Advanced Powder Technology 27 (2016) 461–470

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Advanced Powder Technology

journal homepage: www.elsevier .com/locate /apt

Original Research Paper

Synthesis of amorphous Fe75Si20�xMxB5 (M = Ti, Ta, Zr) via wetmechanical alloying and its structural, thermal and magneticcharacterisation

http://dx.doi.org/10.1016/j.apt.2016.01.0270921-8831/� 2016 The Society of Powder Technology Japan. Published by Elsevier B.V. and The Society of Powder Technology Japan. All rights reserved.

⇑ Corresponding author.E-mail address: Bogdan.Neamtu@stm.utcluj.ro (B.V. Neamt�u).

B.V. Neamt�u a,⇑, H.F. Chicinas� a, T.F. Marinca a, O. Isnard b,c, I. Chicinas� a, F. Popa a

aMaterials Science and Engineering Department, Technical University of Cluj-Napoca, 103-105, Muncii Avenue, 400641 Cluj-Napoca, RomaniabUniversité Grenoble Alpes, Institut NEEL, F-38042 Grenoble, FrancecCNRS, Institut NEEL, 25 rue des martyrs, BP166, F-38042 Grenoble, France

a r t i c l e i n f o a b s t r a c t

Article history:Received 22 July 2015Received in revised form 8 January 2016Accepted 27 January 2016Available online 4 February 2016

Keywords:Fe-based amorphous powdersSoft magnetic materialsMechanical alloyingTransitional metals for Si substitution

Amorphous Fe75Si15B5M5 powders with M = Ti, Ta and Zr, were synthesized by mechanical alloying andtheir thermal stability, crystalline state, particles morphology and magnetic properties were investigatedin this study. It is found that the time to achieve amorphisation during milling increases by the substi-tution of any of the above transition metals 5% for silicon. The particle size distribution is dependenton the mechanical properties of the transitional metals used. Crystallisation temperature of the alloyswas enhanced with up to 24% compared to the reference ternary alloy Fe75Si20B5. Magnetic measure-ments revealed a higher saturation magnetisation in the case of the amorphous alloy containing titaniumthan for the reference amorphous alloy. XRD patterns for the samples heated up to 900 �C reveal the pres-ence of carbides as a result of powder contamination by the process control agent (PCA).� 2016 The Society of Powder Technology Japan. Published by Elsevier B.V. and The Society of Powder

Technology Japan. All rights reserved.

1. Introduction

In the last decades, great progress has been achieved in thesearch for metallic amorphous materials, leading to the develop-ment of many amorphous systems, such as Ti-based, Zr-based,Cu-based, Ni-based and Fe-based alloys [1–4]. Due to their disor-dered structure, amorphous materials may possess unique proper-ties, such as high mechanical strength, high corrosion and wearresistance, low Young’s modulus and large elasticity. These proper-ties impose them as important candidates for a series of industrialapplications in various fields. In the case of soft magnetic materials,the relatively high electrical resistivity of amorphous materialsensures, in AC applications, low eddy current losses, and high mag-netic permeability [1].

Amorphous alloys can be obtained by two main routes: rapidquenching and mechanosynthesis. The rapid quenching routeimplies ultra-fast solidification from liquid state, which can beachieved by several techniques such as: arc-melt casting, suctioncasting, low pressure copper mould casting, high pressure die cast-ing, squeeze casting, unidirectional melting and centrifugal casting[2,5,6]. The mechanosynthesis route (detailed in the following

paragraph) involves the processing in high energy mills of a mix-ture of elemental, pre-alloyed powders or amorphous ribbons[7,8].

Mechanical alloying technique (MA), developed at the begin-ning of the 1970s as a method to obtain oxide dispersion strength-ened superalloys, can be considered as an effective and versatileapproach to produce non-equilibrium structures/microstructuresas amorphous alloys, extended solid solutions, nanocrystallinematerials etc. [9]. The solid state reactions that occur in mechani-cally activated powders (a mixture of elemental or pre-alloyedpowders) during high energy ball milling leads to the alloy forma-tion. Basically, the particles are repeatedly flattened, cold welded,fractured and re-welded leading to new particles characterisedby a sandwich like structure (layered structure consisting in vari-ous combinations of the starting elements). The large number ofstructural defects induced by milling (vacancies, dislocations, andincreased number of grain boundaries), together with a slightincrease of the powder’s temperature during milling, favours theatomic inter-diffusion and finally the alloy formation [9]. In orderto prevent the agglomeration and cold welding of the particles,process control agents (PCA) are being added inside the millingvials. The quantity of PCA introduced in the milling vials is usuallycomprised between 1 and 5 wt.% of the total amount of powderprocessed [9]. The PCA are, most commonly, organics substances,

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which act as surfactants, leading to the achievement of a balancebetween cold welding and fracturing phenomena by reducing thequantity of cold welding processes. Due to the fact that most PCAhave a low melting/boiling point, in almost all the cases, thedecomposition of this PCA during the milling experiment mustbe considered and a contamination of the processed powder withoxygen, carbon, hydrogen, etc. is expected [10].

Preparation of amorphous alloys in powder form leads to theelimination of several drawbacks of amorphous materials (pre-pared as thin ribbons or wires) such as impossibility to create com-plex 3D shapes. This drawback is, somehow, eliminated by afurther processing (compaction) of the amorphous powdersthrough powder technology techniques such as spark plasmasintering. Different studies have highlighted the efficiency ofmechanical alloying to induce amorphisation in various Fe-basedalloys such as: Fe–Si–B [11], Fe–Zr–B [12], Fe–Al–P–B–C [13],Fe–Co–Ni–Zr–B [14], Fe–Cu–Nb–Si–B [15], and Fe–Ni–P–Si [16].

Numerous alloys are derived from Fe–Si–B based mixturethrough addition of small quantities of several elements (Cu, Nb,Zr, Ti, Hf, Ta, W, P, C etc.) such as the FINEMET alloys. In spite oftheir relatively large magnetisation, the main drawback of theFe-based amorphous alloys is the relatively low thermal stabilityof the amorphous phase. It was proved that the addition of severaltransition metals such as Ti, Zr, Nb and Ta leads to the improve-ment of the thermal stability domain of the amorphous phase[17,18].

The goal of the present research on mechanical alloying tech-nique is to develop new amorphous powders of Fe75Si15B5M5

(M = Ti, Ta or Zr) type in order to increase the thermal stability ofthe classical Fe75Si20B5 alloy. Indeed an enhanced thermal stabilityis desired to increase the chances to preserve the amorphous stateof the alloy during compaction stage [19]. In addition, the obtain-ing of a higher crystallisation temperature would allow higher sin-tering temperature leading to compacts with higher density andenhanced magnetic properties.

2. Experimental procedure

Ternary Fe75Si20B5 alloy and 3 alloys deriving from it weresynthesised for 5 at.% substitution of transition metals for Si. Start-ing mixtures corresponding to the composition Fe75Si20B5 andFe75Si15B5M5, where M is titanium, tantalum or zirconium, wereprepared starting from high purity elemental powders: HöganäsNC 100.24 iron powder, silicon powder 99.9% (Alfa Aesar), boronpowder 99.9% (Alfa Aesar), titanium powder 99.5% (Alfa Aesar),tantalum powder 99.6% (Alfa Aesar) and zirconium powder 99.2%(Alfa Aesar). To ensure the starting mixture’s homogeneity thepowders were homogenized in a turbula-type blender for 30 min.Each of the homogenized powders was then processed in a high-energy planetary ball mill (Fritsch Pulverisette 6). The milling bod-ies were manufactured from tempered steel. The diameter of themilling balls was 14 mm, and the ball to powder ratio (BPR) waschosen to be 16:1. This led to a filling factor of 60%. In order toavoid oxidation, argon atmosphere was chosen. Several millingtimes were used ranging from 2 h up to 60 h. The powders werewet milled and the chosen PCA was benzene – C6H6. Each timewhen sampling was done, 1 ml of PCA was added in order to supplythe quantity of benzene that evaporates. The vial rotation speedwas set to 350 rpm. The collection of samples was done under Aratmosphere in a glove-box (Iteco Engineering SGS 30).

The particle morphology was investigated by scanning electronmicroscopy using a Jeol-JSM 5600 LV scanning electron micro-scope. The particle size distribution has been determined using aLaser Particle Size Analyzer (Fritsch Analysette 22-Nanotec), withan analysis field of 10 nm – 2000 lm.

Structural evolution of the ternary and quaternary alloys wasinvestigated by X-ray diffraction (XRD), in the angular range of2h = 20–110�, on an Inel Equinox 3000 powder diffractometer,working with Co Ka radiation (k = 1.7903 Å).

The thermal stability of the alloys was studied by differentialscanning calorimetry (DSC) on a Setaram Labsys apparatus usingas reference high purity alumina powder. The measuring rangewas 20–900 �C, with a heating/cooling rate of 20 �C/min underargon atmosphere.

The M(H) curves were recorded using the sample extractionmethod in a continuous magnetic field up to 7 T. The magnetisa-tion measurements were done at room temperature.

The Curie temperature of the amorphous powders was deter-mined from the thermomagnetic measurements. The powder wassealed in a quartz tube under vacuum and the used heating ratewas 10 �C/min. The value of the applied magnetic field during ther-momagnetic measurements was 0.1 T.

3. Results and discussion

The X-ray diffraction patterns of the Fe75Si20B5, Fe75Si15B5Ti5,Fe75Si15B5Ta5 and Fe75Si15B5Zr5 alloys have been recorded for dif-ferent wet mechanical alloying durations up to 60 h as can be seenin Fig. 1. In the XRD patterns of the starting samples the Braggreflections characteristic for the elemental Fe and Si powders arenoticed. Additionally, for the quaternary alloys, the Bragg reflec-tions of the transitional metal used for substitution can beobserved also. Due to the low boron content as well as, its lowX-ray scattering factor and its amorphous state the Bragg reflec-tions are too weak to be observed. In the case of ternary Fe75Si20B5

alloy a nanocrystalline iron based solid solution is beginning toform after 2 h of wet mechanical alloying. At this milling time acertain amount of Si remains unreacted as can be seen from itsmost intense Bragg peak observed in the diffraction pattern atabout 2h = 33�. Increasing the milling time at 5 h leads to completealloying of the elements and the material consists in a single phase,iron based solid solution. The relatively broad peaks indicate thatthe solid solution is obtained in nanocrystalline state. A certainamount of amorphous phase exists in material and this amorphousphase is situated in the grain boundaries [9]. After 10 h of millingan important broadening of the Fe-based solid solution can beobserved suggesting the decrease of the crystallite size and theincrease of the amount of amorphous phase. The increase ofmilling time up to 20 h is leading to complete amorphisation ofthe alloy. This amorphisation of the material is shown by the pres-ence of a single broadened feature on the XRD pattern. The absenceof any Bragg peak is revealing the absence of the long range atomicorder. Further increase of milling time is leading to another phe-nomenon: a mechanically induced crystallisation. Mechanicallyinduced crystallisation is shown by the appearance of a Bragg peakon the low angle part of the broad feature which consequentlytends to becomes asymmetric.

In the case of quaternary alloys, the nanocrystalline iron basedsolid solution is obtained also after 5 h of mechanical alloyingwhen the diffraction peaks of the starting phases are gone. Energydispersive X-ray analysis performed on such sample reveals thatthe sample is homogeneous at nanometric scale and that the ele-ments are homogeneously distributed in the alloys. In the case ofFe75Si15B5Ta5 alloy, the contribution of the amorphous phase canbe observed in the diffraction pattern for this milling time: a halois visible at the base of the main Bragg reflection of iron based solidsolution. As shown from the line broadening, the crystallite size isdiminishing upon increasing the milling time for all the investi-gated alloys. In the case of Ta and Zr containing alloys, the XRD

Fig. 1. Evolution of the X-ray diffraction pattern for: (a) ternary Fe75Si20B5 alloy, (b) quaternary Fe75Si15B5Ti5 alloy, (c) quaternary Fe75Si15B5Ta5 alloy and (d) quaternaryFe75Si15B5Zr5 alloy. The samples were wet milled for the indicated time up to 60 h. The XRD patterns have been vertically shifted.

Fig. 2. Detailed view of Bragg reflections visible after mechanically inducedcrystallisation for the Fe75Si20B5 and Fe75Si15B5Zr5 alloys milled for 60 h. Forcomparison, the XRD patterns of the corresponding amorphous alloys are givenbelow for each composition. For clarity, the XRD pattern have been shiftedvertically.

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patterns of the samples wet milled for 20 h, present a single broadbump around 52�. As can be noticed, this feature has a slight asym-metry (on the low angle side), revealing a small amount of crys-talline phase embedded in an amorphous matrix. For the Ticontaining alloy, the XRD pattern of 20 h milled sample is charac-teristic for a nanocrystalline material (Bragg reflections are broad-ened and with low intensities). As compared to the ternary alloy, inthe case of all quaternary alloys the complete amorphisationoccurs at longer milling time. This can be assigned, according tothe Fe–Zr, Fe–Ta and Fe–Ti binary phase diagrams, to the verysmall (practically null) solubility of Zr, Ta or Ti in Fe at room tem-perature. For all the investigated quaternary alloys the amorphisa-tion occurs after 40 h of wet mechanical alloying. In their XRDpatterns a single halo is noticed, revealing the absence of the longrange atomic order.

Increasing the milling time reveals the prevention of mechani-cally induced crystallisation by substituting Ti and Ta for Si. Con-trary to the previously mentioned alloys, the substitution of Zrfor Si alloy presents different behaviour. In this case the mechani-cally induced crystallisation is more easily noticed, the bump isreplaced by a convolution of several Bragg reflection. For compar-ison reasons, the XRD patterns of the Fe75Si20B5 and Fe75Si15B5Zr5alloys milled for 60 h alongside with XRD patterns of correspond-ing amorphous alloys are given in Fig. 2. It can be noticed thatmechanically induced crystallisation leads to the formation of a

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certain amounts of crystalline Fe23B6 and a-Fe(Si) phase for bothFe75Si20B5 and Fe75Si15B5Zr5 alloys. Similar results were also previ-ously reported for mechanically induced crystallisation of FINEMETlike alloys [20].

The particle size distribution for the ternary and quaternaryalloys wet milled up to 20 h and respectively 40 h are presentedin Fig. 3. All particles size distribution reveal a Gaussian like profile,which is given by the achievement of a balance between cold weld-ing and fracturing phenomena induced by milling. This Gaussiandistribution is favourable for the future compaction processes ofthe powders. The median particle diameter, D50, of the amorphousreference alloy Fe75Si20B5 is 11.3 lm. One can notice that the sub-stitution of Si for Ti, Ta or Zr significantly influences the D50 param-eter. The D50 parameter variation can be linked to the mechanicalproperties of the used substitution element. For the element withthe largest value of the hardness (Ta) the lowest value of the D50

parameter is obtained, while for the element with the lowest hard-ness (Zr) the D50 parameter is larger. For the reference alloy, theparticle size distribution shows the presence of certain amountof sub-micrometric particles. The presence of such particles wasevidenced by SEM analysis, which shows the occurrence of veryfine particles on the surface of larger particles as can be observedin the inset presented in Fig. 3 for the ternary based alloy.

The particles morphology investigation for the ternary and qua-ternary amorphous alloys is presented in Fig. 4. The SEM investiga-tion showed that powder is formed by particles with polyhedralshape. Also, SEM investigation confirms the influence of the substi-tution element hardness on the particle size as previously evi-denced by particle size distribution analysis. The SEM imagesreveal that the large particles are formed by smaller particles thatare agglomerated and cold welded during milling. SEM investiga-tion showed a relatively smooth surface of the particles in areasnot covered by the smaller particles. The morphology of the parti-

Fig. 3. Particle size distribution for the ternary (Fe75Si20B5) and quaternary (Fe75Si15B5Ti5comparison reasons, the mean particle size (D50) and specific surface of the particles (S

cles presented in this study is dissimilar to particles obtained bychemical routes (generally sponge like shape) or atomisation tech-niques (often sphere like shape and smooth surface).

In order to determine the thermal stability of the reference alloyFe75Si20B5, and compare it to the newly developed quaternaryalloys, Fe75Si15B5M5 (M = Ti, Ta or Zr), DSC measurements wereperformed on the amorphous samples. The DSC heating curves ofthose alloys are presented in Fig. 5. During heating up to 900 �C,several events can be noticed. In the temperature range of 100–250 �C, a broad exothermic peak can be noticed, for all the samples,independent of their chemical composition. This peak is attributedto the internal stresses release phenomenon [19]. This event is typ-ically encountered during heating of the samples subjected to sev-ere plastic deformation which occurs during mechanical milling.The next event visible on the DSC curves appears at different tem-peratures depending on the chemical composition of the amor-phous phase. For the ternary Fe75Si20B5 alloy, around 485 �C, asharp and intense exothermic peak is found. By substituting Ti,Ta and Zr for Si this peak position is found at 540 �C, 590 �C and600 �C respectively. This exothermic peak was assigned to the crys-tallisation phenomenon of the amorphous alloys [21]. The shift ofthe crystallisation phenomenon to higher temperatures by substi-tuting Ti, Ta and Zr for 5 at.% of Si is a positive effect because of theincrease of the thermal stability of the amorphous phase. The val-ues of the crystallisation temperature, Tx, for the substitution alloysare comparable with the values of the similar silicon rich Fe basedamorphous alloys [22–24]. In the case of FINEMET like alloys, P.Kollar et al. evidenced that the crystallisation temperature is at520 �C, for the Fe73.5Si13.5B9Nb3Cu1 composition [22]. H. Okumuraet al. studied an alloy with similar composition. Fe73.5Si16.1B6.4-Nb2.9Cu1.1, and found a lower crystallisation temperature of theamorphous phase which is at 510 �C [23]. Also, Dong et al. reporteda crystallisation temperature of 545 �C for the ternary Fe79Si10B11

, Fe75Si15B5Ta5, Fe75Si15B5Zr5) alloys wet milled up to 20 h and respectively 40 h. For) are given.

Fig. 4. SEM images of ternary and quaternary powder alloys wet milled up to 20 h and respectively 40 h at a magnification of �1000.

Fig. 5. DSC heating curves of amorphous Fe75Si15B5M5 (M = Si, Ti, Ta, Zr) powdersobtained by wet milling. A heating rate of 20 �C/min and Ar atmosphere were used.

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[24]. Although in our case, the values of the crystallisation temper-ature appear to be higher, it is difficult to state which is the highestvalue due to different heating rates used by the previously men-tioned authors.

In Fig. 6 is given the relationship between the amorphous alloycrystallisation temperature, atomic radius, valence bond and elec-tronic configuration as it results from DSC investigations. The lar-ger thermal stability of the of the quaternary amorphous alloyscan be explained by increased degree of dense random-packedstructures and thus more difficult atomic rearrangement as hasbeen proposed elsewhere [25]. This increase can be attainedthrough geometric or electronic effects [2,26]. Both these effectsinfluence the crystallisation phenomenon of this type of alloys.The higher values of the crystallisation temperatures of the quater-

nary alloys compared to the ternary master alloy are the conse-quence of both these effects simultaneously. One empirical rulefor increasing the difficulty of atomic rearrangement is increasingthe number of constituents of the alloy. This will determine ahigher degree of random-packed structures. Also, in order toachieve higher thermal stability, the atomic radius ratio of the ele-ments in the mixture must comply with the Dr/rP 0.12 formula[26]. Both of these circumstances represent geometric effects.Long-range atomic ordering requires higher energies due to thefact that the samples contain atoms of different size, that can beordered in three classes: small atoms (Si, B), intermediate sizeatoms (Fe, Ti and Ta) and by case, large atoms (Zr – 1.602 Å)[27]. The smaller atoms occupy the interstices that are determinedby larger atoms, creating a larger degree of dense random-packedstructures. Due to this fact, atomic rearrangement requires higherenergy; therefore the thermal stability of the amorphous phase isenhanced. However, titanium and tantalum atoms haveintermediate size and they have approximately the same atomicradius (Ti – 1.462 Å, Ta – 1.467 Å) [27]. As a consequence thesamples containing either Ti or Ta should have roughly the samecrystallisation temperature. The high difference concerning thecrystallisation temperature of these two substitutions can beexplained through the electronic effects that are also governingthe crystallisation phenomenon. Indeed the Ta atom is expectedto carry one additional electron in the outer shells. Atomic rear-rangement is linked as well to the bond valence and the outer layerelectronic configuration of the elements. According to the follow-ing Refs. [2,26] higher electron concentration on the outer layer,determines a higher bond valence, and as a result higher crystalli-sation temperature. In this case, Ta has a 5d electronic configura-tion and 4.51 valence bond whereas Ti has a 3d electronicconfiguration with 3.2 valence bond [25]. The larger difference ofthe crystallisation temperature values is the result of these elec-tronic effects. The fact that, Zr has a 4d electronic configuration,and the sample containing zirconium has the highest crystallisa-tion temperature indicates a larger influence of the geometriceffects on the thermal stability than the electronic effects, for thistype of amorphous alloys.

Fig. 6. Relationship between the amorphous alloy crystallisation temperature, atomic radius, valence bond and electronic configuration for Fe75Si15B5M5 powders (M = Ti, Taand Zr).

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The DSC curve of Fe75Si15B5Zr5 alloy in the 400–650 �C temper-ature range is presented in Fig. 7. In this detailed range the glasstransition temperature, Tg, and crystallisation temperature, Tx,can be observed on the heating curve, while on the cooling curvethe Curie temperature, Tc, of the crystallized alloy is visible.

Due to the fact that for this alloy the Tx has the largest value, it ispossible to observe the glass transition temperature, and also, thesupercooled liquid region, DTx. For this alloy, Fe75Si15B5Zr5, the DTxis about 20 �C, and the glass transition temperature is 580 �C. In thecase of the other alloys this event cannot be observed due to alower crystallisation temperature which could overlap it. Furthercompaction of the Fe75Si15B5Zr5 alloy using adequate sinteringtechnologies, such as spark plasma sintering, can allow the obtain-ing of high density compacts in amorphous state. This can beachieved by choosing a sintering temperature comprised in theDTx interval. In this temperature range, the material acts like a vis-

Fig. 7. Detailed view of the DSC curve in the 400–650 �C temperature rangecorresponding to the Fe75Si15B5Zr5 alloy. Tg, Tx and TC refer to the glass transitiontemperature, the crystallisation temperature and Curie temperature respectively.

cous liquid thus enabling the preparation of high density compactsand avoiding simultaneously the alloy crystallisation.

On the DSC cooling curve a change of slope can be noticed, atabout 530 �C. This thermal event is attributed to the Curie temper-ature of the crystallised phase, a-Fe(Si) type in agreement withearlier study [28].

Fig. 8 presents the evolution of the saturation magnetisationversus milling time for the as milled samples. The saturation mag-netisation value of the starting sample is given by the iron powderand is presented for comparison in the same plot. The saturationmagnetisation is dependent upon the chemical composition. Ineach of the four studied compositions, the saturation magnetisa-tion has a different value, despite the fact that all of the four sam-ples contain the same atomic percent of iron. Due to differentatomic mass of the elements, the starting mixtures have a differentFe/alloy weight ratio. The ternary alloy has 87.2 wt.% iron, and thequaternary alloys have as follow: 85.4 wt.% in the case of Ti substi-tution, 75.2 wt.% for the sample containing Ta, and 81.8 wt.% forthe Zr substitution. Different Fe/alloy weight ratio in the startingsample, and the fact that the saturation magnetisation is expressed

Fig. 8. Evolution of the saturation magnetisation of the ternary and quaternaryalloys versus milling time. The measurements were performed at 300 K.

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in A�m2/kg are leading to the observed different saturation mag-netisation values.

The saturation magnetisation has a tendency to decrease versusmilling time for all the samples. It is expected that alloying withnon-magnetic elements will reduce the magnetisation upon dilu-tion. For all the samples investigated, one can notice that for shortmilling times, up to 5 h, the value of the saturation magnetisationexhibits a sharp decrease. This is due to the formation of the alloyoccurring in this milling time interval. During formation of thealloy the metalloids like B and Si are chemically bonding withthe ferromagnetic atoms Fe, effect that induces a reduction of theirmagnetic moments [29–31]. The saturation magnetisation of thealloy has, as expected, a lower value than the starting sample,because of the nonmagnetic elements which interfere with theiron’s structure. Upon further increase of the milling duration,the decrease rate of the saturation magnetisation is lower. This isdue to the fact that the alloy is formed and the decrease is theeffect of the structural defects and internal stresses induced byhigh energy mechanical milling. Also, it was proved that a certainamount of PCA, in this case benzene, can be adsorbed on the pow-der’s surface [32]. Wet mechanical milling leads to very fine parti-cles, which is equivalent to large specific surface area of thepowder and so the influence of the benzene adsorbed on the pow-der surface is more and more accentuated. One can take intoaccount the possible contamination of the alloys by C atomsresulted from the benzene decomposition which can also lead tothe decrease of the magnetisation via Fe–C bonds. The decomposi-tion of the surfactant during milling caused by the severe thermo-mechanical conditions was earlier reported in the case of Ni–Feand Fe–Zr based powder obtained by wet MA [32–33]. M. Pilaret al. obtained Fe–Zr based alloy by wet mechanical alloying andstudied the influence of different PCA such as: hexane, polyethy-lene glycol, cyclopentane, naphthalene and hexanone on the struc-tural and thermal stability of the alloy. They observed that thedecomposition of PCA leads to the dissolution of C atoms in theFe–B phase, without the formation of any crystalline carbide [33].

The Fe75Si20B5 alloy wet milled up to 20 h has a saturation mag-netization of 131 A�m2/kg corresponding to 1.50 lB/Fe atom. Theamorphous Fe75Si15B5Ti5 sample, obtained after 40 h of MA, has asaturation magnetisation of 133 A m2/kg (1.56 lB/Fe atom). Thequaternary alloy containing tantalum has the saturation magneti-sation at 98 A m2/kg (1.30 lB/Fe atom), for the sample wet milledup to 40 h. In the case of Zr for Si substitution the saturation mag-netisation of the amorphous alloy is 112 A m2/kg (1.37 lB/Featom). It can be noticed that the calculated magnetic momentper Fe atom depends on the chemical composition of the amor-phous alloy varying from 1.3 to 1.56 lB/Fe atom. The valuesreported in the literature for FINEMET compositions are around1.67 lB/Fe atom [23,34], which is slightly higher as compared withour present study. For the calculus of the magnetic moment of Featoms, the theoretical composition has been taken into account.The lower value obtained for the magnetic moment is assumedto be given by the following: PCA adsorbed on the particles surfaceand the contamination of the alloys with nonmagnetic atomsresulted from the PCA decomposition. Both phenomena lead tolower weight percent of Fe in amorphous alloys. Also, as isexpected the magnetic moment of iron atom values for all compo-sition of amorphous alloys are lower as compared to the iron mag-netic moment from Fe3Si (2.4 lB for A site and 1.2 lB for B site) andFe2B (1.9 lB) crystalline phases [35]. The increase of the millingdurations leads to the decrease of the value of the saturation mag-netization, as can be seen in Fig. 8. Usually, for prolonged millingtime a contamination with Fe of the processed powder is noticedwhich can lead to an increase of the saturation magnetisation ofthe samples. In this case, the expected increase of the samples

magnetisation is counterbalanced by the dominant decrease ofthe samples saturation magnetisation induced by milling.

The amorphous sample containing Ti has a slightly higher satu-ration magnetisation as compared to the Fe–Si–B base alloy. It wasproved that a certain amount of minor alloying elements can leadto a more dense packed structure and to a higher degree of ran-domness [25]. The slightly higher saturation magnetisation of Ticontaining alloy can be explained by a more dense random packedstructure, which can lead to a higher magnetic interaction betweenferromagnetic atoms.

The thermomagnetic measurements revealed the fact that, thesubstitutions lead to the formation, during alloy crystallisation,of multiple magnetic phases as compared to the ternary referencealloy. The presented thermomagnetic plots are not quantitativemeasurements. They were recorded with the purpose of evidenc-ing the changes in the magnetic behaviour of the samples versustemperature during heating and cooling (Curie temperature ofamorphous alloy or of different phases resulted from the alloycrystallisation). In order to identify the phases that form duringheating up to 900 �C, X-ray diffraction were performed on the heattreated samples. The M = f(T) plots and XRD patterns of the crystal-ized samples are presented in Fig. 9. On the heating curve of theamorphous Fe75Si20B5 alloy (Fig. 9a), the ferromagnetic–paramagnetic transition of this alloy can be noticed as a sudden drop ofmagnetisation at 320 �C. At about 550 �C, an increase of the samplemagnetisation can be observed and is attributed to the samplecrystallisation. According to the literature, the crystallisation ofthe FINEMET like alloys take place in this temperature range, lead-ing to the formation of magnetic a-Fe(Si) solid solution from theamorphous phase [22,23]. During cooling, the curve has twoslopes, which reveal that after full crystallisation of the sample,two magnetic phases are formed: Fe2B and Fe3Si having Curie tem-perature at 730 �C and 626 �C respectively. The Fe2B phase has atetragonal crystal structure and the Fe3Si phase can have a facecentred cubic crystal structure or if it is in ordered state a DO3

superstructure. The presence of the twin Bragg reflexions at about31.6� and 36.7� (Fig. 9b) confirms that the crystallisation symmetryof the obtained Fe3Si in our case is DO3. The assignment of theobserved Bragg reflections to a phase or another is challengingdue to the multiphase character of the material formed and thenon-stoichiometry of the phases obtained after crystallisation.Similar results were also reported in different studies [19,23].

For the sample containing Ti (Fig. 9c), the Curie temperature ofthe amorphous phase has a slightly higher value, 345 �C (8%increase as compared with the ternary alloy). This most probablycan be connected to the larger magnetic moment carried by theFe atoms in this alloy. In addition the change of interatomic dis-tances may also play a role in the strength of the exchange interac-tion of the different alloys as reported by the Néel–Slater curves[36]. However a precise knowledge at the atomic scale of the struc-ture of the amorphous alloys hampers to go further in the analysis.On the cooling curve, there are also only 2 slope changes, one atabout 580 �C which represents the Curie temperature of the Fe3Siphase. The value of the Curie temperature for the obtained Fe3Siis higher as compared to the DO3 super structure obtained in thecase of base alloy crystallisation. According to Fe–Si phase diagram,the Fe3Si compound exists in the range of 10–28 at.% of silicon con-tent [37]. In this range, the Curie temperature varies roughly in the520–710 �C range, having higher values in the iron rich region.However, different Curie temperatures can be noticed at heatingand cooling of the analysed samples which can be explained bythe following aspects: (i) the thermal hysteresis. Highheating/cooling rate will shift (increase/decrease) the temperatureat which the ferromagnetic–paramagnetic transition will beobserved on heating curve and vice versa; (ii) the reordering

Fig. 9. Thermomagnetic measurements, recorded in the temperature range of 30–900 �C, and XRD patterns for the crystallized samples of the (a), (b) ternary Fe75Si20B5 alloyand for the quaternary alloys: (c), (d) Fe75Si15B5Ti5, (e), (f) Fe75Si15B5Ta5 and (g), (h) Fe75Si15B5Zr5. The value of the applied magnetic field during thermomagneticmeasurements was 0.1 T.

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B.V. Neamt�u et al. / Advanced Powder Technology 27 (2016) 461–470 469

phenomena. After primary crystallisation, the powder is furtherheated up to 900 �C. During heating, the reordering phenomenacan occur leading to changes in the vicinity of Fe atoms, thus differ-ent Curie temperature are noticed on heating and cooling curves ofthe thermomagnetic measurements. It is worth mentioning that inthe case of the Ti containing alloy, the Fe3Si compound does notpresent superstructure lines, as proved by XRD investigationshown in Fig. 9d. This indicates the absence of full atomic orderingin such case. The Curie temperature for the Fe2B0.7C0.3 is about425 �C [38]. This phase formation is favoured due to the contami-nation by carbon resulted from the PCA decomposition. Thedecomposition of the PCA during milling has been highlighted inseveral publications underlining the risk of contamination [32,33].

The sample containing Ta has a Curie temperature lower withabout 27% as compared to the Fe–Si–B amorphous phase. Thiswas found to be around 232 �C on the heating curve. The Ta forSi substitution leads to the multiple phase crystallisation as com-pared to the ternary Fe–Si–B alloy. On the cooling curve of thisalloy, there are 3 slope changes, as can be seen in Fig. 9e. At tem-peratures of 680 �C and 548 �C two changes of slope are observed.They correspond to the Curie temperature of an a – Fe(Si) solidsolution (with possible other atoms dissolved in it) and respec-tively Fe3Si phase. In this case, also the Fe2B0.7C0.3 can be foundsimilarly like the sample containing titanium at the same temper-ature. The above mentioned phases were identified in the XRD pat-tern of the sample heated up to 900 �C and are indicated in Fig. 9f.

In the case of the Zr for Si substitution, Fig. 9g, two Curie tem-peratures can be noticed on the heating curve. The first one has avalue of 300 �C while the other one has a higher value and is com-parable with the Curie temperature of the Fe3Si compound(556 �C). The presence of two Curie temperature can be explainedby several hypothesis: (i) the co-existence of two different amor-phous phases in the material; (ii) the presence of a small amountof magnetic crystalline phase. This phase can originate frommechanical induced crystallisation or thermal induced crystallisa-tion during the thermomagnetic investigation itself. On the coolingcurve there are two Curie temperatures, similarly to the referencealloy, one for the Fe2B boride at 653 �C, and the other one for theFe3Si phase at 556 �C. The Curie temperature of the Fe3Si phase isalso visible on the DSC curve presented in Fig. 7. The different val-ues of the Curie temperatures derived from DSC and thermomag-netic measurements can be explained by the different heatingrate and different investigation technique. The Bragg reflectionsof Fe2B boride, Fe3Si phases alongside with ZrB, a-Fe(Si), Fe3Band Fe3B0.7C0.3 can be identified in the XRD pattern of the sampleheated up to 900 �C as presented in Fig. 9h.

It can be noticed that similar phases that crystallise from theamorphous alloys have different Curie temperature. This can begiven by the different chemical composition of each alloy leadingto different vicinity of ferromagnetic atoms. Also, it can beobserved that the small amount of metals that have been intro-duced changed the crystallisation behaviour of the material. Thecrystallisation becomes more complicated and multiple secondaryphases can be encountered.

It is worth to be noticed that mechanically induced crystallisa-tion favours the formation of borides as main phase as compared tothe thermal induced crystallisation that favour the formation ofFe–Si main phases.

4. Conclusions

Quaternary Fe75Si15B5M5 (M = Ti, Ta or Zr) alloys were obtainedby wet mechanical milling using benzene as PCA. Substitution of5 at.% of any of the three transition metals for Si leads to anincrease of the milling time required to attain the amorphous state,

from 20 h up to 40 h. It was found that all samples present Gaus-sian particle size distribution and the D50 parameter is influencedby the mechanical properties of the used substitution element.The thermal domain of the amorphous alloys increases by the sub-stitutions: 12% increase for M = Ti, 21% for M = Ta and 24% in thecase of substitution with Zr. Due to high crystallisation tempera-ture, for the quaternary alloy containing zirconium, the glass tran-sition temperature is visible on the DSC curve. In the case of the Tifor Si substitution the value of saturation magnetisation is largerthan the value of the base ternary alloy, due to a more dense ran-dom packed structure which determines enhance interactionsbetween magnetic moments. Also, the alloy containing Ti has ahigher Curie temperature of the amorphous phase than the refer-ence Fe75Si20B5 alloy. Thermomagnetic measurements revealedthe formations of multiple phases in the cases of the quaternaryalloys as compared to the ternary base alloy. The carbon contami-nation deriving from the PCA’s decomposition is highlighted by theXRD investigations performed on the powder heated up to 900 �C,where several carbides are identified. This can however probablybe much reduced by heating under dynamic pumping or underhydrogen atmosphere.

Acknowledgement

This work was supported by a Grant of the Romanian NationalAuthority for Scientific Research CNCS – UEFISCDI, Project NumberPN II-RU-TE-2012-3-0367.

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