Effects of annealing treatment on the microstructure and mechanical properties of the AlN–SiC–TiB 2ceramic composites prepared by SHS–HIPLijuan Zhou a ,⇑,Hongbo Li b ,Yongting Zheng ba School of Materials Science and Engineering,Shandong University of Technology,ZiBo 255049,PR China bCenter for Composite Materials and Structures,Harbin Institute of Technology,Harbin 150001,PR Chinaa r t i c l e i n f o Article history:Received 14July 2012Received in revised form 10November 2012Accepted 12November 2012Available online 20November 2012Keywords:AlN–SiC–TiB 2ceramic Annealing treatment MicrostructureMechanical propertiesa b s t r a c tAnnealing treatments were carried out at different temperatures (600–1200°C)for various holding times (2–10h)to evaluate the microstructural and mechanical properties changes of the AlN–SiC–TiB 2ceramic composites prepared by SHS–HIP.The experimental results show that the annealing treatment is bene-ficial to improve the mechanical properties.In the meantime,the desolution of SiC-rich solid solution and precipitation of fine TiB 2grains occurred.Due to the precipitation strengthening by the SiC-rich solid solution and the grain boundary strengthening by fine TiB 2particles,the improvement of the mechanical properties was obvious with higher annealing temperatures and longer holding times.The highest HRA,bending strength and fracture toughness were 94.1,517.5MPa and 5.79MPa m 1/2,respectively,after the annealing treatment at 1200°C for 10h.Ó2012Elsevier B.V.All rights reserved.1.IntroductionAlN–SiC solid solution has a wide range of engineering appli-cations due to their excellent mechanical properties at high tem-peratures and has been investigated considerably in the past decades years [1,2].Many preparation methods,such as hot pressing [3],reaction synthesis [4],spark plasma synthesis sinter-ing [5]and combustion synthesis [6],were adopted to fabricate AlN–SiC ceramics.TiB 2is another important material for high temperature applications because of its high melting point,mechanical properties and electrical conductivity,and relatively low coefficient of thermal expansion.So AlN–SiC–TiB 2ceramics have wide potential applications in the field of electrical conduc-tivity at high temperatures.It is known that combustion synthesis (self-propagating high-temperature synthesis,SHS)is a novel method for preparation of AlN–SiC–TiB 2ceramics with many advantages,such as high efficiency and energy saving.However,combustion synthesis is a high temperature and velocity reaction process,which often result in large amount of non-equilibrium phases and high residual stress,and this deteriorates the proper-ties of the combustion synthesized products.So annealing treat-ment at elevated temperatures (800–1200°C)has been introduced to reduce the residual thermal stress and improve the mechanical properties [7,8].In this paper,the effects of annealing treatment on properties of the AlN–SiC–TiB 2ceramics prepared by self-propagating high tem-perature synthesis and hot isostatic pressing (SHS–HIP)were stud-ied.The influences of the annealing temperature and holding time on the composition,microstructure and mechanical properties were discussed in detail.2.Experimental procedureThe AlN–SiC–TiB 2ceramics were prepared by SHS–HIP reported in reference [9].The reaction for preparing AlN–SiC–TiB 2ceramics is as follows:Al þSiC þTiB 2þN 2!AlN þSiC þTiB 2ð1ÞThe samples with Al 35wt%–SiC 35wt%–TiB 230wt%were chosen to carry out the annealing treatment tests.The annealing treatments were carried out on the box-type electric furnace (SXK-8-16,Longjiang electrical furnace works,Harbin,China)in vacuum atmosphere.The specimens were enclosed and vacuumed in quartz glass tube.During the annealing treatment test,four different temperatures (600,800,1000and 1200°C)and five different holding times (2,4,6,8and 10h)were adopted to study the effects of annealing temperature and holding time on the microstructure and mechanical properties of the AlN–SiC–TiB 2ceramics.The composition of the annealing treated samples was analyzed by X-ray dif-fraction (D/max-rB,Rigaku,Japan).The microstructures were observed by scanning electron microscope (SEM,HITACHI S-4700),energy dispersive X-ray spectrometer,and transmission electron microscopy (TEM,Philips CM12/STEM,Holland).The specimen of 3mm Â4mm Â36mm and 2mm Â4mm Â22mm in dimension for bending strength and fracture toughness tests were sliced respectively.The three-point bending strength and fracture toughness were tested on electronic uni-versal test machine with crosshead speed of 0.5mm/min (Instron5569,USA).The Rockwell hardness test was carried out with a ½00ball and 60kg load applied for 30s on the Rockwell hardness test machine (HR-150A,China).0925-8388/$-see front matter Ó2012Elsevier B.V.All rights reserved./10.1016/j.jallcom.2012.11.080Corresponding author.Tel./fax:+865332781357.E-mail address:zhoulijuan@ (L.Zhou).3.Results and discussion3.1.MicrostructureFig.1shows the XRD patterns of theafter the annealing treatment for10h atAs thefigure indicated,oxidationnot occurred and the main phases wereTiB2after the annealing treatment.temperature and nitrogen pressure duringcontent of hexagonal boron nitride(h-BN)ucts.With the increasing of the annealingtion intensity of h-BN phase decreased,TiN phase appeared at the annealing10h,and slightly increases as the annealingto1200°C.The formation of the smallfavorable to the improvement of thermalfracture toughness because of theirand micro-cracking interactionkind of structural ceramic with excellenttion of h-BN phase during preparation anding treatment will not decrease the properties of the products,onthe contrary,it is helpful to the improvement of mechanical properties.Fig.2shows the XRD patterns of AlN–SiC–TiB2ceramics after annealing treatment at1200°C for different holding times.It can be found that the composition of the annealing treated ceramics was almost unchanged at1200°C for8h.When the holding time increased to10h,a small quantity of TiN was formed in the prod-uct,which was consistent with the XRD results shown in Fig.1. Furthermore,the X-ray diffraction peaks of2H–AlN–SiC phase were becoming narrower and sharper with the increasing of the annealing holding time.Meantime,the XRD peaks of the2H–AlN–SiC phase appeared to shift to larger angles,which indicated 3.2.Strengthening and toughening mechanismFig.3shows the BSE images of AlN–SiC–TiB2samples annealed for10h at600and1200°C.AlN–SiC matrix shows gray color and the dispersed white particles are TiB2phase.Fig.3a shows the sam-ple surface annealed at600°C for10h.White columnar TiB2were dispersed inhomogenously in the matrix,and also the sample sur-face presented many pores with different sizes and shapes.When the annealed temperature was1200°C,the dispersion of TiB2 was relatively homogenous andfiner,and the porosity of the sur-face was relatively decreased,as shown in Fig.3b.It is noted that some light gray phase exists between gray ma-trix and white particles,as shown in Fig.4.Line scanning for the light gray region was conducted to determine the phase composi-tion,as shown in Fig.4.The results show that the light gray phase contained large amount of Si element and tiny content of Al,and no Ti element was detected,which can be indicated that the light gray phase was the SiC-rich phase precipitated during the annealing treatment.After the precipitation of SiC,the AlN-rich and SiC-rich solid solutions were formed correspondingly,which was consis-tent with the phase diagram analysis[10].Phase separation leads to grain size refinement,and then results in the improvement of the yield strength.According to Hall–Petch equation,the relationship between the yield strength and grain size is described as follows[11]:ry¼r0þkdÀ1=2ð2Þwhere r y is the yield strength,r0is a materials constant for the starting stress for dislocation movement,k is the strengthening coefficient,and d is the average grain diameter.As the equation indicated,the grain size refinement is in favor of the improvement of yield strength of the materials.Fig.5shows the TEM image and SAED pattern of AlN–SiC–TiB2 ceramic annealed at1200°C for10h.Light gray matrix was AlN–SiC and the black particles embedded in the matrix were TiB2 phase,as shown in Fig.5a.The SAED pattern of the TiB2particles and the calibration were shown in Fig.5b and c,respectively.The average size of these TiB2particles was much less than those in reactant,and also,these particles were completely embedded in the AlN–SiC matrix,which indicated that thefine TiB2particles were precipitated during the annealing treatment.Since combustion synthesis is a high temperature and velocity reaction process,many dislocations,grain boundaries andFig.1.XRD patterns of AlN–SiC–TiB2ceramics annealed at different temperatures for10h.Fig.2.XRD patterns of AlN–SiC–TiB2ceramics annealed at1200°C for different holding times.500L.metastructures can be formed in the products.In addition,boron and titanium could dissolve in the AlN–SiC matrix or in the grain boundaries because of high reaction temperature.What is more,the standard enthalpy of formation (D H )of TiB 2is relatively low (D H =À293kJ mol À1),and the fine TiB 2particles can be precipi-tated from the supersaturated and non-equilibrium AlN–SiC ma-trix during the high temperature annealing treatment.Precipitating of these fine TiB 2particles decreased the lattice distortion,which also released the inner stress and improved the order degree of the lattice.Meantime,composition gradient formed by the precipitation between the matrix and TiB 2phases enhanced the interface bonding strength.In summary,the improvement of the strength and toughness of the AlN–SiC–TiB 2ceramics can be attributed to the particle dis-persion strengthening effect of the precipitated SiC-rich solid solution.And also the fine TiB 2particles precipitated uniformly at the phase interface enhanced the strength of the grain boundaries.nano-particles in AlN–SiC–TiB 2ceramic annealed at 1200°C for 10h:(a)TEM image of AlN–SiC–TiB 2composite;(b)SAED of light gray phase in the AlN–SiC–TiB 2ceramic annealed at 1200°C for 10h:(a)SEM image of the surface;(b)line Fig.3.Backscattered electron images of AlN–SiC–TiB 2samples annealed at (a)600°C,10h;(b)1200°C,10Fig. 6.Typical SEM fractograph in annealing treated AlN–SiC–TiB2ceramic composite.Fig.7.Rockwell hardness of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.Fig.8.Bending strength of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.Fig.9.Fracture toughness of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.annealing temperature reached800and1200°C.The highest HRA,bending strength and fracture toughness were94.1, 517.5MPa and 5.79MPa m1/2,respectively,after the annealing treatment at1200°C for10h.4.ConclusionsThe effects of annealing treatments on the microstructure and mechanical properties of AlN–SiC–TiB2ceramics prepared by SHS–HIP were studied.The annealing treatment can promote the precipitation SiC-rich solid solution in AlN–SiC matrix andfine TiB2particles on the grain boundaries,which improve the mechan-ical properties obviously.With the annealing temperature increas-ing from600to1200°C and holding times increasing from2to 10h,the strengthening effect was more evident.The highest HRA,bending strength and fracture toughness of AlN–SiC–TiB2 ceramics were94.1,517.5MPa and5.79MPa m1/2,respectively, after the annealing treatment at1200°C for10h.AcknowledgmentsThis work has beenfinancially supported by‘‘A Project of Shandong Province Higher Educational Science and Technology Program(J09LD01)’’and‘‘Shandong Province Natural Science Foundation(ZR2010EM045)’’.References[1]L.J.Zhou,Y.T.Zheng,S.Y.Du,H.B.Li,Mater.Sci.Forum546–549(2007)1505–1508.[2]vrenko,J.Desmaison,A.D.Panasyuk,M.Desmaison-Brut,E.Fenard,J.Eur.Ceram.Soc.25(2005)1781–1787.[3]K.Strecker,M.J.Hoffmann,J.Eur.Ceram.Soc.25(2005)801–807.[4]X.M.Yue,G.J.Zhang,Y.M.Wang,J.Eur.Ceram.Soc.19(1999)293–298.[5]M.Hotta,J.Hojo,J.Eur.Ceram.Soc.30(2010)2117–2122.[6]L.Mei,J.T.Li,Acta Mater.56(2008)3543–3549.[7]A.Varma,A.S.Rogachev,A.S.Mukasyan,S.Hwang,Adv.Chem.Eng.24(1998)79–226.[8]D.Sciti,S.Guicciardi,A.Bellosi,J.Eur.Ceram.Soc.21(2001)621–632.[9]L.J.Zhou,Y.T.Zheng,S.Y.Du,Key Eng.Mater.353–358(2007)1517–1520.[10]A.Zangvil,R.Ruh,J.Am.Ceram.Soc.71(1988)884–890.[11]H.Conrad,J.Narayan,Scripta Mater.42(2000)1025–1030.L.Zhou et al./Journal of Alloys and Compounds552(2013)499–503503。
Microstructure and mechanical properties of wroughtMg-4.1Li-2.5Al-1.7Zn-1Sn alloyRuizhi Wu1,2, a, Dayong Li1,b, Xuhe Liu2,c, and Milin Zhang2,d1College of Materials Science & Engineering, Harbin University of Science & Technology, Harbin,P.R. China 1500802Key Laboratory of Superlight Materials & Surface Technology (Harbin Engineering University),Ministry of Education, Harbin, P.R. China 150001a ruizhiwu2006@,b dyli@,c liuxuhe@,d zhangmilin@Keywords: Mg-Li alloy, deformation, microstructure, mechanical properties.Abstract.An Mg-Li-Al-Zn-Sn alloy was prepared by vacuum melting. The actual content of the elements in the alloy was determined using inductively coupled plasma-atomic emission mass spectrometry (ICP-AEMS). The density of the alloy was detected using Archimedes’ method. Extrusion and rolling deformation were carried out on this alloy. Its microstructures and mechanical properties were then studied with an optical microscope (OM), scanning electronic microscope (SEM), X-ray diffractometer (XRD), energy dispersive spectrometer (EDS), and tensile tester. The extruded alloy was composed of α-Mg and Mg2Sn phases and had good strength and elongation properties as well as a good comprehensive performance. After further rolling deformation, an Al-Li phase appeared due to atomic diffusion during the hot rolling process. Strain-hardening and the strengthening effect of the Al-Li phase further improved the strength of the alloy but decreased its elongation capacity.IntroductionSince the Mg-Li alloy was discovered in 1910, it has attracted a lot of attention from researchers because of its low density, high specific strength, stiffness, good processing performance, and dimensional stability. With these properties, it has shown great potential for use in applications in the aerospace, automotive, electronics, and defense industries [1-3].The Mg-Li alloy is the lightest metal material. Its binary phase diagram shows that, when the lithium content is less than 5.7%, the alloy displays an α single-phase that resembles a close-packed hexagonal structure. When the lithium content is more than 10.3%, the alloy shows a β single-phase, which is a body-centered cubic structure. At lithium contents between 5.7-10.3%, the alloy shows a two-phase organization [4]. Studies have shown that the addition of lithium causes the length of the c-axis in the hexagonal close-packed (HCP) structure to decrease, thus bringing about a decrease in the axial ratio c/a and rendering the alloy more able to undergo dislocation slip. This factor improves the deformability of the alloy [5].Al and Zn are two other elements that are commonly used in alloys. Appropriate amounts of Al added to alloys not only increases their strength and hardness but also improves their ductility and corrosion resistance. The addition of Zn can improve the deformation capacity of alloys [6, 7].Xiang Qi et.al[8] studied the influence of Sn on the microstructure and mechanical properties of a Mg-Li-Al-Zn alloy. Their results indicated that the addition of Sn refined the alloy, thereby improving its strength due to the formation of an Mg2Sn strengthening phase. When the Sn content was 1%, the grain size of the alloy was at the minimum size.Extrusion and rolling are the main deformation methods used for Mg-Li alloys. Deformation not only eliminates some casting defects but also causes dynamic recrystallization under certainconditions [9, 10]. This allows for the formation of refined alloys with improved comprehensiveIn this paper, a Mg-Li-Al-Zn-Sn alloy was prepared. After being subjected to two deformation modes, the microstructure and mechanical properties of the alloy obtained were determined.1. Materials and MethodsPure Mg, Li, Al, Zn, and Sn were used as raw materials. The alloy was created by vacuum induction melting using Ar gas as the protection gas in low-carbon steel molds. Homogenization treatments were performed at 300 o C for 24 h. The alloy ingot was extruded from Φ56 mm to 14 mm at 350 o C and denoted as an as-extruded alloy. As a last step, the extruded alloy was reheated and rolled at 260 o C to a final thickness of 3 mm.The chemical composition of the alloy was tested by inductively coupled plasma-atomic emission mass spectrometry (ICP-AEMS). The density of the alloy was determined by Archimedes’ method. The specimens for microscopic examinations were prepared using standard metallographic sample preparation methods. In brief, the specimens were etched with 1 vol% of nitric acid alcohol for 5-10 s.A LEICA DMIRM and JSM-6480 scanning electron microscopes (SEM) were used to observe the surface and fracture morphology of the alloy. A TTR III X-ray diffractometer (XRD) was utilized to identify the different phases in the alloy. An Energy Dispersive Spectrdmeter (EDS) was used to analysis micro-area composition.The tensile specimens were prepared according to the ASTM E8M-04 standard procedure. Tensile tests were carried out on an Instron4505 electronic universal testing machine with a speed of 1.5 mm/min. Five samples for each test were subjected to analysis along the extrusion and rolling directions.2. Results and Discussion2.1 The composition and density of the alloyThe analysis determined that the alloy composition was made of Mg-4.1Li-2.5Al-1.7Zn-1Sn. The density of the alloy was found to be 1.57 g/cm 32.2 The microstructure of the alloyThe microstructure of the extruded Mg-Li-Al-Zn-Sn alloy is shown in Figures 1a and 1b. The alloy was composed of a single α-phase, although some black matter appeared to be distributed in the matrix material. The grain size was small and its shape was equiaxial. These observations are typical of a material that has undergone dynamic recrystallization. Thus, it can be said that dynamic recrystallization occurred during the extrusion process.The microstructure of the alloy after rolling is shown in Figures 1c and 1d. Except for the black material, there also existed some eutectic compounds in the crystals, which may impact the performance of the alloy. The grain size for as-rolled samples was bigger than that for extrusion alloys.Figure 1. The microstructure of the alloy: (a) As-extruded alloy, (b) Magnified as-extruded alloy, (c) As-rolled alloy, and (d) Magnified as-rolled alloy.3.3 Phase analysisbelonged to Mg 22Sn in the 2Sn exists as a phase in the alloy.There were also sections of the rolled alloy that2θ/(°)I nte nsi t y /a .u .Figure 2. The XRD patterns of the alloy.30µmElements Wt./% At./%Mg 33.68 68.77Al 1.07 2.43Zn 1.47 1.36Sn 63.78 27.47A3.4 Mechanical properties of the alloyThe stress-strain curves of the extruded and rolled alloy are shown in Figure 4. Furthermore, Table 1 lists the mechanical properties of the two deformation state of the alloy. It also lists the corresponding performance parameters of commercially available LA141 Mg-Li alloy.Compared with the LA141 alloy, the as-extruded Mg-4.1Li-2.5Al-1.7Zn-1Sn had a higher tensile and yield strength with a considerable elongation capacity (>20%). Its specific strength and modulus are significantly higher than those of the LA141 alloy. These differences may be due to the α-phase (i.e., the α-phase alloy has higher strength than the β-phase alloy). It is also possible that Mg 2Sn, which was extensively distributed throughout the matrix, hinders dislocation slips when the alloy is deformed, thus playing a role in second phase strengthening.After rolling, the strength of the alloy was further increased, and its tensile strength reached 290.26 MPa. The elongation capacity of the alloy decreased but was still above 10%. The increase in strength may be explained in part by several factors, including work-hardening, additional deformation processes, and an increase in internal dislocation density. The latter causes flow stress to increase and improves the strength of alloys. The existence of an Al-Li phase after rolling could also be another reason for the increase in alloy strength. These two strengthening mechanisms, however, contribute to a decline in alloy plasticity. The alloy grain size increased, leading to a decline in its plasticity too.0510152025050100150200250300A s-extruded A s-rolledT e nsi le stre s s/MP a Strain/%Figure 4. Stress-strain curves of the alloy.Table 1. Mechanical properties of the alloyCondition As-extruded As-rolled LA141Tensile strength, MPa 267.51 290.26 144.69Specific tensile strength, cm ×105 170.39 184.87 105.03Yield strength, MPa 161.27 192.19 124.14Specific Yield strength, cm ×105 102.72 122.41 90.12Elastic modulus, GPa — 57.3 42.1Specific Elastic modulus,×106 — 36.49 31.12Elongation, % 21 11 24Density, g/cm 3 1.57 1.57 1.353.5 Fracture microstructure of the alloyThe fracture microstructure of the alloy is shown in Figure 5. The fracture microstructure of the as-extruded alloy was composed of a large number of small dimples. In some individual dimples,mechanism of the as-extruded alloy was ductile in nature. The fracture microstructure of the as-rolled alloy consisted of cleavage planes and a small number of dimples, which indicate that the fracture mechanism of the as rolled alloy can be ascribed to quasi-cleavage fractures.Figure 5. Fracture microstructure of the alloy. (A) As-extruded alloy and (B) As-rolled alloy.Summary1) An ultra-light Mg-Li-Al-Zn-Sn alloy was prepared by vacuum melting, then it was extruded and rolled. The density of the resulting alloy is 1.57 g/cm 3.2) The Mg(4.1)-Li(2.5)-Al(1.7)-Zn-Sn alloy ingot was subjected to two kinds of deformation processes: extrusion and rolling. The as-extruded alloy was found to be composed of the α-Mg and Mg 2Sn phases. After further rolling deformation, however, the alloy was found to consist of the α-Mg, Mg 2Sn, and Al-Li phases.3) Rolling deformations could further improve the strength of the alloy, but this resulted in a decrease of the elongation capacity.AcknowledgmentsThis work was supported by the National Natural Science Foundation of China (No. 51001034), China Postdoctoral Science Foundation(No. 20100481016) and Heilongjiang Postdoctoral Science Foundation.References[1]. R.Z. Wu, M.L. Zhang: Rev. Adv. Mater. Sci. Vol. 24 (2010), p.14[2]. H.Y Wu, Z.W. Gao and J.Y. Lin: J. Alloys Compd. Vol. 474 (2009), p.158[3]. Z.K. Qu, X.H. Liu, R.Z. Wu and M.L. Zhang: Mater. Sci. Eng. A Vol. 527 (2010), p.3284.[4]. L.Y.Wei, G.L.Dunlop and H.Westengen: Mater. Sci. Technol. Vol. 12 (1996), p.741[5]. C.H.Chiu, H.Y.Wu and J.Y.Wang: J. Alloys Compd. Vol. 460 (2008), p.246[6]. R.Z Wu, M.L Zhang: Mater. Sci. Eng. A Vol. 520 (2009), p.36[7]. T.C.Chang, J.Y.Wang and C.L.Chu: Mater. Lett. Vol. 60 (2006), p.3272[8]. Q. Xiang, R.Z.Wu and M.L. Zhang: J. Alloys Compd. Vol. 477 (2009), p.832[9]. T.C. Chang, J.Y. Wang and C.L. Chu: Mater. Lett. Vol. 60 (2006), p.3272[10]. R. Ninomiya and K. Niyake: J. Jpn. Inst. Met. Vol. 10 (2001), p.509[11]. R.Z Wu, Y.S Deng and M.L Zhang: J. Mater. Sci. Vol. 44 (2009), p.4132[12]. D.K. Xu, L. Liu and Y.B. Xu: Scripta Mater. Vol. 57 (2007), p.285(a)Material and Manufacturing Technology IIdoi:10.4028//AMR.341-342Microstructure and Mechanical Properties of Wrought Mg-4.1Li-2.5Al-1.7Zn-1Sn Alloydoi:10.4028//AMR.341-342.31。
International Journal of Minerals, Metallurgy and Materials Volume 25, Number 9, September 2018, Page 1060https:///10.1007/s12613-018-1657-9Corresponding authors: Yi-bo Li E-mail: yibo.li@; Song-sheng Zeng E-mail: zsscsu@© University of Science and Technology Beijing and Springer-Verlag GmbH Germany, part of Springer Nature 2018Microstructure, mechanical properties, and wear resistance ofVC p-reinforced Fe-matrix composites treated by Q&P processPing-hu Chen1,2), Yi-bo Li1,2,3), Rui-qing Li2,3), Ri-peng Jiang3), Song-sheng Zeng4), and Xiao-qian Li1,2,3)1) College of Mechanical and Electrical Engineering, Central South University, Changsha 410083, China2) State Key Laboratory of High Performance Complex Manufacturing, Changsha 410083, China3) Light Alloy Research Institutes, Central South University, Changsha 410083, China4) Valin ArcelorMittal Automotive Steel Co., Ltd., Loudi 417000, China(Received: 10 January 2018; revised: 4 April 2018; accepted: 13 April 2018)Abstract: A quenching and partitioning (Q&P) process was applied to vanadium carbide particle (VCp)-reinforced Fe-matrix composites (VC-Fe-MCs) to obtain a multiphase microstructure comprising VC, V8C7, M3C, α-Fe, and γ-Fe. The effects of the austenitizing temperature and the quenching temperature on the microstructure, mechanical properties, and wear resistance of the VC-Fe-MCs were studied. The re-sults show that the size of the carbide became coarse and that the shape of some particles began to transform from diffused graininess into a chrysanthemum-shaped structure with increasing austenitizing temperature. The microhardness decreased with increasing austenitizing tem-perature but substantially increased after wear testing compared with the microhardness before wear testing; the microhardness values im-proved by 20.0% ± 2.5%. Retained austenite enhanced the impact toughness and promoted the transformation-induced plasticity (TRIP) ef-fect to improve wear resistance under certain load conditions.Keywords: vanadium carbide; Fe-matrix composites; quenching and partitioning process; transformation-induced plasticity effect; micro-hardness; impact toughness; wear resistance1. IntroductionV anadium carbide [1] has advantages of high hardness, excellent high-temperature strength, good corrosion resis-tance, and high chemical and thermal stability even at high temperatures [2–4]; therefore, it is widely used in industrial applications [5–7]. V anadium carbide in alloys such as aus-tenite/martensite steel and cast iron is formed by chemical reaction at temperatures ranging from 1100 to 1500°C [2,8–9], which improves the mechanical properties and wear resis-tance of composites reinforced with vanadium carbide. Fur-thermore, the mechanical properties and wear resistance of vanadium carbide-reinforced composites are greatly en-hanced through optimization of the composition ratios of the carbon and vanadium elements and through application of a heat treatment to regulate the type and size of the micro-structure [10–15], thereby extending the service life of the composites [16].The quenching and partitioning (Q&P) process was first proposed by J. G. Speer in 2006. Plasticity and toughness can be improved through this process, which utilizes a car-bon partition to stabilize austenite [17–18]. The Q&P process has been adopted to adjust the ratio of austenite in the matrix to enhance the mechanical properties and wear behavior of high-strength steels. A hardening layer of the Q&P-treated specimen can form during wear tests because of the transformation-induced plasticity (TRIP) effect [19–24]. The optimized Q&P process and its TRIP effect are benefi-cial to enhancing the mechanical properties and wear resis-tance of the materials. Therefore, these new vanadium car-bide-reinforced Fe-matrix composites (VC-Fe-MCs) can re-place high-chromium and high-manganese wear-resistant cast iron in some industries. However, the study of VC-Fe-MCs on its properties has been reported rarely.In the present paper, the setting of austenitizing tempera-ture and quenching temperature was optimized to enhanceP.H. Chen et al., Microstructure, mechanical properties, and wear resistance of VC p-reinforced Fe-matrix (1061)the mechanical properties and wear resistance of VC-Fe-MCs to promote the application. The effects of heat-treatment parameters on the microstructures, mechani-cal properties, and wear resistance of VC-Fe-MCs were stu-died, and the TRIP effect before and after wear test was discussed.2. Experimental2.1. Materials and sample preparationThe chemical composition of the VC-Fe-MCs is listed in Table 1. The preparation process of VC-Fe-MCs is as fol-lows. Firstly, ferrous scrap was melted in a 200-kg me-dium-frequency induction furnace. Secondly, to increase the absorptivity of vanadium, ferrovanadium was added into liquid iron after being deoxidized. Thirdly, a certain amount of pure titanium was added to accelerate the nucleation of vanadium carbide, after which the melted iron was deox-idized. The final de-oxidation was conducted by adding 0.5wt% pure aluminum. Finally, the liquid cast iron was poured into a sand mold at approximately 1450°C and then formed by cooling in air. A schematic of the pouring process has been reported elsewhere [25]. The size of ingots were 200 mm × 100 mm × 14 mm. Samples for microhardness tests, impact tests, and wear tests were cut to small ingots with volume of 10 mm × 10 mm × 10 mm , 10 mm × 10 mm × 55 mm, and 10 mm × 10 mm × 40 mm, respectively. All samples were prepared for a batch of VC-Fe-MCs ingots.Table 1. Chemical composition of VC-Fe-MCs wt%C Si Mn V Cr Mo Fe2.80 1.53.0 8.1 2.0 1.5 Bal.2.2. Heat-treatment techniqueHeat treatments were carried out using a heat-treatment furnace. The heat-treatment schedules were based on the Q&P process and the TRIP effect. Specimens were austeni-tized at 900, 950, 1000, and 1050°C for 30 min and then quenched in a salt-bath furnace (55vol% KNO3 + 45vol% NaNO2 with a melting point of 130°C and service temper-ature range from 150 to 500°C) at 150, 180, 210, and 300°C respectively for several minutes before being parti-tioned at 320°C for 30 min and cooled to room tempera-ture.2.3. Measurement of mechanical properties and wear re-sistanceMaterial impact tests were carried out on a JBW-300B pendulum-type impact-testing machine with an impact energy of 150 J. Specimens were prepared without a notch. Each group included three specimens; the final reported value was an average of the results of the three specimens. The microhardness of the specimens before and after wear test was measured on a Zwick Roell Indentec Vickers tes-ter (ZHV1-AFC, Germany) with a load of 9.8 N for a dwell time of 15 s. Five points were measured for every specimen, with the reported value being the average of the five values. A wear resistance test was performed using a block-on-cylinder-type wear testing machine (M-2000) at room temperature with a load of 200 N at a rotating speed of 400 r/min. The material of the cylinder was YG8 hard alloy with a hardness of HRA 89 and outer and inner diameters of 40 mm and 16 mm, respectively. The weight loss (mg) of the wear specimens was measured using an analytical bal-ance, and the width and depth of the wear trace were meas-ured by a reading microscope (JC10, China) and an ul-tra-deep field 3D microscope (VHX5000, Japan), which provided a quantitative correlation among weight loss, wear width, wear depth, and wear time.2.4. Morphology characterization and composition anal-ysisThe microstructure of the specimens was characterized by scanning electron microscopy (SEM) on a TESCAN MIRA 3 LMH/LMU electron microscope operated at 15–25 kV. The samples were examined before and after they were etched with a 4vol% nitric acid alcohol solution for 30 s. The fracture morphology of specimens after the impact test was characterized by SEM. The chemical composition of the specimens was analyzed by energy-dispersive spectros-copy (EDS). Phase identification of the samples was carried out using X-ray diffraction (XRD) with a Cu Kα radiation; the samples were scanned at a rate of 0.04°/s over the 2θ in-terval from 10 to 90°.3. Results3.1. Microstructures and XRD analysisFig. 1 displays the microstructure and composition anal-ysis of an as-cast specimen. The microstructure was charac-terized by VC, V8C7, M3C, pearlite, and a small amount of martensite. The size of VC was between 5 μm and 10 μm, and a small amount of V8C7 was observed as long strips. On the basis of XRD composition analysis (Fig. 1(b)), the car-bide phases were identified as VC and V8C7. The crystal structure of VC was a 2 × 2 × 2 super cell of VC with 32 vanadium atoms and 32 carbon atoms. The crystal structure of V8C7 was the unit cell of V8C7 with 32 vanadium atoms1062 Int. J. Miner. Metall. Mater ., Vol. 25, No. 9, Sep. 2018and 28 carbon atoms [26]; the microstructure is shown in Fig. 1(c) and Fig. 1(d). α-Fe in the matrix was mainly pear-lite with a small amount of martensite. The microstructure is shown in Fig. 1(d).Fig. 2 shows the XRD composition analysis of the VC-Fe-MCs specimens austenitized at different temperatures and quenched at different temperatures. Compared with the XRD composition analysis of the as-cast specimen, it can be found that different proportions of the γ-Fe phase (retainedaustenite) were detected by XRD at different austenitizingFig. 1. Microstructure and XRD pattern of the as-cast specimen: (a) morphology and distribution of carbides; (b) XRD composi-tion analysis; (c) and (d) M 3C, V 8C 7, VC, pearlite, and small amount of martensite.Fig. 2. XRD composition analysis of VC-Fe-MCs under different (a) austenitizing temperatures and (b) quenching temperatures.P.H. Chen et al., Microstructure, mechanical properties, and wear resistance of VC p -reinforced Fe-matrix …1063temperatures and quenching temperatures. As shown in Fig. 2(a), the peaks associated with the γ-Fe phase (retained aus-tenite) at 50, 75, and 90° increased in intensity, nevertheless, the peaks associated with the α-Fe phase at 45 and 82° de-creased in intensity. The intensity of the VC peaks in-creased with increasing austenitizing temperature; however, with increasing quenching temperatures, the peak intensity of VC increased after initially decreasing, as shown in Fig. 2(b).3.2. Microhardness and toughnessThe effect of austenitizing temperature on impact tough-ness is shown in Fig. 3(a). The impact toughness initially increased and then decreased, reaching a toughness value greater than the initial value; the maximum value was 10.18 J/cm 2 at an austenitizing temperature of 1000°C. The effect of quenching temperature on impact toughness is shown in Fig. 3(b). Impact toughness increased with increasing quenching temperature. Fig. 4 shows the microhardness of specimens austenitized and quenched under different tem-peratures before and after wear test. The microhardness of the as-cast specimen was HRC 53.1 before and after wear test, whereas the microhardness of specimens changed sub-stantially as the austenitizing and quenching temperatures were varied. Specifically, the microhardness decreased with increasing austenitizing temperature. As displayed in Fig. 4(a), the microhardness of specimens at austenitizing tem-peratures of 1000°C and 1050°C were lower than that of the as-cast specimens before wear test, which were HRC 50.3 and HRC 43.3, respectively. However, after carrying out continuous wear test for 60 min, the values increased by 17.3% and 40.9%, respectively. Fig. 4(b) shows the micro-hardness at different quenching temperatures before and af-ter wear test. Before wear test, the microhardness of all the specimens was between HRC 47.5 and HRC 50.5. The mi-crohardness increased substantially after wear test compared with that before test, and the hardness values improved by20.0% ± 2.5%.Fig. 3.Variation in toughness under different (a) austenitizing temperatures and (b) quenching temperatures.Fig. 4. Variation in microhardness under different (a) austenitizing and (b) quenching temperatures before and after wear testing.1064Int. J. Miner. Metall. Mater ., Vol. 25, No. 9, Sep. 20183.3. Wear propertiesFig. 5(a) shows the weight losses at different austenitiz-ing temperatures. In general, weight loss increased with wear time, and a comparison of the results of the wear expe-riment show that the wear loss of the specimen austenitized at 900°C was lower than the others; the maximum value was only 8.7 mg, whereas that of the as-cast specimen was 16.4 mg after wear for 60 min. The wear loss ∆m at different austenitizing temperatures decreased in the order of ∆m AC >∆m 950 > ∆m 1000 > ∆m 1050 > ∆m 900, as shown in Fig. 5(a) (∆m AC is a wear loss of the as-cast specimen). The effect of quenching temperature on wear loss was shown in Fig. 5(b), which indicated that little difference exists in wear losses at different quenching temperatures. After a wear time of 60 min, the wear loss of the specimen treated at 300°C ranked behind only that of the as-cast specimen. The wear losses of the specimens treated at 150–210°C were virtually equal.The wear width of the austenitized specimens was shown inFig. 5. Variation in weight losses under different (a) austenitizing temperatures and (b) quenching temperatures; variation in widths under different (c) austenitizing temperatures and (d) quenching temperatures; and variation in depths under different (e) austenitizing temperatures and (f) quenching temperatures. AC represents the as-cast specimen.P.H. Chen et al., Microstructure, mechanical properties, and wear resistance of VC p -reinforced Fe-matrix …1065Fig. 5(c). The wear width exhibited a parabolic increase with wear time; however, the wear width of the 900°C-treated spe-cimen exhibited linear growth. After a wear time of 60 min, the relationship of wear width agreed with that of wear loss. The maximum wear width of the as-cast specimen was 4.2 mm, and that of the specimen treated at 900°C exhibited a maximum value of 3.3 mm. The wear width exhibited a pa-rabolic trend increase with increasing wear time when the quenching temperature was between 150°C and 300°C, as shown in Fig. 5(d). The maximum wear width of the speci-mens quenched at 180°C was the lowest (about 3.75 mm). The effect of the austenitizing temperature and quenching temperature on wear depth was shown in Fig. 5(e) and Fig. 5(f). The depth of the as-cast specimen was the largest, with an average value of 130 μm and a maximum value greater than 150 μm, as shown in Fig. 5(e). With an increase in aus-tenitizing temperature, the wear depth increased after in-itially decreasing, whereas the wear depth of the specimen treated at 950° reached the maximum value. As shown in Fig. 5(f), in contrast, the wear-depth trend decreased initially and then increased.4. Discussion4.1. Effect of heat treatment on mechanical properties On the basis of the experimental results, the impacttoughness of the 900°C/300°C-treated specimen (where 900°C and 300°C are the austenitizing and quenching tem-peratures, respectively) was the lowest among the investi-gated samples, whereas that for the 1000°C/300°C-treated specimen was the highest. As shown in Fig. 6, four fracture morphologies were observed in the fracture surface of the 900°C/300°C-treated specimen: zone I (spalling of spherical carbides), zone II (brittle fracture of dendritic carbide), zone III (ductile fracture of Fe-matrix), and zone IV (brittle frac-ture of Fe-matrix). Carbides were mainly distributed in the ductile fracture region (zone III). The size of the spherical carbide particles was between 5 μm and 10 μm, whereas the dendritic carbides were larger than 250 μm, as shown in Fig. 6(a) and Fig. 6(b). The bonding between the spherical car-bide and the Fe-matrix was poor, and a gap was readily ap-parent in the impact fracture. Cracks initiated in zone I un-der the impact condition. The bonding between the dendritic carbide and the Fe-matrix was quite good. The interdiffusion of alloying elements was detected between the dendritic carbide and the Fe- matrix, as shown in Fig. 6(c). However, the dendritic carbide was considered a ceramic particle, and brittle fracture occurred in zone II because the crack dif-fused from zone I to zone II (i.e., dendritic carbide). In addi-tion, according to the mass fraction of the carbon element in the designed materials, the VC-Fe-MCs belong to particlereinforced wear-resistant cast iron, which exhibited poorFig. 6. Fracture morphologies of the 900°C-treated specimen in different regions: (a) different fracture regions; (b) magnified view of ductile fracture and spherical carbide of zone I and zone II in subfigure (a); (c) magnified view of dendritic carbide of zone III in subfigure (a) and EDS composition analysis of line scanning; (d) magnified view of brittle fracture of zone IV in subfigure (a).1066Int. J. Miner. Metall. Mater ., Vol. 25, No. 9, Sep. 2018impact toughness. In the 1000°C-treated specimen, the carbon diffused from α-Fe to γ-Fe. A higher percentage of γ-Fe (austenite phase) could exist at room temperature be-cause the cooling time was longer than that for the 900°C-treated specimen. The spherical carbide particles of zone I appeared to break but not fall off, as shown in Fig. 7. Therefore, the impact toughness of 1000 treated spec ℃i-men was better than the others with the results of SEM morphology and EDS composition, as shown in Fig. 6 andFig. 7.Fig. 7. Fracture morphologies and EDS composition analysis of the 1000°C-treated specimen: (a) showing fracture microstructure for different fracture regions; (b) magnified view of spherical carbides, dendritic carbide and ductile fracture of zone I, zone II and zone IV in (a); (c) magnified view of brittle carbide of zone III in (a); (d) EDS composition analysis of line scanning.4.2. Effect of heat treatment on wear resistanceFig. 8 shows the wear morphology of the as-cast (AC), 900°C-treated, and 1000°C-treated specimens. Debris of the YG8 hard alloy was attached to the surface of the AC spe-cimen, whereas a large amount of granular precipitates dis-persed in the wear region of the 900°C-treated specimen. Debris of the YG8 hard alloy were detected on the wear surface of the tested specimen austenitized at 1000°C. As shown in Fig. 9, four different types of wear morphologies were characterized by SEM: type 1 (carbide); type 2 (debris of the YG8 hard alloy); type 3 (tested specimen); and type 4 (debris of the tested specimen). Therefore, evaluating whether the specimens treated at different austenitizing temperatures and quenching temperatures exhibited good wear resistance by measuring weight loss was difficult; in-stead, wear width and depth were taken into account to eva-luate the wear resistance, as shown in Fig. 5. 5. Strengthening mechanismAccording to previously reported experimental results [15], wear resistance mainly depends on the shape, size, and dis-tribution of vanadium carbide when the microhardness is greater than HRC 58. When the microhardness is less than HRC 58, the wear resistance of the VC-Fe-MCs depends pri-marily on its hardness and microstructure. Texture of the ma-trix is almost transformed completely into the austenite tex-ture when the austenitizing temperature is greater than 800°C. Through a rapid quenching treatment, most of the austenite texture is preserved in the matrix. However, the rest can re-main in an unstable state. In the subsequent process of car-bon partitioning, part of the austenite becomes more stable because of the redistribution of the carbon element. Howev-er, the remainder can change from γ-Fe (austenite texture) to α-Fe (martensite texture) during the final cooling process. The microstructure of the Fe-matrix can be transformedP.H. Chen et al., Microstructure, mechanical properties, and wear resistance of VC p -reinforced Fe-matrix (1067)from γ-Fe (good toughness) into α-Fe (excellent hardness and strength) under a certain friction load [27]. The micro-hardness improves greatly before and after wear test. Differ-ent proportions of retained austenite can be achieved through the Q&P process. Therefore, this process and the TRIP effect not only ensure impact toughness but also improve wear re-sistance during use. On the bases of these experimental re-sults, under friction conditions with either low impact or no impact, an austenitizing temperature of 900°C and a quench-ing temperature of 180°C are the optimal parameters for im-proving the microhardness and wear resistance of VC-Fe-MCs. However, under a high-impact condition, an austenitizing temperature of 1000°C and a quenching temperature of300°C are needed for excellent comprehensive performance.Fig. 8. Wear morphologies of typical specimens: (a) and (b) wear morphologies of an AC specimen; (c) and (d) wear morphologies of a specimen treated at 900°C; (e) and (f) wear morphologies of a specimen treated at 1000°C.1068 Int. J. Miner. Metall. Mater ., Vol. 25, No. 9, Sep. 2018Fig. 9. SEM morphologies (a -c) and EDS analyses (d) of points 1-4 in (b -c) for the 1000°C/300°C-treated specimen after wear test.6. ConclusionsThe microstructure of the Fe-matrix was translated from α-Fe into γ-Fe through a Q&P process, and γ-Fe was stabi-lized at room temperature through a carbon partitioning treatment. Thus, the TRIP effect was observed because γ-Fe (good toughness) was transformed into α-Fe (excellent hardness and strength) under a certain friction load. In addi-tion, the average hardness was improved substantially be-fore and after wear test. Under low-impact conditions, spe-cimens austenitized at 900°C and quenched at 180°C showed favorable wear resistance; however, under a cer-tain impact condition, the specimens exhibited good wear resistance when austenitized at 1000°C and quenched at 300°C.AcknowledgementsThis work was financially supported by the National Natural Science Foundation of China (Nos. 51475480 and U1637601), the Research Funding from the State Key La-boratory of High-Performance Complex Manufacturing (No. ZZYJKT2017-01), Innovation Platform and Talent Plan of Hunan Province (No. 2016RS2015), and the Project of In-novation Driven Plan in Central South University (No. 2015CX002). References[1] L.L. Wu, T.K. Yao, Y.C. Wang, J.W. Zhang, F.R. Xiao, andB. Liao, Understanding the mechanical properties of vana-dium carbides: Nano-indentation measurement and first-principles calculations, J. Alloys Compd., 548(2013), p. 60.[2] B. Zhang and Z.Q. Li, Synthesis of vanadium carbide bymechanical alloying, J. 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表面技术第53卷第5期高温对含氢DLC涂层的微观结构及力学性能的影响贾伟飞1，梁灿棉2，胡锋1,2*（1.武汉科技大学 高性能钢铁材料及其应用省部共建协同创新中心，武汉 430081；2.广东星联精密机械有限公司，广东 佛山 528251）摘要：目的针对含氢DLC涂层热稳定性很差的问题，探究高温下含氢DLC涂层的微观组织变化特征，以及高温对其力学性能的影响。
方法采用等离子体强化化学气相沉积（Plasma Enhanced Chemical Vapor Deposition, PECVD）在S136模具不锈钢表面沉积以Si为过渡层的含氢DLC复合涂层，利用光学显微镜、扫描电镜、拉曼光谱、X射线电子衍射仪、三维轮廓仪研究DLC涂层的微观结构，采用划痕测试仪、往复式摩擦磨损试验机、纳米压痕仪研究DLC涂层的力学性能，并通过LAMMPS软件，利用液相淬火法建立含氢DLC模型，模拟分析经高温处理后涂层的组织变化特征和纳米压痕行为。
结果在400 ℃、2 h的退火条件下，拉曼谱峰强度I D/I G由未退火的0.7增至1.5，涂层发生了石墨化转变，同时基线斜率下降，H元素析出；XPS结果表明，在此条件下涂层中sp2杂化组织相对增加，氧元素增多，涂层粗糙度增大；在600 ℃、2 h退火条件下，DLC发生了严重氧化，LAMMPS模拟结果表明，在400 ℃高温下涂层的分子键长变短，表明sp3杂化组织在高温下吸收能量，并向sp2杂化转变。
纳米压痕模拟结果显示，在400 ℃下退火后，涂层的硬度下降。
结论在400 ℃下退火处理后，涂层中的H元素释放，涂层内应力减小，保证了涂层的强度；在600 ℃退火条件下，过渡层的Si和DLC在高温下形成了C—Si键，使得DLC薄膜部分被保留；LAMMPS 模拟结果表明，在高温下涂层发生了石墨化转变，涂层的硬度减小。
关键词：含氢DLC涂层；退火处理；微观组织；力学性能；LAMMPS模拟中图分类号：TB332 文献标志码：A 文章编号：1001-3660(2024)05-0174-10DOI：10.16490/ki.issn.1001-3660.2024.05.018Effect of High-temperature on Microstructure and MechanicalProperties of Hydrogen-containing DLC CoatingJIA Weifei1, LIANG Canmian2, HU Feng1,2*(1. Collaborative Innovation Center for Advanced Steels, Wuhan University of Science and Technology, Wuhan 430081,China; 2. Guangdong Xinglian Precision Machinery Co., Ltd., Guangdong Foshan 528251, China)ABSTRACT: The thermal stability of hydrogen-containing DLC coating is poor, and the work aims to explore the microstructure changes of hydrogen-containing DLC coating at high temperature and their impact on mechanical properties. The收稿日期：2023-01-09；修订日期：2023-05-18Received：2023-01-09；Revised：2023-05-18基金项目：中国博士后科学基金（2021M700875）Fund：China Postdoctoral Science Foundation (2021M700875)引文格式：贾伟飞, 梁灿棉, 胡锋. 高温对含氢DLC涂层的微观结构及力学性能的影响[J]. 表面技术, 2024, 53(5): 174-183.JIA Weifei, LIANG Canmian, HU Feng. Effect of High-temperature on Microstructure and Mechanical Properties of Hydrogen-containing DLC Coating[J]. Surface Technology, 2024, 53(5): 174-183.*通信作者（Corresponding author）第53卷第5期贾伟飞，等：高温对含氢DLC涂层的微观结构及力学性能的影响·175·hydrogen-containing DLC composite coating with Si as the transitional layer was deposited on the surface of S136 stainless steel by plasma enhanced chemical vapor deposition (PECVD). The microstructure of DLC coating was investigated by optical/scanning electron microscopy, Raman spectroscopy, XPS (X-ray photoelectron spectroscopy) and three-dimensional profiler, the mechanical properties of DLC coating were studied by scratch, reciprocating friction wear and nano-indentation experiment, and the nano-indentation experiment behavior of DLC coating was simulated by LAMMPS to analyze the microstructure characteristics in annealing. The coating was subject to annealing conditions of 400 ℃for 2 hours and 600 ℃for 2 hours. Under the former condition, Raman spectroscopy showed an increase in the intensity ratio of the I D/I G peaks from0.7 to 1.5, indicating graphitization transition, accompanied by a decrease in baseline slope and H element segregation. XPSanalysis revealed an increase in sp2 hybridization and oxygen content in the coating under this condition, as well as an increase in surface roughness. At 600 ℃, severe oxidation of the DLC coating was observed. Under that condition, the matrix stainless steel was also oxidized. Molecular dynamics simulations using LAMMPS suggested a decrease in molecular bond length at 400 ℃high temperature. The three-dimensional profile test showed that the roughness under the unannealed condition was mainly from the large particles produced during deposition. At 400 for 2℃h, the coating had the minimum surface roughness. At this time, some large particles in the coating structure fell off, and the coating was basically completely damaged at 600 for℃ 2 h. The roughness was mainly from the original stainless steel roughness. The scratch test showed that under the condition of 400 for℃2 h, due to the release of the internal stress of the coating and the tighter bonding of the transition layer, the coating had the bestbonding effect with the substrate and was the least likely to fall off. The statistical results of LAMMPS simulation showed that the chemical bonds of the original DLC model tended to become shorter after annealing at high temperature. Relative to the unannealed DLC coating, the mechanical properties of DLC coating were best under 400 for℃ 2 h. Under this condition, the precipitation of mixed H elements in the coating led to the transformation of the original C—H sp3 structure, which occupied a large space to the smaller C—C sp3 and C—C sp2 structure, releasing internal stress in the coating, while ensuring the strength.The nano-indentation experiments showed that the elastic recovery and hardness of the coating were the highest at 400 for℃ 2 h, compared with that at other annealing temperature. The structure of the DLC coating containing hydrogen changed due to the precipitation of H element at 400 ℃. On the one hand, the coating structure changed from sp3 to sp2 due to high temperature, and on the other hand, the precipitation of H element changed the original C—H sp3 to C—C sp3, reducing the internal stress of the coating and improving the mechanical properties. The coating is basically damaged at 600 for 2 h, but the substrate still℃retains part of the coating. This is because the transition layer Si reacts with the coating to improve the heat resistance of the remaining coating. Molecular dynamics simulations using LAMMPS showed that the coating undergoes a graphitization transition at high temperature, leading to a reduction in its hardness.KEY WORDS: hydrogen-containing DLC coating; annealing treatment; microstructure; mechanical properties; LAMMPS simulationDLC（Diamond-Like Carbon，类金刚石碳，简称DLC）涂层材料具有超高硬度、低摩擦因数、优良化学稳定性等特点，广泛应用于机械、电子、生物医学等领域[1-3]。
