Mechanical properties and structure of rapidly solidified bulk Fe89 − xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15 − yB4 (y = 7) alloys

JMADE-01698; No of Pages 14 Materials and Design xxx (2016) xxx–xxx

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Mechanical properties and structure of rapidly solidified bulk Fe89 − xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15 − yB4 (y = 7) alloys Piotr Józef Bardziński a,⁎, Piotr Błyskun b a b

Department of Mechanics, Materials Science and Engineering, Wrocław University of Science and Technology, M. Smoluchowskiego St. 25, 50–370 Wrocław, Poland Faculty of Materials Science and Engineering, Warsaw University of Technology, Wołoska St. 141, 02–507 Warsaw, Poland

a r t i c l e

i n f o

Article history: Received 16 January 2016 Received in revised form 14 April 2016 Accepted 15 April 2016 Available online xxxx Keywords: Compression Bending Hardness Ultrasonic through-transmission Fe16Hf6Si7

a b s t r a c t Authors investigated mechanical properties of microcrystalline Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys obtained by copper mold injection casting to a form of rods, 3 mm in diameter. Structure of the samples was analyzed by Cu K-α powder X-ray diffractometry and the unit cell parameters of identified h-Fe2Hf0.8Ta0.2 (C14 Laves), fcc-Fe16Hf6Si7 (D8a) and bcc-Fe phases were calculated. The study revealed that ascending Si content leads to reduction in values of both unit cell parameters of C14 phase in the alloys studied while causes the expansion of D8a lattice; decrease in secondary dendrite arm spacing in the middle of the rods; stepwise improvement of yield and compressive strength at the expense of plastic deformation region, what was attributed to the precipitation of C14 Laves/D8a in the interdendritic regions of bcc-Fe; decrease in flexural strength with no effect on deflection-at-yield and exponential increase in hardness of the studied alloys. There is also clear correlation between unit cell volume of fcc-Fe 16 Hf 6 Si 7 phase, yield and compressive strength on Si content. For the alloys with x = 0–13, relation between yield strength and Vickers hardness was described with logarithmic function. It was found, that the lattice expansion of bcc-Fe and, in lesser extent, C14 Laves phase is correlated with the value of bending modulus in the whole composition range. © 2016 Elsevier Ltd. All rights reserved.

1. Introduction Proper alloy components selection together with rapid solidification techniques opens the possibility for obtaining micro- and nanocrystalline, quasicrystalline or amorphous structure of metallic materials [1]. In such cooling conditions, increased solid solubility limit [2,3] and reduced time for nucleation and grain growth are responsible for the formation of unique non-equilibrium structures, leading to outstanding mechanical properties [4]. Some iron-based systems can form bulk metallic glass, such as Fe73−xNb4Hf3YxB20 (x = 0–3) [5], Fe61Co10Me7Y2B20 (Me = Y(7), Y(6)Ti(1) or Zr(2.5)Hf(2.5)W(2)) [6] or (Fe0.61Co0.10Zr0.025Hf0.025Ti0.02W0.02B0.20)100−xYx (x = 0, 2 or 4) [7, 8]. Inoue [9] reports that certain bulk nanocrystalline alloys, as Zr65Al7.5Ni10Cu7.5Pd10, exhibit even larger compressive and tensile strength as well as ductility than their amorphous counterparts with the same composition. However, mechanical properties enchancement by fast melt quenching is not limited to ultrafine-grained alloys. According to Minemura et al. [10], high strength and moderate ductility could be achieved in rapidly quenched high-carbon steels. Mechanical ⁎ Corresponding author. E-mail address: [email protected] (P.J. Bardziński).

properties of the alloys can also be improved by the intrinsic formation of certain reinforcing crystalline phases. For instance, tensile strength of gray cast iron was elevated by the presence of primary austenite dendrites in its microstructure [11]. Occurrence of intermetallic phases such as hexagonal C14 Laves phase was reported to have strengthening effect on iron-aluminide alloys [12,13]. Development of alloys reinforced with hard microcrystalline precipitates is driven by the necessity of manufacturing durable and wear-resistant components for special applications. Due to the complex compositions involved, use of expensive elements and sophisticated manufacturing equipment, economic feasibility justifies its applications as a small parts incorporated in larger devices [14]. Wieczerzak et al. [15] obtained abrasive wear-resistant rapidly solidified bulk Fe–Cr–Mo–C alloys, where hard carbides precipitated in interdendritic zones of ferrite dendrites. Ibrahim et al. [16] described Al–Si–Mg alloys with hardening Mg2Si phase, used in aerospace industry. Park et al. [17] obtained rapidly solidified Al–Si–Fe alloys with improved hardness, tensile strength and ductility resulting from a fine dispersion of DO22–(Al,Si)3Ti phase. Sheng et al. [18] found, that in copper mold injection-cast NiAl–Cr(Nb)/Dy alloy, NiAl/Cr2Nb lamellar eutectic microstructure is formed instead of coarse Cr2Nb grains identified in conventionally cast samples. This microstructural feature together with grain refinement leads to improved compressive strength.

http://dx.doi.org/10.1016/j.matdes.2016.04.054 0264-1275/© 2016 Elsevier Ltd. All rights reserved.

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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Recently, other interesting approach in alloy design was the formation of multicomponent alloys with large configurational entropy and more than five elements, so called, high entropy alloys. In this case, tendency to formation of hard intermetallic phases is hindered due to equimolar component quantities that promote solid solubility [19]. For instance, Li et al. [20] found 3-times increase in yield strength under compression for CoCrFeNi high entropy alloy, owed to substantial supercooling. Guo et al. [21] obtained a refractory high entropy alloy forming bcc-MoNbHfZrTi solid solution with high yield strength of 1.7 GPa measured in compression tests performed at room temperature. Bardziński et al. [22] recently described the magnetic properties of amorphous Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 0, 7) alloys obtained in the form of thin melt-spun ribbons. In the current study the authors want to address the mechanical performance of the alloy with y = 7 in the form of rapidly quenched microcrystalline rods with reinforcing hard crystalline phases. As the properties desirable in the applications for transformer cores, such as eddy current loss and magnetostriction reduction together with increase in specific resistivity could be achieved by Si substitution to steel [23,24], it could be important to study the impact of Si substitution for Fe in Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) rapidly quenched rods. Microstructure and mechanical properties of the alloys obtained by rapid solidification, were compared with reference steel abbreviated as Ref in the text. The latter is unalloyed hypoeutectoid steel with ferritic-pearlitic microstructure, equivalent to 10 or 09 A grade (ISO 683–3:2014). It is used as a core material for arcwelding electrodes that complies with EN 499/ISO 2560: E 38 0 RC 11 classification [25]. Initial tests (as demonstrated in Figs. 1 and 2 in Appendix A) shown acceptable weldability of Ref steel with bulk rapidly quenched alloys investigated in the current study. Ref could serve as an intermediate material for joining them with a range of relatively inexpensive steels, such as S235J2G3–S355J2G3 grades [26]. Thus, it is important to compare the mechanical properties of Ref steel with the newly engineered materials described, in order to choose the right composition for a certain application. 2. Experimental details Pure elements with comparable melting points, Hf (UQF, 98.07 at.%, with 1.65 at.% Zr and 0.28 at.% Fe impurities) and Ta (99.9 at.%) and, separately, Fe (Aldrich, 99.98 at.%), Cu (Hutmen, 99.99 at.%), B (Aldrich, N99 at.% with 1 mg/g Fe and Na, 0.5 mg/g Ca, Cu, K, Mg, Mn, Ni, Zn impurities), Si (Kurt Lesker, N99.9 at.%), Gd (Aldrich, 99.99% REO) and, where applicable, La (Aldrich, 99.9% REO) were arc-melted under Zr-gettered high purity Ar atmosphere, until homogenous dispersion of elements was reached, as confirmed by energy-dispersive X-ray spectroscopy (EDX). Then, each crushed ingot was placed in quartz crucible, high frequency induction-melted under high purity Ar and injection cast into a 3-mm wide channel of copper mold that was 100 mm in diameter and 40 mm long. Note that temperature of the melt in the moment of injection was higher by approx. 36 K than the liquidus temperature of each alloy. Resulting samples were rapidly solidified cylindrical rods, approx. 3 mm in diameter and approx. 40 mm long. Apparatus used to obtain the samples along with as-quenched rods were shown in the Fig. 1. Each sample was spark-cut to a 1-mm thick slice, mechanically polished to achieve mirror luster and placed on Si zero-background sample holder. The structure was investigated by powder X-ray diffractometry performed on Bruker D8 Advance diffractometer with Ni-filtered Cu K-α radiation (λ = 1.5406 Å). The data was recorded in 2θ range of 20–100° with scan rate of 0.016° and counting time of 4.63 s per step. Surface morphologies after compressive and flexural fracture were examined by means of optical microscopy. Cross-sections of the samples were spark-cut perpendicular to the rod axis, ground with wet sand-paper with the grit size from 80 to 4000 and polished with a

Fig. 1. Experimental setup for copper mold injection casting (left) together with the samples of the alloys Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15), Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) and Ref steel (right).

3 μm diamond suspension. The microstructure of untreated surfaces was investigated by Hitachi S-3400N Scanning Electron Microscope (SEM) equipped with energy-dispersive X-ray spectroscopy (EDX) analyzer, used to study the elemental composition. The magnification was maintained at 2000 times. Polished sample surfaces were then etched with 5% Nital and studied with Olympus GX51 metallographic microscope in bright-field mode with 1000 times magnification. Samples for compression tests were spark-cut to a length of approx. 4.5 mm, so the length-to-diameter ratio was maintained at about 1.5, to avoid buckling. Both sample ends were plane-parallel and polished to a mirror luster. Cemented tungsten carbide BHH S20S bearing blocks with dimensions: 12.66 × 12.66 × 4.72 mm and HV of about 1500 were placed on each sample end to prevent plastic deformation of testing machine heads, according to guidelines given in E9-09 ASTM standard [27]. Uniaxial compression tests performed on MTS 809 axial/torsional test system were displacement-controlled with a crosshead speed of 5 μm s −1. Test stand for compression tests was shown in the Fig. 2(a). Due to safety reasons, sample is placed inside the protective housing, consisting of a piston and cylinder. Cylindrical samples for three-point flexural test were approx. 35–40 mm long, mean sample diameter ϕ and section modulus W of each sample were given in Table 7. Note that distance between supports lo was equal to 23.19 mm in every case. Fig. 2(b) shows test stand during bending of Ref sample. Bending tests were displacement-controlled with a crosshead speed of 0.5 μm s − 1. The tests were performed on MTS Bionix except for the alloy with y = 7, for which three-point bending was realized on MTS Synergie100 with force sensor measuring up to 500 N. Vickers macrohardness test (HV3) was realized by means of Zwick/ Roell ZHU 187,5 universal hardness tester with 3 kgf load and dwell time of 10 s. Hardness and elastic moduli were also determined by nanoindentation (abbreviated here as NHT), by means of Nanoindentation Tester, CSM Instruments SA. Indentation was performed with linear loading with acquisition rate of 10 Hz, loading as well as unloading rate of 1000 nm/min to a maximum depth of 1450 nm. The indenter was maintained at this depth for 12 s before unload. For contact detection, the 80% slope variation of the depth versus time was adopted. The Oliver and Pharr method [28] was selected to calculate a best fit for the contact area function. In this approach, the first part of the unloading curve, i.e., the normal force plotted against penetration depth was described by a power law relationship. All indentation tests were performed

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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Fig. 2. Experimental setup for (a) uniaxial compression test and (b) three-point flexural test.

on the cast rod cross-section parallel to its length axis, with the samples included in poly(methyl methacrylate) resin, cut and mechanically polished to a mirror luster. Deviation from plane-parallelism of the resin-included cylindrical samples was not exceeding 1°. The rod samples for ultrasonic through-transmission [29,30] measurements were spark cut into cylinders with plane-parallel bases and mean length of 9.25 mm. In the method applied the ultrasonic wave propagates once through the sample that is placed between two wave probes and the time of impulse transmission is measured. In case of longitudinal waves propagation, the probes were coupled to the sample through a layer of oil. For the propagation of transverse waves, the resin was chosen as a coupling medium due to its viscosity sufficiently high to carry the shear vibrations. Considering that the wave propagates through a layer of coupling medium, calibration was carried out, i.e., delay between two successive multiple echoes was subtracted from the single transit time of a wave through a standard sample [31,32]. Schematic representation for this method was shown in the Fig. 3. A pair of Olympus V156 transverse wave probes with a transmitting frequency of 5 MHz and 6.35 mm in diameter was used to measure transverse wave velocity vT. To determine the longitudinal wave velocity vL, two Olympus V543 longitudinal wave probes with a transmitting frequency of 5 MHz and 6.35 mm in diameter were utilized. Calibration was performed on standard sample no. 2, 12.5 mm thick. For transverse waves, wave propagates the thickness of a standard sample once in 4150 ns, delay time between two consecutive multiple echoes was 7640 ns/2 and the resulting time delay correction was 330 ns. For longitudinal waves, propagation time was 2272 ns for the first impulse and 4174 ns/2 for multiple echoes, what gives 185 ns of a time delay correction. Data acquisition was realized with Sofratest SFT4001HPCI doublechannel ultrasonic card coupled with a notebook through a Magma PCI2CardBus interface. The ultrasonic signal was averaged 30 times, zero-level slope crossing method was used for impulse time delay measurements with a resolution of ± 0.5 ns. Physical dimensions along with the mass of the samples were measured in order to calculate the bulk volumetric mass density d from its definition.

3. Results and discussion Structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15), Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys and steel (Ref) was investigated with Cu K-α powder X-ray diffractometry. In Fig. 4, diffractograms of the samples examined in the current study were compared with corresponding PDF (ICDD, 2008) reference patterns. Respective reflections were indexed according to reference. Phase composition of each alloy was presented in Table 1, together with elemental

Fig. 3. Test stand for ultrasonic through-transmission experiment.

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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y = 7 and fcc-Fe16Hf6Si7 phase was identified in the alloys with x = 8–15. The Fe2Hf1−zTaz pseudobinary system is more stable as hexagonal C14 MgZn2-type than cubic C15 MgCu2-type Laves structure [33] and tend to crystallize in C14 structure in the whole compositional range, i.e., for z = 0–1 [34–36]. Well known for its itinerant-electron magnetic properties [34,35,37] especially for z = 0.2, C14 Laves phases are expected to be brittle, exhibit the Young modulus of about 30–50 GPa and Poisson's ratio of 0.235–0.242 [38]. There is only a limited information in the literature regarding the fcc-Fe16Hf6Si7 D8a-type phase isostructural with Mg16Cu6Si7 [39]. To the best of our knowledge, the phase with the exact composition was described only two times, by Lisenko et al. [40] and Lisenko [41]. Rixecker and Haberkorn [39] found that D8a and C14 phases can coexist in annealed Fe16Ta6Si7 and Fe16Nb6Si7 amorphous powders obtained by mechanical alloying what complies with the findings stated in the present study. The group [39] also claim that crystallization kinetics favours the formation of D8a phase over C14, when the initial material is an amorphous solid, rather than when it is cooled down starting from above the liquidus temperature. Rixecker et al. [42] identified Fe16Nb6Si7 with D8a structure in annealed Fe73.5Cu1Nb3Si13.5B9 melt-spun amorphous ribbons, with a composition similar to investigated in the current study. There was additional tLaFeSi phase, described by Raghavan [43], identified in the alloy Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7). In this alloy, despite that it contains 8 at.% Si, no reflections related to D8a phase were identified. As compared with the alloy with x = 8, this could lead to the conclusion that the presence of La have a detrimental effect on the formation of D8a in favor of the C14 Laves phase in the studied compositions. By substituting interplanar spacing dhkl from the Bragg equation into the quadratic equation of lattice planes for the hexagonal lattice one could get the c and a parameters of the unit cell: Fig. 4. Cu K-α X-ray diffractograms for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref); reflections ascribed to identified crystalline phases correspond to the reference patterns from PDF (ICDD, 2008) database [44].

composition of the areas marked in Fig. 5, studied by SEM-EDX. It was found that, besides bcc-Fe and borides identified in all samples except Ref, h-Fe2Hf0.8Ta0.2 was present in the alloys with x = 0–15 as well as

Table 1 Crystalline phases indentified by X-ray diffractometry (XRD) and energy-dispersive X-ray spectroscopy (SEM-EDX) for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0-15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref). All numbers given in atomic percent. Note that data presented from SEM-EDX analysis are not stoichiometric formulae. Note that geometric symbols given in parentheses refers to the respective areas on the micrographs given in the Fig. 5. Composition

XRD

SEM-EDX

Ref x=0

bcc-Fe bcc-Fe; t-Fe3B; h-Fe2Hf0.8Ta0.2

x=8

bcc-Fe; t-Fe3B; bct-Fe2B; h-Fe2Hf0.8Ta0.2; fcc-Fe16Hf6Si7 bcc-Fe; t-Fe3B; bct-Fe2B; h-Fe2Hf0.8Ta0.2; fcc-Fe16Hf6Si7 bcc-Fe; bct-Fe2B; h-Fe2Hf0.8Ta0.2; fcc-Fe16Hf6Si7 bcc-Fe; t-Fe3B; bct-Fe2B; h-Fe2Hf0.8Ta0.2; fcc-Fe16Hf6Si7 bcc-Fe; t-Fe3B; bct-Fe2B; h-Fe2Hf0.8Ta0.2; t-LaFeSi

Fe98.53Si0.59Mn0.51C0.37 [25] Fe97(Cu,Gd,Ta,B) [○]; Fe75Hf20Ta5 [□]; Fe90Hf8Ta2 [△] Fe92(Si7Cu1) [○]; Fe75Si7Hf10Ta3Gd5 [□]

x = 11

x = 13

x = 15

y=7

Fe89.4Si10Cu0.6 [○]; Fe77Si10Gd5Hf5(Ta2Cu1) [□]; Fe65Si14Hf16Ta5 [△] Fe88Si11Cu1 [○]; Fe73Si15Hf4Gd8 [□]; Fe63Si18Hf15Ta4 [△] Fe86Si12Cu1 [○]; Fe73Si12Hf3Ta1Gd10Cu1 [□]; Fe60Si16Hf17Ta5Gd1Cu1 [△] Fe93Si5Cu2 [○]; Fe83Si6Gd5La1Hf5Ta1 [□]; Fe67Si8Hf18Ta7 [△]; Fe42Si30La20Gd8 [◊]

0

112

 2 2 l 22 h1 þ h1 k1 þ k 21 −l 21 h2 þ h2 k2 þ k 22 B C c¼B 2 
2
C @− 
A ð1Þ  2 1 2 h 22 þ h2 k2 þ k 22 − sinθ h1 þ h1 k1 þ k 21 4 sinθ n1 λ n2 λ

0
112 2 2 4 B3 h1 þ h1 k1 þ k1 C a ¼ @
2
2 A : l1 1 þ 4 sinθ c n1 λ

ð2Þ

Scattering angles 2θi are given in radians and reflection orders ni are corresponding to a pair of reflections hikili where i = 1, 2. λ, given in ångström, stands for the wavelength of Cu K-α radiation. By analogy, the equation for a parameter of unit cell in the cubic system can be written as:

a¼

 !1 ðnλÞ2 h 2 þ k 2 þ l 2 2 : 4sin2 θ

ð3Þ

Unit cell parameters of h-Fe2Hf0.8Ta0.2 phase found to be present in a number of studied alloys were calculated according to Eqs. (1) and (2) from the reflections with the following Miller indices: 200 and 112, 110 and 103. The values of c and a parameters are given in Table 2. Ouyang et al. [37] reported c and a parameters of h-Fe2Hf0.8Ta0.2 to be 8.058 and 4.924 Å, respectively. In [35–1387] PDF (ICDD, 2008) database [44] record, based on work of Nishihara and Yamaguchi [45], c and a are found to be 8.065 and 4.930 Å, respectively. According to Table 2 those values are close to the obtained for the C14 phase in Fe89−xHf4Ta1Cu1Gd1SixB4 with x = 0. It was found that Si addition leads to the reduction in values of both unit cell parameters of C14 phase in the alloys studied. Atomic radius of Si is significantly smaller than for Fe. Assuming, that Fe atoms are substituted by Si in C14 Laves

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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Fig. 5. Secondary electron micrographs of untreated cross-sections of rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0 − 15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref); scale bar stands for 20 μm. Note that composition of the areas on the micrographs marked with respective geometric symbols is given in Table 1.

Table 2 Lattice constants: a, c and unit cell volume V of hexagonal Fe 2 Hf 0.8 Ta 0.2 phase in rapidly quenched bulk Fe 89−x Hf 4 Ta 1 Cu 1 Gd 1 Si x B 4 (x = 0–15) and Fe 74 Hf 4 Ta 1 Cu1 Gd1 La ySi 15−yB4 (y = 7) alloys. Composition

a [Å]

c [Å]

V [Å3]

x=0 x=8 x = 11 x = 13 x = 15 y=7

4.901 ± 0.003 4.850 ± 0.008 4.816 ± 0.013 4.828 ± 0.013 4.826 ± 0.013 4.883 ± 0.005

8.022 ± 0.007 7.943 ± 0.003 7.900 ± 0.004 7.913 ± 0.010 7.900 ± 0.024 7.993 ± 0.001

166.86 ± 0.078 161.84 ± 0.556 158.68 ± 0.759 159.73 ± 0.630 159.35 ± 0.359 165.07 ± 0.373

lattice, interplanar spacings are statistically smaller and thus, reflections on the diffractogram are shifted to the higher angles. This could explain decreasing unit cell volume with ascending Si content. As can be seen

Table 3 Lattice constant a and unit cell volume V of fcc-Fe16Hf6Si7 phase in rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 8 − 15) alloys. Composition

a [Å]

V [Å3]

x=8 x = 11 x = 13 x = 15

11.530 ± 0.025 11.535 ± 0.012 11.558 ± 0.022 11.532 ± 0.005

1532.83 ± 9.94 1534.81 ± 4.97 1544.08 ± 8.95 1533.57 ± 1.80

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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Table 4 Lattice constant a and unit cell volume V of bcc-Fe phase in rapidly quenched bulk Fe 89−x Hf 4 Ta 1 Cu 1 Gd 1 Si x B 4 (x = 0 − 15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref). Composition

a [Å]

V [Å3]

Ref x=0 x=8 x = 11 x = 13 x = 15 y=7

2.869 ± 0.001 2.871 ± 0.001 2.868 ± 0.001 2.863 ± 0.002 2.871 ± 0.004 2.867 ± 0.003 2.871 ± 0.002

23.61 ± 0.02 23.67 ± 0.01 23.58 ± 0.06 23.46 ± 0.03 23.66 ± 0.19 23.56 ± 0.07 23.65 ± 0.03

from EDX data given in Table 1, Si is distributed among all phases identified. Eq. (3) was used to calculate the values of a parameter in fccFe16Hf6Si7 (D8a) and bcc-Fe phases and the results are shown in Tables 3 and 4, respectively. In case of D8a phase, hkl indices chosen to compute the unit cell parameter were: 111, 411, 1331, 711, 511, 440, 822, 933, 844, 733, 331, 422 and for bcc-Fe: 200, 110, 211, 220. In the work of Lisenko [41] reported as a source for PDF (ICDD, 2008) [29–0698] record [44], a for Hf-containing D8a phase is equal to 11.480 Å. In the alloy with x = 8 investigated in the current study it is higher by 5 · 10 −2 Å. As the a parameter increases in the Si-richer alloys, what is also reflected in the phase composition given in Table 1,

Fig. 6. Bright-field optical micrographs of the cross-sections of rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0 − 15), Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) and steel (Ref) etched with 5% Nital; scale bar stands for 20 μm.

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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one can conclude that the presence of Si atoms causes the expansion of D8a lattice. Silicon might be incorporated in the interstitial positions of D8a lattice, leading to lattice expansion with increasing amount of Si atoms. Dróżdż et al. [46] investigated the crystallization of melt-spun amorphous Fe-C-Si(B) cast irons, containing 2.75–4.08 at.% Si. The group reported lattice constant a of bcc-Fe phase from 2.855 to 2.860 Å. Carsley et al. studied the structure of mechanically alloyed Fe90Cu10 and found a of bcc-Fe phase equal to 2.877 Å, compared to 2.866 for pure bcc-Fe powder [47]. Those values are close to those obtained in the current study. Bright-field optical micrographs of the sample cross-sections etched with 5% Nital are presented in Fig. 6 and corresponding SEM micrographs of untreated sample surfaces, that revealed elemental contrast are given in Fig. 5. According to average grain sizes measured for the middle and edge part of each rapidly quenched rod cross-section, collected in Table 5, it is obvious that the grain diameter decreases over the length of the radius from the middle to the edge, from 3 to 0.5 μm on average, respectively. It is owed to the fact, that the heat transfer during cooling of the sample takes place only through the interface between the molten metal and the copper walls of the mold, so the interior part of the sample remains in higher temperature for a longer period of time, allowing the crystals to grow. In the alloy with x = 0 one can distinguish nearly hexagonally shaped C14 Laves phase grains (lighter-shaded) uniformly dispersed with round bcc-Fe (darker) with rather smooth edges and comparable size. Microstructure of the alloys with x = 8–15 contains starfish-like bcc-Fe dendrites with six-fold rotational symmetry as well as lamellar eutectic structure. The latter can form due to the near-eutectic composition in Fe–Hf system [48]. As can be seen by comparing the Fig. 5 with the data presented in Table 1 which show compositional effects on microstructure evolution, when the Si:Fe ratio is gradually increased, the grains of Hf-rich phase start to form in the vicinity of iron dendrites. It seems that the addition of 8 at.% Si hinders the formation of bigger grains of C14 Laves phase as in the alloy with x = 0. In fact, in the alloy with x = 8, the Hf-rich phase exists only as a part of the eutectic. This might suggest the presence of the near-eutectic composition in the ternary Fe–Hf–Si system. The Hf-rich phase grains in the alloy with x = 13 contain remnants of bcc-Fe dendrites while in x = 15 three distinct phases, i.e., bcc-Fe and C14 Laves/D8a grains and lamellar eutectic are clearly separated. The microstructure of the alloy with y = 7 consists of the lamellar peritectic in the Fe–La–Si system [49,43] in which other phases are embedded. These include bcc-Fe, C14 Laves and the intermediate phase containing approx. 5 at.% Hf, as listed in Table 1. Characteristic microstructural features obtained for the composition with y = 7 are strikingly similar to those found for as-cast LaFe11.6Si1.4 bulk alloy studied by Liu et al. [50]. This material was described due to giant magnetocaloric effect of its constituent La(Fe,Si)13 phase and the authors [50] stated that the lamellae are the intermediate phase upon formation of La(Fe,Si)13

Table 5 Mean grain diameter of bcc-Fe and Hf-rich C14 Laves/D8a phases together with secondary dendrite arm spacing (SDAS), all given in μm, for two significant areas on the rod crosssection, for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys. Composition

x=0 x=8 x = 11 x = 13 x = 15 y=7

C14 Laves/D8a (light-shaded grains)

bcc-Fe (dark grains)

SDAS

Middle

Edge

Middle

Edge

Middle

Edge

2.0 ± 0.7 0.5 ± 0.1 3.0 ± 1.8 1.7 ± 0.9 2.9 ± 2.7 1.2 ± 1.3

1.1 ± 0.4 0.4 ± 0.1 2.2 ± 1.2 0.7 ± 0.4 1.8 ± 0.6 0.4 ± 0.1

4.3 ± 2.2 2.6 ± 1.8 2.2 ± 1.8 3.2 ± 1.2 2.5 ± 0.8 1.6 ± 1.9

1.6 ± 0.4 0.4 ± 0.1 0.3 ± 0.1 0.6 ± 0.3 0.9 ± 0.6 0.8 ± 0.7

– 1.6 ± 0.9 1.0 ± 0.5 0.7 ± 0.2 0.3 ± 0.1 –

– 0.7 ± 0.4 0.7 ± 0.4 0.7 ± 0.2 0.3 ± 0.1 –

7

after appropriate annealing. Lamellar eutectic growth was also found to be present in hypo- and hyper-peritectic Zn–Ag alloys [51]. For the alloys with x = 8–15, secondary dendrite arm spacing (SDAS) was determined and presented in Table 5. It was found that in the middle of the rods, mean value of SDAS decreases from 1.6 to 0.3 μm as the Si content incerased from 8 to 15 at.%, respectively. On the regions near the mold-melt contact surface, the composition had no significant effect on the SDAS, which turns to be about 0.7 μm for the alloys with x = 8–13. This complies with the finding of Easton et al. [52], who claim that SDAS depends on constitutional undercooling generated by the liquid remaining near the end of solidification. Average grain diameter and SDAS are decreasing from the middle to the edge of the rod due to cooling rate gradient across the rod radius during quenching process [53]. Elemental composition mapping on untreated sample cross-sections was performed with SEM-EDX method and presented in the Appendix B. It revealed that in the alloy with y = 7, La, Cu and Gd atoms precipitated in the bcc-Fe–t-LaFeSi peritectic interface. Similar observation was conducted for the alloys with x = 0–15, i.e., there were significant concentration of Gd and Cu atoms observed on the grain boundaries of both Fe- and Hf-rich phases. Ohkubo et al. [54] gave evidence that fcc-Cu clusters are serving as heterogeneous nucleation sites for bcc-Fe in amorphous Fe89Zr7B3Cu1 alloy. As in our case, the enthalpy of mixing between Cu and M (Hf, Gd) has much higher negative value [55–57] than in the alloys from the Finemet group [58,59], it is more likely that during solidification, Cu atoms are involved in the formation of certain intermetallic compound rather than fcc-Cu clusters. It is thus presumed that GdCu clusters with simple cubic structure [60] are formed and are involved in the bcc-Fe dendrite formation in the alloys investigated in the current study. Although there was only bcc-Fe phase revealed in XRD pattern (see Fig. 4), according to optical micrograph given in Fig. 6 Ref steel has typical ferritic structure (light-shaded grains) with dark uniformly dispersed pearlite platelets typically 10 μm long and Fe3CIII cementite on the grain boundaries. Note that the SEM micrograph depicted in Fig. 5 for Ref was etched with 5% Nital in order to reveal its microstructural features, as the compositional contrast in this sample was negligible. Uniaxial compression tests were performed for all alloys studied and the results are collected in Table 6. Note that εm reported here is a maximum plastic strain. It was defined as: εm ¼ εt −ε0:2

ð4Þ

where: εt and ε0.2 stand for total strain at fracture and offset yield point measured at plastic strain of 0.2%, respectively. The considered mechanical parameters have comparable values with those obtained for other rapidly quenched alloys, as reported in the literature. For example, Inoue [9] found that bulk nanocrystalline Zr65Al7.5Ni10Cu7.5Pd10 subjected to uniaxial compression have Rm of 1820 MPa and εm of 0.005, while the same amorphous as-cast alloy exhibit Rm of 1630 MPa and almost no plastic elongation. Fe73−xNb4Hf3YxB20 (x = 0–3) bulk metallic glass

Table 6 Yield strength R0.2, compressive strength Rm and maximum plastic strain εm determined in static uniaxial compression test for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0 − 15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref). Composition

R0.2 [MPa]

Rm [MPa]

εm [mm mm −1]

Ref x=0 x=8 x = 11 x = 13 x = 15 y=7

716 ± 11 1382 ± 120 1909 ± 290 2036 ± 134 2397 ± 237 2250 ± 373 1690 ± 140

918 ± 64 1939 ± 248 2076 ± 296 2212 ± 43 2532 ± 89 2368 ± 507 1697 ± 140

0.203 ± 0.048 0.287 ± 0.118 0.045 ± 0.007 0.019 ± 0.008 – – –

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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P.J. Bardziński, P. Błyskun / Materials and Design xxx (2016) xxx–xxx

studied by Long et al. [5], exhibit E = 177–190 GPa, Rm of 3200–3490 MPa and εm = 0.02 with brittle fracture. Fracture morphologies were shown in Fig. 8. In case of Ref steel as well as in the alloys with x = 0–11, shear fracture was observed. In the alloys with x = 8 and 11, maximum shear stress plane was inclined by nearly 45° to the direction of loading force. In the samples of the alloys with x = 13 and y = 7 shear-cleavage fracture occurred. Finally, brittle fracture was observed in the alloy with x = 15. Sharp tips of bcc-iron dendrites present in the alloys with x = 8–15 might act as stress concentrators, contributing to brittle fracture. The similar notching effect was described for graphite in cast iron [11]. According to Song et al. [61], yield strength of rapidly solidified Fe-Mn alloys decreased with ascending elongation of needle-shaped ε-martensite phase. Brittleness of high silicon steels could be caused by the presence of impurities on grain boundaries [24]. Chen et al. [62] observed that higher quenching rate correlates with a tendency of certain phases precipitation on grain boundaries in rapidly solidified bulk Mg-6Zn-3Sn-2Al-0.2Ca alloy. Here, higher concentration of Cu and Gd on the grain boundaries together with the presence of hard borides, might also contribute to the higher brittleness of the rapidly solidified alloys studied with respect to Ref steel. As can be seen from comparing engineering stress–strain curves shown in Fig. 7, increasing Si content leads to stepwise improvement of yield R0.2 and compressive Rm strength at the expense of plastic deformation region represented here by maximum plastic strain εm. This result complies with the increase of brittleness of steel with ascending Si content, due to formation of D03 and B2 phases [23,63]. In the current study the precipitation of C14 Laves/D8a in the interdendritic regions, as shown in Fig. 5, could be the major contributor of this effect. According to Fig. 9, there is also clear correlation between unit cell volume V of fcc-Fe16Hf6Si7 phase, R0.2 and Rm on Si content. It is worth noting here that SDAS reported in Table 5 is decreasing accordingly with higher Si content in the alloy. It was found that lowering the dendrite arm spacing increases interaction between them, leading to improved ultimate tensile strength of cast iron [64]. It was confirmed by Hemanth [65] who found that larger dendrite arm spacing is correlated with lower ultimate tensile strength in the case of austempered chilled ductile iron. Similar findings were reported by Kaiser et al. [53] who found that the improvement in HV10, E, ultimate tensile strength and R0.2 while decrease in εm are connected with reduced SDAS in Co66Cr28Mo6 prepared by lost wax investment casting.

Fig. 7. Typical uniaxial compressive engineering stress–strain curves for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0 − 15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref).

Fig. 8. Fracture morphologies after uniaxial compression test of rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref).

Note that Rm and εm given in Table 6 for Ref steel which deforms plastically regardless of the stress applied, refers to the arbitrarily chosen strain at which the compressive stress–strain curve changes the slope

Fig. 9. Dependency between unit cell volume V of fcc-Fe16Hf6Si7 phase, yield strength R0.2 and compressive strength Rm on Si content in rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 8–15) alloys.

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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9

to higher value, while ε N ε0.2. The alloy with x = 0 exhibits εm comparable with Ref while R0.2 was two times higher. Maximum value of R0.2 and Rm was recorded for the alloy containing 13 at.% Si. This behaviour could be owed to its microstructural feature shown in Fig. 5, namely, that the hard Hf-rich C14 Laves/D8a phase grains that precipitated in the interdendritic zones of bcc-Fe dendrites are intersected with 0.3 ± 0.1 μm thick softer bcc-Fe platelets. Substitution of 7 at.% Si by La in the alloy with x = 15 is detrimental to R0.2, Rm and εm, probably owed to the poor mechanical performance of the peritectic structure observed in Fig. 5. Material's properties derived in three-point flexural test were calculated according to relations proposed by Katarzyński et al. [66]. Assuming pure bending (without shear), one can write: W¼

πϕ3 32

ð5Þ

Rb ¼

Fl0 4W

ð6Þ

f 0:2 ¼

Eb ¼

2  10−3 l 20 6ϕ

Fig. 10. Flexural strength Rb-deflection f plots for rapidly quenched bulk Fe89−x Hf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref).

ð7Þ

l 30 1 n F 1 F n−1 F n ∑ þ⋯þ þ 48W n i¼1 f 1 f n−1 f n

ð8Þ

where: ϕ, F, l0, f, W, Rb, f0.2, Eb, stand for mean sample diameter, force applied perpendicular to the half length of the sample, distance between supports, deflection, section modulus, flexural strength, deflection-at-yield and bending modulus, respectively. The values of thus obtained quantities are presented in Table 7. Due to ductile properties of Ref steel, Rb and f values for this material were taken from the point at which Rb plotted as a function of f changes its slope, as can be seen in Fig. 10. For the rapidly solidified samples, all of which exhibited brittle fracture at bending, Rb and f were the maximum stress and deflection the sample could carry, respectively. For the alloys with x = 0–11 values of both Rb and f were larger than for steel, by 25% on average, while Eb was smaller by about 40% for the alloys with x = 0–15. The largest value of Rb, i.e., 1567 MPa was obtained for x = 0. By studying the alloy microstructure given in Figs. 5 and 6, the striking feature is that the grains are largest among the alloys investigated and that their edges are smooth. Since larger grains have less grain boundaries that hinder the dislocation movement, this might lead to improved ductility [53]. Compared to other rapidly quenched alloys reported in the literature, the obtained values are moderate. For instance, Inoue [9] performed three-point flexural test for bulk amorphous Zr65Al7.5Ni10Cu7.5Pd10 and found Rb of 2000 MPa and near zero f. After nanocrystallization, Rb increased to 4400 MPa and f to about 0.5 mm. In general, Rb decreased with ascending Si content, while f0.2 remain unaffected. It was found, according to Fig. 12, that the lattice expansion of

Table 7 Mean sample diameter ϕ, section modulus W, flexural strength Rb, deflection-at-rupture f, deflection-at-yield f0.2 and bending modulus Eb determined in three point flexural test for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref). Composition

ϕ [mm]

W [mm3]

Rb [MPa]

f [mm]

f0.2 [mm]

Eb [GPa]

Ref x=0 x=8 x = 11 x = 13 x = 15 y=7

3.14 3.00 3.03 2.96 2.96 3.00 3.02

3.04 2.66 2.72 2.53 2.53 2.66 2.71

1147 1567 1263 1454 1070 879 791

0.28 0.35 0.32 0.37 0.24 0.23 0.39

0.057 0.060 0.059 0.061 0.061 0.060 0.059

211.32 ± 14.5 138.44 ± 2.00 120.58 ± 1.55 117.88 ± 4.08 132.73 ± 3.10 118.53 ± 4.72 84.77 ± 0.68

bcc-Fe and, in lesser extent, C14 Laves phase is correlated with the value of Eb in the whole composition range. In contrast, the similar dependency for R0.2 and Rm was only observed for the alloys that contain from 11 to 15 at.% Si. Fracture surfaces of bent samples were presented in Fig. 11. For all rapidly quenched samples, brittle fracture was observed, coarsegrained for x = 0, y = 7 and fine-grained for x = 8–15. Results of hardness measurements performed by two methods: Vickers HV3 and nanoindentation NHT were compared in Table 8. Although both reveal similar dependency on Si-content, as presented in Fig. 13, results of NHT measurements seem to be overestimated. Instrumental hardness determined in the nanohardness measurement and converted to HV, is fraught with error resulting from imperfect indenter geometry, especially when small indentation depth of 1.45 μm was used. The other factors include local compositional nonuniformities, indentation performed on grain boundaries with size comparable with the indenter tip and smoothness derogations of the sample surface [67]. Shin et al. who investigated the effect of Si-content on micro-Vickers hardness of Si-steels found [24] (Fig. 6), that the slope of said dependency becomes larger when Si content is higher than ca. 7–9 at.% and this relation is not affected by heat treatment. Despite larger macro- and nano-hardness values overall found in the current study, owed to the presence of h-Fe2Hf0.8Ta0.2 and fcc-Fe16Hf6Si7 phases, the dependency of at.% Si on hardness is very similar to the one described by Shin et al. Taking into account the SEM-EDX data given in Table 1, this could be attributed to the increase in Si-content in bcc-Fe phase. Note that Shin et al. [24] chosen to treat lower- and higher-Si-content parts of hardness relation separately, while in the current study it is proposed to describe them for the currently investigated alloys by a single relation in the whole composition range: HV ¼ HV 0 þ A exp


x B

ð9Þ

with the parameters: a) for Vickers macrohardness HV3 test: HV 0 ¼ 485:0  15:6 HV A ¼ 0:8  0:9 HV B ¼ 2:4  0:4 at:%:

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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Fig. 11. Fracture surfaces of bent samples of rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref).

The coefficient of determination R2 = 0.969 and the reduced χ 2 = 0.418. b) for Vickers nanohardness NHT test: HV 0 ¼ 532:7  19:1 HV A ¼ 28:1  10:5 HV B ¼ 6:1  0:8 at:%: In this case, the fit was more accurate, yielding R2 = 0.997 and the reduced χ2 = 0.101. The data was fitted with damped least-squares [68] regression method without weighting utilizing the Origin package. As can be

seen from Table 8 the alloy with y = 7 has comparable HV3 and NHT hardness with alloy with x = 8, thus it can be described with Eq. (9). As the difference between those two alloys is that 7 at.% Fe is substituted by La and the Si-content is equal, it proves that hardness of the studied alloys is primarily affected by the amount of this element. According to Tabor [69], work hardening and deformation that accompanies the indentation test leads to the presence of correlation between yield strength and hardness measured in such way. In the current study it was found that yield strength R0.2, compressive strength Rm and Vickers hardness increases with comparable monotonicity for the alloys with x = 0–13, as exemplified in Fig. 14. Mutual dependency between R0.2 and hardness, obtained either from macroor nanohardness tests was shown in Fig. 15. Relation between this

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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Fig. 12. Dependency between unit cell volumes V of bcc-Fe and h-Fe2Hf0.8Ta0.2 phases, bending modulus Eb, yield strength R0.2 and compressive strength Rm on Si content in rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) alloys.

11

Fig. 13. Least-squares exponential function fit for Vickers hardness HV determined in Vickers hardness test HV3 and nanoindentation test NHT for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) alloys.

quantities was fitted with damped least-squares [68] regression method and described with logarithmic function:

Utrasonic through-transmission method was chosen to determine the dynamic elastic moduli of the studied alloys and the results are summarized in Table 9. Longitudinal vL and transverse vT velocities of ultrasonic waves obtained and bulk volumetric mass density d were used to calculate the following quantities [70,29,71]:

R0:2 ¼ AlnðHV−BÞ

E ¼ dvT 2

ð10Þ

with the parameters: a) for Vickers macrohardness HV3 test: A ¼ 430  37 MPa= ln HV

3vL 2 −4vT 2 vL 2 −vT 2

ð11Þ

G ¼ dvT 2

ð12Þ

4 B ¼ dvL 2 −d vT 2 3

ð13Þ


2

B ¼ 468  16 HV: 2

2

The coefficient of determination R = 0.695 and the reduced χ = 1.773. b) for Vickers nanohardness NHT test:

vL −2 vT ν ¼
2 vL 2 vT −2

ð14Þ

where: E, G, B, ν stands for elastic modulus, shear modulus, bulk modulus and Poisson's ratio, respectively.

A ¼ 405  11 MPa= ln HV B ¼ 534  6 HV: The fitting accuracy for NHT is better than obtained for HV3, that is, R2 = 0.959 and the reduced χ 2 = 0.240.

Table 8 Vickers hardness determined in Vickers macrohardness test HV3 and nanoindentation test NHT for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref). Composition

HV3

NHT

Ref x=0 x=8 x = 11 x = 13 x = 15 y=7

257 ± 18 501 ± 31 494 ± 20 573 ± 32 649 ± 30 856 ± 53 474 ± 0.52

363 ± 14 564 ± 32 632 ± 23 707 ± 67 783 ± 34 862 ± 12 662 ± 39

Fig. 14. Dependency of Vickers nanohardness NHT, yield strength R0.2 and compressive strength Rm on Si content in rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) alloys.

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

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Fig. 16. Correlation between the elastic E, shear G and bulk B moduli, Poisson's ratio ν, flexural strength Rb and deflection-at-rupture f for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) alloys.

ν comparable with Ref steel, i.e., near 0.28. Values of several mechanical properties of rapidly quenched Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) alloys were normalized and compared in Fig. 16 in order to study their monotonicity. It was revealed that E ∝ G and that both are inversely proportional to B ∝ ν ∝ Rb ∝ f while plotted with ascending Si content. It was found that B, ν, Rb and f exhibit a clear maximum at x = 11. Note that E and G have local minimum at this composition. Nanoindentation test was also used to determine the elastic moduli of the studied alloys and the results were summarized in Table 10. The reduced modulus Ered was defined as follows [73–76]: pffiffiffi S π Ered ¼ pffiffiffiffiffiffi 2 Ap

Fig. 15. Correlations between Vickers macrohardness HV3 (top) as well as nanohardness NHT (bottom) and yield strength R0.2 for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) alloys. Numbers near datapoints indicate the Si content; dashed line represents the fitting with logarithmic function.

The values of B for alloys with x = 0, 8 and 13 are higher by approximately 10 GPa than 167 GPa obtained for Ref steel. Alloy with x = 15 have B higher than steel by only 5 GPa, while x = 11 exhibit B higher by nearly 20 GPa. E determined in the ultrasonic method for x = 0–13 is lower than for steel by about 35 GPa. Alloy with x = 15 have E comparable with Ref. G for all investigated alloys, except with x = 15, is by 15 GPa lower than for Ref. Lowest moduli are reported for the alloy with y = 7. ν of about 0.32, obtained for alloys with x = 0–13 are also common for the iron-based metallic glass with similar composition, reported in the literature [72]. Alloys with x = 15 and y = 7 exhibit Table 9 Time delay error-corrected longitudinal vL and transverse vT velocities of ultrasonic waves, elastic modulus E, shear modulus G, bulk modulus B and Poisson's ratio ν determined in ultrasonic through-transmission method for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref), given together with bulk volumetric mass density d. Composition

d [kg m −3]

vL [m s −1]

vT [m s −1]

E [GPa]

G [GPa]

B [GPa]

ν [−]

Ref x=0 x=8 x = 11 x = 13 x = 15 y=7

7780 8390 8190 8080 7980 7930 7820

6016 5703 5787 5882 5837 5986 5244

3321 2876 2940 2938 2972 3246 2975

219.81 184.52 187.75 186.05 186.79 215.86 174.79

85.81 69.40 70.79 69.75 70.49 83.55 69.21

167.17 180.35 179.89 186.56 177.90 172.74 122.76

0.281 0.329 0.326 0.334 0.325 0.292 0.263

ð15Þ

where Ap is projected contact area and S is the contact stiffness, i.e., slope of applied load-penetration depth curve recorded at unloading. As noted by Oyen and Cook [73], Ered is a series sum of the plane strain moduli of the material and the indenter. It was found that the values of elastic modulus E determined in ultrasonic through-transmission test are comparable with indentation modulus Eind, which was obtained utilizing Poisson's ratio ν given in Table 9. They also comply with the values of the reduced modulus Ered, which, according to Eq. (15) is derived solely from nanoindentation experiment. According to transformed Eq. (2) [28] given by Oliver and Pharr, Eind is an approximation of the Young modulus of the studied sample [77]: Eind ¼



!−1 

−1 1−0:072  1−ν2 Ered − : 1−ν2 1141

ð16Þ

The values of E and Eind found in the current study for rapidly quenched alloys are very close to the ones reported for iron-based

Table 10 Indentation modulus Eind, reduced modulus Ered and indentation relaxation Rind determined in nanoindentation test for rapidly quenched bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7) alloys compared with steel (Ref). Composition

Eind [GPa]

Ered [GPa]

Rind [%]

Ref x=0 x=8 x = 11 x = 13 x = 15 y=7

204.39 ± 27.67 186.12 ± 12.82 176.08 ± 2.49 185.56 ± 5.73 188.89 ± 7.95 197.88 ± 5.95 178.12 ± 8.54

185.52 ± 20.26 176.49 ± 10.30 168.13 ± 2.03 176.66 ± 4.62 178.31 ± 6.39 181.97 ± 4.60 163.95 ± 6.74

−4.52 ± 1.41 −4.56 ± 0.20 −2.72 ± 0.06 −3.39 ± 0.83 −3.21 ± 0.59 −2.93 ± 0.41 −4.84 ± 0.68

Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054

P.J. Bardziński, P. Błyskun / Materials and Design xxx (2016) xxx–xxx

bulk metallic glass, which are in the range of 120–170 GPa [78], 170–200 GPa [79]. As can be seen by comparing the data collected in Tables 7 and 9, Eb and E of Ref steel are comparable while Eb for rapidly solidified alloys are substantially lower than E. Keeping in mind the grain size difference across the sample radius, given in Table 5, it could be explained by sample anisotropy. According to Lucca et al. [80], in the test in which the maximum penetration depth is kept constant for a period of time, the measured change in force is defined as indentation relaxation Rind [81,82]: Rind ¼ 100

F2− F1 : F1

ð17Þ

The Rind, given in percent, is a measure of rheological properties of the material [83]. Data presented in Table 10 shows that alloys with x = 0 and y = 7 exhibit Rind comparable with steel, while this quantity is lower by about 1% for the alloys containing Si and no La. 4. Summary and conclusions In the current study, a range of novel multicomponent iron-based alloy compositions, Fe89−xHf4Ta1Cu1Gd1SixB4 (x = 0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−yB4 (y = 7), was investigated. Rapidly quenched microcrystalline rods with reinforcing hard crystalline phases were synthesized by copper mold injection casting. Besides bcc-Fe and borides identified in all rapidly quenched samples, C14 Laves phase was present in the alloys with x = 0–15 as well as y = 7 and D8a phase was identified in the alloys with x = 8–15. It was noted that fcc-Fe16Hf6Si7 D8a-type phase was described only two times in the known literature [40,41] before this study. La substitution for Si in the alloy with x = 15 was found to have a detrimental effect on the formation of D8a in favor of the C14 Laves phase. In the alloy with x = 0 one can distinguish nearly hexagonally shaped C14 Laves phase grains uniformly dispersed with round bcc-Fe with rather smooth edges and comparable size. Microstructure of the alloys with x = 8–15 contains starfish-like bcc-Fe dendrites with six-fold rotational symmetry as well as lamellar eutectic structure. Microstructural features of the alloy with y = 7 include lamellar peritectic in the La–Fe–Si system in which other phases are embedded. Average grain diameter and SDAS decreased from the middle to the edge of the rod due to cooling rate gradient across the rod radius during quenching process. The study revealed that ascending Si content leads to: • reduction in values of both unit cell parameters of C14 phase in the alloys studied while causes the expansion of D8a lattice. • decrease in secondary dendrite arm spacing (SDAS) from 1.6 to 0.3 μm by average in the middle of the rods. • stepwise improvement of yield R0.2 and compressive Rm strength at the expense of plastic deformation region, what was attributed to the precipitation of C14 Laves/D8a in the interdendritic regions of bcc-Fe. • decrease in flexural strength Rb with no effect on deflection-at-yield f0.2. • increase in hardness of the studied alloys, what was described by exponential function. There is also clear correlation between unit cell volume V of fccFe16Hf6Si7 phase, R0.2 and Rm on Si content. For the alloys with x = 0–13, relation between yield strength and Vickers hardness was described with logarithmic function. It was found, that the lattice expansion of bcc-Fe and, in lesser extent, C14 Laves phase is correlated with the value of bending modulus Eb in the whole composition range. It was revealed that E ∝ G and that both are inversely proportional to B ∝ ν ∝ Rb ∝ f while plotted with ascending Si content. It was found that B, ν, Rb and f exhibit a clear maximum at x = 11. E and G exhibit local

13

minimum at this composition. Maximum values of R0.2 and Rm were recorded for the alloy containing 13 at.% Si, to be approx. 2397 and 2532 MPa, respectively. The largest value of Rb, i.e., approx. 1567 MPa was obtained for the alloy with x = 0. The values of E and Eind found in the current study for rapidly quenched alloys are very close to the ones reported for iron-based bulk metallic glass, which are in the range of 120–200 GPa [78,79]. The optimum combination of strength and ductility can lead to select the alloy with x = 0 as a potential candidate for structural applications. It can be characterized with twice higher strength parameters than the reference steel with moderate level of plastic deformation. Acknowledgements The work of Piotr Józef Bardziński was co-financed from the European Union funds under the European Social Fund. Authors want to acknowledge the outstanding technical assistance of Mr Piotr Gutkiewicz PhD, Mr Janusz Szymkowski PhD, Mr Piotr Kotowski PhD, Mr Miłosz Siczek PhD, Mr Tomasz Baraniecki PhD, Mr Bartosz Babiarczuk MSc and Ms Anna Zięty MSc. The Origin package (OriginLab, Northampton, MA) was used under the terms of the University's License. Appendix A. Supplementary data Supplementary data to this article can be found online at http://dx. doi.org/10.1016/j.matdes.2016.04.054. References [1] A. Inoue, A. Takeuchi, Recent progress in bulk glassy, nanoquasicrystalline and nanocrystalline alloys, Mater. Sci. Eng. A 375–377 (2004) 16–30. [2] T.-H. Lee, Y. Kawamura, A. Inoue, S.-S. Cho, T. Masumoto, Mechanical properties of rapidly solidified Al-Si-Ni-Ce P/M alloys, Scr. Mater. 36 (1997) 475–480. [3] E. Karaköse, M. Keskin, Structural investigations of mechanical properties of Al based rapidly solidified alloys, Mater. Des. 32 (2011) 4970–4979. [4] A. Inoue, T. Masumoto, H. Tomioka, Microstructure and mechanical-properties of rapidly quenched L20 and L20 + L12 alloys in Ni-Al-Fe and Ni-Al-Co systems, J. Mater. Sci. 19 (1984) 3097–3106. [5] Z. Long, Y. Shao, F. Xu, H. Wei, Z. Zhang, P. Zhang, X. Su, Y effects on magnetic and mechanical properties of Fe-based Fe-Nb-Hf-Y-B bulk glassy alloys with high glass-forming ability, Mater. Sci. Eng. B 164 (2009) 1–5. [6] J. Olszewski, J. Zbroszczyk, K. Sobczyk, W. Ciurzynska, P. Bragiel, M. Nabialek, J. Swierczek, M. Hasiak, A. Lukiewska, Thermal stability and crystallization of iron and cobalt-based bulk amorphous alloys, Acta Phys. Pol. A 114 (2008) 1659–1666. [7] M. Hasiak, K. Sobczyk, J. Zbroszczyk, W. Ciurzynska, J. Olszewski, M. Nabialek, J. Kaleta, J. Swierczek, A. Lukiewska, Some magnetic properties of bulk amorphous Fe–Co–Zr–Hf–Ti–W–B–(Y) alloys, IEEE Trans. Magn. 44 (2008) 3879–3882. [8] K. Błoch, Magnetic properties of the suction-cast bulk amorphous alloy: (Fe0.61Co0.10Zr0.025Hf0.025Ti0.02W0.02B0.20)96Y4, J. Magn. Magn. Mater. 390 (2015) 118–122. [9] A. Inoue, Bulk amorphous and nanocrystalline alloys with high functional properties, Mater. Sci. Eng. A 304–306 (2001) 1–10. [10] T. Minemura, A. Inoue, Y. Kojima, T. Masumoto, Microstructure and mechanical properties of non-equilibrium austenite in Fe–C–(Mo, W) systems rapidly quenched from melts, Trans. Iron Steel Inst. Jpn. 22 (1982) 934–941. [11] J. Zhang, F. Ren, Influence of primary austenite dendrite on the microstructure and mechanical properties of gray cast iron, in: J. Bu, Y.H. Kim (Eds.),Adv. Eng. Mater. III, PTS 1–3, Volume 750–752 of Adv. Mater. Res 2013, pp. 450–453. [12] L. Machon, G. Sauthoff, Deformation behaviour of Al-containing C14 laves phase alloys, Intermetallics 4 (1996) 469–481. [13] D.D. Risanti, G. Sauthoff, Microstructures and mechanical properties of Fe–Al–Ta alloys with strengthening Laves phase, Intermetallics 19 (2011) 1727–1736. [14] G. Haour, P. Boswell, Rapidly solidified materials — a status report, Mater. Des. 8 (1987) 10–12. [15] K. Wieczerzak, P. Bala, M. Stepien, G. Cios, T. Koziel, Formation of eutectic carbides in Fe–Cr–Mo–C alloy during non-equilibrium crystallization, Mater. Des. 94 (2016) 61–68. [16] M. Ibrahim, S. Alkahtani, K. Abuhasel, F. Samuel, Effect of intermetallics on the microstructure and tensile properties of aluminum based alloys: role of Sr, Mg and Be addition, Mater. Des. 86 (2015) 30–40. [17] W.-W. Park, B.-S. You, N.J. Kim, Microstructure and mechanical properties of rapidly solidified Al–Si–Fe–X base alloys, Mater. Des. 17 (1996) 255–259. [18] L. Sheng, W. Zhang, J. Guo, Z. Wang, H. Ye, Microstructure evolution and elevated temperature compressive properties of a rapidly solidified NiAl–Cr(Nb)/Dy alloy, Mater. Des. 30 (2009) 2752–2755.

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Please cite this article as: P.J. Bardziński, P. Błyskun, Mechanical properties and structure of rapidly solidified bulk Fe89−xHf4Ta1Cu1Gd1SixB4 (x=0–15) and Fe74Hf4Ta1Cu1Gd1LaySi15−..., Materials and Design (2016), http://dx.doi.org/10.1016/j.matdes.2016.04.054