Intrinsic surface hardening and precipitation kinetics of Al 0.3 CrFe 1.5 MnNi 0.5 multi-component alloy Ming-Hao Chuang a , Ming-Hung Tsai a , Che-Wei Tsai a , Nai-Hao Yang a , Shou-Yi Chang b , Jien-Wei Yeh a , Swe-Kai Chen c , Su-Jien Lin a,⇑ a Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu 30013, Taiwan b Department of Materials Science and Engineering, National Chung Hsing University, Taichung 40227, Taiwan c Center for Nanotechnology, Materials Science, and Microsystems, National Tsing Hua University, Hsinchu 30013, Taiwan article info Article history: Received 16 August 2012 Received in revised form 26 September 2012 Accepted 27 September 2012 Available online 5 October 2012 Keywords: High-entropy alloy Multi-component Precipitation Surface hardening Wear abstract An Al 0.3 CrFe 1.5 MnNi 0.5 multi-component alloy with a very effective surface hardening ability attributed to intrinsic q phase precipitation and applicable to complex tool components was developed. Under a con- ventional aging treatment in a normal atmosphere at 550 °C for 2 h, the alloy with the surface precipita- tion hardening layer of 74 lm thick exhibited markedly enhanced surface hardness from HV 338 to HV 840 and efficiently improved wear resistance to 1.4 times the values of SUJ2 and SKD61 steels, while high fracture toughness close to that of ductile SKD61 steel was effectively retained. Precipitation thermody- namics and growth kinetics of the surface hardening layer were also investigated. The growth of the surface hardening layer was much faster than that of the precipitation in the bulk matrix; it did not follow typical long-distance diffusion kinetics but behaves more similar to a self-induced or reaction- accelerated short-range decomposition with a thickness increase proportional to the cube of aging time. On the surface, a lower heterogeneous nucleation energy and a reduced strain energy (total 55 kJ/mol) than the regular nucleation energy in the bulk matrix (78 kJ/mol) dominated the rapid formation and growth of the intrinsic surface precipitation with significant strain relaxations. Ó 2012 Elsevier B.V. All rights reserved. 1. Introduction Tribological damages of machinery tools and components during practical use at high speeds and high temperatures cause materials severe wear losses, consequently leading to most of the failures of machines and of the significant increase in maintenance expense [1]. The developments of wear-resistant materials and the researches on surface modifications have therefore continually been emphasized [2–4]. Particularly, surface hardening that will efficiently improve the wear, corrosion and fatigue resistance of tool components without a toughness loss has been intensively studied, including from (1) a mechanical aspect by work hardening (such as shot peening and roller burnishing) [3–5], (2) a thermal aspect by surface quenching to form a martensitic layer (such as flame hardening and laser surface melting) [5], (3) a diffusional aspect by introducing a surface carbide or nitride layer (such as carburizing or nitriding) [4,6], and (4) an additional aspect by protective hard coatings [7]. However, most of the surface hardening techniques are complicated, expensive, time consuming, and even inapplicable to complex-shape tool components [2–4]. In another aspect, conventional age hardening owing to second-phase precipitations from oversaturated solid solutions under atmospheric heat treat- ments provides a simple way for strengthening bulk materials [8]. For surface hardening, nevertheless, to confine the precipitations only within surface regions, ex. extrinsically by introducing high concentrations of defects with intensive irradiation damages [9,10] or cold work [5,11–12], needs additional procedures, cost and time. In 2004, a new concept of alloy design with multi-principal elements/components, i.e. the development of high entropy alloys (HEAs), was proposed by Yeh [13]. Previous studies indicated that, because of the incorporations of differently-sized atoms and conse- quent severe lattice distortions and sluggish atom diffusion, HEAs showed the advantages of high hardness, good wear and oxidation resistance and also high elevated-temperature strength and soften- ing resistance [14–17]. Especially, attributed to high mixing entro- pies and markedly reduced free energy, HEAs would stabilize in a form of simple solid solutions with significantly increased solubility and show a high capability of age hardening [18,19]. In particular, 600–800 °C aged Al x CrFe 1.5 MnNi 0.5 alloys presented a markedly in- creased hardness (from HV 300 to HV 800) and improved wear resistance owing to a q-phase precipitation ((Cr, Fe, and Mn)-rich precipitates with the solid solutions of Al and Ni) [20,21]. Unfortu- nately, along with the hardening, the fracture toughness of the 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.09.133 ⇑ Corresponding author. Tel.: +886 3 5719543; fax: +886 3 5722366. E-mail address: [email protected](S.-J. Lin). Journal of Alloys and Compounds 551 (2013) 12–18 Contents lists available at SciVerse ScienceDirect Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Transcript
Journal of Alloys and Compounds 551 (2013) 12–18
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds
journal homepage: www.elsevier .com/locate / ja lcom
Intrinsic surface hardening and precipitation kinetics of Al0.3CrFe1.5MnNi0.5
multi-component alloy
Ming-Hao Chuang a, Ming-Hung Tsai a, Che-Wei Tsai a, Nai-Hao Yang a, Shou-Yi Chang b, Jien-Wei Yeh a,Swe-Kai Chen c, Su-Jien Lin a,⇑a Department of Materials Science and Engineering, National Tsing Hua University, Hsinchu 30013, Taiwanb Department of Materials Science and Engineering, National Chung Hsing University, Taichung 40227, Taiwanc Center for Nanotechnology, Materials Science, and Microsystems, National Tsing Hua University, Hsinchu 30013, Taiwan
a r t i c l e i n f o
Article history:Received 16 August 2012Received in revised form 26 September 2012Accepted 27 September 2012Available online 5 October 2012
An Al0.3CrFe1.5MnNi0.5 multi-component alloy with a very effective surface hardening ability attributed tointrinsic q phase precipitation and applicable to complex tool components was developed. Under a con-ventional aging treatment in a normal atmosphere at 550 �C for 2 h, the alloy with the surface precipita-tion hardening layer of 74 lm thick exhibited markedly enhanced surface hardness from HV 338 to HV840 and efficiently improved wear resistance to 1.4 times the values of SUJ2 and SKD61 steels, while highfracture toughness close to that of ductile SKD61 steel was effectively retained. Precipitation thermody-namics and growth kinetics of the surface hardening layer were also investigated. The growth of thesurface hardening layer was much faster than that of the precipitation in the bulk matrix; it did notfollow typical long-distance diffusion kinetics but behaves more similar to a self-induced or reaction-accelerated short-range decomposition with a thickness increase proportional to the cube of aging time.On the surface, a lower heterogeneous nucleation energy and a reduced strain energy (total 55 kJ/mol)than the regular nucleation energy in the bulk matrix (78 kJ/mol) dominated the rapid formation andgrowth of the intrinsic surface precipitation with significant strain relaxations.
� 2012 Elsevier B.V. All rights reserved.
1. Introduction
Tribological damages of machinery tools and components duringpractical use at high speeds and high temperatures cause materialssevere wear losses, consequently leading to most of the failures ofmachines and of the significant increase in maintenance expense[1]. The developments of wear-resistant materials and theresearches on surface modifications have therefore continuallybeen emphasized [2–4]. Particularly, surface hardening that willefficiently improve the wear, corrosion and fatigue resistance of toolcomponents without a toughness loss has been intensively studied,including from (1) a mechanical aspect by work hardening (such asshot peening and roller burnishing) [3–5], (2) a thermal aspect bysurface quenching to form a martensitic layer (such as flamehardening and laser surface melting) [5], (3) a diffusional aspectby introducing a surface carbide or nitride layer (such as carburizingor nitriding) [4,6], and (4) an additional aspect by protective hardcoatings [7]. However, most of the surface hardening techniquesare complicated, expensive, time consuming, and even inapplicableto complex-shape tool components [2–4]. In another aspect,
ll rights reserved.
+886 3 5722366.
conventional age hardening owing to second-phase precipitationsfrom oversaturated solid solutions under atmospheric heat treat-ments provides a simple way for strengthening bulk materials [8].For surface hardening, nevertheless, to confine the precipitationsonly within surface regions, ex. extrinsically by introducing highconcentrations of defects with intensive irradiation damages[9,10] or cold work [5,11–12], needs additional procedures, costand time.
In 2004, a new concept of alloy design with multi-principalelements/components, i.e. the development of high entropy alloys(HEAs), was proposed by Yeh [13]. Previous studies indicated that,because of the incorporations of differently-sized atoms and conse-quent severe lattice distortions and sluggish atom diffusion, HEAsshowed the advantages of high hardness, good wear and oxidationresistance and also high elevated-temperature strength and soften-ing resistance [14–17]. Especially, attributed to high mixing entro-pies and markedly reduced free energy, HEAs would stabilize in aform of simple solid solutions with significantly increased solubilityand show a high capability of age hardening [18,19]. In particular,600–800 �C aged AlxCrFe1.5MnNi0.5 alloys presented a markedly in-creased hardness (from HV 300 to HV 800) and improved wearresistance owing to a q-phase precipitation ((Cr, Fe, and Mn)-richprecipitates with the solid solutions of Al and Ni) [20,21]. Unfortu-nately, along with the hardening, the fracture toughness of the
M.-H. Chuang et al. / Journal of Alloys and Compounds 551 (2013) 12–18 13
fully-precipitating alloys was significantly reduced from 216 to only1–2 J/cm2, limiting the practical applications of the alloys. Thus inthe present study, a new route to intrinsically confine the precipita-tion within the surface of Al0.3CrFe1.5MnNi0.5 multi-componentalloy has been developed by using simple atmospheric aging treat-ments at a lower temperature range and applying the volume misfitbetween the q-phase precipitates and the matrix. The microstruc-ture and mechanical performances of the intrinsic surface precipita-tion hardening (denoted as SPH) layer were characterized, and theformation thermodynamics and growth kinetics of the SPH layerwere studied.
2. Experimental
Comb-shape ingots of Al0.3CrFe1.5MnNi0.5 multi-component alloy were pre-pared by vacuum induction-melting the constituent elements of designed compo-sition and casting in an argon atmosphere, followed by a homogenizationtreatment at 1100 �C for 4 h. The as-homogenized ingots were then cut into platespecimens of 10 � 10 � 5 mm in size for structural and mechanical characteriza-tions, cylindrical specimens of 8 mm in diameter and 3 mm in height for wear tests,and bar specimens of 28 � 5 � 5 mm for bending tests. The alloy specimens weresubsequently aging-treated, at 500 �C for 4–24 h or at 550–700 �C for 0.5–4 h, forprecipitations. Cross-sectional microstructures of the specimens were observedby an optical microscope and a scanning electron microscope (SEM, JEOL JSM-5410, Japan) with an X-ray energy dispersive spectrometer (EDS, Oxford, England).The crystal structures were analyzed by an X-ray diffractometer (XRD, RigakuME510-FM2, Japan). The macro- and microhardness of the plate specimens afterdifferent heat treatments were measured using a Vickers hardness tester and amicrohardness tester (Mitutoyo HV-115 and HM-115, Kawasaki, Japan) under loadsof 5 kg and 25 g, respectively. The wear resistance of the cylindrical specimens wasexamined by a pin-on-belt abrasive test against a 400 mesh-Al2O3 cloth under aload of 1 kg and at a sliding speed of 0.5 m/s, and presented as the total sliding dis-tance over the volume loss of the specimens (m/mm3). Three-point bending tests ofthe bar specimens were conducted, with a span distance of 21 mm, by a universaltesting machine (Instron 4468, USA) for determining the strength and fracturetoughness of the specimens. For comparison, wear and bending tests of SUJ2 (AISI52100) bearing steel and SKD61 (AISI H11) hot-work tool steel as well as of thepresent aging-treated (at 550 �C for 2 h) alloy specimen with an SPH layer removed(i.e. bulk matrix only, by polishing the specimen surface) were performed.
3. Results and discussion
3.1. Formation and structural characterizations of SPH layer
Effective surface precipitation hardening of Al0.3CrFe1.5MnNi0.5
multi-component alloy was applicable to complex tool compo-nents, as seen in the photographs of comb-shape alloy specimensafter aging treatment at 550 �C for 2 h in Fig. 1. In the optical
Fig. 1. Comb-shape Al0.3CrFe1.5MnNi0.5 alloy specimens after aging treatment at 550 �C fomicrograph around a specimen notch, showing the formation of a uniform and continmicrographs of the alloy specimens after aging treatment at 550 �C for 1–4 h, respective
micrograph around a specimen notch in Fig. 1(c), a uniform andcontinuous SPH layer just along the surface contour of the speci-men was observed to form. Further from the micrographs of the al-loy specimens after aging treatment at 550 �C for 1–4 h shown inFig. 1(d)–(g), respectively, the growth of the intrinsically formedSPH layer was found to proceed with aging time but showed a non-linear acceleration behavior, possibly in an exponentially increasedrate (to be discussed below in the section of thermodynamics andkinetics).
Fig. 2(a) shows the SEM image of the Al0.3CrFe1.5MnNi0.5 alloyspecimen after aging treatment at 550 �C for 2 h. Though a distinctinterface between the formed SPH layer and the bulk matrix wasobserved, however as determined by the EDS mapping analysesof elemental distribution presented in Fig. 2(b), it was very inter-esting that all the constituent elements uniformly distributed ineither the bulk matrix or the formed SPH layer; basically all theconstituents formed uniform solid solutions attributed to theirhigh mixing entropies [20], and no obvious macroscopic segrega-tion or difference in composition was found. The SPH layer wasformed intrinsically, not a product of oxygen penetration and/orsurface oxidation since no increase in the content of oxygen wasdetected. In the image shown in Fig. 2(c), dendritic (circled ‘‘D’’)and interdendritic (circled ‘‘I’’) regions were clearly identified inboth the bulk matrix and the formed SPH layer. The EDS analysesof elemental contents listed in Table 1 indicated that the dendriteswere very slightly rich in Al and Cr, and interdendrites rich in Mnand Ni. Even at the microscale either in the dendritic or interden-dritic region, the composition of the formed SPH layer was veryclose to that of the bulk matrix, within a small deviation less than10 at.%, suggesting that the intrinsic formation of the SPH layer wasvery possibly a consequence of short-range decomposition or sep-aration rather than long-distance diffusion or segregation of spe-cific elements.
Fig. 3 shows the XRD patterns of Al0.3CrFe1.5MnNi0.5 alloy spec-imens after different heat treatments. According to literature[20,21], two sets of diffraction peaks that appeared in the patternof as-homogenized specimen corresponded to body-centered cubic(bcc) and face-centered cubic (fcc) structures; the (Al, Cr)-rich den-drites consisted of the bcc phase because Al and Cr were bcc stabi-lizers, while the (Mn, Ni)-rich interdendrites the fcc because Mnand Ni were fcc stabilizers [20,21]. After aging treatment at550 �C for 2 h, many other diffraction peaks additionally appearedat 2h = 40–50� in the pattern of the specimen with a formed SPHlayer, whereas the specimen with the SPH layer removed (i.e. bulk
r 2 h: (a) and (b) photographs of whole and mounted-polished specimens, (c) opticaluous SPH layer along the surface contour of the specimen; (d)–(g) cross-sectionally, showing the growth of the formed SPH layer.
Fig. 2. Microstructures and elemental analyses of Al0.3CrFe1.5MnNi0.5 alloy specimens after aging treatment at 550 �C for 2 h: (a) SEM image of bulk matrix and formed SPHlayer, (b) EDS mapping analyses of elemental distribution in (a); (c) SEM image showing dendritic (circled ‘‘D’’) and interdendritic (circled ‘‘I’’) regions in the bulk matrix andthe formed SPH layer for EDS analyses of elemental contents.
Table 1Elemental contents of dendritic and interdendritic regions in the bulk matrix and the formed SPH layer of Al0.3CrFe1.5MnNi0.5 alloy specimens.
Fig. 3. XRD patterns of Al0.3CrFe1.5MnNi0.5 alloy specimens after different heattreatments (As-homo: as-homogenized, 550 �C for 2 h (Matrix): after agingtreatment at 550 �C for 2 h (bulk matrix only, the formed SPH layer removed), w.SPH: with the formed SPH layer retained; h: fcc phase, s: bcc phase, arrow-pointed: q phase diffraction peaks).
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matrix only) showed a pattern the same as the as-homogenizedone. A comparison of the patterns indicated the appearance ofthe additional peaks being attributable to the formation of theSPH layer, which was confirmed by the pattern of the fully precip-itating specimen after aging treatment at 700 �C for 2 h. The pre-cipitation was reported to be a q phase (JCPDS No. 36-1373) thattransformed from the dendritic bcc phase and was composed ofCr5Fe6Mn8 (tetragonal structure, lattice constants a = 9.09 Å andc = 9.99 Å) with the solid solutions of Al and Ni [20,21]; meanwhile,the interdendritic fcc structure remained unchanged. It was ex-pected that, in the bcc matrix as some Al and Ni of strong binding
were separated from the other elements to form B2-type precipi-tates (bcc structure), the remaining Cr, Fe and Mn spontaneouslyformed the q phase precipitates (similar to up-hill spinodaldecomposition). Because of short-range separation rather thanlong-distance diffusion, the composition of the SPH layer deter-mined above by EDS analyses did not vary much from that of thebulk matrix.
3.2. Mechanical performance of SPH layer
Table 2 lists the macrohardness of Al0.3CrFe1.5MnNi0.5 alloy spec-imens after different heat treatments as well as the microhardnessof the dendritic and interdendritic regions in the bulk matrix andthe formed SPH layer of the alloy specimens. It was very clear that,under a conventional aging treatment in a normal atmosphere at550 �C for 2 h, the alloy with the SPH layer of about 74 lm thickshowed significantly enhanced surface hardness. Because the qphase precipitation was expected to occur within the bcc dendrites,the hardness of the dendritic region dramatically increased from HV386 to HV 1090, while the hardness of the interdendritic fcc regionremained at about HV 250, even similarly for the fully precipitatingspecimen after aging treatment at 700 �C for 2 h. Although the over-all macrohardness of the SPH layer was not available due to aninsufficient thickness for accurate measurement, it was believedto be at the high level of HV 840 (according to the similar hardnessof the SPH layer and the fully precipitating specimen in both thedendritic and interdendritic regions), much superior to the valueof the unhardened bulk matrix, HV 338. In comparison with othersurface treatment routes and hardening techniques, such as shotpeening, flame hardening and carburizing of typical alloys [4,6],the intrinsic SPH of the present Al0.3CrFe1.5MnNi0.5 alloy alterna-tively provides a more efficient and effective as well as cheap andsimple way for uniform surface hardening (up to HRC 70) of anycomplex tool components including the comb-shape one with deeptrenches (Fig. 1, teeth length of 5.4 mm and spacing of 1.5 mm).
Table 2Overall macrohardness of Al0.3CrFe1.5MnNi0.5 alloy specimens after different heat treatments and the microhardness of dendritic and interdendritic regions in the bulk matrix andthe formed SPH layer of the alloy specimens.
a The overall macrohardness of the SPH layer is not available due to insufficient thickness.
M.-H. Chuang et al. / Journal of Alloys and Compounds 551 (2013) 12–18 15
Fig. 4 further plots the wear resistance, another indication of tri-bological application, of SUJ2 and SKD61 steels and the presentAl0.3CrFe1.5MnNi0.5 alloy specimens after different heat treatments.It showed that the unhardened bulk matrix (the formed SPH layerremoved) of the alloy specimen after aging treatment at 550 �C for2 h, even with a low hardness of HV 338 only, presented high wearresistance close to that of the robust SKD61 (HV 520) and evenSUJ2 (HV 723) steels. With the formed SPH layer retained, the wearresistance was more efficiently improved to 15 m/mm3, equivalentto the high value of the fully precipitating specimen and about 1.4times the values of the SUJ2 and SKD61 steels, suggesting the highpotential of the alloy for practical use.
Even with high hardness and wear resistance, on the other handpresumable brittleness, the Al0.3CrFe1.5MnNi0.5 alloy specimenwith an SPH layer effectively retained high toughness close to thatof ductile SKD61 steel, as seen in the stress–strain curves and frac-tographs of bending tests in Fig. 5. The stresses, r, and strains, e,measured by the three-point bending tests were estimated fromthe following equations that [22]
r ¼ 3PL
2bt2 ð1Þ
e ¼ 6td
L2 ð2Þ
where P is the load of bending test, L the span distance (21 mm), band t the width and depth (both 5 mm) of the specimens, respec-tively, and d the deflection of the specimens. From the bendingstress–strain curves and fractographs of the SUJ2 and SKD61 steelsand the alloy specimens after different heat treatments, it was real-ized that the alloy with a simple aging treatment at 550 �C for 2 hshowed a considerable ductility (bending strain �0.31), as high asthat of the SKD61 steel and much better than those of the SUJ2 steel
Fig. 4. Wear resistance of SUJ2 and SKD61 steels and Al0.3CrFe1.5MnNi0.5 alloyspecimens after different heat treatments (As-homo: as-homogenized, 550 �C for2 h (Matrix): after aging treatment at 550 �C for 2 h (bulk matrix only, the formedSPH layer removed), w. SPH: with the formed SPH layer retained).
(bending strain �0.083) and the fully hardened specimen (bendingstrain �0.016 mm). The bending strengths of the SUJ2 and SKD61steels as well as of the alloy specimens in the as-homogenized, sur-face hardened and fully hardened states were then determined as3700, 2460, 1770, 1850 and 550 MPa, respectively. The abovemechanical analyses revealed that the effectively strengthenedAl0.3CrFe1.5MnNi0.5 alloy with substantial increases in surface hard-ness and wear resistance presented a negligible loss in fracturetoughness relative to the original value of the as-homogenizedstate.
3.3. Formation thermodynamics and growth kinetics of SPH layer
Based on the good mechanical performances of presently devel-oped Al0.3CrFe1.5MnNi0.5 alloy, it is of great interest to clarify theformation thermodynamics and growth kinetics of the intrinsicSPH layer for properly controlling heat treatments, especially asthe layer presents a nonlinear acceleration growth much rapiderthan the precipitation in a bulk matrix. Thus, the thicknesses ofthe SPH layers formed on the alloy specimens after aging treat-ments at 500, 550 and 600 �C for different durations were plotted(vs. aging time) in Fig. 6(a), both in logarithmic scales. Undoubt-edly, with increasing aging temperature and time, a thicker SPHlayer formed, and the specimen was then fully hardened. However,it was interestingly found that the formation rate of the SPH layervaried during the aging process; for the 550 �C-aged specimens asexamples, the thickness slowly increased to 31 lm in 1 h but rap-idly from 74 to 203 lm with time from 2 to 3 h. A correlation be-tween the thickness, x, of the formed SPH layer and the aging time,t, is simply given to describe the growth as that [23]
x ¼ ðktÞn i:e: log x ¼ n log kþ n log t ð3Þ
where k is defined herein as a rate constant, and n rate exponent, allof which at the beginning and late stages of the SPH layer formationat 500–600 �C are determined as listed in Table 3, by linearly fittingthe data. Very obviously, the growth of the SPH layer did not followtypical diffusion-controlled kinetics, x ¼ ðktÞ1=2 (square root declinein rate, R ¼ dx=dt ¼ k0t�1=2). Instead, the initial formation was moresimilar to a first-order, uniform interface-controlled reaction, x ¼ kt(constant in rate, R ¼ dx=dt ¼ k), while the subsequent growthchanged to a non-uniform self-induced or reaction-acceleratedcombustion-like behavior, x ¼ ðktÞ3, with thickness proportional tocube of time (quadratic increase in rate, R ¼ dx=dt ¼ k00t2). This find-ing confirms the above prediction that the SPH layer formation ispossibly a consequence of short-range decomposition or separationrather than long-distance diffusion or segregation. By further plot-ting the logarithmic rate constants, logk, vs. 1/temperature, 1/T, inFig. 6(b) and using the following equation [23], the logarithmicconstants logk0 and activation energy, Q, were determined as alsolisted in Table 3.
k ¼ k0 exp � QRT
� �i:e: log k ¼ log k0 �
Q2:3RT
ð4Þ
Fig. 5. Stress–strain curves and fractographs of three-point bending tests of SUJ2 and SKD61 steels and Al0.3CrFe1.5MnNi0.5 alloy specimens after different heat treatments(As-homo: as-homogenized, 550 �C for 2 h: after aging treatment at 550 �C for 2 h with the formed SPH layer retained).
Fig. 6. (a) Thicknesses (vs. aging time) of SPH layers formed at the surface of Al0.3CrFe1.5MnNi0.5 alloy specimens after aging treatments at different temperatures for differentdurations and (b) rate constants (vs. 1/temperature) of the SPH layer formation at the surface at different temperatures.
Table 3Rate exponents (n), logarithmic rate constants (logk), logarithmic constants (logk0) and activation energy (Q) of SPH layer formation at the surface and q phase precipitation inthe bulk matrix of Al0.3CrFe1.5MnNi0.5 alloy specimens.
Precipitation Temperature (�C) Stage n logk logk0 Q (kJ/mol)
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The activation energy of the SPH layer formation at the beginningand late stages were then obtained as 55 and about 245 kJ/mol,respectively.
By comparison the rate of q phase precipitation in the bulk ma-trix of Al0.3CrFe1.5MnNi0.5 alloy was much slower than that of theSPH layer formation at the surface; the precipitation in the bulkmatrix would dominate the hardening of the alloy only at temper-atures higher than 650 �C. Because of the difficulty in measuringthe sizes of the precipitates in the bulk matrix, the hardening rateof the alloy reported in our previous study [20] was adopted for
estimating the precipitation rate. By assuming that the hardeningof the alloy is linearly proportional to the size of the precipitatesand that the ultimate size approaches the scale of a Vickers indentmark, around the order of 100 lm, a predicted relation of the pre-cipitate size, d, to the alloy hardening is given as that
d100 lm
� DHðH1 � H0Þ
i:e: d � DHðH1 � H0Þ
� 10�4 m ð5Þ
where DH is the increase in hardness, H1 and H0 the ultimate andoriginal values, respectively. Fig. 7(a) plots the predicted sizes of the
Fig. 7. (a) Predicted sizes (vs. aging time) of q phase precipitates formed in the bulk matrix of Al0.3CrFe1.5MnNi0.5 alloy specimens after aging treatments at differenttemperatures for different durations (at 800 �C, the size of the precipitates was assumed as very small as one micrometer at a very short aging time of one minute) and (b) rateconstants (vs. 1/temperature) of the q phase precipitations in the bulk matrix at different temperatures.
M.-H. Chuang et al. / Journal of Alloys and Compounds 551 (2013) 12–18 17
q phase precipitates in the bulk matrix of the alloy specimens afteraging treatments at 600–800 �C for different durations. Only onestage was presented; the late stage was omitted because of the sat-uration of hardening due to precipitate collision. Similarly by usingEqs. (3) and (4), and regarding the size of the precipitates, d, as thethickness, x, the rate constants, k, rate exponents, n, the logarithmicconstants logk0, and the activation energy, Q, of the precipitation inthe bulk matrix were determined, as listed in Table 3 and plotted inFig. 7(b). Based on Eqs. (3) and (4), it should be noted that thechange in assumed ultimate size of the precipitates, 100 lm, willnot alter the calculations of rate exponents and activation energy.Similar to the SPH layer formation at the beginning stage, the pre-cipitation in the bulk matrix followed an interface-controlledbehavior, x � kt (constant in rate, R ¼ dx=dt ¼ k) but had a higheractivation energy of 78 kJ/mol.
It is concluded that, with the same rate exponent, n, of 1, both theSPH layer formation at the beginning stage and the precipitation inthe bulk matrix present an interface-controlled mechanism, con-firming both the precipitations consequent on short-range decom-position or separation rather than long-distance diffusion orsegregation. In addition, from Eq. (4), it is realized that, with theclose constants k0 (a term of frequency and accommodation factors,independent of temperature [23]), 10�4.7 and 10�4.4, the activationenergy, Q, will dominate the rate of precipitations. At the surface,the smaller activation energy, 55 kJ/mol, than that in the bulk ma-trix, 78 kJ/mol, accordingly yields a high rate as observed. Twoimportant factors are believed to facilitate the decrease in activationenergy and the rapid formation of the intrinsic SPH layer, includinglower heterogeneous nucleation energy and strain energy at the
Fig. 8. (a) Predicted variation of critical activation energy (energy barrier, DG⁄, vs. nucleunucleation and strain relaxation, along with the schematic illustrations of incoherent vwithout consideration of strains, PBM/SPH: precipitations in bulk matrix/surface precitemperature) of SPH layer formation at the surface and precipitations in the bulk mtemperatures.
surface than that of regular nucleation in the bulk matrix. To theknowledge of phase transformation [24], the change in free energyduring the nucleation or formation of a new phase is given by
DG ¼ �DGvV þ cAþ DGsV � DGd ð6Þ
in which�DGv is the reduction in volumetric (chemical) free energyper unit volume, c the increase in interfacial energy per unit area,DGs the increase in strain energy per unit volume, and �DGd thereduction in free energy owing to nucleation on any defects suchas dislocations, boundaries, free surfaces, etc. [24,25]; V and A arethe volume and surface area of nucleus of the new phase, respec-tively. In total, critical activation energy, DG⁄, i.e. an energy barrierto be overcome, is needed for stabilized nucleation, as schematicallyillustrated in Fig. 8(a).
First, the lowered activation energy of heterogeneous nucleationat the free surface, from DG�hom to DG�het by�DGd (the subscripts homand het denote the homogeneous nucleation in the bulk matrix andthe heterogeneous nucleation at the surface, respectively) [24,25],will simply favor the preferred, rapid formation of the SPH layer.Second, the increase in strain energy, DGs, which is originated fromthe incoherent volume misfit between the precipitates and matrixwill enlarge the energy barrier, DG⁄, for forming a stable nucleus[24]. In the fully precipitating Al0.3CrFe1.5MnNi0.5 alloy, an increaseof 0.85% in density owing to the volume shrinkage of q phase precip-itates transforming from the bcc matrix was measured; as well, adifference of�1.74% in volume between the precipitates and matrixwas calculated according to the numbers of atom per unit cells andthe lattice constants measured from XRD analyses. Both the differ-ences in density and volume indicated a considerable volume misfit
s radius, r) for precipitations in bulk matrix and at surface owing to heterogeneousolume misfit originated strains (hom/het: homogeneous/heterogeneous nucleationpitation hardening with strain energy) and (b) comparison of rate constants (vs.atrix of Al0.3CrFe1.5MnNi0.5 alloy specimens during aging treatments at different
18 M.-H. Chuang et al. / Journal of Alloys and Compounds 551 (2013) 12–18
of 0.85–1.74% between the two phases, accordingly a high strain en-ergy and a marked increase in activation energy, from DG�hom toDG�PBM by DGPBM (Strain), for the q phase precipitation in the bulkbcc matrix. A smaller increase in activation energy, from DG�het toDG�SPH by DGSPH (Strain), for the SPH layer formation at the surfaceis expected, since the reductions in strain energy along with phasetransformations at surfaces have been attributed to strain relax-ations through surfaces [26–33]. Suezawa showed that the strainenergy for spinodal decomposition and martensitic transformationnear free surfaces was reduced by 20% and 10%, respectively [28]. Inconsequence, order–disorder transition [26–28], spinodal decom-position [28–30], and martensitic transformation [31–33] preferen-tially occurred at free surfaces. Conclusively, the SPH layerformation at the surface, which is believed to be a short-rangedecomposition or separation rather than long-distance diffusionor segregation, with smaller activation energy will have a larger rateconstant than the q phase precipitation in the bulk matrix does dur-ing aging treatments at the same temperature as compared inFig. 8(b). At higher temperatures, the activation energy relative tothe temperature term in Eq. (4) becomes trivial, and the rate con-stant is alternatively dominated by the constant k0 and will ap-proach a level of about 10�4.7–10�4.4 as determined above.
It is also interesting to clarify the reason that at the late stagethe growth rate of the SPH layer dramatically increases. Based onEqs. (3) and (4), expectably, smaller activation energy, Q, for theSPH layer formation at the beginning stage, 55 kJ/mol, than thatat the late stage, 245 kJ/mol, would plainly yield a larger rate con-stant, k, and a higher formation rate, R. However, the anticipation iscontrary to the above observation because, kinetically, the rateconstant, k, is determined not only by the activation energy, Q,but also by the constant k0. A much larger constant k0 at the latestage, 1010.3, than that at the beginning, 10�4.7, accordingly pro-vides a larger rate constant and leads to a rapider growth of theSPH layer. The constant k0 is a term of frequency factor and thenumber of accommodation site [23]. It means, because of the antic-ipated easy strain relaxation through the free surface, at the latestage as the SPH layer continues to grow, more suitable low-energysites for atom migration and subsequent phase transformation arevery possibly formed. From another perspective, the growth rate atthe late stage could be regarded as that R ¼ dx=dt ¼ k00t2 ¼ k000x2=3
which shows a function of time, t, or correspondingly a functionof the ratio of product transformed, i.e. the thickness of the SPHlayer, x. Both the larger constant k0 implied more accommodationsites and the introduction of product transformation ratio into thegrowth rate at the late stage suggest a self-induced or reaction-accelerated combustion-like growth behavior, i.e. an autocatalyticreaction, attributable to the easy stress relaxation.
4. Conclusion
In this study, effective surface hardening of Al0.3CrFe1.5MnNi0.5
multi-component alloy was developed. Under a conventionalatmospheric aging treatment, a uniform and continuous surfaceprecipitation hardening layer of q phase (Cr5Fe6Mn8) was intrinsi-cally transformed from a dendritic bcc matrix. Attributed to thehardening layer, the alloy presented a markedly improved surfacehardness of HV 840, bending strength of 1850 MPa, and wear resis-tance of 1.4 times the values of SUJ2 and SKD61 steels, along withretained high fracture toughness close to that of ductile SKD61steel. In comparison with conventional hardening techniques, theintrinsic surface precipitation provides a simple, efficient and
effective route for surface hardening of complex tool components,and the surface-hardened alloy shows high potential for tribologi-cal use. Studies of formation thermodynamics and growth kineticsindicated that the very rapid growth of the surface precipitationlayer did not follow typical long-distance diffusion kinetics butwas more similar to a self-induced or reaction-accelerated short-range decomposition. Important factors including low heteroge-neous nucleation energy and reduced strain energy at surfaceduring nucleation as well as easy strain relaxation through surfaceduring growth dominated the rapid formation and growth ofthe surface precipitation layer.
Acknowledgement
The authors gratefully acknowledge financial support for thisresearch by the National Science Council, Taiwan, under GrantNo. NSC 99-2221-E-007-069-MY3.
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