel matrix to the ceramic reinforcement improved. This effect led to an increase in compressive strength and was more pronounced with increasing ceramic contents up to 40 vol%. Furthermore, the steel was strongly deformed during high-energy ball milling and recrystallized during SPS. The resulting fine-grained steel matrix improved the compressive strength further. However, the grain refinement was diminished by large amounts of the reinforcing phase. Due to the fine-grained steel matrix and the reduced number of ceramic clusters, therefore, the compressive strength was improved without a distinct reduction in deformability – except for an Mg-PSZ content of 40 vol%. When a content of 40 vol% Mg-PSZ was used as the reinforcement, the percolation threshold was exceeded. Thus, the composite became very brittle and failure occurred – primarily in the Mg-PSZ phase. In contrast, an Mg-PSZ content of 5 vol% seemed to be low, but increased σ1% by approx. 100 MPa. Due to the fine-grained steel matrix, Mg-PSZ made a significant contribution to compressive strength, and by not exceeding the percolation threshold, the composite 10C250 exhibited a high σ1%, a high degree of deformability and a large degree of stress-induced and deformation-induced phase transformation. It was assumed that the ceramic content could have been increased further without exceeding the percolation threshold.
Table 5 1% offset compressive yield strength σ1% and calculated contribution of reinforcement phase as a function of the volume fraction of steel Vs. Material [rpm]
σ1% [MPa]
Strength σs of Fine-Grained Unreinforced Steel Matrix [MPa]
Increase in Strength (1-VS)∙σMg-PSZ due to Mg-PSZ [MPa]
S 5C100 5C250 10C250 40C100 40C250
400 540 636 681 776 1328
400 488 568 528 483 513
0 76 96 206 486 1020
logarithmic strain curve, the degrees of work hardening of 5C100 and 5C250 were similar, and they exhibited the same type of damage. As shown in Fig. 7, it was primarily the steel matrix that deformed. In Fig. 7d, deformation bands are clearly visible (marked by the white ellipse), and these deformation bands were the sites where α′-martensite was formed. The α′-martensite nuclei constrained the dislocation movement in a manner similar to grain boundaries and, therefore, intensified the work-hardening effect [30]. Furthermore, delamination occurred at the steel/ceramic interfaces (marked by white arrows in Fig. 7). Due to the high deformability of these composites, a large amount of α′-martensite was formed (Table 6). Furthermore, the ceramic material underwent stress-induced transformation (Table 6). In contrast, the deformation behavior of the composites with 40 vol % Mg-PSZ varied. Due to the exceedance of the percolation threshold in these composites, the damage was concentrated within the ceramic material (Fig. 7e), and it was mainly large ceramic particles and clusters that fractured. Due to the sintering temperature, which was quite low for zirconia, the Mg-PSZ clusters exhibited low strength and failed at a low degree of deformation of the composite. In some regions, deformation of the steel phase was visible. The cracks in the ceramic particles and clusters were stopped by the ductile steel matrix. Despite the failure of the Mg-PSZ particles, stress-induced phase transformation took place (Table 6). Furthermore, the strength of the interface between the steel phase and the ceramic phase was weak due to stress caused by their differences in thermal expansion during cooling. Thus, delamination occurred at these interfaces. In 40C250, some deformation bands formed during deformation within the steel phase. Due to a smaller number of clusters and large particles, their fracture was not the main type of damage, as it was mainly the steel/ceramic interfaces that failed. Furthermore, fewer Mg-PSZ particles experienced stress-induced phase transformation (Table 6). However, the strength of the interfaces seemed to be higher than the strength of the ceramic clusters. Hence,
Data availability The raw/processed data required to reproduce these findings cannot be shared at this time, as the data also forms part of an ongoing study. Acknowledgements The authors would like to thank the German Research Foundation (DFG) for supporting the investigations, which were part of the Collaborative Research Center TRIP-Matrix Composites (SFB 799, subproject A06). Furthermore, the authors would like to thank the Institute of Ceramic, Glass and Construction Materials and the Institute of Iron and Steel technology at TU Bergakademie Freiberg for the chemical characterization of the Mg-PSZ and steel powders. 6
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Fig. 7. Deterioration in the microstructure of the composites after 60% compressive deformation, a) 5C100, b) and d) 5C250, c) 10C250; and after failure, e) 40C100, f) 40C250 (light grey: ceramic particles, dark grey: steel, black: cracks and holes). Table 6 Phase contents due to deformation up to 60% or until fracture; amount of monoclinic content refers only to the ceramic material, and not the entire composite. Material
α′-Martensite [%]
Stress-Induced Monoclinic Content [%]
S 5C100 5C250 10C250 40C100 40C250
50 51 47 42 5 0
– 22 57 77 33 10
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