SF

Ra

b

c

a

ARRAA

KALPF

1

afaltoeaaruaiaG

t

0h

Applied Surface Science 264 (2013) 176– 183

Contents lists available at SciVerse ScienceDirect

Applied Surface Science

jou rn al h om epa g e: www.elsev ier .com/ locate /apsusc

tudy on feasibility of producing an amorphous surface layer ofe49Cr18Mo7B16C4Nb3 by pulsed Nd:YAG laser surface melting

eza Mojavera, Faezeh Mojtahedib, Hamid Reza Shahverdib,∗, Mohammad Javad Torkamanyc

Office of Applied Researches and Technology, Tarbiat Modares University, P.O. Box 14115-143, Tehran, IranDepartment of Materials Engineering, Tarbiat Modares University, P.O. Box 14115-143, Tehran, IranIranian National Center for Laser Science and Technology, P.O. Box 14665-576, Tehran, Iran

r t i c l e i n f o

rticle history:eceived 5 June 2012eceived in revised form 3 September 2012ccepted 30 September 2012vailable online 8 October 2012

eywords:morphousaser surface melting

a b s t r a c t

This work aims to investigate whether an amorphous surface layer can be obtained when as-castFe49Cr18Mo7B16C4Nb3 alloy is submitted to pulsed Nd:YAG laser surface melting. The experiments wereconducted in the various laser scanning speeds. The microstructures of laser treated zones were inves-tigated by X-ray diffraction XRD and Field Emission Scanning Electron Microscope (FESEM) and theirmicrohardness were measured, too. The chemical composition of different points of each sample wasanalyzed by energy-dispersive X-ray spectroscopy EDS. Although the estimated cooling rates in surfacelayers were higher than the required cooling rate to achieve full amorphization, but the present experi-ments were unable to retain complete glassy microstructure on surface and a mixture of amorphous (low

ulsed Nd:YAG lasere-based glass former

volume fraction) and ultrafine grained phases were produced in surface of samples. Based on the find-ings, it was understood that the overlapping of successive pulses and element redistributions occurredin pulsed laser melting could severely restrict amorphization. The influence of laser scan speed and laserpower on heat input, melting ratio, compositional changes and cracking in laser treated zone were dis-cussed separately. It is suggested that the limited range of laser variables in pulsed Nd:YAG laser meltingmay help to produce a sound amorphous phase of as-cast Fe49Cr18Mo7B16C4Nb3 alloy.

. Introduction

The amorphous alloys usually have the high hardness, high wearnd corrosion resistance which make them excellent in their sur-ace properties [1–4]. Among the metallic glass formers, Fe-basedlloys are attractive because of their low cost [4,5]. So, during twoast decades, many studies on the glass transition and crystalliza-ion process, structure, physical properties and corrosion behaviorf Fe-based metallic glass alloys have been conducted [6–11]. How-ver, compared to other amorphous alloys like Zr- and Pd-basedlloys, Fe-based amorphous alloys have a lower glass-formingbility (GFA). Consequently, the higher critical cooling rates areequired for their amorphization [6,8,10–12], ranging from val-es around 102 K/s in comparison with values about 1–10 K/s forlloys with very good glass-forming ability [8]. This limitation hasnspired some attempts to develop some new Fe bulk metallic glasslloys which have a larger supercooled liquid region and a higher

FA [10,11,13,14].

While the characteristics mentioned above provide a few limi-ations in producing Fe-based bulk metallic glass parts, for example

∗ Corresponding author. Tel.: +98 21 82893307; fax: +98 21 82883101.E-mail address: shahverdi@modares.ac.ir (H.R. Shahverdi).

169-4332/$ – see front matter © 2012 Elsevier B.V. All rights reserved.ttp://dx.doi.org/10.1016/j.apsusc.2012.09.167

© 2012 Elsevier B.V. All rights reserved.

obtained thickness could not be more than a few millimeters [8],but these alloys are expected to easily vitrify in forms of atom-ized powders and surface layers where production routes of themare associated with very high cooling rates. Amorphous layers ofa few Fe system alloys might be obtained by different techniquesincluding high velocity oxy-fuel, air plasma spraying [15] and theselayers in different states, fully or partly amorphous or nano-grainsized microstructures (resulted from devitrifying-heat treatment)have even been entered in market [16].

In laser surface treatment LST (laser melting\alloying\cladding),a high energy laser beam can be irradiated on the surface of a crys-talline bulk in a very short time by laser scanning at high speedswhich could afford the cooling rates exactly in the bounds of thequench rates necessary for vitrification and nano-crystalization[17–22]. On the other hand, from the literature it can be found thatnot all attempts to provide an amorphous layer on metallic bases bylaser surface treatment were successful [23–27]. The main reasonsfor the failure were explained as: (1) the redistribution of elementsin weld metal [24,25]; (2) devitrification resulted from the reheat-ing cycles induced by the successive laser tracks [20,21,23,26]; (3)

epitaxial solidification which favors the crystal growth [18]; (4)inter-diffusion or dilution of some phases of substrate in the layer[24,25]. As well, contributions of some parameters of LST includ-ing laser scan speed, laser power [24–27], overlapping [21,23] and

R. Mojaver et al. / Applied Surface Science 264 (2013) 176– 183 177

Table 1Chemical composition of as-cast alloy.

Elements Fe Cr Mo C B Nb

mi

acpfwttpcla

mbhiFsp

2

Fm(spiedlw

lMww46ippttt

TS

D

wt% 57 19 14 2 3.5 4.5

ulti-track scanning [25] on surface amorphization have beennvestigated.

Generally speaking, it could be inferred that in order to gain anmorphous layer by LST, the criterion of cooling rate is not suffi-ient and all involving parameters in complicated reactions of laserrocessing and metallic glass formation should be considered. Thisact might become more importance in pulse laser melting PLMhere more variables like pulse shaping, pulse energy, pulse dura-

ion and frequency inevitably complicate the operation [28,29]. Onhe other hand, in pulsed laser melting coupling the high peakower with overlapping pulses produce a good melting ratio, a highooling rate and a low heat input than those of the continuous waveaser melting that makes it theoretically a good tool for producingn amorphous layer [30].

Fe52−xCr18Mo7B16C4Nbx (at %) alloys are a group of Fe-basedetallic glass alloys which glass forming ability and corrosion

ehaviors of their amorphous bulk in different amounts of Nbave earlier been studied [31–34]. In present study, the feasibil-

ty of producing an amorphous layer on surface of the as-caste49Cr18Mo7B16C4Nb3 by pulsed laser melting in different scanpeeds has been investigated and it has been aimed at correlatingulsed laser parameters with surface microstructure.

. Experimental

Multicomponent alloy with nominal composition ofe49Cr18Mo7B16C4Nb3 (at%) was prepared. The high purityaterials Cr (>99.9 wt%), Mo (>99.9 wt%), C (>99.5 wt%), Nb

>99.9 wt%) and B (>99.9 wt%) as well as pre-alloyed Fe–C wereelected and melted in a vacuum induction furnace. The melt wasoured into a permanent mold. The chemical composition of the

ngot was investigated by inductively coupled plasma spectrom-try (ICP) and the result is given in Table 1. A few plates withimensions of 30 mm × 60 mm × 2 mm were cut from the ingot for

aser processing. Before laser treatment the surface of the platesere ground and cleaned with acetone to remove impurities.

A pulsed Nd:YAG laser Model IQL-10 with a maximum meanaser power of 400 W was used as laser source for the experiments.

ore detail about the laser processing setup can be found else-here [30]. The treatments were carried out using pulsed laserith the focused beam diameter of 0.9 mm, pulse frequency of

0 Hz and pulse duration of 2 ms at various scanning speeds of 4, 5,.67 and 8 mm/s, A1–A4. A summary of these parameters is shown

n Table 2. These conditions were chosen in the way that laserrocessing occurred in conduction mode, but with a little change in

arameters, it might change to keyhole mode. During the process,he surface was protected by pure Ar gas emerged coaxially withhe laser beam. Two types of samples were prepared: single trackreated plates which were used for study on microstructures by

able 2ummary of pulsed laser melting.

Samplename

Average power(W)

Frequency(Hz)

Pulse duration(ms)

Beam diameter(mm)

A1 170 40 2 0.9

A2 170 40 2 0.9

A3 170 40 2 0.9

A4 170 40 2 0.9

, 1.2 mm; pulse peak power, 2.125 KW.

Fig. 1. Back scattered SEM image of microstructure of the as-cast alloy.

Field Emission Scanning Electron Microscope (FESEM) and multitrack treated plates which were used for X-ray diffraction (XRD)studying, using Co K� radiation. In the later case, the treatmentwas conducted in the way that percentage of overlapping betweenneighboring tracks was less than 10%.

The elements distribution in parent alloy and remelted poolswas analyzed by energy-dispersive X-ray spectroscopy (EDS) afterAr-ion etching (Ar-ion accelerating voltage-3 kV, ion current 30 �A,diameter and raster of ion beam were 120 �m and 1 mm, respec-tively and etching time was 240 s). The sample surface layerwas removed about 20 nm in thickness during cleaning proce-dure which was necessary for appropriate light elements analyses,mainly B and C. EDS analyses were done at both accelerating vol-tages of the electron beam of 10 kV and 30 kV. The results showedthat distribution of B, C and Nb were much clear in 10 kV.

In addition, the Vickers microhardness profiles were extractedfrom cross-sections of the processed samples by applying a 200 gfload for 15 s. The reported dates are the average of nine measure-ments for each sample.

3. Results and discussion

3.1. Microstructure of base metal

Fig. 1 shows the microstructure of parent alloy. Investigation byXRD, EDS and X-ray map showed that five phases could be iden-tified in the parent alloy, labeled A–E. The phases called D and Ewhich are distributed in matrix of phase C are very small ones andthey are not visible at the image represented in Fig. 1. So, they are

not labeled on this figure. Authors could not find any referencesin literature to characterize phases certainly. However, accord-ing to compositions (represented at Table 3) and XRD of parentalloy (Fig. 2) and by considering the average atomic weight and

Laser scanspeed (mm/s)

Overlappingfactor (%)

Accumulativeoverlapping index F

Effective peak powerdensity (KW/mm2)

4 89 6.5 12.3635 86 5.3 10.086.67 81 4.1 7.7958 78 3.5 6.657

178 R. Mojaver et al. / Applied Surface Science 264 (2013) 176– 183

Fig. 2. XRD patterns of the as-cast alloy and the top surface of fo

Table 3Chemical composition of phases in as-cast alloy.

Elements wt%

Fe Cr Mo C B Nb

A 75 19 4 2 0 0B 12 16.6 38.1 0 26 7.3C 50.86 32.5 11.14 5.5 0 0

cAbAmmee

3

saeopiasAXIbwotp

pt

3c

a

and/or remelted by nearly six later pulses. Also, numbers labeled indifferent portions in Fig. 3(b) presents the possible times of remelt-ing or partial melting of each zone due to the high number of

D 43.2 45.7 3.8 0 7.3 0E 4 5.9 0 12.1 0 78

omparing them to results reported in [13,26,35]: phase labeled as could be known as �-Fe. The quadrangular phase B seems theoride compound which is enriched in Mo (M23B6 or Mo2FeB2).lso, there is another small phase distributed in matrix, D, thatight be a boride compound (may be M7B3). The Needle-likeatrix phase C might be other morphology of �-Fe which is highly

nriched in Cr and Mo. The last phase, E, is the tiny phase highlynriched in Nb which could be identified as the carbide (NbxCy).

.2. XRD analyses of surface layers

The results of XRD analyses of top surface of four laser treatedamples and as-cast one are represented in Fig. 2. It can be seen thatlthough by increasing the scan speeds from 4 mm/s to 8 mm/s, thexperiments failed to retain complete amorphous microstructuren surfaces of samples. In XRD patterns of A1 and A2, the crystallineeaks of borides and carbides are still remained with a drop in their

ntensities. Also, three characteristic peaks of �-Fe at 2�: 44◦, 67◦

nd 86◦ become very broad that indicates on fine features of re-olidified microstructures. With increasing scan speed, in samples3 and A4, there are no longer peaks of borides and carbides in theirRD patterns and it seems that these phases completely dissolved.

n addition, two peaks of �-Feat 2�: 67◦ and 86◦ turn out to beroad humps while the peak around 2�: 44◦ becomes much wideith a plateau in its head. Comparing these characteristics with

ther results [23,26], it can be claimed that the microstructure ofop layer of A4 (and probably A3) contains a mixture of amorphoushase (low volume fraction) and ultrafine crystalline phases.

In order to understand the reason of failure to achieve full amor-hous phase, it is necessary to consider some aspects of PLM andransportation phenomena in melt zones.

.3. Pulse laser melting parameters, cooling rates and

haracterization of laser treated zones

In overlap pulsed laser melting because the energy pumped ton area not from a single pulse but also from overlapping pulses on

ur samples after laser treatment in different scan speeds.

the same spot, different parts of weld experienced various thermalcycles [28]. It is known that altering scan speed changes percentageof overlapping and also heat input in pulsed laser process [28,29].

Fig. 3(a) illustrates a schematic of successive pulses (Qf ≈ 90%)and sidetracks overlapping (≈10%) in top view and a schematic ofpulse overlapping in direction of parallel to weld track is shownin Fig. 3(b). Boundaries shown in figures indicate on fusion linesof successive pulses; it can be seen that overlapping of pulsescan divide the melt pool into the zones which each of them maymelt or partially melt a few times during irradiation. The frequentnumbers in Fig. 3(b) depicts that melt pool of nth pulse did nothave enough time to completely solidification and it was reheated

Fig. 3. (a) Schematic of highly overlapped successive pulses and side-overlappedtracks and (b) schematic of overlap pulse laser in section parallel to track direction.The digit in each portion indicates on the frequency of remelting of that portion.

rface Science 264 (2013) 176– 183 179

ord

Q

wfoct

E

w

F

rolaahuf

mmsrcaa

wa

cifthpi[a

cmiostefwrs

am

Table 4Estimated cooling rates in top surface of samples based on Eq. (3) as a function ofTsubstrate .

Assumed substratetemperature (K)

Cooling rates (K/s)

A1 A2 A3 A4

473.15 2085.26 2606.57 3477.16 4170.51573.15 1938.97 2423.71 3233.23 3877.94

R. Mojaver et al. / Applied Su

verlapping. These numbers show that top surface of melt pooleceived the higher energy and solidification of these parts waselayed for a longer time.

The percentage of overlapping can be calculated as (Eq. (1)):

f =(

1 − (v/f )(D + vt)

)× 100 (1)

here v is the process travel speed (mm/s), t the pulse duration (s), the pulse frequency (1/s) and D refers to the typical spot diameterf laser beam on the work piece [30]. Also, another parameter Falled cumulative overlapping index has been introduced (Eq. (2))o assess peak power density much precisely in PLM [28]:

ffective peak power density EPPD = peak power × (F)area of spot on work piece

here

= 1 + n{1 − ((n + 1)v|2fD)}n =⌈

Df

v

⌉(2)

The amounts calculated for Qf, F and EPPD for five samples areeported in Table 2. It can be seen that by increasing scan speed,verlapping factor and effective peak power density decreased. Theower EPPD (heat input) results to the higher cooling rate. So, aver-ge cooling rate of melt zone increased from A1 to A4. However,mount of overlapping factor of all samples might be consideredigh. As shown in Fig. 3(b), each considered region of melt zonenderwent several melting and solidifying cycles, e.g. the top sur-ace of A1 almost eight times remelted and resolidified.

The primary condition for formation glassy structure is achieve-ent to a critical cooling rate. The physics of simple laser surfaceelting is highly complex [22,26,27] and considering the circum-

tances in present experiments, exact estimation of local coolingates becomes much more complex. Nevertheless, in many arti-les [18,26,36] in laser surface melting, the Rosenthal solution for

moving point heat source proposed by Steen (Eq. (3)) used as anpproximately prediction of cooling rates:

dT

dT= −2�k

( vp

)�T2 (3)

here k is the thermal conductivity (≈22.5 W/m K [26] for thislloy), v is the scan speed and P is the pulse peak power.

In order to use (Eq. (3)), some assumptions should be made foralculation the parameter �T. Earlier, some researchers found thatn PLM, the maximum temperature usually are achieved in the sur-ace of melt pool and also, there is a very steep gradient betweenop and bottom of melt pool. Also, it might be expected that the pre-eating of successive pulses leads to establish an almost constanteak temperature in surface and another constant temperature in

nterface of melt zone and substrate during pulsed laser melting37]. Here, �T can be considered as a difference between Tsurface

nd Tsubstrate.Due to conducting the present experiments in upper limit of

onduction mode (refer to Section 2), it was assumed that theaximum temperature in top surface might be around the boil-

ng temperature of steel alloys, nearly 3273.15 K (Tsurface). On thether hand, there is not any certain criterion for assuming the sub-trate temperature. However, one can consider a few acceptableemperatures as Tsubstrate which in consequence they provide a vari-ty of possible cooling rates occurred in PLM of this study. Hereour temperatures including 473.15, 573.15, 673.15 and 773.15 Kere considered as the possible Tsubstrate. The estimated cooling

ates from (Eq. (3)) for all samples based on the different assumed

ubstrate temperatures are given in Table 4.

It can be seen that the achieved cooling rates in top surface ofll samples with different assumed substrate temperatures wereuch bigger than the critical cooling rate needed for amorphization

673.15 1798.00 2247.50 2998.17 3596.00773.15 1662.35 2077.94 2771.97 3324.71

of this alloy (between 600 and 103 K/s [2,6,12–15]). On the otherhand, XRD analyses show incomplete amorphization.

The microstructures of laser treated zones of samples A1 and A4are represented in Figs. 4 and 5, respectively. It can be observed thatin all samples some boride particles partially dissolved and they stillleft in melt zone, especially in interface with substrate, which indi-cates that the heat input was not enough to melt/dissolve the courseparticles. Comparing microstructure of A4 to one of A1 shows thatmore primary borides left in middle and top of the melt pool of A4that it is in agreement with decreasing the heat input by increas-ing scan speed. These remained borides can restrict amorphizationin two ways. Firstly, they can act as nucleation sites and secondly,they deplete the liquid from elements B and Mo.

In addition, while based on the gradual decreasing of cooling ratefrom surface to bottom, it is expected that microstructure becamesteadily finer in depth direction but one can distinguish some dis-tricts with coarser feature in the middle of melt pool or beneaththe surface. This phenomenon can be explained by overlapped suc-cessive pulses. In high overlapping factor, a large portion of theearlier melt pool generated by the former pulse will be remeltedby the next pulse [28–30]. Now it is possible that some tiny cellscannot be remelted, especially particles with high thermal andthermodynamic stability like borides and carbides, and during nextresolification, they can be the heterogeneous nucleation sites andenhance crystallization and growth in the middle or top zones ofmelt pool. Also, the more size growth of phases in these districtscould happen because of reheating after solidification.

The frequency of remelting (and reheating) of different portionsin a constant Qf increases from interface to surface, as demonstratedin Fig. 3(b). That is in contrast to the trend of cooling rate and maydiscontinue amorphization.

Fig. 6 shows the chemical analysis of matrix (single phase) indifferent points of three samples of A1, A3 and A4.

The results reveal that there was a great composition changein all samples. Many variables should be considered effective oncomposition change in PLM. Debory et al. have had the comprehen-sive study on vaporization rate and concentration change duringconduction mode laser welding and derived equations to calcu-late these rates. According to mass balance, it was expressed thatthe composition change from evaporation is proportional to theratio of vaporization rate and melting rate given by �vA where� is the density of weld and A refers to spot area (vaporizationrate G1 = �C1 A�v) [38–40]. Since vaporization rate is a function oflaser power, it can be concluded that scanning speed has impact oncomposition change through melting rate directly and indirectlyvia laser power by changing F cumulative overlapping index. It isfound from Fig. 6 that amount of composition variation from nom-inal composition of liquid greatly increased with decreasing scanspeed, from A4 to A1. The concentration of elements Mo, C andCr increased in vicinity of melt pool/substrate interface while bymoving toward surface their weight percentage became close to

their amount in nominal composition of liquid. The trend for Fewas completely reverse; its percentage showed a drop in points 1and 2 (marked on figures) but it reached to its amount in nominal

180 R. Mojaver et al. / Applied Surface Science 264 (2013) 176– 183

Fig. 4. (a–c) Back scattered FESEM images of different positions of laser treated zoneof sample A1 (scan speed = 4 mm/s).

Fig. 5. (a–c) Back scattered FESEM images of different positions of laser treated zoneof sample A4 (scan speed = 8 mm/s)

R. Mojaver et al. / Applied Surface S

Fig. 6. The composition changes in different positions of laser treated zone of sam-ples (a) A1 (scan speed 4 mm/s), (b) A3 (6.67 mm/s) and (c) A4 (8 mm/s). The locationof points 1–3 are shown in Fig. 5c and Fig. 6c.

cience 264 (2013) 176– 183 181

composition of liquid in point 3. On the other hand, percentage ofNb was constant at all points and equal to its amount in nominalcomposition of liquid. For B, it can be seen that in all samples itslightly rose in bottom of pool but it completely went out in zonesnear the surface of weld (point 3). Enrichment of microstructuresfrom Mo, C, Cr and B in interface can be resulted from partiallydilution of phases in substrate and diffusion of these elements tolower positions of remelted zone. Increasing heat input can boostthe total of dilution and diffusion. This fact agrees with great com-position change in A1 and more chemical homogeneity in A4 byvirtue of decreasing scan speed.

Besides, the high overlapping factor leads to the fact that liq-uid is stabilized in longer period of time. This extended liquid stageand multi-time resolidification may lead to much elemental distri-bution and segregation.

The deficiency of B in the regions near to the surface in A1 and A3can be explained by its low boiling temperature and high diffusioncoefficient [41]. Also, it is notable that B completely left the lasertreated zone in sample A4. In addition to the higher melting ratein A4, the reason might be the lower depth of remelted zone in A4compared to other samples. The diagram in Fig. 7 represents thevariation of depth of laser treated zone versus scanning speed. Itcan be found that depth of remelted zone significantly decreasedwith increasing scan speed. It equals with the shorter path for B todiffuse out of melt.

This chemical variation from nominal composition of liquid inlaser treated zone can greatly reduce glass forming ability of alloyand suppresses amorphization [5–7].

Considering the above discussion, it can be inferred that increas-ing scan speed in PLM can favor vitrification via increasing coolingrate and chemical homogeneity and decreasing overlapping factor.On the other hand, it results in lower heat input that as a result,process cannot melt primary borides and carbides (a barrier foramorphization) and also, cannot produce an adequate thickness oflayer on surface. Moreover, the higher speed leads to higher melt-ing rate that in consequence enhances the vaporization rate of lightelements.

In order to solve the mentioned above problems, it seems thatincreasing energy of pulse and scanning speed simultaneously mayoptimize the process. But it should be noticed that amplifying pulseenergy might change the process from conduction mode to keyholemode. Formation of keyhole greatly increases the evaporation and

as a result glass forming ability of liquid is lessened.

Another point that has to be considered is related to sound-ness of the processed zone. Although the process was not able to

Fig. 7. Variation of depth of laser treated region versus laser scanning speed.

182 R. Mojaver et al. / Applied Surface S

Fs

pamtpolhlsiecmmgimtoifP

iFtolts

4

lTims

-

[

[

[

[

[

[

[[

[

[

[

[

[

ig. 8. Average microhardness of laser treated samples in different laser scanningpeed.

roduce a full amorphous layer, but based on the images of FESEMnd broadening of XRD peaks, it can be said that an ultrafine grainedicrostructures have obtained on surfaces of samples. The forma-

ion of this ultrafine grained microstructure may improve someroperties of surface like hardness. The measured microhardnessf four samples is shown in Fig. 8 and as expected, the hardness ofaser treated microstructure increased with scan speed. The highardness of surface layers can render superior surface properties

ike high wear resistance which may make them interesting forome applications. On the other hand, laser surface melting resultedn cracking of the resolidified regions in all four samples. The gen-ration of cracks is a frequent result of laser surface melting. Theracking might be caused by the thermal stresses of pulsed laserelting combined with the high hardness and brittleness of theicrostructure which is the mixture of amorphous and ultrafine

rained phases [35]. Torkamany and co-workers [42] studied thenfluence of different parameters of pulsed Nd:YAG laser on sheet

etal bending. If the bending angle is considered equivalent withhermal strains induced in a sheet when is completely fixed, basedn their results, it can be found that the amount of thermal strainsncrease with increasing both scan speed and laser power. There-ore, other limit to choose adequate scan speed and laser power inLM should be considered to avoid cracking in amorphous layer.

The present results shows that the range of laser variablesn pulsed laser melting to produce a sound amorphous phase ofe49Cr18Mo7B16C4Nb3 is narrow, i.e. it seems not easy to optimizehe parameters of process. The reasons can be related to formationf a complicated heat and mass transportation in melt pool by over-apped laser pulses and also the nature of glass forming ability ofhe alloy. As well, some features of as-cast microstructure like largeize of borides can be known responsible, too.

. Conclusion

An attempt has been made to produce a glassy metallicayer of Fe49Cr18Mo7B16C4Nb3 by pulsed laser surface processing.he result was not successful. From investigation the possiblenvolved variables and by considering the glass forming ability and

icrostructural features of the as-cast alloy, the following conclu-ions may be drawn:

In pulsed laser process, overlapping factor can greatly affect thestructure and composition of laser treated zone. A high over-lapping factor even in presence of high cooling rate can restrict

[

[

cience 264 (2013) 176– 183

amorphization. So, to gain an amorphous phase by pulsed lasermelting, this factor should be keep as low as possible.

- Increasing laser scanning speed increases cooling rate and com-position homogeneity in melt pool which help amorphization.On the other hand, it decreases heat input that can lead to sur-vive some primary borides from completely melting and also, areduction in thickness of obtained layer.

- Increasing laser power to supply enough heat input and meltingratio is necessary up to a limit that it does not lead to form keyholeand enhance evaporation of light elements.

- It should be considered that increasing laser scanning speed andlaser power develops cracking in laser treated zone. Therefore,another boundary should be chosen for these parameters basedon this fact.

- Considering the glass forming ability of Fe49Cr18Mo7B16C4Nb3,it can be suggested that the range of laser variables in pulsedNd:YAG laser melting to produce a sound amorphous phase isnarrow, i.e. it seems not easy to optimize the parameters of pro-cess.

References

[1] W.H. Wang, C. Dong, C.H. Shek, Bulk metallic glasses, Mater. Sci. Eng. Rep. 44(2004) 45–89.

[2] A. Inoue, Stabilization of metallic supercooled liquid and bulk amorphousalloys, Acta Mater. 48 (2000) 279–306.

[3] A. Inoue, B. Shen, A. Takeuchi, Fabrication, properties and applications of bulkglassy alloys in late transition metal based systems, Mater. Sci. Eng. A 441 (2006)18–25.

[4] B. Prakash, Abrasive wear behavior of Fe, Co and Ni based metallic glasses, Wear258 (2005) 217–224.

[5] S.J. Pang, T. Zhang, K. Asami, A. Inoue, Synthesis of Fe–Cr–Mo–C–B–P bulkmetallic glasses with high corrosion resistance, Acta Mater. 50 (2002) 489–497.

[6] L. Ma, A. Inoue, On glass-forming ability of Fe-based amorphous alloy, Mater.Lett. 38 (1999) 58–61.

[7] S.J. Pang, T. Zhang, K. Asami, A. Inoue, Bulk glassy Fe–Cr–Mo–C–B alloys withhigh corrosion resistance, Corros. Sci. 44 (2002) 1847–1856.

[8] M. Stocia, J. Eckert, S. Roth, L. Schultz, Preparation of bulk amorphousFe–Cr–Mo–Ga–P–C–B alloys by copper mold casting, Mater. Sci. Eng. A 375–377(2004) 399–402.

[9] K.A. Lee, Y.C. Kim, J.H. Kim, C.S. Lee, J. Namkung, M.C. Kim, Mechanical propertiesof Fe–Ni–Cr–Si–B bulk glassy alloy, Mater. Sci. Eng. A 449–451 (2007) 181–184.

10] Q. Chen, D. Zhang, J. Shen, H. Fan, J. Sun, Effect of yttrium on the glass-formingability of Fe–Cr–Mo–C–B bulk amorphous alloys, J. Alloys Compd. 427 (2007)190–193.

11] K.Q. Qiu, J. Pang, Y.L. Ren, H.B. Zhang, C.L. Ma, T. Zhang, Fe-based bulk metallicglasses with a larger supercooled liquid region and high ductility, Mater. Sci.Eng. A 489 (2008) 464–467.

12] K. Hildal, N. Sekido, J.H. Perepezko, Critical cooling rate forFe48Cr15Mo14Y2C15B6 bulk metallic glass formation, Intermetallics 14(2006) 898–902.

13] D.J. Branagan, M. Carig Marshall, B. Meacham, Glass stability, glass forming, andmicrostructural refinement, United States Patent 7553382 B2 (2009).

14] D.J. Branagan, Method of modifying iron based glasses to increase crystalliza-tion temperature without changing melting temperature, United States Patent7186306 B2 (2007).

15] D.J. Branagan, Method of forming a hardened surface on a substrate, UnitedStates Patent 7785428 B2 (2010).

16] http://www.nanosteelco.com/, The Nanosteel Company, Inc., US, 2012.17] D.G. Morris, Glass-forming condition during laser surface melting, Mater. Sci.

Eng. 97 (1988) 177–180.18] F. Audebert, R. Colaco, R. Vilar, H. Sirkin, Production of glassy metallic layers by

laser surface melting, Scripta Mater. 48 (2003) 281–286.19] D.T.A. Matthews, V. Ocelik, J.Th.M. de Hosson, Tribological and mechanical

properties of high power laser surface-treated metallic glasses, Mater. Sci. Eng.A 471 (2007) 155–164.

20] G. Yang, X. Lin, F. Liu, Q. Hu, L. Ma, J. Li, W. Huang, Laser solid forming Zr-basedbulk metallic glass, Intermetallics 22 (2012) 110–115.

21] H. Yoshioka, K. Asami, A. Kawashima, K. Hashimoto, Laser-processed corrosion-resistant amorphous Fe–Ni–Cr–P–B surface alloys on a mild steel, Corros. Sci.27 (1987) 981–995.

22] C. Cui, J. Hu, Y. Liu, K. Gao, Z. Guo, Formation of nano-crystalline and amorphousphases on the surface of stainless steel by Nd:YAG pulsed laser irradiation, Appl.Sur. Sci. 254 (2008) 6779–6782.

23] Y. Yang, Y. Song, W. Wu, M. Wang, Multi-pass overlapping laser glazing ofFeCrPC and CoNiSiB alloys, Thin Solid Films 323 (1998) 199–202.

24] I. Manna, J. Dutta Majumdar, B. Ramesh Chandra, S. Nayak, N.B. Dahotre, Lasersurface cladding of Fe–B–C, Fe–B–Si and Fe–BC–Si–Al–C on plain carbon steel,Surf. Coat. Technol. 201 (2006) 434–440.

rface S

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

[

R. Mojaver et al. / Applied Su

25] A. Basu, A.N. Samant, S.P. Harimkar, J. Dutta Majumdar, I. Manna, N.B. Dahotre,Laser surface coating of Fe–Cr–Mo–Y–B–C bulk metallic glass composition onAISI 4140 steel, Surf. Coat. Technol. 202 (2008) 2623–2631.

26] V.K. Balla, A. Bandyopadyay, Laser processing of Fe-based bulk amorphous alloy,Surf. Coat. Technol. 205 (2010) 2661–2667.

27] B. Chen, Y. Li, Y. Cai, R. Li, S. Pang, T. Zhang, Surface vitrification of alloys bylaser surface treatment, J. Alloys Compd. 511 (2012) 215–220.

28] F. Malek Ghaini, M.J. Hamedi, M.J. Torkamany, J. Sabbaghzadeh, Weld metalmicrostructural characteristics in pulsed Nd:YAG laser welding, Scripta Mater.56 (2007) 955–958.

29] J. Sabbaghzadeh, M.J. Hamedi, F. Malek Ghaini, M.J. Torkamany, Effect of processparameters on the melting ratio in overlap pulsed laser welding, Metall. Mater.Trans. B 39B (2008) 340–347.

30] M.J. Hamedi, M.J. Torkamany, J. Sabbaghzadeh, Effect of pulsed laser parameterson in-situ TiC synthesis in laser surface treatment, Opt. Lasers Eng. 49 (2011)557–563.

31] S. Ahmadi, H.R. Shahverdi, S.S. Saremi, Effects of Nb alloying on Nano-crystallization kinetics of Fe55−xCr18Mo7B16C4Nbx (x = 0–3) bulk amorphousalloys, J. Mater. Sci. Technol. 27 (8) (2011) 735–740.

32] S. Ahmadi, H.R. Shahverdi, S.S. Saremi, Nanocrystallization of alpha-Fe crystalsin Fe52Cr18Mo7B16C4Nb3 bulk amorphous alloy, J. Mater. Sci. Technol. 27 (6)

(2011) 497–502.

33] H. Sun, Q. Man, Y. Dong, B. Shen, H. Kimura, A. Makino, A. Inoue, Effect of Nbaddition on the glass-forming ability, mechanical and soft-magnetic proper-ties in (Co0.942Fe0.058)72−xNbxB22.4Si5.6 bulk glassy alloys, J. Alloys Compd. 504S(2010) S31–S33.

[

cience 264 (2013) 176– 183 183

34] H. Zohdi, H.R. Shahverdi, S.M.M. Hadavi, Effect of Nb addition on corrosionbehavior of Fe-based metallic glasses in Ringer’s solution for biomedical appli-cation, Electrochem. Commun. 13 (2011) 840–843.

35] M.S. Bakare, K.T. Voisey, K. Chokethawai, D.G. McCatney, Corrosion behav-ior of crystalline and amorphous forms of the glass forming alloyFe43Cr16Mo16C15B10, J. Alloys Compd. 527 (2012) 210–218.

36] H. Alemohammad, S. Esmaeili, E. Toyserkani, Deposition of Co–Ti alloy on mildsteel substrate using laser cladding, Mater. Sci. Eng. A 456 (2007) 156–161.

37] J. Sabbaghzade, M. Azizi, M.J. Torkamany, Numerical and experimental investi-gation of seam welding with a pulsed laser, Opt. Lasers Tech. 40 (2008) 289–296.

38] P.A.A. Khan, T. Debroy, Alloying element vaporization and weld pool tempera-ture during laser welding of AISI202 stainless steel, Metall. Mater. Trans. B 15(1984) 641–644.

39] K. Mundra, T. Debroy, Calculation of weld metal composition change in high-power conduction mode carbon dioxide laser-welded stainless steels, Metall.Mater. Trans. B 24 (1993) 145–155.

40] H. Zhao, T. Debroy, Weld metal composition change during conduction modelaser welding of aluminum alloy 5182, Metall. Mater. Trans. B 32 (2001)163–172.

41] M.J. Torkamany, P. Parvin, M. Jandaghi, J. Sabbaghzadeh, in: Na Xiaodong (Ed.),Estimation of Composition Change in Pulsed Nd:Yag Laser Welding, InTech,

2010, pp. 194–218.

42] M. Hoseinpour Gollo, H. Moslemi Naeini, G.H. Liaghat, M.J. Torkamany,S. Jelvani, V. Panahizade, An experimental study of sheet metal bendingby pulsed Nd:YAG laser with DOE method, Int. J. Mat. Form. 1 (2008)137–140.